A Facile Method To Fabricate Double Gyroid as a Polymer Template

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A Facile Method To Fabricate Double Gyroid as a Polymer Template for Nanohybrids Hsiao-Fang Wang,† Lv-Hong Yu,† Xin-Bo Wang,*,†,‡ and Rong-Ming Ho*,† †

Department of Chemical Engineering, National Tsing Hua University, Hsinchu 30013, Taiwan School of Materials Science and Engineering, Harbin Institute of Technology at Weihai, Weihai, Shandong 264209, China

‡

S Supporting Information *

ABSTRACT: Here, we suggest a facile method to acquire double gyroid (DG) phase from the self-assembly of chiral block copolymers (BCPs*), polystyrene-b-poly(L-lactide) (PS−PLLA). A wide region for the formation of DG can be found in the phase diagram of the BCPs*, suggesting that helical phase (H*) from the self-assembly of BCPs* can serve as a stepping stone for the formation of the DG due to an easy path for order−order transition from two-dimensional to three-dimensional (network) structure. Moreover, the order− order transition from metastable H* to stable DG can be expedited by blending the PS−PLLA with compatible entity. Unlike the conventional way for blending BCP with homopolymer, PS−PLLA blends are prepared by using styrene oligomer (S) to fine-tune the morphologies of the blends at which the molecular weight ratio of the S and compatible PS block (r) is less than 0.1. Owing to the use of the low-molecular-weight oligomer, the increase of BCP chain mobility in the blends significantly reduces the transformation time for the order−order transition from H* to DG. Consequently, by taking advantage of degradable character of the PLLA, nanoporous gyroid SiO2 can be fabricated using hydrolyzed PS−PLLA blends as a template for sol−gel reaction followed by removal of the PS matrix.

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INTRODUCTION Block copolymer (BCP) can self-assemble into various ordered phases.1 By introducing chirality in BCP self-assembly, new structure phases can be found in the self-assembly of BCP composed of chiral block (designated as chiral block copolymer, BCP*).2 A novel helical phase (H*) can be formed from the self-assembly of polystyrene-b-poly(L-lactide) (PS− PLLA) in addition to the conventional phases, such as spheres in a BCC lattice (S), hexagonally packed cylinders (HC), double gyroid (DG), and lamellae (L).2b The H* phase consists of hexagonally packed, interdigitated PLLA helical microdomains in a PS matrix with the space group of P622. Also, phase transition from H* to DG can be found after longtime annealing, suggesting that the H* phase is a long-lived metastable phase.2b The theoretical and simulation studies have been carried out for the formation of H* phase in BCPs*, consistent with the experimental results, revealing a significant chirality effect on BCP self-assembly.3 The theoretical results suggest that it is possible to acquire equilibrium stable H* phase from BCP* melts once the degree of chirality reaches a critical value, and the chirality is critically sensitive to the thermodynamic preference for cholesteric twist of the chiral microdomain. A drive for twist at the scale of chain segments indeed reshapes the structure and symmetry of the ordered phases on much larger length scales. Among all of the nanostructures formed by the self-assembly of BCPs, the gyroid is one of the most appealing morphologies for practical applications because of its unique geometry, composed of a © 2014 American Chemical Society

matrix and two continuous, but independent, interpenetrating networks in three-dimensional space; as a result, it is also referred to a double gyroid phase.4 After selective degeneration of the minority block, the gyroid nanostructure could be exploited to create fully interconnected nanochannels.5 Because of high porosity and large specific surface area, the nanoporous polymers, resulting from the BCP gyroid, are very promising for use in a variety of applications, such as photonic crystals,6 catalysts,7 ceramic membranes,8 membrane reactors,9 and hybrid solar cells.5c Unfortunately, the gyroid morphology which has the kind of multicontinuous structure occurs over rather limited areas of the phase portrait, in narrow slivers on either side of the diagram between cylinders and lamellae.1d Blending homopolymer to BCP is a well-known technique to obviously enrich microdomain morphology and control domain size.10 For the blends of BCP with a selective homopolymer (i.e., compatible with one of constituent blocks in the BCP), the phase behavior for the binary blends depends strongly on the compatibility, resulting from the affinity of the homopolymer with the constituent blocks and the corresponding molecular weight of the homopolymer to that of the selected block in the BCP. In the case of blends of A homopolymer and A−B BCP, the ratio (r) of the molecular weight of homopolymer (Mh) to that of the corresponding (or Received: September 22, 2014 Revised: October 17, 2014 Published: November 3, 2014 7993

