ARTICLE pubs.acs.org/Biomac
Colloidal Ionic Assembly between Anionic Native Cellulose Nanofibrils and Cationic Block Copolymer Micelles into Biomimetic Nanocomposites Miao Wang,† Anna Olszewska,‡ Andreas Walther,*,†,|| Jani-Markus Malho,† Felix H. Schacher,§ ‡ € Janne Ruokolainen,† Mikael Ankerfors,# Janne Laine,‡ Lars A. Berglund,^ Monika Osterberg, ,† and Olli Ikkala* †
Molecular Materials, Department of Applied Physics and ‡Department of Forest Products Technology, Aalto University (formerly Helsinki University of Technology), FIN-00076 Aalto, Espoo, Finland § Laboratory of Organic Chemistry and Macromolecular Chemistry, and Jena Center for Soft Matter, University of Jena, D-07743 Jena, Germany # Innventia AB, SE-11486 Stockholm, Sweden ^ Wallenberg Wood Science Center and Department of Fiber and Polymer Technology, Royal Institute of Technology, SE-10044 Stockholm, Sweden ABSTRACT:
We present a facile ionic assembly between fibrillar and spherical colloidal objects toward biomimetic nanocomposites with majority hard and minority soft domains based on anionic reinforcing native cellulose nanofibrils and cationic amphiphilic block copolymer micelles with rubbery core. The concept is based on ionic complexation of carboxymethylated nanofibrillated cellulose (NFC, or also denoted as microfibrillated cellulose, MFC) and micelles formed by aqueous self-assembly of quaternized poly(1,2-butadiene)block-poly(dimethylaminoethyl methacrylate) with high fraction of the NFC reinforcement. The adsorption of block copolymer micelles onto nanocellulose is shown by quartz crystal microbalance measurements, atomic force microscopy imaging, and fluorescent optical microscopy. The physical properties are elucidated using electron microscopy, thermal analysis, and mechanical testing. The cationic part of the block copolymer serves as a binder to NFC, whereas the hydrophobic rubbery micellar cores are designed to facilitate energy dissipation and nanoscale lubrication between the NFC domains under deformation. We show that the mechanical properties do not follow the rule of mixtures, and synergistic effects are observed with promoted work of fracture in one composition. As the concept allows wide possibilities for tuning, the work suggests pathways for nanocellulose-based biomimetic nanocomposites combining high toughness with stiffness and strength.
’ INTRODUCTION Nature exhibits composite materials with excellent mechanical properties, such as mollusk shells, bone, silk, antler, tooth enamel, or wood. Therein hierarchically ordered self-assemblies with a delicate interplay between the hard and soft domains are commonly observed13 where typically a large fraction of hard and reinforcing aligned segments are surrounded by a soft biopolymer matrix, such as proteins or polysaccharides. It is fascinating to realize how biology creates synergistic material properties combining high stiffness, strength, and toughness from mostly relatively simple components. Interestingly, the soft phase often only acts as an advanced binder, allowing crack guidance and deflection and nanoscale lubrication between the spatially closely located reinforcing segments, thus promoting toughness. Nanoscale friction reduction under high load r 2011 American Chemical Society
enables dissipative relative movement of the reinforcing domains, thus suppressing catastrophic fracture. For instance, proteins and fibroins provide soft separation of inorganic sheets in pearl of nacre, and lignin and hemicellulose are the biological lubricants between the reinforcing native cellulose microfibrils in wood.3 The resulting combination of stiffness, strength, and toughness with lightweight structure has spurred the quest to transfer such biological building concepts into synthetic materials within the field of biomimetics.4,5 The main challenge is to construct composites with high fractions of reinforcements that are Received: December 23, 2010 Revised: April 16, 2011 Published: April 26, 2011 2074
dx.doi.org/10.1021/bm101561m | Biomacromolecules 2011, 12, 2074–2081
Biomacromolecules hierarchically ordered into hard/soft mesostructured materials. Most importantly, we need to find pathways, components, and compositions that allow synergetic properties beyond the common rule of mixtures. A particular challenge is to identify routes for dissipative polymer layers between the reinforcements. In terms of efficient materials synthesis, we also have to develop strategies that allow their preparation with energy-efficient and fast processes. The most elegant approach is based on assembly of the components as encoded by their architecture to overcome the laborious sequential deposition techniques or multistep templating approaches. We recently demonstrated such a progress for technologically relevant nacre mimics based on the self-assembly of coreshell (hard/soft) polymer-coated nanoclay sheet-like building blocks.6,7 Herein, we aim to develop such a concept of hard/soft biomimetic ordered composites to stiff and strong native cellulose nanofibrils (NFC = nanofibrillated cellulose or also denoted as microfibrillated cellulose, MFC). The micro- and nanoscale native cellulose fibers have gained increasing attention due to their mechanical robustness combined with low weight, biodegradability, and renewability.8,9 In fact, the high-performance NFC, as derived from wood, is among the globally most abundant materials that do not interfere with the food chain and could therefore be one cornerstone for future biobased high technology materials applications. It has been shown earlier that cellulose I crystals within the native cellulose can reach a Young’s modulus of up to 136 GPa, and also a strength in the range of a few GPa might be achieved.10,11 These singularly attractive properties could already be translated into very strong nanopapers.12 Besides, a number of research groups explored nanocellulose composites with thermoset resins,13,14 poly(vinyl alcohol),15 amylopectin,16 poly(ethylene oxide),17 or poly(lactide),18 very often, however, only focusing on rather small amounts of reinforcements without selfassemblies and alignments, contrasting the concepts found in biological materials. Interestingly, most of the polymers used for composite applications are so far either water-soluble or postinfiltrated into NFC fiber sheets, requiring multistep processes. This exemplifies a main obstacle for a broader applicability of NFC, which is the poor solubility in non-aqueous dispersion and the often insufficient interaction with commonplace plastics. The aggregation of nanofibrillated cellulose in blends of hydrophobic polymers is a considerable problem.19 A study covering the full range of compositions was reported for composites made of NFC with amylopectin and glycerol.16 Therein, the mechanical properties evolved in a typical predictable and systematic fashion as a function of the ratio of the components: with decreasing amount of NFC, the stiffness and strength decreased, but the strain-to-failure of the materials increased, thus leading to a higher work of fracture. Consequently, the next challenge is to maintain the stiffness and strength, while increasing the strain-to-failure. This clearly requires synergies and calls for new concepts toward biomimetic materials. A control of the nanoscale hard/soft architecture and lubrication at the same time represents a promising approach.13,5,6,20 The first indication that such a method could provide success also in nanocellulose was demonstrated by in situ coating of biotechnologically produced bacterial cellulose (BC) with hydroxyethylcellulose (HEC) that was added to the culture medium of the organisms.21 Therein, the biologically grown BC with HEC coating displayed better mechanical properties than the pure BC or a postprepared blend of HEC and BC. This example shows that a well-defined soft compartmentalized shell can provide substantial synergies.
