Article pubs.acs.org/cm
Formation of Periodically-Ordered Calcium Phosphate Nanostructures by Block Copolymer-Directed Self-Assembly Rui-Qi Song,† Tobias N. Hoheisel,† Hiroaki Sai,†,⊥ Zihui Li,† Joseph D. Carloni,† Suntao Wang,‡ Randall E. Youngman,§ Shefford P. Baker,† Sol M. Gruner,‡,∥,¶ Ulrich Wiesner,*,† and Lara A. Estroff*,†,∥ †
Department of Materials Science and Engineering, Cornell University, Ithaca, New York 14853, United States Department of Physics, Cornell University, Ithaca, New York 14853, United States § Science & Technology Division, Corning Incorporated, SP-AR-02-4, Corning, New York 14830, United States ∥ Kavli Institute at Cornell for Nanoscale Science, Ithaca, New York 14853, United States ¶ Cornell High Energy Synchrotron Source, Cornell University, Ithaca, New York 14850, United States ‡
S Supporting Information *
ABSTRACT: Structuring ionic solids at the nanoscale with block copolymers (BCPs) is notoriously difficult due to solvent incompatibilities and strong driving forces for crystallization of the inorganic material. Here, we demonstrate that elucidating pathway complexity in the BCP-directed selfassembly of an ionic solid, amorphous calcium phosphate (ACP), is a key component in obtaining nanostructured, bulk composite materials in which the nanostructure is the result of thermodynamically controlled BCP self-assembly, i.e., exhibiting sequences of bulk morphologies as known from typical equilibrium BCP phase diagrams. Specifically, we identify three critical pathway “decision points” for the evaporation-induced self-assembly of composites from ultrasmall, organosilicatemodified amorphous calcium phosphate nanoparticles (osm-ACP-NPs) and poly(isoprene)-block-poly(2-(dimethylamino)ethyl methacrylate) (PI-b-PDMAEMA) block copolymers. Using this strategy enabled us to obtain composites with hexagonal, cubic network, and lamellar BCP morphologies, in addition to mesoporous, cellular materials and macrophase separated materials. The osm-ACP-NPs are synthesized via a two-step sol−gel process in which (3-glycidyloxypropyl)trimethoxysilane (GLYMO) quenches the reaction, limits the particle size, and functionalizes the NP surface. Dynamic light scattering evidences a transition from BCP unimers to micellar aggregates with increasing amounts of sol solution, which is reflected by a corresponding switch from BCP-type morphologies to micellar/cellular morphologies of the nanocomposites. Nanostructured organic−inorganic composites with a continuous osm-ACP-NP matrix phase have indentation moduli (measured by nanoindentation) that are an order of magnitude larger than unstructured composites with similar compositions. Insights provided by this study have relevance to understanding the effects of pathway complexity in the assembly of organic−inorganic composites and may enable access to a broad range of hybrid nanostructures with potential applications in areas including dental repair and hard tissue engineering.
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metals to ceramics to transition metal oxides.8−11 However, despite multiple reports of the use of BCPs as macromolecular additives to control the solution growth of ionic solids, such as calcium phosphates and carbonates,12−18 to the best of our knowledge, there are no demonstrations of nanostructured, bulk composite materials of BCPs and ionic solids where the nanostructure is the result of self-assembly controlled by equilibrium BCP thermodynamics, i.e., exhibiting sequences of bulk morphologies as known from typical equilibrium BCP phase diagrams. Here, by unraveling the pathway complexity during solution assembly, we report the successful use of block
INTRODUCTION Mammalian hard tissues (e.g., bone and teeth) are composed of carbonated apatite (Ca10(PO4)6−y(CO3)x+(3/2)y(OH)2−2x) nanocrystals embedded within a primarily collagen matrix.1−3 The mechanical properties of these tissues are directly related to their nano-, meso-, and macroscale structures.4 Despite much effort, creating synthetic analogues to these macroscopic, nanostructured composites of ionic solids within organic matrices has proven challenging.5 A promising approach to form composites with controlled nanostructure is to use amphiphilic block copolymers (BCPs) as structure directing agents where inorganic nanoparticles (NPs) are loaded into specific domains.6 Exploiting microphase separation of BCPs,7 composites with a range of ordered nanostructures have been obtained for a variety of inorganic nanoparticles ranging from © XXXX American Chemical Society
