Enthalpy Relaxation and Morphology Evolution in Polystyrene-b-poly

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Article Cite This: Macromolecules 2018, 51, 7368−7376

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Enthalpy Relaxation and Morphology Evolution in Polystyrene‑b‑poly(methyl methacrylate) Diblock Copolymer Mingchao Ma,†,‡ Yage Huang,†,‡ and Yunlong Guo*,†,‡,§ †

The State Key Lab of Metal Matrix Composites, ‡University of Michigan−Shanghai Jiao Tong University Joint Institute, and § School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China

Macromolecules 2018.51:7368-7376. Downloaded from pubs.acs.org by UNIV OF TOLEDO on 09/26/18. For personal use only.

S Supporting Information *

ABSTRACT: Desirable patterns formed by microphase separation in block copolymers make these materials attractive for miniaturization of functional devices in which formation of nanostructures is highly demanded. The microphase separation forms microdomains under three-dimensional confinement for each component and thus substantially changes physical properties such as structural relaxation of the polymer blocks. In this article, we report enthalpy relaxation and morphology evolution in polystyrene-b-poly(methyl methacrylate) (PS-b-PMMA) during isothermal physical aging, measured by calorimetry and atomic force microscopy (AFM). Under nanoconfinement formed by microphase separation, the PS blocks have higher relaxation rate and amount of relaxed enthalpy compared with the corresponding homo-PS, while the PMMA blocks have lower relaxation rate and relaxed enthalpy than the homo-PMMA. The in situ morphology evolution shows that the characteristic distribution profiles of the AFM phase angle become lower and wider with increasing aging time for both the blocks, and the average distance between the central phase angles of the two blocks increase during aging. As large AFM phase angle of PMMA representing higher modulus, it demonstrates that the PS blocks under hard confinement exhibit greater enthalpy change with enhanced relaxation rate from the bulk. In contrast, PMMA blocks under soft confinement show the opposite trend.

1. INTRODUCTION Upon rapid cooling to a temperature below the glass transition temperature (Tg), a polymeric material enters its glassy state; subsequently, many physical properties such as specific volume,1−3 modulus,4−6 enthalpy,7−9 and fluorescence response10−12 gradually change with time. This evolution process is known as structural relaxation or physical aging.13−15 Structure relaxation is of significant importance for polymers as they are widely used as functional or structural materials and has attracted remarkable research interest in the past several decades.1−13,16,17 Although many efforts have been devoted on physical aging of homopolymers, polymer blends, and polymer composites, structural relaxation of copolymers has not been well understood. For instance, Tant and Wilkes measured both the stress relaxation and enthalpy recovery in styrene− butadiene and carbonate−siloxane block copolymers but did not make a direct comparison for relaxation behavior in thermal and mechanical response.18 Araki and co-workers investigated physical aging of an ethyl methacrylate (EMA)−methyl methacrylate (MMA) random copolymer via dynamic viscoelastic measurements.19 In their study, the copolymer sample contained 3.9 mol % of EMA. It was found that relaxation rate was obviously faster than that of the homo-PMMA when aged below Tg of PEMA. Moreover, Cameron et al. reported physical aging of styrene−maleic anhydride (MA) random copolymer with various component fractions.20 They observed a single © 2018 American Chemical Society

enthalpy relaxation process of the bicomponent system and found that the relaxation speed was not monotonous with respect to the component fraction, i.e., the relaxation rate showing a complicated trend with increasing MA content. Yun and co-workers examined the mechanical response of PS/PMMA blends, PS-b-PMMA copolymers, and the corresponding homo-PS and homo-PMMA during physical aging by means of microcantilever deflection measurements.21 In their results, the temperature dependence of the deflections of copolymercoated cantilever after aging showed one broad peak due to overlapped Tgs of the PS and PMMA components; hence, the individual aging response of the blocks was not identified. Although a complicated aging phenomenon was reported, the majority of existing research in the literature did not systematically investigate the physical aging of polymer blocks and answer a question about how the aging behavior deviates from the homopolymer or polymer blends. The application of block copolymers could be bottlenecked due to lack of the knowledge of structure relaxation, which dramatically alters properties of polymers in their service life. Today block copolymers are considered as a key element in material solution for emerging technologies.22−31 For instance, block copolymers can be used as the templates or scaffolds in Received: June 21, 2018 Published: September 14, 2018 7368

