Cooperative and Sequential Phase Transitions in it-Poly(propylene

Article pubs.acs.org/Macromolecules

Cooperative and Sequential Phase Transitions in it-Poly(propylene oxide)‑b‑poly(ethylene oxide)‑b-it-poly(propylene oxide) Triblock Copolymers Weichao Shi,† Alaina J. McGrath,† Youli Li,† Nathaniel A. Lynd,*,†,⊥,# Craig J. Hawker,*,†,‡,§ Glenn H. Fredrickson,*,†,∥,§ and Edward J. Kramer†,∥,§ †

Materials Research Laboratory, ‡Department of Chemistry and Biochemistry, §Department of Materials, and ∥Department of Chemical Engineering, University of California at Santa Barbara, Santa Barbara, California 93106, United States ⊥ McKetta Department of Chemical Engineering, University of Texas at Austin, Austin, Texas 78712, United States # Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States S Supporting Information *

ABSTRACT: Sequential phase transitions were investigated in detail for a set of newly designed poly(propylene oxide)-b-poly(ethylene oxide)-b-poly(propylene oxide) (PPO−PEO−PPO) triblock copolymers, where the PPO blocks have isotactic, (R)- or (S)-, or atactic configurations. Microphase separation and thermal composition fluctuations were studied near the order−disorder transition temperature (133 ± 1 °C), and the interaction parameters χ(T) were determined. A unique feature of these materials is the cooperative and sequential nature of phase transitions, which leads to the inversion of the crystallization/melting sequence of PEO and isotactic PPO blocks. The PEO blocks crystallize quickly (ca. 50 °C) within a soft, confined lamellar template, which is followed by slow crystallization of the isotactic PPO blocks (ca. 40 °C) under hard confinement. However, in the melting process, PEO crystals melt first (ca. 60 °C) affording soft confinement which enables isotactic PPO crystals to undergo recrystallization and lamellar thickening. This cooperative behavior allows the PPO crystals to survive at higher temperatures (68 °C). The c-axes of PEO and isotactic PPO crystal lattices are both found perpendicular to the lamellar planes. We believe this study provides insight into the fundamental interrelationship between microphase separation and confined crystallization that can occur in semicrystalline materials capable of self-assembly across multiple length scales.

■

INTRODUCTION A general understanding of AB and ABA block copolymer selfassembly has been established in past decades.1,2 In general, block copolymers remain nearly homogeneous when the incompatibility between the two blocks is small. When the repulsive interactions between the dissimilar blocks increases, the composition fluctuations increase in magnitude, which eventually lead to the formation of ordered structures at the mesoscopic length scale;3 this process is known as the order− disorder transition (ODT). Depending on the composition and temperature of the block copolymer, several morphologies may be achieved, including body-centered-cubic spheres, hexagonally packed cylinders, gyroid bicontinuous structures, and lamellae. Block copolymer microphase separation is not the only possible transition in the temperature range of interest. For example, crystallization and glass transition events may occur interactively with microphase separation.4−7 Upon crystallization or glass transition (for at least one block), the corresponding crystallized or glassy block becomes immobile; thus, if the crystallization or glass transition temperature of one block is higher than the ODT temperature, then microphase © 2015 American Chemical Society

separation may be kinetically hindered or retard the formation of ordered morphologies. Conversely, when a crystallization or glass transition event takes place at a lower temperature than the ODT temperature, crystallization or glass formation may occur within the well-ordered nanodomains and can be directed or influenced by microphase separation. The thermodynamics and kinetics of crystallization in nanoconfinement are significantly different than in the bulk.8−25 There are three main parameters to consider for crystallization within confined nanodomains based on block polymer microphase separation. (1) Segregation strength:8 when crystallization occurs in discrete microphase-separated spheres/cylinders, it is known that the nucleation and growth of crystals are fully confined in separate nanodomains when the segregation strength of the constituent blocks is very high. However, crystal growth can break out of the discrete nanodomains at low or moderate segregation strengths. (2) Modulus:9−15 the crystallization of linear polymers is a liquid− Received: February 13, 2015 Revised: April 16, 2015 Published: April 27, 2015 3069

DOI: 10.1021/acs.macromol.5b00326 Macromolecules 2015, 48, 3069−3079

Article

Macromolecules

used as surfactants,45−48 modifiers for thermosets,49,50 and matrices for polymer electrolytes.51

solid transition; upon crystallization, the modulus increases several orders of magnitude. If the equilibrium crystal size is larger than the nanodomain spacing, mechanical interactions between the growing crystal and the confining environment occurs: if the environment is soft, the crystal can grow toward its equilibrium size. In contrast, if the environment has a high modulus, crystal size and/or growth orientation may be compromised under this hard confinement. Glassy and crystalline polymer domains as well as inorganic and metal templates are examples of systems that provide hard confinement. Notably, soft confinement permits the transport of nearby crystallizable components into the growth front, but hard confinement does not. (3) Topology:16−25 the confinement space can be further classified as one-, two-, or threedimensional. Lamellar structures typically represent onedimensional confinement in the thickness direction but no confinement in the other two dimensions. Cylindrical and spherical domains provide two- and three-dimensional confinement templates, respectively. It should be noted that crystallization thermodynamics and kinetics suppressed with increasing confinement dimensions.8 Crystal orientation may also vary depending on the geometrical nature and degree of confinement.21−24 The effect of nanoconfinement in crystalline−amorphous and crystalline−crystalline diblock and triblock copolymers has been extensively investigated.8−38 In most double-crystalline block copolymers, the crystallization and melting temperatures of the two blocks are quite different, so that phase transitions and confinement constraints occur sequentially with temperature variation.25−35 The typical pattern is that microphase separation at a high temperature TODT imposes soft confinement to both crystalline components, while the crystalline component with the highest crystallization (or melting) temperature (below TODT) subsequently provides hard confinement to the second crystalline component. A much more complex scenario arises when the two crystalline blocks have comparable crystallization (or melting) temperatures. Little is understood to date about the cooperativity and sequence of phase transitions and the concomitant evolution of nanoconfinement in such systems36 Recently, a few copolymers with distinguishing chiral crystalline blocks have been prepared (such as poly(L-lactide) and poly(D-lactide)). In such cases, the role of tacticity of the blocks further complicates, yet enriches, the possible crystallization, phase transformation, and morphological behaviors.39−43 To investigate these effects, we employed a series of novel poly(propylene oxide)-b-poly(ethylene oxide)-b-poly(propylene oxide) (PPO−PEO−PPO) triblock copolymers. Well-defined PPO−PEO−PPO triblock copolymers were synthesized with isotactic end-blocks, designated (R)-it-PPOb-PEO-b-(R)-it-PPO (R-it-PEP) and (S)-it-PPO-b-PEO-b-(S)it-PPO (S-it-PEP), where R and S designate the enantiomer of the propylene oxide monomer polymerized from a PEO macroinitiator.44 The atactic triblock copolymer, at-PPO-bPEO-b-at-PPO (at-PEP), was prepared in an identical manner using racemic propylene oxide.44 With these architecturally and stereochemically defined materials in hand, the microphase separation and confined binary crystallization/melting behavior of PPO and PEO, as well as the relationship of crystallite orientation within the microphase-separated domains, were investigated. We believe this study will also add significantly to the field of amphiphilic at-PPO-b-PEO and related block copolymers, commercially known as Pluronics which are widely