Received: November 4, 2015 Revised: January 1, 2016
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DOI: 10.1021/acs.chemmater.5b04266 Chem. Mater. XXXX, XXX, XXX−XXX
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Chemistry of Materials copolymers to structure-direct the assembly of amorphous calcium phosphates (ACP). The resulting synthetic strategy ultimately yields an unprecedented level of control over the morphology, resulting in tailored mechanical properties of the final bulk materials: periodically ordered, organic−inorganic nanocomposites. ACP has gained increasing attention recently as a precursor to carbonated apatite formation in both biological and synthetic systems19−21 as well as a component in ACP/polymer composites for hard tissue engineering.22−24 Two main approaches have been developed to prepare ACP/polymer composites. The first involves embedding ACP fillers in methacrylate resin matrices and often leads to uncontrolled agglomeration and low strength and toughness as a result of poor interfacial interaction between ACP particles and resin matrices.24 The second approach is based on commercial BCPs employed for the synthesis of mesoporous bioactive glasses (MBGs; composition: SiO2CaOP2O5). MBGs often suffer from brittleness, small mesopore sizes of a few nanometers, extreme deficiencies in calcium phosphate content compared to silica, and compositional heterogeneity.25−27 The synthesis of ACP/polymer nanocomposites with periodically ordered nanostructures, as well as feature sizes larger than 10 nm, has not been demonstrated. New assembly strategies are thus highly desirable offering better control over the nanostructure of ACP/polymer composites. While the formation mechanisms of biominerals are still being elucidated, it is clear that small changes to the assembly conditions can lead to large changes in the final structure and properties.28−30 Such “pathway complexity” is well-known in biological systems31−33 and has recently been addressed in allorganic synthetic systems.34−38 Much research in recent years has focused on developing bioinspired synthetic routes to nanostructured organic−inorganic composites,17,18,21,39 however, a lack of understanding of pathway complexity during assembly still hinders progress in the targeted synthesis of bulk materials with desired nanostructures. The current work is motivated by the following consideration: If the pathway complexity involved in the self-assembly of ACP NPs with amphiphilic BCPs can be elucidated, then we expect to gain access to a large array of composites with wellcontrolled nanostructures, thus enabling exquisite control of materials properties including tailoring of mechanical properties of the resulting composites. Toward this goal, we studied the solution behavior of poly(isoprene)-block-poly(2(dimethylamino)ethyl methacrylate) (PI-b-PDMAEMA) as a function of solvent composition, and developed a synthetic strategy in which organosilicate-modified ACP (osm-ACP) NPs swell hydrophilic domains of PI-b-PDMAEMA resulting in composites with well-defined periodic BCP lattices with lamellar, cubic network, and hexagonal symmetries (Figure 1).
Figure 1. Schematic of pathway complexity in the solvent evaporation induced BCP-directed assembly of osm-ACP NPs. (1) Assembly into macro- vs microphase separated BCP hybrids depends on NP size; (2) formation of cellular/micellar vs periodically ordered BCP hybrid morphologies is a function of the amount of added sol solution. (3) Formation of a specific periodically ordered nanostructure via a microphase separation pathway of osm-ACP/BCP composites depends on the amount of osm-ACP NPs added to the same BCP. From left to right the five bottom graphics represent macrophase separated, lamellar, cocontinuous cubic, hexagonal, and 3-D cellular/ micellar network nanostructures. In these composites, osm-ACP NPs are selectively incorporated into the PDMAEMA domains. The osmACP NP/PDMAEMA domains are represented in bright blue and the PI domains in red.