DOI: 10.1021/acs.macromol.8b01323 Macromolecules 2018, 51, 7368−7376

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blocks, the Tas were set at least 33 °C lower than Tg,PMMA block or at least 5 °C higher than Tg,PS block. As such, the aging response of one block (PMMA at low Tas or PS at high Tas) was almost totally suppressed in our aging time up to 38400 s due to large departure from Tg or thermodynamic equilibrium. The aging response reflected in heat capacity overshoot is attributed from one block in the copolymers. Therefore, by selecting pertinent Tas, structural relaxation in individual polymer blocks can be measured by calorimetry. The enthalpy relaxed during aging process (ΔHa) was calculated by integrating of the heat capacity discrepancy between aged and unaged data over a temperature range, as express by eq 1:

nanofabrication by utilizing their hierarchical nanostructures formed by microphase separation.22−24 In recent years, the conjugated diblock copolymers play a rising role in photovoltaics.29−31 That is, the microphases of the two blocks serve as electron donor and acceptor with capacity of transporting charge in high efficiency. Additionally, drug delivery based on copolymer micelles shows its promising potential in biomedical research and development.25−28 Many studies have been conducted to develop smart and functional copolymers in meeting material requirements by some specific techniques. A general strategy for creating and tuning nanostructures in block copolymers utilizes the diversiform microphase separation through thermal annealing or organic vapor treatment. Microphase separation forms block domains on a nanometer scale with multitype geometries such as cylinders, thin layers, gyroids, and other complicated structures. The blocks in these microdomains are nanoscopically confined by other component blocks in the macromolecular chains, and vice versa. Therefore, block copolymers provide a model system to investigate the nanoconfinement effect on glassy dynamics of polymers. The structural relaxation of block copolymer under such confinement, including the aging behavior of each block and the influence of interface interactions, still remains to be elucidated. In this article, we report enthalpy relaxation of a diblock copolymer under microphase separation. The aging response of each block was individually captured and compared with the corresponding homopolymer. The morphology evolution during aging was monitored by an atomic force microscope (AFM) to compare the response in structure evolution and enthalpy recovery.

ΔHa =

∫T

T2

1

(Cp ,aged(T ) − Cp ,unaged(T )) dT

(1)

where cp,aged and cp,unaged are the heat capacity of the aged and unaged samples, respectively. T1 and T2 are reference temperatures far below and far above Tg where the measured heat capacity data of unaged and aged sample are equal. Morphology Evolution Measurement. Morphology evaluation of PS-b-PMMA copolymer during physical aging was monitored by a Bruker multimode 8 AFM equipped with a heating stage. To fabricate a sample, first, a solution of 1 wt % PS-b-PMMA copolymer in chloroform was prepared. By use of this solution, copolymer films were produced from spin-coating at 800 rpm for 50 s on a silicon wafer with native oxide top layer. The thickness of resultant film was 350 nm, determined by AFM. After spin-coating, the film was annealed at 100 °C under vacuum overnight to remove the remaining solvent and residual stress. The film was taken out and the morphology was directly checked by AFM. It was found that the film was pretty flat, and no phase separation was detected. To examine the morphology change during aging, the samples underwent the same temperature histories used in DSC testing. That is the film was heated to 180 °C and was held at that temperature for 10 min, followed by a quench to the desired aging temperature. The morphology of the sample under various aging times was observed by in situ AFM measurements.