■

EXPERIMENTAL SECTION

Materials. The syntheses of at-PEP, R-it-PEP, and S-it-PEP were reported previously.44 The architectural information for all of the samples used in this study is shown in Table 1.

Table 1. Characterization Data for the Samples Used in This Study sample PEO R-it-PPO at-PEP R-it-PEP S-it-PEP

Mn (PEO)

PDI (PEO)a

20000

1.30

20000 20000 20000

1.30 1.30 1.30

Mn (PPO)b

PDI (PEP)a

PPO vol % (NMR)c

3300 8600 8800 8600

1.20 1.28 1.27 1.26

0 100 32.5 33.6 32.5

a

Determined by GPC in chloroform calibrated to PEO standards (g/ mol), PDI = Mw/Mn. bTotal PO polymerized determined by 1H NMR spectroscopy. cVolume fraction of the PPO block calculated through 1 H NMR spectroscopy where the mass densities of amorphous PEO and PPO are 1.124 and 0.998 g/cm3, respectively. Differential Scanning Calorimetry (DSC). The heat flows for the crystallization and melting processes were recorded on a TA Instruments Q-2000 DSC using the Tzero technique. The rate of heating and cooling was 10 °C/min over a temperature range of −20 to 100 °C. Transmission Electron Microscopy (TEM). Ultrathin (∼100 nm) sections were cut at room temperature for fully crystallized R-itPEP. The thin slices were collected on copper grids (CF-300 from Electron Microscopy Sciences) from a hexane surface (nonsolvent for PPO and PEO). Staining was carried out in ruthenium tetroxide vapor for 15 min to enhance the contrast for TEM images. Large-Amplitude Oscillatory Shear (LAOS). LAOS was carried out on a rheometer (TA Instruments, Inc.) with two parallel plates. The system temperature was controlled by a heating oven with a nitrogen atmosphere. The sample was loaded at 100 °C and compressed to 1 mm thick film. The shear rate was 1 rad/s, and shear amplitude was 200% strain. LAOS started from 100 °C with a slow cooling rate of 0.1 °C/min and stopped at room temperature. Small- and Wide-Angle X-ray Scattering (SAXS and WAXS). Both SAXS and WAXS experiments were carried out on a house-made X-ray generator with a rotating Cu target. The wavelength of the beam was 0.154 nm, and a 2D position sensitive detector was placed perpendicular to the beam direction. For the SAXS experiments, the sample-to-detector distance was 1362 mm. The instrument geometry was calibrated using a silver behenate standard. For the WAXS experiment, the sample-to-detector distance was 245.5 mm, which afforded a scattering wavevector (q) ranging from 6 to 20 nm−1. The calibration was based on Si powder with a featured peak position at 28.443°. Thermal treatment used in the scattering studies was carried out by a stepwise temperature ramp procedure. Isothermal conditions were maintained for 20 min between temperature increments. For the crystallization and melting temperature range of primary interest (25− 70 °C), measurements were obtained in 1 deg increments. At 70−100 °C, the temperature increment was 10 °C for each measurement; above 100 °C, the tests were carried out by stepwise cooling with a temperature step of 2 °C. The SAXS and WAXS patterns were obtained with 20 min exposures at each temperature. All scattering data were corrected for air scattering and absorption by the sample. The SAXS results were further corrected by subtracting the contribution from thermal density fluctuations. Following Ruland’s equation, the high-q range intensity was fitted by I(q)q4 = Kp exp(−σ2q2) + IBq4, where the wave vector is q = 4π sin(θ/2)/λ (θ is the scattering angle). Kp is the Porod constant, σ 3070


Article

Macromolecules is related to the interface smearing width, and IB is the scattering from thermal density fluctuations. The background intensity was subtracted over the q-range to give the corrected intensity, IC(q).52,53 The normalized 1D correlation function is defined as

g (r ) =

∫0

∞

system is homogeneous at high temperatures, the block connectivity leads to depletion of interchain segment correlation (correlation hole effect).55,56 The resultant weak composition inhomogeneity can be detected by SAXS, which showed a broad, low-intensity peak at q* (Figure 1c). The composition fluctuations grew stronger as the temperature decreased, which produced a narrower, higher intensity peak. When the temperature decreases below the ODT temperature, additional higher order peaks appeared due to the formation of long-range ordered structures (Figure 1b). For all three triblock copolymers used in this study, an additional weak shoulder/ peak emerged at (4/3)1/2 × q*. Thus, we suspect that the gyroid morphology forms immediately below the ODT temperature and remains stable until 110 °C. When the temperature is decreased further (Figure 1a), the weak shoulder/peak disappears and a small peak appears at 2 × q*, consistent with a lamellar structure. From the SAXS curves, we conclude that the original double gyroid transforms into a lamellar morphology at lower temperatures. A practical method to identify the various regimes of phase behavior from disorder to the ODT and beyond is to plot the reciprocal of the maximum scattering intensity (Im−1) against the reciprocal of the absolute temperature (T−1). Following the random phase approximation (RPA),2,57 the structure factor for the linear triblock copolymer is expressed as

IC(q)q2 cos(qr ) dq/Q

2 where Q = ∫ ∞ 0 IC(q)q dq is the scattering invariant. The intensities at −1 low (near 0.10 nm ) and high q (1.2−1.4 nm−1) were fitted by Guinier’s law and Porod’s law, respectively.54 The thickness of the lamellar crystallites (or the amorphous layer) was obtained by the intersection between the slop of the triangle at g(r) = 0 and the horizontal line at the first minimum. The thicknesses of the crystalline and amorphous regions are identified by crystallinity, which is calculated from DSC.