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Table 1. Characterization of Synthesized PS−PLLA BCPs*, Polystyrene Homopolymer (SH and SL), and Styrene Oligomer (SO) code

Mn,PS (g/mol)

Mn,PLLA (g/mol)

f PLLAv

total Mn (g/mol)

ĐM

morphology

PS22-PLLA16 PS30-PLLA25 PS45-PLLA29 PS38-PLLA17 PS22 (SH) PS9 (SL) PS2.6 (SO)

22000 30400 45200 38000 21800 9000 2600

15700 24600 29000 17000

0.37 0.40 0.34 0.31

37700 55000 74200 55000

1.37 1.32 1.35 1.22 1.27 1.03 1.01

helix helix helix helix

compatible) block (Mb) is the key parameter (r = Mh/Mb) for the mixing state between the homopolymer and the BCP, giving the final microdomain morphology.10a,11 When r < 1, the added homopolymer swells the corresponding block chains (wet brush regime) to change the interfacial mean curvature, resulting in a morphological transition. On the other hand, when r ∼ 1, the added homopolymer is mixed into the corresponding microdomains without swelling the block chains and localized in the middle regions of the microdomains (dry brush regime).10b Thus, interfacial mean curvature and microdomain morphology do not change while domain size is greatly expanded if the homopolymer of the major component with r ∼ 1 is added. Generally, macrophase separation will appear at r value much greater than 1. Moreover, increasing the homopolymer concentration and/or the homopolymer molecular weight would cause the microdomains to swell less uniformly, resulting in segregation of the homopolymer toward the middle of the microdomains. Matsen found that large homopolymers are considered to be immiscible with the nanostructure and hence phase separate, whereas small ones tend to be miscible and can swell the nanostructure, resulting in the lattice spacing to diverge. Before macrophase separation occurs, the minority-component region of a nanostructure can only accommodate a limited amount of homopolymer; the majority-component regions can swell indefinitely with the addition of homopolymer, eventually leading a homopolymer-rich disorder phase.11 In addition, the formation of H* can be also obtained by blending PS homopolymer (HS) with different molecular weights (Mn) into the PS−PLLA BCP* for self-assembly. The H* can be formed in the blends with low-Mn HS due to the enhancement of helical steric hindrance.12 Here, we aim to acquire DG phase by directly blending homopolymer or oligomer with the BCP*. Note that the formation of DG phase requires extensive synthetic efforts due to the narrow composition window for the formation of DG. As reported previously in our laboratory, the phase transition from metastable H* to stable DG could be found after reasonably long-time thermal annealing.2 By taking advantage of the forming metastable H* from BCP*, the DG-forming region in the phase diagram can be enlarged due to the lower Gibbs free energy with helical conformation thermodynamically and the lower energy barrier for the H* → DG kinetically. Moreover, with the introduction of low-molecular-weight oligomer (r value as low as 0.05), the transformation time from H* to DG can be significantly reduced and the packing frustration of forming DG can be relieved. By combining the phase transition from H* to DG (H* → DG) and blending styrene oligomer with high mobility, the DG-forming region can be significantly enlarged and expedited in PS−PLLA BCPs* on the basis of thermodynamic and kinetic considerations. An easy process of fabrication of DG can be thus obtained, and the acquired DG

can be then used as a template after hydrolysis of PLLA blocks. As a result, the nanoporous gyroid SiO2 can be fabricated using this template for sol−gel reaction followed by removal of the PS matrix, giving appealing potentials for applications as antireflection materials with high transmission and low reflection properties as well as superior environmental resistance.5d