2. EXPERIMENTAL SECTION Materials. The PS-b-PMMA diblock copolymer was purchased from Polymer Source Inc. and used without any further purification. The weight-average molecular weight of the PS block (Mw,PS block) and the PMMA block (Mw,PMMA block) are 259000 and 629000 g/mol, respectively. Thus, the mole fraction of the blocks PS is 29%. The polydispersity index (PDI) of 1.08 was determined by size exclusion chromatography (SEC). PS-b-PMMA was chosen to investigate structural relaxation in blocks under confinement because both blocks are representative amorphous polymers; thence, aging should be explicit. In addition, the Tgs of the two blocks depart from each other by ∼20 °C so that the aging responses are not overlapped and can be separated on some aging temperatures. The homopolymers PS and PMMA used in our control experiments were purchased from Sigma-Aldrich and American Polymer Standards Corporation, respectively, and used as received. The Mw of homo-PS is 254 kg/mol and PDI is 2.83, whereas the Mw of homoPMMA is 620 kg/mol and PDI is 1.05. For a direct comparison, both of the homopolymers have similar molecular weight with the corresponding polymer blocks. The materials above are the same as what we used in our previous work.32 Enthalpy Relaxation Test. Enthalpy relaxation measurements were performed by a Shimadzu DSC-60 plus differential scanning calorimeter (DSC) under a nitrogen atmosphere. The instrument was calibrated by indium. In aging testing, a sample of 3−5 mg was first sealed in an aluminum pan; subsequently, it was heated to 180 °C (well above the Tgs of the both polymer blocks) with a constant rate of 20 °C/min and held for 10 min to remove residual stress. The sample was then quenched to an aging temperature (Ta) with a cooling rate of 30 °C/min and annealed thereafter for a specific aging time (ta). The aging effect was characterized by the thermogram when reheating the sample from Ta to 180 °C at 20 °C/min. The details of this test protocol were described in our previous work.32 The glass transition temperatures of the PS and PMMA blocks are 109 and 127 °C, respectively. To obtain uncoupled aging response of the two

3. RESULTS AND DISCUSSION Figure 1 depicts physical aging behavior of PS-b-PMMA when the temperature is far (at least 33 °C) lower than Tg,PMMA block. The aging temperatures in Figures 1a and 1b are 89 and 94 °C, corresponding to Tg,PS block − 20 °C and Tg,PS block − 15 °C, respectively. No apparent heat capacity peaks of PMMA block can be distinguished in Figure 1 on our experimental aging time scale up to 38400 s. Hence, we consider that all gradual changes in heat capacity curves are contributed by aging of the PS block. When Ta increases from 89 to 94 °C, the maximum of heat capacity (i.e., cp curves at longest ta) decreases a little bit, while the aging response in short ta at 94 °C is apparently greater than that at 89 °C. These observations reflect the fact that the higher thermodynamic driving force initiating structural relaxation at higher Ta becomes sluggish in further aging when the material is getting closer to equilibrium. When a Ta is far lower than Tg,PMMA block, the aging response of block copolymer in these DSC traces merely comes from the PS block. As such, we can directly compare the response from PS block and the corresponding homo-PS with same molecular weight. Selected representative aging results of homo-PS with same temperature departure from Tg are shown in Figure 2. Tg,PS is equal to 106 °C so that two aging temperatures are 86 and 91 °C. For homo-PS, the maximum value and the shape of the heat capacity curves at longest ta are nearly the same, indicating that the relaxed energies during physical aging are very close, as are the associated average aging rates. This result is consistent with the findings in the literature; i.e., the aging rate in a 7369

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Figure 1. Heat capacity of PS-b-PMMA diblock copolymer under lower aging temperatures: (a) Ta = Tg,PS block − 20 °C; (b) Ta = Tg,PS block − 15 °C.

Figure 3. Heat capacity of PS-b-PMMA diblock copolymer under high aging temperatures: (a) Ta = Tg,PMMA block − 13 °C; (b) Ta = Tg,PMMA block − 8 °C.

identical thermogram under aged and unaged conditions in the vicinity of the Tg,PS block. When the aging temperature increases from 114 to 119 °C, the maximum of heat capacity curve at the same ta decreases. The sample reaches thermodynamic equilibrium after ta = 9600 s, demonstrated by the overlapped cp curves when ta ≥ 9600 s, when Ta = Tg,PMMA block − 8 °C. Therefore, the material approached the upper bound of enthalpy span for thermal energy relaxation under such condition. In contrast, the aging process was kept ongoing as ta up to 38400 s, when Ta = Tg,PMMA block − 13 °C. Similarly, the results of aging tests of homo-PMMA are shown in Figure 4. The glass transition temperature of homoPMMA is 3 °C higher than Tg,PMMA block, i.e., 130 °C, so the aging temperatures for homo-PMMA are 117 °C and 122 °C, 13 °C and 8 °C lower than its Tg, respectively. From these results, we find that the time for homo-PMMA to reach its equilibrium is shorter than that of PMMA block in the copolymer. To characterize the enthalpy change and aging speed, we calculated the enthalpy from the heat capacity curves. For homopolymers, an integral of cp results in the enthalpy, but for the copolymer sample, a calculation method needs to be established to carry out the enthalpy relaxation of each individual block from the apparent cp data as shown in Figures 1 and 3. Our aim is to compare the difference between one block of copolymer and its corresponding homopolymer. The apparent cp curves of the copolymer sample were calculated by