■

RESULTS AND DISCUSSION Microphase Separation and Thermal Fluctuations near the ODT. We began our study by examining phase transitions for at-PEP by small-angle X-ray scattering (SAXS) across a range of temperatures (100−210 °C). The SAXS data are shown in Figure 1. For block copolymers, although the

S ̅ −1 =

S(q) − 2χ W (q)

(1)

χ is the interaction parameter, which is assumed to be temperature-dependent according to χ = A + B/T. From eq 1, the reciprocal of the correlation hole scattered intensity is predicted to correlate linearly with the reciprocal of the absolute temperature. Figure 2a demonstrates that this relationship describes the scattered intensity above 180 °C, which is known as the mean-field linear regime where disordered correlation hole scattering is dominant. Below 180 °C, the inverse scattered intensity deviates from the mean-field prediction. This marks the onset of composition fluctuations which strengthen in magnitude as the temperature is decreased toward the ODT. This regime is known as the fluctuation regime or non-mean-field regime.3,58 The ODT corresponds to the discontinuous decrease in Im−1. For all three block copolymers studied herein (at-PEP, R-it-PEP, and S-it-PEP), the ODTs occur at 133 ± 1 °C. It should be emphasized that within experimental error the triblock copolymers have nearly identical compositions and overall molecular weights (Table 1). These results demonstrate that the tacticity of the PPO end blocks has negligible influence on the microphase separation behavior in the absence of crystallinity. Furthermore, the interaction parameter can be calculated for each block copolymer by fitting the structure factor mentioned above and correcting for the finite-width molecular weight distribution.57 The results are referenced to a common repeat unit volume of 0.1 nm3 and shown in Figure 2b. As expected, the interaction parameter increases with decreasing temperature for all samples. While we acknowledge the uncertainty associated with these measurements (mainly derived from background correction and low signal-to-noise ratio at high temperatures), we conclude that R-it-PEP, S-it-PEP, and at-PEP have very similar interaction parameters, which indicates that the tacticity of the PPO does not significantly impact the interaction parameter between PEO and PPO. Since the RPA is

Figure 1. SAXS curves for at-PEP near the ODT: (a) shows the lamellar and gyroid structures in the ordered state; (b) indicates the gyroid to disordered transition; (c) reveals the correlation hole scattering from the disordered structures at high temperatures. 3071


Article

Macromolecules

In this study, the PEO and it-PPO blocks were crystallized by slowly lowering the temperature so that the crystallites grew inside the microphase-separated lamellar domains. This transition could be observed by SAXS and wide-angle X-ray scattering (WAXS) experiments. For at-PEP, SAXS indicated that the crystallization of PEO occurred at 48 °C (Figure 3a). At this temperature, a second, broader peak emerged at lower q, indicating larger domain spacing after crystallization. Additionally, the scattered intensity after crystallization was significantly stronger than that due to microphase separation alone. The scattering curves did not change significantly below 40 °C (Figure 3b). By WAXS, the intensity of the peak at 13.6 nm−1 ((120) plane) was monitored so that information about the crystallization of PEO in at-PEP could be obtained (Figure 3c); the intensity of the crystallization peak grew quickly from 50 to 47 °C and continued to increase slowly as the temperature was further decreased (Figure 3d). This behavior indicates that the majority of the PEO chains crystallize rapidly immediately below 50 °C. In comparison, PEO homopolymer crystallized rapidly between 52 and 54 °C in the bulk. This indicates that when crystallization occurs within a microphase separated 1-D lamellar template with soft confinement, there is little influence on the crystallization temperature. Similar results have been reported in previous publications.8,16 The scattered intensity was observed to increase significantly after PEO crystallization in at-PEP due to an increase in electron density contrast. Amorphous PEO has an electron density of 369 e/nm3, while crystalline PEO has an electron density of 405 e/nm3. Additionally, the electron densities for amorphous and crystallized PPO are 332 and 364 e/nm3. These values correlate with the mass density of PEO (1.124 g/cm3: amorphous; 1.235 g/cm3: crystalline) and PPO (0.998 g/cm3: amorphous; 1.097 g/cm3: crystalline).61 The resultant calculated difference in electron densities between crystallized PEO and amorphous PPO (405 − 332 = 73 e/nm3) is twice the difference between amorphous PEO and amorphous PPO (369 − 332 = 37 e/nm3). This explains the appearance of higher intensity small-angle scattering due to the size, shape, and arrangement of PEO crystallites interspersed with the remaining soft, microphase-separated PPO domains. The scattering peaks after crystallization were much broader than the sharp scattering peaks due to microphase separation alone, which indicates that additional defects and a wider distribution in lamellar thickness accompanies PEO crystallization. The microphase-separated lamellae are therefore readily deformed when PEO crystallization occurs next to the soft neighboring at-PPO domains, which can result in defect or branch formation during growth or variation in the PEO crystal thickness.8 The change in lamellar thickness during PEO crystallization in at-PEP can be quantitatively calculated from the 1-D correlation function determined from the SAXS data based on a two-phase model of amorphous layers of PEO and PPO.54,62−64 Before crystallization (see Figure 4), the long period (L) was 19 nm, and the corresponding PEO and PPO layer thicknesses were 13 and 6 nm, respectively. From these calculations, we estimated the volume fraction of PPO in at-PEP to be 0.324, which is consistent with the composition of PPO determined by 1H NMR (ca. 0.30). After crystallization of PEO, the long period was calculated to be 21 nm, which consisted of a 13 nm thick high electron density layer and an 8 nm thick low electron density layer. Since the PPO blocks were amorphous during PEO crystallization, we would expect the PPO layer thickness

Figure 2. (a) Plot of Im−1 with respect to T−1 for the at-PEP block copolymer. Im corresponds to the maximum intensity in the SAXS measurement at each temperature. The four regimes (mean field, nonmean field, gyroid, and lamellae) are classified accordingly. The ODT is revealed by the discontinuous drop in Im−1. The extracted χ values for at-PEP, R-it-PEP, and S-it-PEP are shown in (b). The scattering data in the mean-field region were included in the fit to χ. The stereochemistry of the PPO blocks resulted in negligible variation in χ.