ARTICLE
Scheme 1. Concept of Complexation of Quaternized Poly(1,2-butadiene)-block-poly(dimethylaminoethyl methacrylate) Micelles with Anionic Nanofibrillated Cellulose (NFC) and Their Subsequent Packing into Biomimetic Hard/Soft Composites Due to Ionic Complexation of the Fibrillar and Spherical Colloidal Objects (Actual Block Copolymer Structure Used Herein Is Depicted on the Left-Hand Side)
In this work, the hypothesis was to overcome the nanoscale dispersion problem and to generate hard/soft biomimetic composites by ionically complexing two oppositely charged colloidal level objects, that is, negatively charged nanofibrillar cellulose and positively charged block copolymer micelles with soft cores. We hypothesized that the latter could promote dissipation of fracture energy upon deformation, lubrication between the NFC domains, and potentially even allow sacrificial bonding interactions.22,23 In the previous literature, in the molecular (smaller) length scale, oppositely charged polyelectrolytes can easily mutually complex or adsorb onto surfaces through electrostatic attraction, a phenomenon widely used for the electrostatic layer-by-layer growth of lamellar composite materials.24 In our approach, we use an amphiphilic diblock copolymer that carries one cationic, water-soluble function and another hydrophobic part with softness and deformability as a functionality. Such polymers are commercially available or can be synthesized in a large range of compositions and polymer sequences, thus giving rise to a platform concept. Herein, we chose a poly(1,2-butadiene)-blockpoly(dimethylaminoethyl methacrylate) block copolymer (PBb-PDMAEMA), whose amine group can be transformed into a strong polyelectrolyte via quaternization (methylation) of its nitrogen functions (PB-b-PDMAEMAq). Such amphiphilic block copolymers self-assemble in polar solvents, forming colloidal nanoparticles (micelles)25 where the hydrophobic segments collapse into a core,26 surrounded by a stabilizing corona of hydrophilic chains. We aimed at a beneficial lubricating effect of the hydrophobic rubbery polybutadiene phase because of its low glass transition temperature (Tg ≈ 15 C) and hence soft character at room temperature. Upon combination of the cationic PB-b-PDMAEMAq block copolymer micelles and carboxymethylated, anionically charged NFC nanofibrils, a colloidal complex is expected where the cationic polyelectrolyte serves as an anchor for the hydrophobic polybutadiene phase inside the micelles. The concept is illustrated in Scheme 1. In the following, we will present detailed results on the adsorption process and complex formation toward micelle/NFC colloidal complexes, as well as of the structure of the resulting nanocomposites after packing. Additionally, we will highlight the possible access to synergies in the mechanical performance of the biomimetic composites via optimized ratios of micelles to NFC and developing lubrication therein. 2075
dx.doi.org/10.1021/bm101561m |Biomacromolecules 2011, 12, 2074–2081
Biomacromolecules
’ EXPERIMENTAL SECTION Materials. Cellulose nanofibrils were disintegrated from bleached sulfite pulp by using a high-pressure fluidizer (Microfluidizer M-110EH, Microfluidics Corp., Newton MA) at Innventia AB, Stockholm, Sweden (formerly STFI-Packforsk) essentially following the procedures described in ref 27. The anionic NFC was prepared from bleached sulfite softwood dissolving pulp (Domsj€o ECO Bright). The pulp was carboxymethylated in a chemical pretreatment step and then run once through the microfluidizer.28,29 The initial NFC dispersion concentration was 1.87 wt %. Poly(1,2-butadiene)-b-poly(2-dimethylaminoethyl methacrylate) (PB-b-PDMAEMA) was prepared by anionic polymerization in THF with molecular weight Mn = 35 200 g/mol with Mw/Mn = 1.06 with 80 wt % of PB and 20 wt % of PDMAEMA blocks.30 N,N-Dimethylformamide (DMF) (g99.9%), tris(hydroxymethyl)aminomethane (Tris) at pH 8.3, and Nile Red dye were purchased from Sigma Aldrich (Germany). Quaternization and Preparation of Block Copolymer Micelles in Water. The PB-b-PDMAEMA block copolymer was dissolved in DMF to give a 2 g/L solution, which was deoxygenated by bubbling with nitrogen gas. A 5-fold excess of methyl iodide, compared to the nitrogen content of PDMAEMA was added with a syringe to the polymer solution, and the quaternization reaction was kept in a water bath at 50 C for 24 h.31 After quaternization, the polymer solution was filled in a dialysis tube (Spectra/Por 3500 Da MWCO, Canada) and dialyzed against Milli-Q water in a 10 L beaker for 36 h. The water was refreshed once during dialysis. This process yields spherical micelles with a PB core and a cationic polyelectrolyte corona of quaternized PDMAEMA, termed PB-b-PDMAEMAq (to be shown later). The polymer is selected to have a large weight fraction of the soft hydrophobic polybutadiene block coupled to a comparatively short cationized binder segment. The presence of large soft domains was expected to be beneficial to facilitate energy dissipation in the resulting composite materials. The amphiphilic balance of the diblock copolymer was designed to approach the so-called crew-cut regime, but still kept in a ratio so that the formation of spherical micelles is favored in contrast to worm-like ones or vesicles.26