Figure 2. Heat capacity of homo-PS under very low aging temperatures: (a) Ta = Tg,PS − 20 °C; (b) Ta = Tg,PS − 15 °C.

certain temperature range not far below Tg is not sensitive to the aging temperature.33,34 Figure 3 illustrates the heat capacity curves of the diblock copolymer when aging temperatures are Tg,PMMA block − 13 °C and Tg,PMMA block − 8 °C, i.e., 114 °C and 119 °C, respectively. The aging temperature is at least 5 °C higher than Tg,PS block; thus, we speculate that aging response here comes from the PMMA block, and this speculation is confirmed from the

cp,copolymer =

1 dQ copolymer /dt m dT /dt

(2)

where cp,copolymer is the heat capacity of the copolymer, m is the mass of whole copolymer sample, dQ copolymer/dt is the heat flow signal from DSC, and dT/dt is the heating rate. Considering the diblock copolymer sample, the heat flow originates from two blocks as 7370

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Figure 4. Heat capacity of homo-PMMA under high aging temperatures: (a) Ta = Tg,PMMA − 13 °C; (b) Ta = Tg,PMMA − 8 °C.

dQ copolymer /dt = dQ PS block /dt + dQ PMMA block /dt

Figure 5. Relaxed enthalpy of PS in copolymer and homopolymer under very low aging temperatures. The error bar represents the standard deviation of data of three duplicates, and β is the aging rate from linear fit. (a) Ta = Tg,PS/PS block − 20 °C; (b) Ta = Tg,PS/PS block − 15 °C.

(3)

where dQPS block/dt and dQPMMA block/dt are the heat flow signals of PS block and PMMA block, respectively. Choosing PMMA block enthalpy relaxation as an example, the aging temperature is higher than glass transition temperature of PS block, so the aging response is considered only coming from the PMMA block. The enthalpy change of PMMA block during aging, ΔHa,PMMA block, was calculated by the following equation: ΔHa,PMMA block =

∫T

T2

Enthalpy relaxation results of PS block and homo-PS are illustrated in Figure 5. The aging rate β represents the slope of the linear fit of the data. From the results, not only the ΔHa but also the aging rates of PS block are much higher than the values of homo-PS, so that PS block has a larger magnitude and velocity in enthalpy relaxation under similar thermal conditions. For the PS block shown in red, β decreases with increasing Ta, and the maximum value of ΔHa also decreases. These trends are in agreement with previous findings in homopolymers that showed lower driving force but larger energy space in structural relaxation at low Tas.4,35−37 For homo-PS, β and the maximum ΔHa may increase slightly, but on the whole there was no apparent change under these temperatures; it implies that the homo-PS is probably on the aging rate plateau (that is, β stays in high level in a Ta range before dropping down when Ta gets closer to Tg), as shown in the literature.33 Note that the PDI of the homo-PS is much larger than the PDI of the copolymer. To assess whether the high PDI remarkably affects aging behavior, we employed another homoPS as a reference material. This homo-PS was purchased from Shanghai Ziqibio Co., Ltd., and used as received. The Mw of this homo-PS is 213.6 kg/mol, and the PDI is 1.01. The Mw is ∼18% lower than PS block, but the PDI is pretty good. The enthalpy relaxation and associated aging rate of this homo-PS are shown by the green symbols and lines in Figure 5. The results are consistent with the former homo-PS, confirming that the aging behavior determined from these homo-PS samples is reproducible and reliable. Figure 6 illustrates the enthalpy recovery of the PMMA block and homo-PMMA. With aging temperature increases from Tg − 13 °C to Tg − 8 °C, the maximum value of relaxed enthalpy decreases greatly. The PMMA block and homoPMMA both reach their equilibrium states after a short aging

(Cp ,coplymer,aged(T ) − Cp ,coplymer,unagd(T ))

ij 1 dQ PMMA block,aged /dt jj jj MPMMA dT /dt 1 jm k 1 dQ PMMA block,unagd /dt yzzz Mcopolymer − dT zz × z m MPMMA dT /dt { dQ PMMA block,aged/dt T2 i 1 jjj = jj T1 j mPMMA block m k dQ PMMA block,unaged /dt yz 1 zzdT − zz z mPMMA block dT /dt { 1