only valid in the linear regime at higher temperatures, the interaction parameter fitting is carried out above 180 °C; the extracted χ values deviate from mean-field behavior at lower temperatures due to the onset of composition fluctuations. The interaction parameter between PEO and PPO based on the RPA fit to the disordered state scattering is 34.95 χ = −0.0236 + T According to the calculated interaction parameter, χN (where N is the total number of segments) is estimated to be 27 at the ODT, which qualitatively agrees with the value predicted by self-consistent field theory (χN = 24) for triblock copolymers.59 The slight difference may be due to fluctuation effects (as shown in Figure 2a) or the molecular weight distribution associated with the polymers. Moreover, the value of χN at room temperature is calculated to be ∼40, which represents microphase separation in the weak segregation regime. Sequential Crystallization and Melting. During crystallization, chains fold back and forth to form polymer crystal lamellae while entanglements and defects are repelled outward, leaving an interface with the remaining amorphous material. Within the crystallites, the polymer chains arrange into a helix with highly ordered lattice periodicity. PEO and PPO crystals occupy monoclinic and orthorhombic lattices, respectively.22,60 3072


Article

Macromolecules

Figure 3. SAXS (a, b) and WAXS (c) curves during the PEO crystallization in at-PEP block copolymer. (d) Normalized temperature-dependent evolution of the (120) plane at 13.6 nm−1 indicated by an arrow in (c).

To summarize thus far, in at-PEP, thin PEO crystals form within a narrow temperature window, which leads to a slightly depressed melting temperature during heating. With an understanding of the crystallization behavior of atPEP, the dual crystallization behavior of R-it-PEP was then examined. From the SAXS curves, the reflections consistent with a microphase-separated lamellar structure from microphase separation transformed into higher intensity, broad peaks below 45 °C (Figure 5a), indicating crystallization of PEO (broad peaks were located at 0.29, 0.6, and 0.9 nm−1 for the first-, second-, and third-order peaks, respectively). Below 40 °C (Figure 5b), the primary peak position did not change, but a weak shoulder emerged at 0.4 nm−1. Additionally, the previous second- and third-order reflections disappeared and were replaced by a weak feature at 0.8 nm−1. Interestingly, this phenomenon indicated that a second lamellar structure was formed below 40 °C. From the WAXS data (Figure 5c,d), we found that the intensity of the (120) reflection at 13.6 nm−1, which is the peak corresponding to the PEO crystal, grew between 50 and 47 °C; this behavior is similar to that observed in at-PEP. In contrast, the (200) peak at 12.0 nm−1, derived from PPO crystals, grew significantly between 45 and 35 °C and remained constant below 35 °C. Combining the SAXS and WAXS data, we conclude that R-it-PEP triblock copolymers crystallize in two steps: the crystallization of PEO blocks occurs rapidly below 50 °C, and subsequently, R-it-PPO blocks crystallize slowly below 45 °C. During PEO block crystallization, the neighboring R-it-PPO blocks were a viscous liquid. Therefore, PEO crystallized within a soft lamellar template, similar to the at-PEP case. However, when the R-it-PPO blocks crystallized, the neighboring PEO blocks had already crystallized and provided a hard confinement template, causing the crystallization temperature of the R-

to remain 6 nm. However, we attribute the increase in the lower electron density, amorphous phase to the accumulation of the remaining amorphous PEO into the PPO domain due to defects, entanglements, and chemical junctions to the PPO blocks. The existence of the interface region will be confirmed later by TEM. The melting behavior of at-PEP was studied by slowly heating in 1 °C increments (Supporting Information, Part A). From the SAXS curves, the broad peaks with high intensity transformed back into narrow peaks with lower intensity and higher wave vectors. The WAXS data demonstrated that the crystal structure melted rapidly between 55 and 60 °C. The change in the 1-D correlation functions was consistent with the recovery of the microphase-separated lamellar structure upon melting of the PEO blocks. In comparison, PEO homopolymers show melting behavior between 60 and 65 °C (Supporting Information, Part B), which is slightly higher temperature range than at-PEP. The heat flow for the crystallization of PEO homopolymer, determined by DSC, was 175.1 J/g; this measurement allowed us to estimate the crystallinity (81%), based on the standard heat of fusion (213.7 J/g) for PEO. From this estimate, the PEO homopolymer crystal thickness (>25 nm, 31.4 nm × 81%) was determined to be more than twice the PEO crystal thickness in the at-PEP triblock copolyer (13 nm). From the heat flow due to PEO melting (104.6 J/g), normalized by the weight fraction (69 wt %), the crystallinity of PEO was found to be 71% in the at-PEP triblock copolymer. The PEO crystal thickness is thinner in atPEP than in the PEO homopolymer due to two primary factors: thermodynamically, increased interface energy resulting in lower equilibrium melting temperatures and dynamically, the central PEO block is tethered by two end blocks, which inhibits entanglements from relaxing in the crystallization/melting process. 3073