Preparation of Nanofibril Dispersion for QCM-D and AFM. The NFC dispersion was diluted to 1.67 g/L and ultrasonicated using a microtip (Branson Digital Sonifire) for 10 min at 25% amplitude. After ultrasonication, the dispersion was centrifuged at 10 400 rpm for 45 min. The clear supernatant fraction was used to prepare ultrathin films on substrates. The end concentration of anionic fibril dispersion was 1.2 g/L, and it was further diluted to the 0.2 g/L at the desired pH. Preparation of Nanocomposites. Typically, 100 μL of Nile Red solution in acetone with a concentration of 0.5 g/L was added to 10 mL of PB-b-PDMAEMAq solution in order to stain the micelles with a fluorescent dye. The hydrophobic dye migrates into the core of the micelles. After sonication for 30 min, the solution was left open overnight to evaporate the acetone. The pH of the micellar solution was adjusted to 8.3 using 1 mM Tris buffer before mixing with NFC. Different mass ratios of NFC (pKa ≈ 4.8)28 and PB-b-PDMAEMAq block copolymer (1:1, 2.5:1, 5:1, and 10:1) were prepared via mixing with an Ultra-Turrax (IKA, T25 Basic) at 11 000 rpm for 2 min. Subsequently, entrapped air bubbles were removed by centrifugation at 2500 rpm for 5 min. Finally, composites were prepared by removing the water using vacuum filtration with 1.2 μm Millipore filter membranes.
Quartz Crystal Microbalance with Dissipation Monitoring (QCM-D). The adsorption of NFC and PB-b-PDMAEMAq on silica crystal was studied by means of quartz crystal microbalance with dissipation (QCM-D) with the E4 instrument from Q-Sense AB, V€astra Fr€olunda, Sweden. The instrument is equipped with an axial flow chamber, which allows the simultaneous measurements of both the frequency changes and the dissipation changes (frictional losses due to viscoelastic properties of the adsorbed layer) at the fundamental resonance frequency of 5 MHz and
ARTICLE
its overtones of 15, 25, 35, 45, 55, and 75 MHz. The third overtone (15 MHz) has primarily been used in the evaluation of the data. Right after starting the measurements, the pump was turned to the low flow rate of 0.1 mL/min. All experiments were repeated at least two times. The basic principles of the instrument have been described previously.32,33 Without adsorbate, the crystal oscillates at its resonance frequency, f0, and upon adsorption, the resonance frequency decreases to f. For rigidly and uniform adsorbed film, the change in frequency, Δf, is proportional to the adsorbed mass per unit surface, Δm (mg/m2), according to the Sauerbrey equation34 Δm ¼
CΔf n
where n is the overtone number and C is a device sensitivity constant. However, in the case of thicker films which show pronounced viscoelastic effects, the Sauerbrey equation is no longer valid. A model35 for the true sensed mass for viscoelastic layers has been used to calculate the mass adsorption: ! Ff 2 d2 _ m ¼ m0 1 þ ^J ðf Þ 3 where true sensed mass (m0) can be calculated by assuming that the shear compliance ^J (f) is independent of the frequency in the accessible frequency range. By plotting the equivalent mass against the square of the resonance frequency (f2), the true sensed mass is given as the intercept.36 Dynamic Light Scattering. Dynamic light scattering (DLS) (Malvern Zetasizer NanoZS) was used to determine the hydrodynamic radius of the micelles in solution at a concentration of 0.2 g/L in Tris buffer (1 mM, pH = 8.3). Atomic Force Microscope Imaging (in Air and Water). AFM imaging in tapping mode was used to characterize the PB-b-PDMAEMAq micelles in never-dried state adsorbed onto freshly cleaved mica and to characterize the morphology of the adsorbed block polymers on the NFC surface in air. A Nanoscope IIIa multimode scanning probe microscope (Digital Instruments Inc., Santa Barbara, CA) was used. For imaging in air, a silicon cantilever (NSC15/AIBS) delivered by MicroMasch, Tallinn, Estonia, with a driving frequency around 300360 kHz was used. The radius of the tip according to the manufacturer was less than 10 nm. A liquid cell was used for imaging the micelles in never-dried state. Soft silicon cantilever with a driving frequency between 7 and 11 kHz (Noncontact ultrasharp silicon cantilevers NSCS12, NT-MDT) was used. The image analysis was done with help of Scanning Probe Image Processor (SPIP) software (version 4.5.3, Image Metrology, Lyngby, Denmark). Only flattening was performed; no other image processing was applied. Optical Microscopy. A Leica DM4500 microscope (Leica Microsystems GmbH, Germany) was used to detect the binding of micelles on NFC in both transmission and fluorescence mode. Field-Emission Scanning Electron Microscopy (FE-SEM). A JEOL JSM-7500 FA field-emission scanning electron microscope was used to characterize the morphology of composites and determine the cross section thickness of composites. A thin Pt/Au layer was sputtered onto the specimens before imaging. Transmission Electron Microscopy (TEM). TEM imaging was performed with a JEOL JEM-3200FSC TEM operating at liquid nitrogen temperature. Zero-loss filtered images were acquired at an acceleration voltage of 300 kV with a Gatan Ultrascan 4000 camera in bright-field mode. The double bonds of the polybutadiene segments of the micelles were stained using OsO4 prior to imaging. For cryogenic TEM, a plasma-treated holey carbon grid was placed into a FEI Vitrobot, and a small droplet of NFC dispersion was applied onto it. After blotting, the grid was plunged into a liquid cryogen of 2076
dx.doi.org/10.1021/bm101561m |Biomacromolecules 2011, 12, 2074–2081
Biomacromolecules
ARTICLE