requires chemical compatibility between the NP surface and the BCP segments to maximize enthalpy gains.40 Failure to offer such compatibility generally leads to an uneven distribution of NPs in BCP domains, segregation, and ultimately loss of nanostructure control. In order to obtain control over the size and surface chemistry of ACP NPs, we developed a sol−gel synthesis41 of ACP using an organosilicate, GLYMO, to functionalize the surface and limit growth. This sol−gel synthesis introduces additional components, in particular water and GLYMO, into the experiment, thus increasing the complexity, as compared to other NP-BCP systems. As is known for all-organic systems,35,36,42 as the number of components in the synthetic “recipe” increases, it is critical to evaluate the added pathway complexity. In particular, the solubility of BCP segments/ blocks, as well as the self-assembly behavior, can vary dramatically with solvent composition. Protic/hydrophilic solvents, such as water and ethanol employed in the sol−gel synthesis of ACP particles, may cause severe solubility problems for large molar mass hydrophobic blocks. Despite numerous studies on amphiphilic BCP/inorganic NP assemblies, not enough is known about solvent effects on assembly
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RESULTS AND DISCUSSION General Considerations. From previous work on BCPdriven self-assembly, several key experimental variables in determining the assembly pathway are already known: First, the diameter of the NPs must be smaller than the end-to-end distance of the block into which the particles are being incorporated (i.e., the hydrophilic domain), equivalent to 61/2 times the radius of gyration of the chain.40 Larger particles will disrupt the equilibrium morphologies of the BCP nanostructures and lead to size-induced phase segregation. Second, using BCPs to spatially direct NP assembly in specific BCP domains B
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Chemistry of Materials pathways.11,43−45 Therefore, before beginning hybrid assembly experiments, we studied the solvent-dependent structure of the PI-b-PDMAEMA diblock copolymer. In doing so, we elucidate a key variable, the water content of the solution, in determining the self-assembly pathway of the organic−inorganic composites. Upon the basis of our experiments, three critical decision points along the pathway for evaporation-induced structure formation were identified (Figure 1): The first is based on nanoparticle size dependent micro- versus macro-phase separation (1), the second involves solution composition dependent microphase separation in BCP micelle/osm-ACP NP composites (2), and the third is based on bulk composition dependent microphase separation in BCP/osm-ACP NP composites (3). As will be demonstrated below, when the inorganic sol content is low, PI-b-PDMAEMA exists as unimers in solution, i.e., with chains molecularly dispersed. Incorporation of osm-ACP NPs into PDMAEMA domains yields hybrid materials without interfering with microphase separation of PI-b-PDMAEMA. In this case, ordered block copolymer-type nanostructures, such as lamellae, hexagonal arrays, and even complex bicontinuous cubic morphologies can be achieved by changing either BCP composition or BCP/inorganic sol ratio. When the content of osm-ACP sol in the polymer solution is increased, solvation of the PI block is decreased due to the larger amount of added water-containing ethanol, which is a nonsolvent for the PI blocks. This change in solvent composition leads to a transition of PI-b-PDMAEMA from unimers to micelles, which in return switches to a different solvent evaporation induced structure formation process. These micelles provide a pathway to the synthesis of cellular/micellar 3-D network nanostructures with majority inorganic domains, irrespective of the polymer/osm-ACP NP ratio. Once identified, these assembly pathways allow precise tailoring of composite nanostructure and properties. Parent Polymer Synthesis and Characterization. Two PI-b-PDMAEMA polymers, referred to as IA-1 and IA-2, were synthesized by anionic polymerization (see Methods section).46,47 IA-1/2 had polydispersities of 1.10/1.08 and Mw of 26.8/34.2 kg/mol, of which 28.4/45.9 wt % was PDMAEMA. The small-angle X-ray scattering (SAXS) patterns of films cast from neat IA-1/2 are shown in Figure S1. The profile obtained from IA-2 was consistent with a lamellar lattice, with a spacing centered at 23 nm, as expected for its weight fraction.48 From the weight fraction of IA-1, a hexagonal morphology was expected. The SAXS pattern for IA-1 exhibited a primary peak corresponding to a microphase separated structure with a spacing centered at 32.3 nm as well as weak and broad higher order reflections. Synthesis and Characterization of osm-ACP Sols. Sol− gel methods have been used widely for the preparation of calcium phosphate materials.41,49 Most efforts, however, focused on optimizing the reaction conditions for the transformation of ACP to hydroxyapatite, rather than controlling NP size or surface functionality.50,51 The success of our strategy depends on using (3-glycidyloxypropyl)trimethoxysilane (GLYMO, see Figure 1) to quench the ACP NP surface and prevent further growth (see Methods section). As a modifier of ACP NPs, GLYMO can serve multiple functions: First, GLYMO-modified ACP NPs are considerably more stable against agglomeration than bare ACP particles as a result of the lower surface energy. By binding to the nanoparticle surface via phosphate-silicate and organo-phosphate interactions, GLYMO also suppresses further particle
growth, thus keeping osm-ACP NP size smaller than the critical limit relative to the size of the PDMAEMA blocks (vide supra). Furthermore, the presence of organosilicates on the surface facilitates dispersion of the osm-ACP NPs in organic solvents while simultaneously rendering particles compatible with the hydrophilic PDMAEMA blocks.8 GLYMO thus acts as a stabilizing agent of colloidal osm-ACP NP dispersions as well as a solubilizing agent in the polymer. The osm-ACP NPs, with a calculated phosphorus to silicon atomic ratio of 2.2:1, were characterized by several techniques. The absence of peaks in the wide-angle X-ray diffraction pattern (Figure S2) suggested their amorphous character, while Fourier transform infrared spectroscopy (Figure S3) showed the presence of organic groups in addition to PO and SiO bands. Solid-state single pulse 31P MAS NMR spectra displayed resonances typical of organophosphorus species and orthophosphate groups (Figure S4),52−55 while solid-state 29Si CPMAS NMR spectroscopy was dominated by T1, T2, and T3 sites, indicating an expected partially cross-linked organosilicate network (Figure S5).56,57 In order to evaluate the importance of GLYMO as a modifier in suppressing particle growth and agglomeration, we used atomic force microscopy (AFM) to measure particle size (Figure 2). In the presence of GLYMO, AFM line scans
Figure 2. AFM analysis of osm-ACP particle size. (a) AFM height image (500 × 500 nm2) of osm-ACP sol NPs used for hybrid synthesis on a silicon substrate. (b) Height profiles of line sections shown in (a) revealing NP heights below 2 nm.