Article

Macromolecules

blocks, which implies that the crystallization of R-it-PPO was fully confined within the nanodomains and no mechanical deformation occurred (Figure 5b). Because of the double semicrystalline nature of R-it-PEP, there are four local environments in the system, including crystalline and amorphous PEO and two states of R-it-PPO (Figure 5e). Because of the excellent electron density contrast between crystallized PEO and the other three local states, a coarse lamellar structure of R-it-PEP can be observed by TEM without any staining, showing alternating bright and dark layers (Figure 6a). After careful staining with RuO4 vapor, more intricate structures were revealed (Figure 6b). RuO4 exclusively diffused into the amorphous domains and enhanced contrast with the crystallized domains. Using this technique, we determined that the dark layers were the amorphous PEO and R-it-PPO interfacial regions. Additionally, there were two bright layers: a thick bright layer identified as crystallized PEO and thin layers that correspond to crystallized R-it-PPO. We conclude that the dual crystallization of PEO and R-it-PPO blocks leads to a hierarchically ordered lamellar structure. Having determined from TEM that the final ordered structure of these doubly crystalline materials was doublelamellar, the SAXS curves were investigated in detail. It is difficult to directly analyze the 1-D correlation functions of the SAXS curves based on the standard two-phase model. However, we can obtain more information by carefully considering the two-step crystallization behavior and compare it to the crystallization behavior of at-PEP and R-it-PPO homopolymer. The first crystallization event of R-it-PEP is identical to that of at-PEP with the PEO blocks crystallizing first within the temperature range of 45−40 °C. Amorphous PPO forms 6 nm thick layers and PEO forms a 15 nm thick layer that consists of 1 nm amorphous, 13 nm crystalline, and another 1 nm of amorphous sandwich layers. The crystallization of the R-it-PPO end-blocks should lead to similar amorphous/crystalline/amorphous sandwich structures within the lamellar nanodomains. The precise dimensions cannot be extracted directly from the scattering data, although there is evidence of additional correlation peaks in the SAXS curves (Figure 5b) and the 1-D correlation function displays a distorted shape (Supporting Information, Part D). The domain sizes can be estimated based on the crystallinity of the R-it-PPO homopolymer which was estimated to be 80%, based on the crystalline and amorphous thicknesses extracted from the 1-D correlation function. From the DSC thermograms, we determined the heat flow associated with crystallization of Rit-PPO homopolymer to be 52.2 and 37.1 J/g for R-it-PPO in R-it-PEP, which is a 71% reduction in crystallinity.44 Therefore, the absolute crystallinity of the R-it-PPO blocks, crystallizing within the 6 nm nanodomains, is estimated to be approximately 57%. Using this information, we estimated the crystal thicknesses of the R-it-PPO domain to be 3.4 nm and each of the two interface layers to be 1.3 nm (Figure 5e). On the basis of their smaller size, we expected the R-it-PPO crystals to melt before PEO during the heating process; however, the experimental results show the opposite behavior. To explain this, it is important to consider that R-it-PPO showed melting and recrystallization behavior during the heating process, while PEO did not (Supporting Information, Figure S5). This unexpected melting behavior was captured by a series of WAXS and SAXS experiments. In Figure 7c,d the WAXS data show that the intensity of the (120) plane, corresponding to the

Figure 4. 1-D density correlation functions at 50 °C (a) and 40 °C (b), which are calculated based on the SAXS data. L is the long period length from the two layers with the largest contrast (L = L1 + L2). (c) Domain size variation at different temperatures. (d) Schematic illustration of the structure before and after PEO crystallization.

it-PPO block to be depressed. For R-it-PPO homopolymer (Mn = 3.3 kg/mol, which is slightly smaller than the R-it-PPO blocks in the triblock copolymer), the (200) reflection drastically increased in intensity during crystallization (between 45 and 41 °C) (Supporting Information, Part C). The shifting of the primary peak position in the SAXS curves is a direct consequence of soft or hard confinement. The primary peak position shifted toward lower values after PEO crystallization in the triblock copolymers (Figures 3a and 5a), indicating that the PEO crystals slightly expanded the primary spacing in the system. In contrast, the primary peak position remained nearly constant during crystallization of the R-it-PPO 3074


Article

Macromolecules

Figure 5. SAXS curves for R-it-PEP during slow cooling show crystallization of the PEO block (a) and the subsequent crystallization of the R-it-PPO block (b). A portion of the WAXS curves are shown in (c). The normalized peak evolution of PEO (from the (120) plane at 13.6 nm−1) and R-itPPO (from the (200) plane at 12.0 nm−1) is plotted in (d). The sequential crystallization process is schematically shown in (e), where the color change with crystallization indicates the change in electron density contrast.

PEO crystals, remains nearly constant below 50 °C and significantly drops near 60 °C. This behavior is very similar to the melting of PEO in at-PEP. For crystallized R-it-PPO blocks, the intensity of the (200) plane decreases slightly below 60 °C, followed by a slight increase and decay to zero intensity at 68 °C, which is higher than the melting temperature of PEO. Because of the sequential melting behavior, the electron density contrast changes significantly, which is evident by an intensity variation in the SAXS curves (Figure 7a,b). At room temperature, there is high contrast between PEO and R-itPPO crystals (electron density difference is 405 − 364 = 41 e/ nm3). After melting PEO, the majority of contrast is between amorphous PEO and crystalline R-it-PPO, which is very low (369 − 364 = 5 e/nm3); the maximum intensity drops from 2000 to 100 counts. At higher temperatures, the R-it-PPO crystals melt, which leads to intermediate contrast between amorphous PEO and amorphous R-it-PPO (369 − 332 = 37 e/ nm3). Accordingly, when the R-it-PPO crystals fully melt, the intensity increased significantly from 100 to 1000 counts. The parallel experiment with R-it-PPO homopolymer shows clear evidence of the recrystallization phenomenon during the differential scanning calorimetry (DSC) heating process with an exothermic event (recrystallization) occurring before the endothermic events (complete crystal melting) (Supporting Information, Part E).44 In the SAXS curves, the primary peak position significantly shifts toward smaller values as the

temperature is increased. From the 1-D correlation function we can estimate that the R-it-PPO crystal thickness increases from 8 nm at room temperature to 11 nm at 60 °C. On the basis of the above experimental results, we conclude that the R-it-PPO blocks show strong recrystallization behavior, which leads to a higher melting temperature than the PEO blocks in R-it-PEP. Notably, the early melting of PEO crystals changes the environment for the R-it-PPO crystals from hard confinement to soft confinement. We believe this modulus inversion offers a mobile neighboring layer, allowing for recrystallization and crystal thickening to proceed. The melting procedure is schematically shown in Figure 8. Precise determination of the layer thicknesses is not straightforward due to complex scattering from four phases. At room temperature, the R-it-PPO crystal thickness is 3.4 nm, which is sandwiched between two amorphous interfacial layers of 1.3 nm. The PEO crystal thickness is 13 nm and sandwiched between two amorphous interface layers of 1 nm. When the PEO crystals have melted, the PEO layer thickness changes from 15 to 13 nm. The long period can also be extracted from the SAXS curves, and by subtracting the amorphous PEO layer from the long period, we can obtain the total thickness of the crystalline R-it-PPO and two surrounding amorphous layers. Within the temperature range of recrystallization, we assume that the thickness of the amorphous R-it-PPO layers in the block copolymer remains 1.3 nm. We base this assumption on 3075