Figure 1. (a) Cryo-TEM micrograph showing anionic NFC nanofibrils in aqueous medium. (b) Distribution of the hydrodynamic diameter of PB-bPDMAEMAq micelles in water obtained by DLS. (c) Liquid AFM height images of PB-b-PDMAEMAq micelles adsorbed onto mica imaged in water. The concentration of the block copolymer was 0.2 mg/L. The z-range is 15 nm. (d) Height profile corresponding to the section analysis shown in (c). propane/ethane (1/1) and transferred into the TEM microscope under cryogenic conditions and imaged close to liquid nitrogen temperature. Thermal Gravimetric Analysis (TGA). A Mettler M3 TG50 thermobalance gravimetric analysis (Mettler Intrumente AG., Switzerland) was used to determine the thermal stability of composites. All samples were heated from 40 to 500 C at a heating rate of 10 C/min in air. Mechanical Testing. Mechanical tests were performed on a DEBEN minitester equipped with a 200 N load cell. All measurements were carried out at room temperature with a humidity of 3060%. The specimen had dimensions of 2 cm 2 mm 70200 μm. The exact specimens’ width and cross section thickness were determined by optical or scanning electron microscopy. At least five specimens were tested for each sample. A nominal strain rate of 0.5 mm/min was used. Digital speckle photography (DSP) was used to monitor the displacement of the sample upon applying stress. The speckle pattern was sprayed onto the sample surface. The Aramis Software of GOM GmbH (Germany) was used for digital image correlation and data computation.
’ RESULTS AND DISCUSSION The Starting Materials: Quaternized PB-b-PDMAEMAq Micelles and Carboxymethylated NFC. Figure 1 summarizes
the characterization of the two colloidal level components used for the ionic complexes to allow the biomimetic composites, that is, carboxymethylated, anionically charged nanofibrillated cellulose (NFC) and cationic PB-b-PDMAEMAq block copolymer micelles. The aqueous NFC is displayed in the cryo-TEM image in Figure 1a. The micrograph depicts micrometer-long NFC nanofibrils with an average diameter at 7.5 nm (standard deviation = 2.3 nm) and a length distribution in the range of a few micrometers. The carboxymethylation of the pulp during the processing leads to a degree of substitution of DS ≈ 0.1 in the nanofibrils28 and a charge density of ca. 575 μequiv/g. The quaternized PB-b-PDMAEMAq30 diblock copolymers form the micelles by self-assembly during the dialysis from DMF into water (see Figure 1bd). The size and shape of the micelles were
assessed by combining scattering and imaging techniques. The z-average hydrodynamic diameter of the PB-b-PDMAEMAq aggregates was determined using dynamic light scattering (DLS) to be 65 nm (Figure 1b) with a low polydispersity index of 1.07. The spherical nature of the micelles was further confirmed by liquid AFM imaging of never-dried micellar dispersion adsorbed onto freshly cleaved mica in water (Figure 1c). A rather uniform size of spherical micelles can be observed. The measured diameters from AFM images were found to be 80100 nm, and the height was around 1114 nm (Figure 1d). Taking into account the overestimation of the diameter due to tip deconvolution of the AFM tip, the dimensions point to some flattening of the micelles for the soft character of the polybutadiene core block with a low glass transition temperature, Tg ≈ 15 C. Adsorption of Quaternized PB-b-PDMAEMA Micelles onto Anionic NFC. In order to control the composite structure, we have to understand the binding between the cationic PB-bPDMAEMAq micelles and the anionic NFC nanofibrils. The two colloidal objects exhibit opposite charges, and thus we expect a tight binding via ionic complexation, similar to interpolyelectrolyte complexes at the smaller molecular length scale. In both molecular and colloidal cases, the main driving force for the complexation is the entropic gain due to the counterion release. Since anionic NFC is a weak polyelectrolyte with a pH-dependent ionization of the carboxylic acid groups, an increase of pH favors the dissociation and thus complexation. Therefore, we bind the PB-b-PDMAEMAq micelles onto the NFC nanofibrils in a Tris-buffered solution at pH = 8.3. This is sufficiently above the pKa value of typical carboxylic acids (3.05.5)37 and that of our NFC (4.8)28 and ensures a near quantitative degree of ionization to promote the complexation. Note that depending on the ratio of PB-b-PDMAEMAq micelles to NFC fibrils, different binding scenarios can be imagined. At low content of PB-b-PDMAEMAq micelles, the micelles can act as multifunctional binders between several NFC fibrils. On the other hand, if the micelles are considerably in excess, individual NFC fibrils are 2077
dx.doi.org/10.1021/bm101561m |Biomacromolecules 2011, 12, 2074–2081
Biomacromolecules
Figure 2. Adsorption of quaternized PB-b-PDMAEMA micelles onto anionic NFC. (a) Fluorescent optical microscopy images in aqueous medium upon mixing an aqueous dispersion of NFC with PB-bPDMAEMAq micelles in a ratio of 10/1 w/w. (b) QCMD data (third overtone) showing the change in frequency (0) and dissipation (b) upon first adding aqueous NFC dispersion (t = 0 min) and thereafter aqueous PB-b-PDMAEMAq micelles (t = 134 min). The inset shows a magnified section of the adsorption process of the PB-b-PDMAEMAq micelles on the NFC surface. (c) AFM height image of adsorbed PB-bPDMAEMAq micelles onto the NFC film imaged in air (z-range 15 nm). (d) Height profile of the section analysis indicated by the line in image (c).