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EXPERIMENTAL SECTION

Synthesis of PS−PLLA Block Copolymers. The PS−PLLA diblock copolymers were prepared by sequential living polymerization processes, which have been described in our previously published results.13 The number-average molecular weight and polydispersity (ĐM) of PS block were determined by GPC. The ĐM of PS−PLLA was determined by GPC, and the numbers of L-LA repeating unit versus styrene repeating unit were determined by 1H NMR analysis. The volume fraction of PLLA, f PLLAv, was calculated according to the assumption that the densities of PS and PLLA are 1.02 and 1.248 g/ cm3, respectively. The chemical characterization results for all PS− PLLA diblock copolymers synthesized are listed in Table 1. Sample Preparation. Binary blends of neat PS−PLLA and styrene homopolymer (SM, SL) or oligomer (SO) were prepared by solution casting at 10 wt % CH2Cl2. The total polymeric weight of each mixture was around 70 mg, and the solution was sealed with aluminum foil with a pin-punched hole. The time required for the solvent to evaporate was controlled to be 2 weeks. After that, the specimen was further dried in a vacuum oven to remove residual solvent for 3 days followed by annealing temperature at 60 °C. Subsequently, the samples were thermally treated at 180 °C (temperature with microphase-separated melt) for 3 min to eliminate PLLA crystalline residues. Then, the PLLA blocks of the PS−PLLA bulk samples were removed by hydrolysis, using a 0.5 M basic solution that was prepared by dissolving 2 g of sodium hydroxide in a 40/60 (by volume) solution of methanol/water. After 3 days of hydrolysis, the hydrolyzed samples were rinsed using a mixture of DI water and methanol and then used as templates for the following sol−gel reaction. Sol−Gel Procedure. SiO2 precursor mixture was introduced into the PS template by immersing the template in TEOS/HCl(aq)(0.1 M)/ methanol mixture (weight fraction of TEOS/HCl(aq)(0.1 M)/ methanol = 10/1/25) with stirring at room temperature for 24 h and then treated under controlled humidity at 40 °C for 48 h. After drying, PS/SiO2 nanohybrids were prepared; then the PS/SiO2 nanohybrids were calcinated at 550 °C for long enough time to remove PS matrix to give nanoporous SiO2. Small-Angle X-ray Scattering (SAXS). Small-angle X-ray scattering (SAXS) experiments were conducted at the synchrotron X-ray beamline X27C at the National Synchrotron Radiation Research Center (NSRRC) in city of Hsinchu, Taiwan. The wavelength of the X-ray beam was 0.155 nm. A MAR CCD X-ray detector (MAR USA) was used to collect the two-dimensional (2D) SAXS patterns. A onedimensional (1D) linear profile was obtained by integration of the 2D pattern. The scattering angle of the SAXS patter was calibrated using silver behenate, with the first-order scattering vector q* (q* = 4λ−1 sin θ, where 2θ is the scattering angle) being 1.076 nm−1. Transmission Electron Microscopy (TEM). Bright-field TEM images were obtained using the mass−thickness contrast with a JEOL 7994

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JEM-2100 LaB6 transmission electron microscope (at an accelerating voltage of 200 kV). Field-Emission Scanning Electron Microscopy (FESEM). FESEM observations were performed on a JEOL JSM-6700F using accelerating voltages of 1.5−3 keV. Before observations, the samples were sputter-coated with 2−3 nm of platinum to avoid the charge effect (the platinum coating thickness was estimated from a calculated deposition rate and experimental deposition time).