×

Mcopolymer

dT =

∫T

T2

∫

=

∫T

T2

(Cp ,PMMA block,aged(T ) − Cp ,PMMA block,unaged(T )) dT

1

(4)

where cp,copolymer, aged and cp,copolymer,unaged are the heat capacity of aged and unaged sample, respectively. T1 and T2 are reference temperatures far below and far above Tg where the measured heat capacities of unaged and aged sample are equal. Mcopolymer and MPMMA block are the molar mass of the copolymer chain and the PMMA block, respectively. mPMMA block is the mass of PMMA block of the sample. Applying the equations above, we obtained the enthalpy change of the PMMA block. The enthalpy relaxation of PS block was determined by a similar method. 7371

DOI: 10.1021/acs.macromol.8b01323 Macromolecules 2018, 51, 7368−7376

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may reflect the influence of nanoconfinement and interfacial effect. On the other side, the aging may also impact microstructure of the block copolymer. To investigate aging effect on microstructure, the morphology evolution of PS-b-PMMA was monitored by in situ AFM measurements. Next, we will show representative AFM images at various aging times and then compare the results of enthalpy relaxation and morphology evolution. We also make discussion on the aging results of PS-b-PMMA and those for thin polymer films, nanoparticles, nanocomposites, etc., in which the polymeric materials showed deviation of out-of-equilibrium dynamics from the bulk when confined on the nanoscale.35 The samples for AFM measurement are PS-b-PMMA thin films. To prevent importing additional confinement from the film thickness, the samples should be sufficiently thick. It is known that the properties of ultrathin polymeric films are different from bulk materials.43,44 The thickness threshold for nanoconfinement is usually defined as ∼100 nm,45−47 although some work set this critical value as 200 nm.48,49 To avoid the thickness effects and establish the relationship between the morphology and the enthalpy recovery, the thicknesses of our films are ∼350 nm, controlled by the speed in spin coating, initial solution concentration, and the solvent. With sufficient thickness, our AFM samples are considered having bulk morphology, and the morphology change corresponds to structure evolution during physical aging. The AFM phase images of PS-b-PMMA copolymer in aging test at 89 °C, as a representative example in a series of experiments, are shown in Figure 7. The color scales from −6° to 6° are kept the same for the four panels. The microphase separation is already set up prior to aging experiments. Figure 7 reveals that the shape and size of the microdomains did not change a lot during aging, but the colors of the patterns, i.e., the average AFM phase angle which represents the intensity of phase separation, changed especially for ta = 38400 s. To explore and analyze the evolution of the distributed phase angle of each polymer block, numerical data of phase angle from AFM measurements were exported. For the purpose of a direct comparison, the raw data were shifted so as to set the average phase angle of the whole diblock copolymer at each aging time to be zero. Thus, the intensity of microphase separation can be expressed by the distance between the average phase angles of the two blocks, while the structural relaxation of individual block could be featured by the change of geometric characteristics (such as the height or width) of the distribution function. The phase angle histograms of the blocks (in color) and the whole sample (black) are summarized in

Figure 6. Relaxed enthalpy of PMMA in copolymer and homopolymer under high aging temperatures. The error bar represents the standard deviation of data of three duplicates, and β is the aging rate from linear fit. (a) Ta = Tg,PMMA/PMMA block − 13 °C; (b) Ta = Tg,PMMA/PMMA block − 8 °C. All data points were used to do the linear fit of the equilibrium state for PMMA block, but only last five points were used for homo-PMMA. The black dashed line is a guide for the eyes.

when Ta = Tg − 8 °C. Thus, the limited enthalpy span at relatively high Ta restricts the relaxation duration in both materials, but the average relaxation speed differs. From Figure 6a, the value of ΔHa of homo-PMMA is higher than the ΔHa of PMMA block; this result is in contrast with the enthalpy relaxation of PS block and homo-PS. The aging rate of homo-PMMA is greater than PMMA block. From Figure 6b, the ΔHequilibrium of homo-PMMA is larger than PMMA block, implying that the enthalpy span of homo-PMMA is larger than PMMA block. As above-mentioned, the PS and PMMA blocks should age under confinement induced from microphase separation as the sample underwent high-temperature equilibration at 180 °C (slightly above the order−disorder transition temperature, TODT38−42) before quenching to the aging temperature. Hence, the difference between polymer blocks and their homopolymers in enthalpy recovery