presumed to be coated by a layer of micelles in the aqueous medium. We expect this to be important in the final packed assembled state to promote nanoscale lubrication between the NFC domains to allow toughening of the composites. Fluorescence microscopy allows visualization of the binding of the PB-b-PDMAEMAq micelles to NFC upon incorporation of a hydrophobic fluorescent dye (Nile Red) into the nonpolar PB core of the micelles (Figure 2a). The fluorescence micrograph in aqueous medium depicts fibrils and fibrillar aggregates, indicating a strong coordination of the dye-stained micelles with the NFC fibrils. The high contrast between dye-stained fibrils and nonfluorescing background demonstrates a strong affinity as free micelles would diminish the contrast. Due to the diffraction limit, it is not possible to resolve the individual nanoscale fibrils with their true diameter. Since the NFC is in large excess compared to the micelles (10/1 w/w), some clustering of nanofibrils into elongated aggregates can be observed at this particular ratio. An undesirable aggregation into large spherical agglomerates does not occur, most likely owing to the stiffness of the NFC fibrils. To better understand the colloidal complexation of PB-b-PDMAEMAq with anionic NFC, in situ adsorption experiments were performed at pH 8.3 using a quartz crystal microbalance with dissipation monitoring (QCM-D) (see Figure 2b). To enable NFC model films, the slightly anionic QCM-D crystal resonator substrates were first spin-coated using a polyethylene imine (PEI) primer layer. Next, aqueous NFC was added at t = 0, and Figure 2b shows a decrease in the QCM frequency (50 Hz), which indicates that NFC quickly adsorbs onto the PEI primer layer of the oscillator. The simultaneous increase in dissipation value (17 106) indicates that the NFC film is highly swollen and contains a substantial amount water. Next, addition of aqueous PB-b-PDMAEMAq micelles causes a further
ARTICLE
decrease of frequency (41 Hz), explained by the absorption of the PB-b-PDMAEMAq onto the NFC layer on the oscillator. Interestingly, the dissipation value decreases upon addition of the block copolymer (7 106), suggesting that water is removed from the NFC network during the adsorption process. Therefore, the film becomes more rigid. The Johannsmann’s model35 was used to calculate the adsorbed mass for NFC and PB-b-PDMAEMAq, leading to 3.5 and 3.0 mg/m2, respectively. The adsorption of PBb-PDMAEMAq on the NFC surface is fast (see the inset of Figure 2b), and the saturation value is reached in less than 1 min. This rapid process underscores the high affinity of the cationic block copolymer micelles to the negatively charged NFC surface. After adsorption, the film was dried and imaged by AFM in air. Figure 2c shows roughly globular structures on the NFC, where the diameter is in the range of 80150 nm and height in the range of 814 nm. Such structures point toward adsorbed micelles on the NFC network layer. Although the glass transition temperature of PB is low (ca. 15 C) and the micelles have soft cores, we do not observe evidence of opening and film formation of the micelles on the surface. This result has major implications for the composite structure and behavior. Structure of the NFC/PB-b-PDMAEMAq Nanocomposites. Next we proceed to the composites upon water removal using simple vacuum filtration. Our interest focuses on the biomimetic compositions with a large fraction of the reinforcing NFC domains as separated by a small weight fraction of soft dissipative block copolymer domains. Therefore, we chose the NFC/PB-b-PDMAEMAq compositions in a range of 10/1, 5/1, 2.5/1, and 1/1 w/w. The composites are translucent on the macroscale (Figure 3a). This indicates the absence of macroscopic phase separation of NFC and block copolymer, which would lead to opaque appearance. The attractive ionic binding between the oppositely charged nanofibrillar and spherical colloids obviously prevents the macroscopic phase separation. Note that a simple mixing of polybutadiene and NFC would lead to macroscopic phase separation, poor interfacial bonding, and little interaction between both components. Scanning electron microscopy shows a nanofibrillar network structure (Figure 3b) and a layered internal structure on the mesoscale (Figure 3c). The NFC/PB-b-PDMAEMAq colloidal complexes are orientated in a random plane fashion to form a web-like fibrous network along the substrates during the vacuum filtration. The fracture surface displays some vertical holes and protrusion between the various nanocellulose layers, thus indicating that the intralaminar NFC interaction is stronger than the interlaminar one. Although the composites may have a certain limited void content, previous work indicates that this is not critical to mechanical performance.16 This is a consequence of the random in-plane orientation distribution of NFC. We can find similar layered structures for all studied NFC/PB-b-PDMAEMAq composites. In order to investigate the structural details down to the nanoscale, we used cross-sectional TEM, where the contrast of the PB domains of the block copolymer was increased via selective staining of the double bonds with OsO4. The TEM micrograph for a NFC/PB-b-PDMAEMAq block copolymer nanocomposite is presented in Figure 3d,e, showing approximately periodic structure with alternating NFC and PB-b-PDMAEMAq domains with a periodicity of around 160200 nm. Interestingly, the micrographs also indicate rather equally spaced darker spheres within the PB-b-PDMAEMAq domains. Since the darker parts in the micrograph originate from the stained PB double bonds, we conclude that the PB-b-PDMAEMAq micelles are observed also in the composites upon drying and are aligned in a pearl necklace fashion on NFC strands (see Figure 3e 2078
dx.doi.org/10.1021/bm101561m |Biomacromolecules 2011, 12, 2074–2081
Biomacromolecules
ARTICLE
Figure 3. Structures of NFC/PB-b-PDMAEMAq composites: (a) photograph of composites having the NFC/PB-b-PDMAEMAq fractions of 2.5/1 and 5/1 w/w. Note that the composites are colored due to the staining by Nile Red dye, used to follow the complexation. (b,c) SEM images of the top surface and fracture surface of nanocomposites for the NFC/PB-b-PDMAEMAq sample with the weight fraction 2.5/1 w/w. (d,e) Cross-sectional TEM image of the same composite and section analysis.