In general, the DG morphologies occur over rather limited composition window of the phase portrait, in narrow range on either side of the diagram between cylinders and lamellae. Also, this phase is found at limited temperature window from the low to intermediate segregation value (i.e., χN ≤ 40, where χ is the interaction parameter and N is the degree of polymerization). The self-consistent-field calculations have been carried out for the formation of DG in BCP, consistent to the experimental results, revealing that DG is found to be stable relative to other complex phases in a range of composition and segregation strength.14 The calculations of Matsen and Bates suggested that DG ceases to be stable at strong segregations (χN ≥ 60) due to packing frustration effects.15 Similarly, several others have also predicted that the nonclassical phases will cease to be stable at strong segregations.16 Accordingly, it might require extensive synthetic efforts to form the DG. Interestingly, the DG in PS− PLLA BCPs* can be found not only in weak segregation region but also in the intermediate and strong segregation values after appropriate thermal treatments that drives the order−order transition from H* to DG. Consistently, there were several reports revealing that the formation of DG can be found in the strong segregation region.17 For example, the phase behavior of polystyrene-b-polyisoprene (PS−PI) with high molecular weight was examined, and the DG can be found in the strong segregation region.17a However, the reasons why these results appear to conflict with self-consistent-field theory (SCFT) remain unclear. By contrast, in the case of PS−PLLA, we speculate the forming DG in intermediate and even strong segregation might be attributed to the helical conformation with larger persistence length, resulting in the alleviation of packing frustration (see below for details). Note that the formation of DG in PS−PLLA BCPs* is apparently still present in a narrow temperature region. The formation of the DG in intermediate and strong segregation is the result from the H* through order−order transition at which the H* in the PS−PLLA is a metastable phase, appearing in relatively wide region from intermediate to strong segregation due to the slow kinetics associated with long, highly entangled chains. By taking advantage of the H* → DG, the DG-forming temperature window can be enlarged in the phase diagram of PS−PLLA due to the formation of the H*. As a result, the reasons why the DG-forming region can be enlarged by the formation of H* need to be elucidated. In the microscopic level, the packing energy of the polymer chains strongly depends on the rigidity of chain conformation, which is related to its persistence length. As the persistence length of a constituted polymer chain increases, the packing energy decreases. As calculated from theoretical prediction, the persistence length of the chiral PLLA chain (11.9 Å) is larger than that of the racemic PLA chain (10.6 Å).18 The larger persistence length for the chiral PLLA chain is attributed to the formation of a helical chain conformation for a chiral polymer chain, as evidenced by vibrational circular dichroism (VCD) (see Figure S1 in Supporting Information).13 Consequently, self-assembled phase with lower packing energy can be achieved due to the highly oriented chains at the interface that lowers the packing enthalpy (step 1 in Figure 2). Moreover, the packing frustration can also be alleviated, resulting from longer persistence length of the PLLA helical chain conformation than the racemic polymer that can sustain excessive stretching for PLLA blocks oriented toward the center of DG nodes without chain stretching to reduce the entropic penalty (step 3 in Figure 2). In the mesoscopic level, we speculate that H* can

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RESULTS AND DISCUSSION Order−Order Transition from H* to DG in PS−PLLA. To systematically study the phase behaviors of order−order transition from helical phase (H*) to double gyorid phase (DG), a series of PS−PLLA BCPs* with different compositions (i.e., various volume fractions of PLLA block) were synthesized to establish the BCP* phase diagram (Figure 1).2 Notably, the

Figure 1. Phase diagram of the PS−PLLA BCPs* with respect to overall molecular weight and composition. All the bulk samples were prepared by solution casting from CH2Cl2 solution (10 wt %) and then dried in a vacuum oven at 60 °C for 3 days. The drying samples were first thermally treated at 180 °C for 3 min to eliminate PLLA crystalline residue and then were quenched to room temperature. Note that the H* phase will transform to DG phase after long-time annealing or adding a small amount of styrene oligomer.

calculation of segregation strength (χN) is based on our previous results, where χ is the Flory−Huggins interaction parameter and can be described by χ(T) = 154.9/T − 0.211 and N is the degree of polymerization.2b Bulk samples of PS− PLLA were prepared by solution casting from CH2Cl2 solution (10 wt % of PS−PLLA) at room temperature for several days. The cast samples were placed in drybox at 40 °C to remove residual solvent. Subsequently, the samples were thermally treated at 180 °C (temperature with microphase-separated melt) for 3 min to eliminate PLLA crystalline residues. After thermal treatment, the samples were quenched to ambient condition for SAXS experiments and then sectioned by microtoming (thickness ∼50 nm) for TEM observations. Staining was accomplished by exposing the samples to the vapor of a 4% aqueous RuO4 solution for 2 h. A variety of selfassembled phases, in particular H* and DG, can be found in the PS−PLLA samples with various compositions and different molecular weights. Note that order−order transition from H* to DG (H* → DG) can be found after long-time annealing (for more than 1 week), suggesting that H* is a long-lived metastable phase. 7995