Article

Macromolecules

the fact that the amorphous interface layers maintain a constant thickness during recrystallization of the R-it-PPO homopolymer. Accordingly, the R-it-PPO crystal thickness can be evaluated during the melting process: the crystal thickness increases from 3.4 to 9 nm near 60 °C. The R-it-PPO chains of 4.3 kg/mol are estimated to produce 25.6 nm crystals with extended chain conformation; once-folded chains result in 12.8 nm crystals, and twice-folded chains result in 8.5 nm crystals. Therefore, on average, the R-it-PPO blocks can transform from crystal lamellae (with multiply folded chains) to twice folded chains. The sequential crystallization behavior of S-it-PEP is similar to R-it-PEP. Detailed data can be found in the Supporting Information (Parts F and G). Similar crystallization behavior indicates that the opposite chirality of the PPO blocks does not have any significant effect on the crystallization of either PEO or PPO. Crystal Orientation. In the experiments and discussions above, we focused our attention on crystal dimensions. The crystal orientation with respect to the microphase-separated lamellar orientation is the focus of the remainder of this study. Typically polymer chains fold back and forth within nanodomains to give crystal lamellae with a specific lattice arrangement. However, the crystal lattice orientation is not always in the same direction as the crystal thickness, which is sometimes tilted.13,21−24 After thermal treatment in the bulk, the block copolymers do not show preferred macroscopic orientations, which are made apparent by the presence of isotropic halos in the 2-D SAXS patterns. With the crystals inside the lamellar nanodomains, the 2-D WAXS patterns should also show isotropic rings in this case. The lack of macroscopic lamellar orientation limits the study of the crystal orientation after thermal treatment; therefore, we used large-amplitude oscillatory shear (LAOS)

Figure 6. TEM images of R-it-PEP without staining (a) and with RuO4 staining (b). The scale bars are 100 nm (a) and 20 nm (b). A coarse lamellar structure is revealed in (a) due to the low electron density contrast. The intricate double lamellar structure is revealed after staining (b).

Figure 7. SAXS curves (a, b) for R-it-PEP during the slow heating process, showing the melting of the PEO and R-it-PPO blocks. A part of the WAXS curves are shown in (c). The normalized peak evolution of PEO (from the (120) plane at 13.6 nm−1) and R-it-PPO (from the (200) plane at 12.0 nm−1) is plotted in (d). 3076


Article

Macromolecules

Figure 9. 2-D SAXS and WAXS patterns for at-PEP (left, a and c) and R-it-PEP (right, b and d) triblock copolymers. The circular regions in (c) and (d) are artificial masks screening the defects in the detector. (e) Schematic illustration showing the geometry of the LAOS and scattering experiments. The “zero” of the azimuthal angle is on the right side of the horizontal direction and angles are measured positively in the counterclockwise direction. The it-PPO crystal has an orthorhombic lattice, with a = 0.104, b = 0.464, and c = 0.692 nm.60 The PEO crystal is monoclinic with a = 0.805, b = 1.304, c = 1.948 nm, and β = 125.40°.22 The c-axes of both it-PPO and PEO are perpendicular to the microphase-separated lamellar planes.

that the lamellar layers are aligned parallel to the shear plane. In the corresponding 2-D WAXS pattern, the (120) reflection at 13.6 nm−1 is more intense near 0 and 180°. This indicates that the c-axis of the crystal lattice in reciprocal space lies in the vertical direction. The reflection at 16.5 nm−1 may be the result of the overlapped peaks from the (004), (1̅24), (112), (2̅04), (032), (014), (1̅32), and (2̅12) reflections.22 The intensity is strongest at −25, 25, 155, and 205°. Furthermore, we calculated the tilt angle of each reflection plane, referenced to the c-axis in reciprocal space. The (004), (1̅24), (2̅04), and (014) planes correspond to tilt angles of 35.40, 36.78, 37.71, and 38.76°, and the (112), (032), (1̅32), and (2̅12) planes correspond to tilt angles of 66.67, 66.95, 67.12, and 67.19°, respectively. Considering the c-axis in the 90 and 270° direction, we identified four intense spots near −25, 25, 155, and 205° as overlapped peaks from the (112), (032), (1̅32), and (2̅12) planes. Crystalline PEO has a monoclinic lattice, where the cand a-axes have a contact angle of 125.4°. Specifically, the PEO chain stems are perpendicular in real space; the a-axis has tilt angles of 54.6° with respect to the normal of the microphaseseparated lamellar plane, and the b-axis is along the X-ray beam direction. Notably, from the 2-D WAXS images, the (120) plane of the PEO crystals presented a broad band near 0 and 180° and the bandwidth expands ±50°, which indicates that

Figure 8. Schematic illustration of the sequential melting behavior of R-it-PEP. The PEO crystal melts first (a, b), which is accompanied by recrystallization and thickening of the R-it-PPO block (c, d). The system returns to a microphase-separated lamellar structure after the R-it-PPO crystals melt.

to induce long-range orientation in the microphase separated lamellae. The lamellar orientation was then used as a reference for the crystal orientation. The 2-D SAXS pattern shown in Figure 9 for LAOS-oriented at-PEP shows two bright arches at 90 and 270°, which indicates 3077


Macromolecules most PEO stems are perpendicular while others are tilted with respect to the microphase-separated lamellar plane. Only the perpendicular orientation is schematically represented in Figure 9e. A similar analysis can be made for R-it-PEP. The lamellar orientation can be clearly determined from the two strong spots at 90 and 270° in the 2-D SAXS pattern. In the 2-D WAXS pattern, we obtained similar information from the PEO crystals in R-it-PEP as we do from the PEO crystallites in at-PEP. In addition, R-it-PPO crystallites showed two strong reflections from the (200) and (110) planes at 12.0 and 14.8 nm−1, respectively. The maximum intensity was located at 0 and 180° for both of the reflection planes, and therefore, we infer that the c-axis is in the vertical direction in reciprocal space. Based on the orthorhombic crystal lattice for PPO crystals, each stem is vertical to the lamellar plane in the real space unit cell; a- and baxes are in the lamellar plane, and b-axis is along the X-ray direction.