revealed heights typically between 1 and 2 nm (Figure 2b) with some NPs reaching up to 5 nm (Figure S6), varying slightly between batches. In contrast, large particles with heights often exceeding 20 nm formed in the absence of GLYMO (Figure S6), suggesting that growth and agglomeration became severe in the absence of GLYMO. In order to estimate whether the osm-ACP NPs are small enough to be incorporated into PDMAEMA domains without loss of BCP structure control, we calculated the end-to-end distance of the PDMAEMA blocks for IA-1 and IA-2 to be 2.2 and 3.1 nm, respectively. Comparison of the calculated block dimensions with AFM derived particle sizes suggests that using either IA-1 or IA-2 as structure directing agent for the osmC
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different size regimes, depending on solvent composition. At the lowest volumes of cosolvent, IA-1 forms unimers with a maximum size below 10 nm. Upon increasing the amount of cosolvent, a transition regime with large (200−300 nm), poorly defined aggregates was observed. Upon further increase of the cosolvent amount, well-defined smaller aggregates (micelles) were observed with narrow size distributions and mean diameters of 40 to 50 nm. In these micelles, it is expected that the poorly solvated PI blocks accumulate in the core, while the solvated PDMAEMA blocks are located in the corona. The liquid micelle core (Tg(PI) = −60 °C) facilitates formation of well-defined aggregate sizes at high cosolvent/sol content. From these results, it can be expected that osm-ACP nanocomposite formation will proceed via different assembly pathways as the amount of osm-ACP sol added to the BCP in solution is increased. osm-ACP/PI-b-PDMAEMA (IA-1) Composites. By adding different amounts of the osm-ACP sol to solutions of IA-1 in THF, a series of osm-ACP/PI-b-PDMAEMA composite materials was synthesized, denoted as ACPIA-1 (see Methods section and Table 1). As-made ACPIA-1 composites were characterized by a combination of SAXS and TEM measurements (Figure 4a−e, Figure S9a−c). In the SAXS pattern for the sample with the lowest sol content (ACPIA-1/1), the main peak is located at a q value corresponding to 36.1 nm and higher order reflections appear at integer multiples 2, 3, 4, and 5 of this first-order maximum. This spacing sequence is indicative of a lamellar structure, which is corroborated by TEM (Figure 4b). The darker regions in the unstained TEM micrograph arise from PDMAEMA domains incorporating osm-ACP NPs. Adding osm-ACP sol NPs to IA-1 thus leads to a phase transition from the morphology of the parent polymer (vide supra) to a lamellar morphology of the composite. This progression of morphologies suggests that the inorganic sol NPs effectively swell the PDMAEMA domains to a volume fraction leading to a lamellar structure.48 With increasing sol content, the morphology of the resulting composite (ACPIA-1/2), as observed by TEM, is characterized by bright oblong objects (PI) in a dark matrix (PDMAEMA + osm-ACP NPs). This structure suggests the formation of an inverse micellar nanostructure with a continuous inorganic matrix (Figure 4c). The corresponding SAXS pattern exhibits a first-order reflection with a peak position at 39.7 nm, as well as at least one higher order reflection indicative of some degree of long-range order (Figure 4a); however, two peaks are insufficient to assign a lattice on the basis of the SAXS pattern. Further increasing the osm-ACP loading, composite