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advantage of the easy path for the H* → DG, the DG-forming temperature window in phase diagram can be enlarged due to the lower Gibbs free energy with helical conformation thermodynamically and the lower energy barrier for the H* → DG kinetically. Moreover, in comparison with the composition window for forming DG phase from achiral block copolymers with equivalent segregation strength, such as PS−PI,1d a larger composition window for the formation of DG will be expected in PS−PLLA BCP*, further implying that the H* can intrinsically serve as a stepping stone for the formation of the DG. Note that the solubility parameter of PLLA, PS, and PI are 20.3, 18.6, and 17.0 MPa0.5, respectively, giving a pair of model systems with equivalent segregation strength (i.e., PS− PLLA and PS−PI) for comparison. As illustrated in Figure 1, DG morphologies can be observed in PS−PLLA BCPs* with f PLLAv = 0.31−0.38 after long-time annealing, whereas the DG forming region can be found with f PIv = 0.36−0.39 in achiral PS−PI. We speculate that the wide composition window is attributed to the same reasons that helical conformation provides the lower Gibbs free energy state for the formation of DG morphology and helical locus gives the lower energy barrier for the phase transition from H* to DG. As a result, the forming DG region (both temperature and composition windows) can be significantly enlarged with consideration of the Gibbs free energy during the transition path. A Facile Method To Fabricate DG from Blending. To efficiently acquire DG from the order−order transition, two polystyrenes with high and low molecular weights (denoted as SH and SL) and a styrene oligomer (SO) were synthesized to blend with H*-forming PS−PLLA BCPs*. The characterization of representative samples and corresponding morphological results are summarized in Table 1. As shown in Figure 3a, RuO4-stained PS matrix appears dark whereas PLLA helical microdomains appear bright, revealing the formation of H* for PS30-PLLA25.2b The corresponding 1D SAXS profile shows well-defined reflections at the relative q values of 1:√4:√7:√13, revealing that the helices are hexagonally packed in the PS matrix. Note that for blends with homopolymer of the major component (i.e., PS) at r ∼ 1 the

Figure 2. Illustration of suggested Gibbs free energy versus transition path from H* and HC to DG.

serve as a stepping stone with helical twisting locus due to the matching of geometric dimensions (step 2 in Figure 2). For the order−order transition from hexagonally packed cylinders (HC) to DG, the transition is initiated by the formation of one 5-fold junction that pinches off breaking a cylinder and leaving a 3-fold junction (i.e., interconnect the middle cylinder with neighboring three cylinder). The two free ends of the cylinder are highly unfavorable and should quickly form two 4fold junctions that would each break off, producing a 3-fold junction.19 As the process repeats, HC evolves into DG. By contrast, in the case of the order−order transition from H* to DG, the formation of H* facilitates the easier interconnection from 2D to 3D network due to the 2D geometry with helical twisting locus which is similar to the reduction of activation energy barrier by catalysts. Also, the helical locus not only provides the lower packing enthalpy due to the enhancement on the incompatibility between PS and PLLA polymer chains but also gives the less entropic penalty due to the convenient chain packing without extension of polymer chain, providing the easy path for interconnection from 2D arrays to 3D networks to form the DG phase. As a result, by taking

Figure 3. TEM micrographs of (a) PS30-PLLA25, (b) blends with 4.3 vol % SM (Mn = 21 800 g/mol), (c) blends with 4.3 vol % SL (Mn = 9000 g/ mol), and (d) corresponding 1D SAXS profiles. 7996