Figure 4. TGA analysis of the thermal degradation of NFC/PB-bPDMAEMAq composites. The labels indicate the mass ratio of NFC to PB-b-PDMAEMAq in the nanocomposites (10/1 is omitted for clarity). For clarity, the arrows indicate the changes induced by the added PB-bPDMAEMAq weight fraction in the composites.
and schematically in Scheme 1). This observation also coincides with the AFM imaging of the intact micelle layer in the QCM-D adsorption studies. The micellar shape is retained although the low glass transition temperature of the PB part could, in principle, allow for sufficient dynamics and a disintegration of the micelles. Note that the dark domains coincide well with the size range of the micelles, especially considering that staining with the bulky OsO4, flattening, and microtome cutting can increase the observed size in the dried state. The micelles are well-separated and do not form a continuous film. Since the periodicity perpendicular to the lamellae is closer to the diameter of a micelle than to individual NFC fibrils, it is likely that the micelles dominate the formation of the assembly at this length scale. Although the structure appears relatively regular, small-angle X-ray scattering only displayed weak correlation peaks (not shown). This is, however, understandable, considering the polydisperse nature of the NFC nanofibril diameter and that of the micelles.
The thermal stability and the actual composition of the NFC/ PB-b-PDMAEMAq composites can be inferred from thermogravimetric analysis (TGA, Figure 4). The complexation of the block copolymer micelles leads to changes in the thermal stability of the nanocomposites. As comparison, the onset of degradation for the pure NFC starts around 260 C and exhibits a second degradation step just above 400 C. This is in the typical range of cellulose paper with a degradation range of 230360 C.3841 The PB-bPDMAEMAq block copolymer also shows a two-step degradation profile with a low temperature and less pronounced first step at ca. 220 C and a second major degradation step just above 400 C. The NFC/PB-b-PDMAEMAq composites show a shift of the first degradation step to slightly higher temperature. All curves lie above the TGA trace of pure NFC. This behavior is unexpected and is beyond additive behavior of mixtures, where a simple additive TGA curve of both components would result in a shift of the first decomposition to smaller temperature as compared to NFC, thus right between NFC and PB-b-PDMAEMAq. Consequently, the polyelectrolyte attraction between NFC and PB-bPDMAEMAq can be identified as the origin of this delayed degradation. In the intermediate decay region (300400 C), a continuous shift of the plateau level from pure NFC to pure PB-bPDMAEMAq can be observed. This difference may be used to estimate the weight fractions of the PB-b-PDMAEMAq in the composite materials (see Table 1). Accordingly, such analyzed fractions of NFC versus PB-b-PDMAEMAq correspond closely to the feed fraction. This indicates that the high affinity of the cationic binder segments to the anionic cellulose nanofibrils prevents both a macrophase separation as well as a physical separation of the micelles through the large pores of the filter. Consequently, the strategy allows well-defined hard/soft nanocomposites with ordered mesoscale structures and predictable composition. Mechanical Properties. After demonstrating the binding and microstructure of the complexes between NFC nanofibers and PB-b-PDMAEMAq micelles, we next turn to the mechanical properties. The stressstrain curves, work of fracture, Young’s modulus, and strain-to-break of the composites series are indicated in Figure 5 and summarized in Table 1. For pure NFC 2079
dx.doi.org/10.1021/bm101561m |Biomacromolecules 2011, 12, 2074–2081
Biomacromolecules
ARTICLE
Table 1. Tensile Properties of the NFC/PB-b-PDMAEMAq Composites Where the Fraction of Rubbery Block Copolymer Increases toward the Bottoma mass fraction NFC (%)
NFC NFC/PB-b-PDMAEMAq
Young’s modulus
tensile strength
strain-to- failure
work of fracture
sample code
feed
expb
(GPa)
(MPa)
(%)
(MJ/m3)
10:1
100 91
94
12.5 (1.0) 10.4 (1.2)
202.2 (21.0) 166.2 (13.7)
6.8 (1.4) 4.5 (0.9)
9.8 (3.0) 5.5 (1.4) 2.5 (0.8)
5:1
83
88
11.1 (1.2)
144.8 (16.2)
2.4 (0.5)
2.5:1
71
77
7.3 (1.0)
165.2 (21.8)
9.1 (0.7)
10.5 (2.1)
1:1
50
54
6.2 (0.2)
95.8 (1.2)
4.6 (0.5)
3.3 (0.5)
Data refer to mean values of pure NFC and NFC/PB-b-PDMAEMAq composites and corresponding standard deviations in brackets of at least five samples. b Estimated from the plateau area of the TGA data.
a
Figure 5. Stressstrain curves for NFC/PB-b-PDMAEMAq composites (a). Mass ratios of NFC and PB-b-PDMAEMAq are indicated. (b) Work of fracture as a function of the composition. The corresponding curves and values of the pure reinforcement NFC are shown as comparison.