ACP NPs should be appropriate for the assembly of welldefined BCP morphologies. In order to validate the importance of GLYMO modification on the suitability of the ACP NPs for successful structure formation with the BCPs, first the large, unmodified sols were mixed with IA-1 (decision point 1 in Figure 1). Macroscopic phase segregation between BCP and ACP NPs was observed in such control samples (Figure S7) establishing the failure of BCP nanostructure control in the absence of GLYMO as organic NP modifier. Segregation is due to the lack of chemical compatibility between the ACP NPs and PI-b-PDMAEMA, as well as the large difference between the sizes of the PDMAEMA block and the ACP NPs. Solution Behavior of PI-b-PDMAEMA in Mixed Solvents. We used dynamic light scattering (DLS) to investigate the influence of cosolvent volume on the structure of 0.03 g of PI-b-PDMAEMA BCPs in 4 mL of tetrahydrofuran (THF), a typical polymer concentration used for the evaporation-induced self-assembly of the hybrids. Size distributions were determined by DLS as a function of the amount of added cosolvent (1:80 v/v water to ethanol) (Table S1), thereby simulating the composite synthesis conditions (see Methods section). A plot of the solvodynamic radii of IA-1 versus the cosolvent volume (Figures 3 and S8) revealed three
Figure 3. Influence of cosolvent on the solution behavior of PI-bPDMAEMA. Plot of solvodynamic radius of IA-1 solution (from dynamic light scattering (DLS) measurements, see Figure S8) as a function of cosolvent volume (1:80 v/v water to ethanol) added to 0.03 g of IA-1 in 4 mL of THF. Introducing water-containing ethanol to IA-1 solutions in THF induces a transition from unimers to micelles with increasing cosolvent volume. The insets show schematic illustrations of unimers, large aggregates, and micelles (from left to right). In the insets, the PI blocks are depicted in red and the PDMAEMA blocks in blue.
Table 1. Composition, Structure, and Mechanical Properties of Hybrid Samples composite
osm-ACP sol [mL]a
polymerb
structurec
wt % NPd
ACPIA-1/1 ACPIA-1/2 ACPIA-1/3 ACPIA-2/1 ACPIA-2/2 ACPIA-2/3 ACP-HP-1 ACP-HP-2
0.60 3.5 5.0 0.45 0.60 3.5 0.85 7.0
IA-1 IA-1 IA-1 IA-2 IA-2 IA-2 PDMAEMA PDMAEMA
lamellar dis. netw. dis. netw. cubic hexagonal dis. netw. N.A. N.A.
20.8 32.8 46.0 16.8 23.8 36.2 16.3 35.1
max. displacement [nm]e
hardness [MPa]e
Eindent [MPa]e,f
500 ± 19
58 ± 4
880 ± 40
180 ± 27 2700 ± 190 320 ± 11
340 ± 80 2 ± 0.2 121 ± 5
6000 ± 1000 58 ± 9 2000 ± 200
a
Aliquot used for the synthesis of osm-ACP-polymer hybrids. bFor the synthesis of the ACP-IA composites, 0.03 g polymer in 4 mL THF were used, for the synthesis of the ACP-HP composites, 0.06 g in 8 mL THF were used. cDetermined by TEM and SAXS. dDetermined by TGA. eDetermined by nanoindentation. fSee SI for the calculation of Eindent. D
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Figure 4. Characterization of as-synthesized osm-ACP NP/PI-b-PDMAEMA composites. (a) SAXS patterns of composites from IA-1. Tick marks for the ACPIA-1/1 curve show the expected peak positions for lamellar structure. (b−e) TEM images of ACPIA-1/1 (b), ACPIA-1/2 (c), ACPIA-1/ 3 (d), and plasma-etched ACPIA-1/3 sections (e). (f) SAXS patterns of composites from IA-2. Tick marks for the ACPIA-2/1 curve show the expected positions for the two primary peaks for double gyroid symmetry, while those for the ACPIA-2/2 curve show the expected peak positions for hexagonal symmetry. (g) 2D SAXS pattern of ACPIA-2/1 and (h) corresponding reflection spot analysis showing consistency with three scattering grains, two of which fit well with BCC-type cubic symmetry. (i−l) TEM images of ACPIA-2/1 (i,j), ACPIA-2/2 (k), and ACPIA-2/3 (l). All scale bars in TEM images are 200 nm. For (i) and (j), sample slices were stained with osmium tetroxide to selectively stain polyisoprene.