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interfacial mean curvature and microdomain morphology will not change but the microdomain size will be greatly expanded. Figure 3b shows the TEM micrograph for the blends of PS30PLLA25 with 4.3 vol % (4 wt %) SH (Mh = 22 000 g/mol, r = 0.99), suggesting the preservation of H* (namely, there is no occurrence of phase transformation). Moreover, no macrophase separation can be observed under TEM, and significant shifting of reflection peaks to lower q in the SAXS result is found, suggesting that the introduction of the SH should be homogeneously localized in-between the PS blocks, resulting in the enlargement of the PS microdomains from 48.6 to 50.9 nm (swollen by about 5% of its original size), as determined from the primary peaks of the SAXS profiles. Upon reduction of the molecular weight of PS for blending to SL (Mh = 9000 g/ mol, r = 0.40), the H* remains while 4.3 vol % SL is introduced as evidenced by TEM (Figure 3c). The corresponding 1D SAXS profile (Figure 1d) shows reflections at the relative q values of 1:√4:√7:√13, further demonstrating the formation of hexagonally packing H* in the blends. However, no shifting of reflection peaks in SAXS profile which is calculated as 48.6 nm (the same as the neat PS− PLLA) can be found by introducing PS with low molecular weight (SL). We speculate that blending of SH is locally solubilized in the middle of PS microdomains so as to give rise to the swelling of the PS microdomains, resulting in an expansion of PS microdomain size. By contrast, the SL can be homogeneously distributed into the PS microdomains in the PS−PLLA, giving a higher tendency to allocate the PS near the microphase-separated interface than that in the blends with SH. However, the amount of SL (4.3 vol %) is still too low to significantly expand the microdomain size. As a result, the insignificant increase in the enthalpic penalty due to the incompatibility between the PLLA and PS polymer chains will not lead the phase transformation from H* → DG. Also, we speculate that the r value is still too high to provide enough mobility for the formation of thermodynamic equilibrium state (i.e., DG phase) under the processing conditions conducted, further demonstrating that the H* is indeed a long-lived metastable phase. To truly provide enough chain mobility for the formation of DG phase, styrene oligomer (SO) with much lower molecular weight than that of the PS block in the PS−PLLA (for instance, Mh = 2600 g/mol, r = 0.12) is thus used for blending. Different TEM projections can be obtained in the blends with the SO. Instead of twisted morphological projections, different projections of DG can be found in the blends (Figure S2, Supporting Information); as exemplified in Figure 4b, the projection result is in line with the simulated [110] projection of DG (inset). Consistently, the reflection peaks in the corresponding 1D SAXS profile occur at the relative q values of √6:√8:√16:√26:√38:√50 (Figure 4c), further demonstrating the formation of DG. Notably, in contrast to the blends with SL, stable DG phase instead of metastable H* phase can be found in the blends with SO. We speculate that the SO not only homogeneously distributes into the PS microdomain to give a higher tendency to allocate the SO near the microphaseseparated interface than that in the blends with higher Mn of homopolymer but also significantly enhances the chain mobility of the PS−PLLA chains, promoting the transformation from the metastable H* to stable DG. The phase behavior of the allocation of the SO near the interface is similar to the BCP composites with inorganic nanoparticles in which the nanoparticles with small size tend to aggregate near the microphase-

Figure 4. TEM micrographs of (a) PS30-PLLA25, (b) blends with 4.3 vol % SO (Mn = 2600 g/mol), and (c) corresponding 1D SAXS profiles. Insets show the corresponding simulated TEM projection.

separated interface on the basis of theoretical predictions.20 Subsequently, the occurrence of the local segregation of the SO with high mobility near the microphase-separated interface gives rise to significant chain mobility for the formation of stable DG with the reduced transformation time. Consequently, by taking advantage of high chain mobility of styrene oligomers, the blending of SO not only relieves the packing frustration of forming DG, which occurs from trying to fill space with Wigner−Seitz unit,11 but also expedites the transition process due to the increase of chain mobility (the contribution of styrene oligomer). Effect of r on DG Formation in PS−PLLA/SO Blends. To further examine the effect of r on the formation of DG in the PS−PLLA/SO blends, H*-forming PS−PLLA samples with different molecular weights (Figures 5a and 6a) were used to blend with the SO (Mh = 2600 g/mol). For the PS−PLLA with large molecular weight (PS45-PLLA29), the TEM micrographs

Figure 5. TEM micrographs of (a) PS45-PLLA29, (b) blends with 4.3 vol % SO (Mn = 2600 g/mol), and (c) the corresponding 1D SAXS profiles. Insets show the corresponding simulated TEM projections. 7997

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Figure 6. TEM micrographs of PS22-PLLA16 (a) blends with 4.3 vol % SO (Mn = 2600 g/mol) ( f PLLAv = 0.35) and (b) blends with 6.5 vol % SO (f PLLAv = 0.35), (c) blends with 8.7 vol % SO ( f PLLAv = 0.34), and (d) the corresponding 1D SAXS profiles. Insets show the corresponding simulated TEM projections.