Figure 7. Phase figure of PS-b-PMMA diblock copolymer under very low aging temperature (Tg,PS block −20 °C). The sample was heated to 180 °C and held for 10 min first, which is like the aging test. The color scales −6° to 6° were set the same for all the panels. The diagrams are from same position whose size is 2 × 2 μm with different aging time: (a) Unaged, (b) 4800 s, (c) 24000 s, and (d) 38400 s. 7372

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Figure 8. Phase angle histograms of PS-b-PMMA diblock copolymer under very low aging temperature (Tg,PS block − 20 °C). The black cityscapes are the phase angle statistics of whole sample. The red and blue bars are the statistics of PS block and PMMA block from AFM results, respectively. The sum of two blocks contribution is equal to the whole sample distribution histograms. (a) Unaged, (b) ta = 4800 s, (c) ta= 24000 s, and (d) ta= 38400 s.

Ta = 89 °C. Here we notice that the magnitude of morphology change at 114 °C is larger than that at 89 °C. As shown in Figure 9, at the end of our aging testing, the deviation of all geometric parameters from unaged status (dashed lines) is larger when Ta = 114 °C. Another thing which should be noticed is that the phase angle or the phase separation did not change during the first 24000 s, but it apparently changed in the next 14400 s under isothermal conditions. This demonstrates that the time scale of structural relaxation on morphology evolution and enthalpy recovery might be different.50 Initially, the aging effect does not cause obvious phase variation. However, when aging time reaches some critical value, the microphase of the polymer block will have a noticeable change. The change of peak position in opposite directions of the two blocks indicates the enhancement of microphase separation with physical aging. The wider distribution of the phase angle of the blocks may reflect an increasing variety of chain conformations of each block toward a larger degree of disorder or disorganization. From the AFM images, the phase angle of the PMMA block is larger than that of the PS block since PMMA has higher elasticity.36,38,39,51 This result is in line with intuition, considering the Ta is closer to the Tg,PS block than Tg,PMMA block. Hence at the aging temperatures in this study, the PMMA and PS blocks respectively provide their chain partners relatively hard and soft confinement. The polymer−polymer interfaces influence the dynamics of a substantial volume of the chain blocks in the microdomains. The non-neutral polymer−polymer interface likely impacts structural relaxation of the adjacent domains in a 2-fold effect which accelerates releasing speed of the trapped-in thermal energy for one domain and depresses the speed in the other domain. Given the diverse relaxation rates in enthalpy recovery of PS and PMMA blocks, we consider the interface plays a key role in such change of the aging speed. The PS-bPMMA copolymer with phase separation, to some extent,

Figure 8. The distribution functions of the two blocks overlap in an intermediate area. The percentage of dark phase in Figure 7 is in line with the PS molar fraction of PS-b-PMMA copolymer (f PS = 0.29), so the dark phase in Figure 7 represents the microdomain of PS, while the bright domain should be PMMA. As shown in Figure 8, the overall phase angle data of the copolymer are identical to the summation of the two blocks. Therefore, digitizing the phase angle distribution of each polymer block from the AFM image and tracking the change in geometry account for the physical aging effect on microstructure evolution in the copolymer. To describe the alteration of phase angle with increasing aging time, three parameterspeak height, peak position, and full width at half-maximum (FWHM)were selected to illustrate the change of distribution function. The same AFM experiment and analysis were also performed for Ta = 114 °C. Figure 9 depicts these parameters versus aging time under the two Tas. For the PMMA block which is shown in Figure 9a−c, the peak height, peak position, and FWHM almost stay constants when aging time is less than or equal to 24000 s in all cases. However, when aging time reaches 38400 s, all the three parameters obviously alter from the previous values. From Figure 9a−c, the peak height decreases and the FWHM increases, which demonstrates that the characteristic distribution function of PMMA phase angle becomes lower and wider. The peak position increase represents the increase of average phase angle of PMMA domains in the AFM image. Figure 9d−f presents the results of the PS block. When aging time is up to 24000 s, no obvious alteration occurs which is similar to the result of PMMA block. Analogous to the response of the PMMA block, the characteristic distribution peaks of PS block also become lower and wider when the aging time increases to 38400 s. In contrast to PMMA block, the negative peak position of PS moves obviously toward the direction of phase angle decrease when Ta = 114 °C but almost no shift when 7373