ACPIA-1/3 showed a similar structure with increased wall thickness of the inorganic-containing matrix (Figure 4d). Indeed, according to TGA measurements, the inorganic content of ACPIA-1 composites 1−3 increased from 21 to 33 to 46 wt % (Table 1). Plasma etching of 50 nm thick sections of ACPIA-1/3 preserved the cellular network structure leading to interconnected mesopores with pore sizes >10 nm (Figure 4e). A selected area electron diffraction (SAED) pattern of the resulting material showed no diffraction peaks indicating the preservation of an amorphous inorganic phase after plasma treatment (data not shown). The range of composite structures formed by ACPIA-1 can be understood by recalling the cosolvent effect on the solution behavior of IA-1 (Figure 3). We attribute the structural transformation from a BCP-type (lamellar) morphology to cellular/micellar structures with increasing inorganic sol content to the cosolvent induced transition of PI-b-PDMAEMA from unimers to micelles. When the polymer in solution is in the form of unimers, BCP-type morphologies are formed upon solvent evaporation (ACPIA-1/1). Because water-containing
ethanol is a nonsolvent for the PI block and a good solvent for PDMAEMA, PI-b-PDMAEMA micelles form upon increasing the volume of added sol. This change in BCP solution behavior leads to a switch in the assembly pathway (decision point 2 in Figure 1). Now, during solvent evaporation ACP nanoparticles segregate between micelles and associate with the PDMAEMA block to reduce interfacial energy, thus leading to the observed cellular structures (ACPIA-1/2,3). Upon further evaporation and subsequent annealing, the PI micelle cores can deform into polyhedra, releasing the PDMAEMA block to close the interstitial voids where interfaces have high energy. Beyond the transition from unimers to micelles, further increases in added sol volume only lead to an increase in cellular wall thickness (see ACPIA-1/3), while the cellular structure itself stays preserved. osm-ACP/PI-b-PDMAEMA (IA-2) Composites. Next, polymer IA-2 with a higher PDMAEMA fraction (46 wt %) and lamellar morphology in the bulk was used to obtain a different set of BCP-directed composites and morphologies (Figure 4f−l, Figure S9d-f). For the lowest sol concentration E
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difference between IA-1 and IA-2. While for both polymers the molar mass of the PI block is similar, the PDMAEMA molar mass of IA-2 is about twice that of IA-1. We therefore hypothesize that this difference leads to smaller sized micelles in solutions for IA-2 as a result of the higher solubilization power of the PDMAEMA block, thus leading to smaller characteristic structural length scales in the micelle driven selfassembly process. Mechanical Properties as Revealed by Nanoindentation. Selected osm-ACP/PI-b-PDMAEMA composites as well as two homopolymer (HP) PDMAEMA-NP composites (ACPHP-1,2) were investigated by nanoindentation in order to elucidate the effect of composition and structure on mechanical properties. The molar mass of PDMAEMA HP was around 54 kg/mol (PDI = 1.04, see Methods). The HP-NP composites did not display mesostructural order according to TEM (Figures S12 and S13). As is evident from Figure 5 and Table 1 (see SI for details), the largest differences in mechanical properties between periodically ordered and disordered composites were observed for samples at low osm-ACP NP content (16−17 wt %). Sample ACPIA-2/1 with 16.8 wt % NPs showed an indentation modulus of 880 ± 40 MPa, while
(ACPIA-2/1), the SAXS pattern was the most complex one of all patterns observed for ACP/PI-b-PDMAEMA composites in this study (lowermost curve of Figure 4f−h). TEM images taken at different tilt angles show long-range order with projections with 4-fold as well as 3-fold symmetries, consistent with a cubic lattice (Figure 4i,j). In contrast to all other hybrids, the 2D SAXS pattern of ACPIA-2/1 showed Bragg spots rather than a powder pattern (Figure 4g). This pattern suggested the existence of large 3dimensionally ordered grains of the material within the scattering volume of the X-ray beam path in the hybrid, consistent with low-magnification TEM images (Figure S10). By analyzing the Bragg reflection spots,58 we identified two microdomains with a body-centered cubic symmetry that can give rise to the appearance of the X-ray patterns (Figure 4h, S11); a third microdomain has two spots that do not index on the same body-centered cubic lattice, possibly due to inhomogeneous plastic compression of the domain.58,59 This analysis puts significant constraints on the possible lattices consistent with the observed structure. The intensity distribution map on our Bragg spots shows that the peaks corresponding to (211) and (220) are the most intense ones, followed by peaks corresponding to (110), (321), (310), and higher indices (Figure S11). While the presence of (110) and (310) peaks excludes the assignment to a pure double gyroid morphology expected from the trend in volume fraction, these peaks are relatively weak compared to the primary peaks at (211) and (220), suggesting that the appearance of these forbidden peaks is the