(Figure 5b) of the blends with r = 0.05 and overall volume fraction of PLLA at 0.33, resulting from the blending of 4.3 vol % styrene monomer, suggest the formation of a network morphology. The projection results are in line with the simulated [110] projection of DG. The corresponding 1D SAXS profile reflection peaks can be found to occur at the relative q values of √6:√8:√16:√32:√50 (Figure 5c), which further demonstrates the formation of DG phase. To examine the upper limit for the r value on the order−order transition from H* → DG, BCP* with lower Mn (i.e., PS22-PLLA16 with r = 0.12) was used for blending. On the basis of TEM images, there is no significant morphological change by adding 4.3 vol % SO to PS22-PLLA16, suggesting that the upper limit for the r value for the blends of PS−PLLA should be at r < 0.12 while the added amount of SO (Mn = 2600 g/mol) is at 4.3 vol %. Interestingly, with further increasing the addition of SO to 6.5 vol % (6 wt %), both of the TEM image (Figure 6b) and SAXS results (Figure 6d) suggest the formation of biphasic morphologies with H* and DG. Surprisingly, as shown in Figure 6c, a well-defined DG can be observed in the blends of PS22-PLLA16 once the adding amount of SO reaches 8.7 vol % (8 wt %). Consistently, the corresponding 1D SAXS profile reflection peaks occur at the relative q value of √6:√8 (Figure 6d). On the basis of reflection peaks of SAXS profiles, we speculate that a complete transition from H* to DG can be achieved by further increasing the added amount of the SO due to the enhancement of mobility. As a result, in addition to the r ratio of the molecular weight of oligomer to that of the compatible block, the added amount of oligomer also plays an important role in the blending of BCPs and oligomer. Namely, at the same processing time, the extent of mobility can be enhanced by increasing added amount of oligomer to significantly reduce the transformation time for the order− order transition from H* to DG. As a result, the introduced amount of oligomer also plays an important factor in the formation of DG due to the efficient enhancement of chain mobility by the increase of introduced SO.

Table 2. Characterization of Blends of PS−PLLA with SO code

Mn,PS−PLLA (g/mol)

f PLLAv (BCP*)

f PLLAv (blends)

r value

PS22-PLLA16 PS30-PLLA25 PS45-PLLA29

37700 55000 74200

0.37 0.40 0.34

0.35 0.38 0.33

0.12 0.08 0.05

Nanohybrids from Templated Sol−Gel Reaction Using Nanoporous PS from Blends. It should be noted that DG is one of the most appealing morphologies for practical applications because of its unique three-dimensional geometry with high porosity and large specific surface. After the selective degeneration of the minority block in degradable BCPs, the nanoporous polymer can be used as a template for templated syntheses such as templated sol−gel reaction. As a result, nanoporous metal oxides can be fabricated after removal of the polymer template. Note that a variety of metal oxides including the oxides of Al, Si, Ti, Zn, and Zr can be synthesized through the sol−gel process.21 As reported previously in our laboratory, nanoporous SiO2 inorganic materials with low refractive index (as low as 1.1) and low absorbance can be fabricated by templated sol−gel reation.5d As demonstrated, an easy method for the preparation of gyroid-forming nanostructured polymer can be achieved by blending PS−PLLA with styrene oligomer. It will be critical to utilize the gyroid-forming PS−PLLA blends for the templated syntheses, in particular templated sol−gel reaction, with respect to practical applications. As shown in Figures 7a and 7c, DG morphology can be identified for selfassembled PS45-PLLA29/SO blends on the basis of TEM observations and corresponding 1D SAXS results. After hydrolysis in a mild aqueous base, the PLLA blocks of the PS−PLLA bulk samples were removed completely. Figure 7c displays the 1D SAXS profile of the PS−PLLA samples after hydrolysis; the diffraction peaks appear at the relative q values of √6:√8:√14:√16:√32:√50, suggesting that the PS/air sample retains as gyroid morphology after hydrolysis of the PLLA blocks in PS−PLLA. After the sol−gel reaction as 7998

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Figure 7. TEM micrographs of (a) PS45-PLLA29 blends with SO with RuO4 staining, (b) PS/SiO2 gyroid nanohybrids without staining, and (c) corresponding 1D SAXS profiles of PS−PLLA BCP after quenching from microphase-separated ordered melt, nanoporous PS template after removal of the PLLA blocks in PS−PLLA BCP by hydrolysis, PS/SiO2 gyroid nanohybrids, and nanoporous gyroid SiO2.