paper, the Young’s modulus can reach 12.5 GPa (Table 1), which is a commonly found value for nanocellulose paper.16 Upon small addition of PB-b-PDMAEMAq (compositions 10/1 and 5/1), the moduli become slightly decreased to values of 10.4 and 11.1 GPa (Table 1). Their stressstrain curves are very similar, indicating an only minor influence of the amount of PB-bPDMAEMAq (see Figure 5a). Still, both curves are remarkably different than the pure NFC, thus pointing to a different interaction between the nanofibrils. Further increase of PB-bPDMAEMAq toward the composition 1/1 leads to a decrease of the modulus to 6.2 GPa. The modulus decreases as expected as a function of added softer PB domains. However, we find strong nonlinearities when looking at stress at break, maximum extension, and in particular at toughness as a function of the composition (Figure 5b). Toughness can be defined as the work of fracture as obtained from the area under the stressstrain curve.12 We identify a synergetic behavior of the mixture at a ratio of 2.5/1. The extension at break and also the toughness are reproducibly 23 times higher than those of the other compositions. Even compared to the pure NFC, we still find a more than 20% higher strain-to-failure. The suppressed brittleness of this synergetic mixture also allows reaching a high strength. We tentatively ascribe the synergy to the PB-b-PDMAEMAq micelles between the reinforcing NFC domains, where the cationic micellar shells serve as anchors to create assembled domains, while the soft PB part dissipates energy and allows friction reduction between the reinforcing NFC domains under stress. It has been demonstrated in literature that the crack propagation in the nanofibrillated TEMPO-oxidized celluloses with hydroxypropylcellulose matrix is similar to fracture-toughened intercalated
clayepoxy nanocomposites,42 in which cracks follow a torturous path. Within the context of our work, we suggest that the soft PB phase serves two purposes. First, it enables a more effective crack deflection at the hard/soft (NFC/PB-b-PDMAEMAq) interface as compared to polymers with high Tg, and therewith it extends the crack path.2 Additionally, the lubrication of the nanoscale fibers by synthetic block copolymer with low Tg can favor a fiber pullout versus a fiber fracture and henceforth contribute a second attractive benefit for toughening.
’ CONCLUSIONS We have implemented a new, simple, and versatile platform concept to complex native nanofibrillated cellulose with block copolymers via noncovalent interactions. This allows assembled biomimetic nanocomposites that contain a large fraction of NFC reinforcements separated by rubbery domains. The concept is demonstrated using anionically charged carboxymethylated NFC as ionically complexed using quaternized poly(1,2-butadiene)block-poly(dimethylaminoethyl methacrylate) block copolymer (PB-b-PDMAEMAq) micelles, possessing a cationic polyelectrolyte corona and a rubbery micelle core. The complexation between the colloidal nanofibrils and colloidal micellar spheres prevents macroscopic phase separation and gives rise to composites with an alternating nanoscale hard/soft architecture. As a consequence of the improved mesoscale order, we identify a synergetic performance of the components beyond the simple rule of mixture. For a ratio of NFC/PB-b-PDMAEMAq of 2.5/1 w/w, significantly larger strain and toughness could be achieved than in the other studied compositions, potentially due to the nanoscopic control 2080
dx.doi.org/10.1021/bm101561m |Biomacromolecules 2011, 12, 2074–2081
Biomacromolecules over fracture energy dissipation and the interfibrillar lubrication between the NFC domains by using low Tg amphiphilic block copolymers. Since block copolymer structures can be widely tuned for different molecular weights, block constituents, and architectures, we expect great possibilities to pursue biomimetic nanocomposites with significantly increased fracture toughness, while maintaining stiffness and strength.
’ AUTHOR INFORMATION Corresponding Author
*E-mail:
[email protected], olli.ikkala@tkk.fi. )
Present Addresses
DWI at the RWTH Aachen University, D-52056 Aachen, Germany.
’ ACKNOWLEDGMENT We thank Prof. Tom Lindstr€om, Prof. Lars Wagberg, Prof. Gero Decher, and Prof. Ludwik Leibler for discussions. We thank the companies involved in project Designcell/Woodwisdom.net and the financial support, National Technology Agency of Finland, Vinnova, Wallenberg Wood Science Center and Academy of Finland. We thank Christoffer Johans for measuring DLS. ’ REFERENCES (1) Bhushan, B. Philos. Trans. R. Soc. London, Ser. A 2009, 367 1445–1486. (2) Fratzl, P.; Gupta, H. S.; Fischer, F. D.; Kolednik, O. Adv. Mater. 2007, 19, 2657–2661. (3) Meyers, M. A.; Chen, P.-Y.; Lin, A. Y.-M.; Seki, Y. Prog. Mater Sci. 2008, 53, 1–206. (4) Podsiadlo, P.; Kaushik, A. K.; Arruda, E. M.; Waas, A. M.; Shim, B. S.; Xu, J.; Nandivada, H.; Pumplin, B. G.; Lahann, J.; Ramamoorthy, A.; Kotov, N. A. Science 2007, 318, 80–83. (5) Munch, E.; Launey, M. E.; Alsem, D. H.; Saiz, E.; Tomsia, A. P.; Ritchie, R. O. Science 2008, 322, 1516–1520. (6) Walther, A.; Bjurhager, I.; Malho, J.-M.; Pere, J.; Ruokolainen, J.; Berglund, L. A.; Ikkala, O. Nano Lett. 2010, 10, 2742–2748. (7) Walther, A.; Bjurhager, I.; Malho, J.-M.; Ruokolainen, J.; Berglund, L. A.; Ikkala, O. Angew. Chem., Int. Ed. 2010, 49, 6448–6453. (8) Siro, I.; Plackett, D. Cellulose 2010, 17, 459–494. (9) Eichhorn, S. J.; Dufresne, A.; Aranguren, M.; Marcovich, N. E.; Capadona, J. R.; Rowan, S. J.; Weder, C.; Thielemans, W.; Roman, M.; Renneckar, S.; Gindl, W.; Veigel, S.; Keckes, J.; Yano, H.; Abe, K.; Nogi, M.; Nakagaito, A. N.; Mangalam, A.; Simonsen, J.; Benight, A. S.; Bismarck, A.; Berglund, L. A.; Peijs, T. J. Mater. Sci. 2010, 45, 1–33. (10) Iwamoto, S.; Kai, W. H.; Isogai, A.; Iwata, T. Biomacromolecules 2009, 10, 2571–2576. (11) Yano, H.; Sugiyama, J.; Nakagaito, A. N.; Nogi, M.; Matsuura, T.; Hikita, M.; Handa, K. Adv. Mater. 2005, 17, 153–155. (12) Henriksson, M.; Berglund, L. A.; Isaksson, P.; Lindstr€ om, T.; Nishino, T. Biomacromolecules 2008, 9, 1579–1585. (13) Nakagaito, A. N.; Yano, H. Appl. Phys. A 2004, 78, 547–552. (14) Nakagaito, A. N.; Yano, H. Appl. Phys. A 2005, 80, 155–159. (15) Millon, L. E.; Wan, W. K. J. Biomed. Mater. Res., Part B 2006, 79B, 245–253. (16) Svagan, A. J.; Azizi Samir, M. A. S.; Berglund, L. A. Biomacromolecules 2007, 8, 2556–2563. (17) Brown, E. E.; Laborie, M.-P. G. Biomacromolecules 2007, 8, 3074–3081. (18) Okubo, K.; Fujii, T.; Thostenson, E. T. Composites, Part A 2009, 40, 469–475.