ACKNOWLEDGMENTS

■

REFERENCES

(1) Bates, F. S.; Fredrickson, G. H. Annu. Rev. Phys. Chem. 1990, 4, 525−557. (2) Leibler, L. Macromolecules 1980, 13, 1602−1617. (3) Bates, F. S.; Rosedale, J. H.; Fredrickson, G. H. J. Chem. Phys. 1990, 92, 6255−6270. (4) Hamley, I. W.; Castelletto, V. Prog. Polym. Sci. 2004, 29, 909− 948. (5) He, W.-N.; Xu, J.-T. Prog. Polym. Sci. 2012, 37, 1350−1400. (6) Kim, J. K.; Han, C. D. Adv. Polym. Sci. 2010, 231, 77−145. (7) Zhu, L.; Cheng, S. Z. D.; Calhoun, B. H.; Ge, Q.; Quirk, R. P.; Thomas, E. L.; Hsiao, B. S.; Yeh, F.; Lotz, B. Polymer 2001, 42, 5829− 5839. (8) Loo, Y.-L.; Register, R. A.; Ryan, A. J. Macromolecules 2002, 35, 2365−2374. (9) Taden, A.; Landfester, K. Macromolecules 2003, 36, 4037−4041. (10) Young, W.-S.; Brigandi, P. J.; Epps, T. H., III Macromolecules 2008, 41, 6276−6279. (11) Chen, H.-L.; Li, H.-C.; Huang, Y.-Y.; Chiu, F.-C. Macromolecules 2002, 35, 2417−2422. (12) Cai, T.; Qian, Y.; Ma, Y.; Ren, Y.; Hu, W. Macromolecules 2009, 42, 3381−3385. (13) Chu, C.-Y.; Chen, M.-H.; Wu, M.-L.; Chen, H.-L.; Chiu, Y.-T.; Chen, S.-M.; Huang, C.-H. Langmuir 2014, 30, 2886−2895. (14) Suzuki, Y.; Duran, H.; Steinhart, M.; Butt, H.-J.; Floudas, G. Macromolecules 2014, 47, 1793−1800. (15) Michell, R. M.; Lorenzo, A. T.; Müller, A. J.; Lin, M.-C.; Chen, H.-L.; Blaszczyk-Lezak, I.; Martín, J.; Mijangos, C. Macromolecules 2012, 45, 1517−1528. (16) Loo, Y.-L.; Register, R. A.; Ryan, A. J.; Dee, G. T. Macromolecules 2001, 34, 8968−8977. (17) He, Y.; Zhu, B.; Kai, W.; Inoue, Y. Macromolecules 2004, 37, 3337−3345. (18) Martínez-Tong, D. E.; Vanroy, B.; Wübbenhorst, M.; Nogales, A.; Napolitano, S. Macromolecules 2014, 47, 2354−2360. (19) Opitz, R.; Lambreva, D. M.; de Jeu, W. H. Macromolecules 2002, 35, 6930−6936. (20) Wang, H.; Keum, J. K.; Hiltner, A.; Baer, E. Macromolecules 2010, 43, 3359−3364. (21) Zhu, L.; Calhoun, B. H.; Ge, Q.; Quirk, R. P.; Cheng, S. Z. D.; Thomas, E. L.; Hsiao, B. S.; Yeh, F.; Liu, L.; Lotz, B. Macromolecules 2001, 34, 1244−1251. (22) Zhu, L.; Cheng, S. Z. D.; Calhoun, B. H.; Ge, Q.; Quirk, R. P.; Thomas, E. L.; Hsiao, B. S.; Yeh, F.; Lotz, B. J. Am. Chem. Soc. 2000, 122, 5957−5967. (23) Huang, P.; Zhu, L.; Guo, Y.; Ge, Q.; Jing, A. J.; Chen, W. Y.; Quirk, R. P.; Cheng, S. Z. D.; Thomas, E. L.; Lotz, B.; Hsiao, B. S.; Avila-Orta, C. A.; Sics, I. Macromolecules 2004, 37, 3689−3698. (24) Huang, P.; Zhu, L.; Cheng, S. Z. D.; Ge, Q.; Quirk, R. P.; Thomas, E. L.; Lotz, B.; Hsiao, B. S.; Liu, L.; Yeh, F. Macromolecules 2001, 34, 6649−6657. (25) Nojima, S.; Ohguma, Y.; Namiki, S.; Ishizone, T.; Yamaguchi, K. Macromolecules 2008, 41, 1915−1918.

CONCLUSION Novel phase transitions of at-PEP, R-it-PEP, and S-it-PEP triblock copolymers with an unconventional thermal sequence have been shown to lead to hierarchical, lamellar nanostructures. During the crystallization process, PEO blocks crystallized in a soft-confining environment provided by a microphase-separated lamellar morphology formed at higher temperatures. Once crystallized, PEO provided a hard-confining environment that constrains the lateral crystallization of itPPO blocks. During the heating process, the PEO crystals melted first at lower temperature in the hard-confined environment provided by the PPO crystals. Subsequently, melting of the PPO crystals occurred in the soft-confining environment provided by surrounding layers of amorphous PEO, which was accompanied by recrystallization and crystal thickening of PPO. These results provide fundamental insights into the melting and crystallization behavior of doublesemicrystalline triblock copolymers that exhibit sequential phase transitions under both soft and hard confinement, resulting in hierarchically structured and orientationally correlated mesostructures. ASSOCIATED CONTENT

* Supporting Information S

Part A: melting behavior of at-PEP triblock copolymer; Part B: melting behavior of PEO homopolymer; Part C: crystallization behavior of R-it-PPO homopolymer; Part D: density correlation function for the crystallization of R-it-PEP; Part E: melting behavior of R-it-PPO homopolymer; Part F: crystallization behavior of S-it-PEP block copolymer; Part G: melting behavior of S-it-PEP block copolymer. This material is available free of charge via the Internet at http://pubs.acs.org.

■

■

This work was supported by the MRSEC Program of the National Science Foundation under Award DMR 1121053 (W.S., Y.L., N.A.L., C.J.H., G.H.F., E.J.K.). This work was partially supported by the National Institutes of Health as a Program of Excellence in Nanotechnology (HHSN268201000046C) (A.J.M., C.J.H., and N.A.L.). Beamline 7.3.3 of the Advanced Light Source is supported by the Director of the Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy, under Contract DE-AC02-05CH11231.