result of grain compression upon drying of the hybrid from its solvent as evidenced in earlier studies.59 Finally, we note some degree of inhomogeneity in the sample as seen in the underlying halo for the 2D SAXS pattern at around q ≈ 0.35 nm−1. Indeed, SAXS patterns taken edge-on against the film showed strong reflections along the film normal direction, indicative of surface reconstructed structures mixed with the double gyroid often seen in bulk films (Figure S9d).60 Adding slightly more osm-ACP sol to polymer IA-2 induced a transition to a different BCP-type morphology. In the SAXS pattern of ACPIA-2/2 (Figure 4f, middle curve), the first order peak is located at a q-value corresponding to 24.7 nm, with at least two more broad higher order reflections. While the SAXS peaks are too few and broad for an unambiguous lattice assignment, a hexagonal lattice as indicated in the figure was corroborated by the observation of a hexagonal honeycomb-like structure in TEM (Figure 4k). The morphological transitions from parent copolymer (IA-2, lamellar) to ACPIA-2/1 (cubic cocontinuous) to ACPIA-2/2 (hexagonal) is consistent with the observed structure evolution of BCP-aluminosilicate NP mixtures.48 These results suggest that we have achieved nearequilibrium, thermodynamic structure control in the composite self-assembly process. In analogy to the discussion of hybrid structure for polymer IA-1 at low sol addition, the suggested self-assembly pathway upon solvent evaporation involves polymer unimers in solution. In contrast, cellular network nanostructures were formed from IA-2 at higher inorganic sol NP content (ACPIA-2/3), suggesting a switch in assembly pathway (Figure 4l) similar to the results observed for polymer IA-1. Interestingly, the characteristic length scale of this network nanostructure, as derived from the position of the first order reflection in the SAXS pattern, is 23.6 nm. It is substantially smaller than what was found for composites ACPIA-1/2 and ACPIA-1/3. This length scale variation is most likely related to the compositional
Figure 5. Load−displacement curves obtained from nanoindentation of two pairs of composites with similar NP content: (a) BCP-NP hybrid ACPIA-2/1 (16.8 wt % NP, orange solid lines) and HP-NP hybrid ACP-HP-1 (16.3 wt % NP, black solid lines) as well as (b) BCP-NP hybrid ACPIA-2/3 (36.2 wt % NP, orange solid lines) and HP-NP hybrid ACP-HP-2 (35.1 wt % NP, black solid lines). F
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this end, it was demonstrated that as a result of continuous NP network formation, nanostructured composites with inorganic NP loading as low as ∼10 vol % had an order of magnitude higher indentation modulus as compared to a disordered composite with the same loading. Furthermore, if ordered structures at higher particle loadings could be obtained, further gains in stiffness over that of the cellular structures, as well as control over stiffness anisotropy, should be possible by taking advantage of the geometrical stiffening afforded by incorporating percolated NPs into ordered 3-D nanostructures. We hope that insights provided from this study will open up access to bulk materials with well-controlled nanostructures of crystalline HA, with potential implications in areas including dental repair and hard tissue engineering.
sample ACP-HP-1 with almost identical loading (16.3 wt % NPs) displayed a modulus of only 58 ± 9 MPa, more than an order of magnitude lower. When moving to higher particle loadings this difference diminished: While ACPIA-2/3, containing 36.2 wt % NPs, showed an indentation modulus of 6000 ± 1000 MPa, sample ACP-HP-2, with similar loading (35.1 wt %), showed a modulus of 2000 ± 200 MPa, only a factor three smaller. These observations can be explained by different percolation thresholds of NPs in periodically nanostructured versus disordered polymer-NP composites. As calculated by Balazs and co-workers, this threshold is expected at 22 vol % for HP-NP hybrids, but at only 9 vol % for BCP-NP hybrids where NPs are confined in nanodomains of the mesophase separated structure.61 Converting the gravimetric NP fraction into a volumetric fraction (see SI for details) suggests that for the higher NP loading case both ordered and disordered composites (22 vol % for ACPIA-2/3, 24 vol % for ACP-HP-2) are above or at their respective NP percolation thresholds, thus exhibiting the same order of magnitude indentation moduli. In contrast, for the low NP loading, unstructured ACP-HP-1 (10 vol %) is below its percolation threshold while structured sample ACPIA-2/1 is at its threshold (9 vol %), consistent with a more than an order of magnitude difference in the indentation moduli of these two samples. Results demonstrate that both NP content and composite nanostructure have a strong influence on mechanical properties of BCP-NP hybrid materials. In particular in the case of BCPs with low glass transition temperature (Tg) blocks (i.e., blocks liquid at room temperature) like PI-b-PDMAEMA (Tg(PI) = −60 °C, Tg(PDMAEMA) = 0 °C), nanostructure can lead to moduli near 1 GPa via confinement-induced NP percolation in composites with NP loadings as low as 9 vol %.