ARTICLE
(19) Wu, Q.; Henriksson, M.; Liu, X.; Berglund, L. A. Biomacromolecules 2007, 8, 3687–3692. (20) Bonderer, L. J.; Studart, A. R.; Gauckler, L. J. Science 2008, 319, 1069–1073. (21) Zhou, Q.; Malm, E.; Nilsson, H.; Larsson, P. T.; Iversen, T.; Berglund, L. A.; Bulone, V. Soft Matter 2009, 5, 4124–4130. (22) Hartmann, M. A.; Fratzl, P. Nano Lett. 2009, 9, 3603–3607. (23) Fantner, G. E.; Hassenkam, T.; Kindt, J. H.; Weaver, J. C.; Birkedal, H.; Pechenik, L.; Cutroni, J. A.; Cidade, G. A. G.; Stucky, G. D.; Morse, D. E.; Hansma, P. K. Nat. Mater. 2005, 4, 612–616. (24) Decher, G. Science 1997, 277, 1232–1237. (25) Chen, Z.; He, C.; Li, F.; Tong, L.; Liao, X.; Wang, Y. Langmuir 2010, 26, 8869–8874. (26) Walther, A.; Goldmann, A. S.; Yelamanchili, R. S.; Drechsler, M.; Schmalz, H.; Eisenberg, A.; M€uller, A. H. E. Macromolecules 2008, 41, 3254–3260. (27) P€a€akk€ o, M.; Ankerfors, M.; Kosonen, H.; Nyk€anen, A.; Ahola, € S.; Osterberg, M.; Ruokolainen, J.; Laine, J.; Larsson, P. T.; Ikkala, O.; Lindstr€ om, T. Biomacromolecules 2007, 8, 1934–1941. (28) Wagberg, L.; Decher, G.; Norgren, M.; Lindstr€om, T.; Ankerfors, M.; Axn€as, K. Langmuir 2008, 24, 784–795. (29) Aulin, C.; Johansson, E.; Wagberg, L.; Lindstr€om, T. Biomacromolecules 2010, 11, 872–882. (30) Schacher, F.; M€ullner, M.; Schmalz, H.; M€uller, A. H. E. Macromol. Chem. Phys. 2009, 210, 256–262. (31) Mori, H.; Walther, A.; Andre, X.; Lanzend€orfer, M. G.; M€uller, A. H. E. Macromolecules 2004, 37, 2054–2066. (32) Rodahl, M.; H€o€ ok, F.; Krozer, A.; Brzezinski, P.; Kasemo, B. Rev. Sci. Instrum. 1995, 66, 3924–3930. (33) H€o€ok, F.; Rodahl, M.; Brzezinski, P.; Kasemo, B. Langmuir 1998, 14, 729–734. (34) Sauerbrey, G. Z. Phys. 1959, 155, 206–222. (35) Johannsmann, D.; Mathauer, K.; Wegner, G.; Knoll, W. Phys. Rew. B 1992, 46, 7808–7815. (36) Naderi, A.; Claesson, P. M. Langmuir 2006, 22, 7639–7645. € (37) Taipale, T.; Osterberg, M.; Nyk€anen, A.; Ruokolainen, J.; Laine, J. Cellulose 2010, 17, 1005–1020. (38) Nystr€ om, G.; Mihranyan, A.; Razaq, A.; Lindstr€om, T.; Nyholm, L.; Stroemme, M. J. Phys. Chem. B 2010, 114, 4178–4182. (39) Johnson, R.; Zink-Sharp, A.; Renneckar, S.; Glasser, W. Cellulose 2009, 16, 227–238. (40) Nadagouda, M. N.; Varma, R. S. Biomacromolecules 2007, 8, 2762–2767. (41) Braun, B.; Dorgan, J. R. Biomacromolecules 2009, 10, 334–341. (42) Zerda, A. S.; Lesser, A. J. J. Polym. Sci., Part B: Polym. Phys. 2001, 39, 1137–1146.
2081
dx.doi.org/10.1021/bm101561m |Biomacromolecules 2011, 12, 2074–2081