Figure 8. FESEM micrograph of nanoporous gyroid SiO2 from PS/SiO2 gyroid nanohybrids after calcination for the removal of PS template. Inset shows the enlarged image.

described in the Experimental Section, the PS/SiO2 gyroid hybrid nanostructure can be found. As shown in Figure 7b, the [100] projected image of the PS/ SiO2 gyroid nanohybrids without staining can be observed, which is similar to Figure 7a but with the inversion of mass− thickness contrast due to the high atomic number of silicon. Moreover, on the basis of the calculation from the primary peak for the gyroid PS template, the interdomain spacing is enlarged to approximately 56.3 nm (increasing by about 4% of its original size after hydrolysis), which compares with the original value of 54.0 nm obtained in the PS−PLLA, indicating that the

methanol used for templated sol−gel reaction may swell the PS matrix during the hydrolysis to cause the variation in size as reported previously.5d Calcination is a common thermal treatment that can be used to increase the crystallinity of metal oxides, resulting in the densification of SiO2. As a result, precise control of the crystallization behavior during calcination is necessary to give the removal of the PS template without damage to the templated morphologies. Figure 7c displays the 1D SAXS profile of the nanoporous gyroid SiO2 materials. The corresponding 1D SAXS profile shows the relative q values of √2:√6:√8:√14:√16. It is noted that there are variations in 7999

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the scattering profile of nanoporous gyroid SiO2; the relative q values of √2:√6:√8:√14:√16 are clearly different than the intrinsic PS−PLLA block copolymer. In particular, a new peak labeled with √2 can be identified. As reported in our previous studies, the appearance of the √2 reflection is attributed to the network shifting of the double gyroid. Network shifting will break the inversion symmetry of the DG phase, resulting in two incoherent SG networks. The variations in the reflections suggest a phase transformation from double gyroid to single gyroid-like nanostructure during the removal of the PS matrix.5d,6c Figure 8 shows the FESEM micrograph of the nanoporous gyroid SiO2 prepared from the PS/SiO2 gyroid nanohybrids after removal of the PS matrix by calcination. Moreover, the texture of the nanoporous gyroid SiO2 materials can be clearly evidenced with enlarged image, as shown in the inset of Figure 8. As a result, the formation of inorganic DG can be successfully achieved by using the templates from such an easy process of DG fabrication.

Radiation Research Center (NSRRC) is thanked for its assistance in the synchrotron SAXS experiments. We also thank Miss Hui-Chun Lee of Department of Chemical Engineering of NTHU for her help in sol-gel experiments.

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CONCLUSIONS Here, we suggest a facile method to fabricate DG phase for the PS−PLLA from blending. A wide region for the formation of DG can be found in the phase diagram of BCPs*, suggesting that H* from the self-assembly of BCPs* can serve as a stepping stone for the formation of the DG due to the easy path for order−order transition from two-dimensional to threedimensional (network) structure through twisting. Moreover, the order−order transition from metastable H* to stable DG can be expedited by blending the PS−PLLA with compatible entity. Unlike the conventional way of block copolymer and homopolymer blends, by controlling the added amount of blends with extremely low r value, the formation of stable DG phase can be easily achieved by blending metastable H*forming PS−PLLA (f PLLAv from 0.31 to 0.38) with SO. Because of the high miscibility and low glass transition of the oligomers with the BCPs*, an easy process to fabricate DG phase can be obtained. Consequently, by taking advantage of degradable character of the PLLA, nanoporous gyroid SiO2 can be fabricated by using hydrolyzed PS−PLLA/SO blends as a template for sol−gel reaction followed by removal of the PS matrix. This may provide a facile way to prepare large-scale, well-ordered nanoporous gyroid inorganic materials.

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ASSOCIATED CONTENT

S Supporting Information *

Figures S1 and S2. This material is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

AUTHOR INFORMATION

Corresponding Authors

*Tel +886-3-5738349; Fax +886-3-5715408; e-mail rmho@mx. nthu.edu.tw (R.-M.H.). *Tel +86-631-5687157; Fax +86-631-5687305; e-mail [email protected] (X.-B.W.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank the National Science Council of the Republic of China, Taiwan, for financially supporting this research under Grant NSC 102-2633-M-007-002 and Grant MOST103-2221-E-007-132-MY3. The National Synchrotron 8000

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