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METHODS
Synthesis and Characterization of PI-b-PDMAEMA and PDMAEMA Polymers. Poly(isoprene)-block-poly(2(dimethylamino)ethyl methacrylate) (PI-b-PDMAEMA) block copolymers IA-1 and IA-2 were synthesized via sequential anionic polymerization as described elsewhere.11,46 PDMAEMA homopolymer was synthesized via anionic polymerization (see SI for further details).62 Gel permeation chromatography was used to determine the molar mass of the PI block and the polydispersity index (PDI) of the polymers. 1H NMR spectroscopy was used to determine the final chemical compositions and the overall molar mass of the polymers. The PDMAEMA homopolymer had a molar mass of 53.8 kg/mol and a PDI of 1.04 with a small lower molar mass impurity at 7 kg/mol and may contain up to 5 wt % of inorganic component, most likely sodium chloride. Synthesis and Characterization of osm-ACP Nanoparticles. Fresh solutions of calcium nitrate tetrahydrate (50 mM, Sigma-Aldrich, stored in a desiccator) and of triethyl phosphite (50 mM, SigmaAldrich) were prepared in anhydrous ethanol. Filtered (200 nm PTFE) aliquots (4 mL each) were mixed in a vial in a room temperature water bath, and double-distilled water (N2-degassed, 0.08 mL) and ammonium hydroxide (0.02 mL, 30%) were added into the solution. After stirring vigorously for 20 min at room temperature, GLYMO (0.02 mL, Sigma-Aldrich) was added. After stirring vigorously for 30 min, the filtered (200 nm PTFE) sol was blended with the polymer solution. For AFM measurements, sols were diluted ten times by ethanol. Filtered (200 nm PTFE) sol aliquots were deposited on clean silicon wafers by spin-coating. After drying in air, the sol NPs were characterized using a Veeco Nanoscope III AFM in tapping mode. DLS Measurements. DLS measurements on BCP solutions (0.03 g BCP in 4 mL of THF) were performed at 25 °C on a Brookhaven Instrument Company BIC 200SM static/dynamic light scattering system, using a HeNe Laser (λ = 632.8 nm). Synthesis of osm-ACP-Polymer Nanocomposites. In a typical synthesis, an aliquot of osm-ACP sol (see Table 1) was added to a solution of 0.03 g of PI-b-PDMAEMA polymer in 4 mL THF. After stirring for 1 h, the mixture was transferred to a Teflon Petri dish 3 cm in diameter and heated beneath a hemispherical glass dome. After 2 days of heating at 50 °C, all the solvent had evaporated. Finally, the composite product was annealed at 120 °C under vacuum for another 2 days and then slowly cooled to room temperature. The homopolymer-NP hybrids were synthesized analogously using 0.06 g of PDMAEMA homopolymer in 8 mL THF and an aliquot of osmACP sol (see Table 1). The resulting composite films were typically around 0.5 mm thick, and several millimeters across (2−6 mm on a side). Characterization of Nanocomposite Samples by TEM. Before TEM characterization, samples were microtomed with a Leica EM UC7/FC7 cryo-ultramicrotome at −60 °C. Thin sections of 50−70 nm thickness were collected on a water/DMSO eutectic solution and transferred to 300 mesh copper grids without carbon film. TEM was performed on a FEI T12 Twin at the accelerating voltage of 120 kV.
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CONCLUSIONS We have established a method for synthesizing nanocomposites from osm-ACP NPs and amphiphilic BCPs with well-defined lamellar, cocontinuous cubic, and hexagonal morphologies, consistent with BCP equilibrium thermodynamics and the selective swelling of the hydrophilic block with NPs. Two key advances enabled access to this structure control: The first was the sol−gel based synthesis of ultrasmall (