■

■

Article

AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected] (N.A.L.). *E-mail [email protected] (C.J.H.). *E-mail [email protected] (G.H.F.). Notes

The authors declare no competing financial interest. Edward J. Kramer passed away on December 27th, 2014. 3078


Article

Macromolecules (26) Hamley, I. W.; Castelletto, V.; Castillo, R. V.; Müller, A. J.; Martin, C. M.; Pollet, E.; Dubois, P. Macromolecules 2005, 38, 463− 472. (27) Castillo, R. V.; Müller, A. J.; Lin, M.-C.; Chen, H.-L.; Jeng, U.-S.; Hillmyer, M. A. Macromolecules 2008, 41, 6154−6164. (28) Kundu, C.; Dasmahapatra, A. K. Polymer 2014, 55, 958−969. (29) Myers, S. B.; Register, R. A. Macromolecules 2008, 41, 6773− 6779. (30) Li, S.; Myers, S. B.; Register, R. A. Macromolecules 2011, 44, 8835−8844. (31) Nojima, S.; Akutsu, Y.; Washino, A.; Tanimoto, S. Polymer 2004, 45, 7317−7324. (32) Chiang, Y.-W.; Huang, Y.-W.; Huang, S.-H.; Huang, P.-S.; Mao, Y.-C.; Tsai, C.-K.; Kang, C.-S.; Tasi, J.-C.; Su, C.-J.; Jeng, U.-S.; Tseng, W.-H. J. Phys. Chem. C 2014, 118, 19402−19414. (33) Lin, M.-C.; Chen, H.-L.; Su, W.-B.; Su, C.-J.; Jeng, U.-S.; Tzeng, F.-Y.; Wu, J.-Y.; Tsai, J.-C.; Hashimoto, T. Macromolecules 2012, 45, 5114−5127. (34) Higa, T.; Nagakura, H.; Sakurai, T.; Nojima, S. Polymer 2010, 51, 576−5584. (35) Lin, M.-C.; Wang, Y.-C.; Chen, J.-H.; Chen, H.-L.; Müller, A. J.; Su, C.-J.; Jeng, U.-S. Macromolecules 2011, 44, 6875−6884. (36) Nojima, S.; Fukagawa, Y.; Ikeda, H. Macromolecules 2009, 42, 9515−9522. (37) Temple, K.; Kulbaba, K.; Power-Billard, K. N.; Manners, I.; Leach, K. A.; Xu, T.; Russell, T. P.; Hawker, C. J. Adv. Mater. 2003, 15, 297−300. (38) Müller, A. J.; Arnal, M. L.; Balsamo, V. Lect. Notes Phys. 2007, 714, 229−259. (39) Nguyen-Dung, T.; Sasaki, S.; Masunaga, H.; Shimizu, N.; Igarashi, N.; Sakurai, S. Polymer 2014, 55, 2562−2569. (40) Jiang, S.; He, C.; Men, Y.; Chen, X.; An, L.; Funari, S. S.; Chan, C.-M. Eur. Phys. J. E 2008, 27, 357−364. (41) Shin, D.; Shin, K.; Aamer, K. A.; Tew, G. N.; Russell, T. P.; Lee, J. H.; Jho, J. Y. Macromolecules 2005, 38, 104−109. (42) Castillo, R. V.; Arnal, M. L.; Müller, A. J.; Hamley, I. W.; Castelletto, V.; Schmalz, H.; Abetz, V. Macromolecules 2008, 41, 879− 889. (43) Ruokolainen, J.; Mezzenga, R.; Fredrickson, G. H.; Kramer, E. J.; Hustad, P. D.; Coates, G. W. Macromolecules 2005, 38, 851−860. (44) McGrath, A. J.; Shi, W.; Rodriguez, C. G.; Kramer, E. J.; Hawker, C. J.; Lynd, N. A. Polym. Chem. 2015, 6, 1465−1473. (45) Wang, R.; Tang, Y.; Wang, Y. Langmuir 2014, 30, 1957−1968. (46) Wang, Q.; Li, L.; Jiang, S. Langmuir 2005, 21, 9068−9075. (47) Jiang, R.; Jin, Q.; Li, B.; Ding, D.; Shi, A.-C. Macromolecules 2006, 39, 5891−5896. (48) Meznarich, N. A. K.; Juggernauth, K. A.; Batzli, K. M.; Love, B. J. Macromolecules 2011, 44, 7792−7798. (49) Larrañaga, M.; Gabilondo, N.; Kortaberria, G.; Serrano, E.; Remiro, P.; Riccardi, C. C.; Mondragon, I. Polymer 2005, 46, 7082− 7093. (50) Cano, L.; Builes, D. H.; Tercjak, A. Polymer 2014, 55, 738−745. (51) Wang, S.-H.; Hou, S.-S.; Kuo, P.-L.; Teng, H. ACS Appl. Mater. Interfaces 2013, 5, 8477−8485. (52) Talibuddin, S.; Wu, L.; Runt, J. Macromolecules 1996, 29, 7527− 7535. (53) Chen, H. L.; Li, L. J.; Lin, T. L. Macromolecules 1998, 31, 2255− 2264. (54) Strobl, G. The Physics of Polymers, 2nd ed.; Springer: New York, 1978; p 151. (55) de Gennes, P. G. Scaling Concepts in Polymer Physics; Cornell University Press: Ithaca, NY, 1979. (56) Bates, F. S. Macromolecules 1985, 18, 525−528. (57) Sakurai, S.; Mori, K.; Okawara, A.; Kimishima, K.; Hashimoto, T. Macromolecules 1992, 25, 2679−2691. (58) Fredrickson, G. H.; Helfand, E. J. Chem. Phys. 1987, 87, 697− 705. (59) Matsen, M. W. Macromolecules 2012, 45, 2161−2165.

(60) Wu, B.; Jeff Harlan, C.; Lenz, R. W.; Barron, A. R. Macromolecules 1997, 30, 316−318. (61) Brandrup, J.; Immergut, E. H.; Grulke, E. A. Polymer Handbook, 4th ed.; John Wiley & Sons: New York, 2005. (62) Hamley, I. W.; Fairclough, J. P. A.; Terrill, N. J.; Ryan, A. J.; Lipic, P. M.; Bates, F. S.; Towns-Andrews, E. Macromolecules 1996, 29, 8835−8843. (63) Saffar, A.; Ajji, A.; Carreau, P. J.; Kamal, M. R. Polymer 2014, 55, 3156−3167. (64) Shi, W.; Yang, J.; Zhang, Y.; Luo, J.; Liang, Y.; Han, C. C. Macromolecules 2012, 45, 941−950.

3079


Cooperative and Sequential Phase Transitions in it-Poly(propylene

Recommend Documents