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work are significantly different. The difference is reflected in two aspects. First the coefficients of thermal expansion of PS and PMMA are quite close in a wide temperature range across Tg. Hence, the two polymeric blocks have similar and nonnegligible volume shrinkage during the quench and aging process, no isochoric confinement was formed due to lack of rigid material. Second, the microphases in diblock copolymers are physically connected by covalent bonds. When the material was quenched from 180 °C and was aged below Tg, the discrepancy of potential shape and volume change in adjacent microdomains caused by structural relaxation should be inhibited by the deformation of the chains at interface which bridge the microdomains. That is, the intramolecular force of the macromolecules at interface varies with temperature and aging time; the force is eventually balanced by intermolecular forces from other chains and acted on the bridging chains via entanglements. As such, the intrinsic elasticity induces hard or soft confinement of polymer blocks; in each block the internal stress level differs from the bulk material, which may have an important influence in aging behavior of the diblock copolymers. We speculate that the enhanced and reduced aging rate should be attributed to the hard and soft confinement, respectively. This phenomenon will certainly be further investigated to attain a better understanding of the relationship between 3D confinement environment and the behavior of structural relaxation.

4. CONCLUSIONS The physical aging of a PS-b-PMMA copolymer and the corresponding homo-PS and homo-PMMA with same molecular weight was investigated by thermal analysis and in situ morphology tracking via AFM. The PS blocks in copolymer show higher aging rate and ΔHa compared with homo-PS at relatively low Tas. On the contrary, the PMMA blocks in copolymer exhibit lower aging rate and ΔHa than the homo-PMMA when Ta is higher than Tg,PS block. In morphology evolution measurement, we find that with increasing aging time, the characteristic distribution peaks of the phase angle become lower and wider for both two blocks. Furthermore, the time scales of structural relaxation on morphology evolution and enthalpy recovery might be different.

Figure 9. Statistics of phase angle of PS-b-PMMA diblock copolymer under two temperature conditions. (a−f) The data from the fitted curves of phase angle histograms. The error bar represents the standard deviation of four duplicates. (a−c) PMMA block. (d−f) PS block. (a, d) The peak value of probability density. (b, e) The position (phase angle) of maximum probability density. (c, f) Full width at half-maximum of phase angle histograms.

could be treated as polymer nanocomposite. We compare the results in this study with the findings in dynamics of polymer under confinement. Our results are consistent with existing work in the literature for other polymeric systems;17,52−56 it was reported that polymers under hard confinement (silica or gold nanocomposites) accelerated the aging rate more than the neat polymers, similar to what we found in PS blocks. Also, the aging rate is discovered increasing with increasing ratio of hard confinement.52,53,56 That will be checked for diblock copolymer in our next work by investigating physical aging behavior of PS-b-PMMA copolymer with different molar fractions of the two blocks. We noticed that there are several studies discussing the negative pressure effect on relaxation or dynamic behavior of glass-forming liquids quasi-isochorically confined in nanopores.57,58 The isochoric condition is generated by a rigid material, e.g., silica, which surrounds the soft glass-former. When the temperature deviates from the processing temperature on that the glass-former was placed into the nanopores, negative or positive pressure is consequently produced between the soft and rigid materials because of the difference of coefficient of thermal expansion. If the temperature change is not quite large, volume variation of the rigid material could be negligible; hence, the inner glass-former is considered under isochoric or “hard” nanoconfinement. Comparing with these material systems, we think the “soft” and “hard” confinement conditions exerted on the PMMA and PS blocks in present

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b01323. Figures S1−S3 (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (Y.G.). ORCID

Yunlong Guo: 0000-0002-4490-2140 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors acknowledge the National Science Foundation of China for financial support through the General Program 2157408. Y.G. is very grateful to the National Youth 1000 Talent Program of China, the Shanghai 1000 Talent Plan, and 7374

DOI: 10.1021/acs.macromol.8b01323 Macromolecules 2018, 51, 7368−7376

Article

Macromolecules

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the support by the Scientific Research Foundation for the Returned Overseas Chinese Scholars, State Education Ministry of China. The authors also acknowledge the start-up fund of Y.G. from both University of Michigan−Shanghai Jiao Tong University Joint Institute and School of Materials Science and Engineering at SJTU. We thank Prof. Xinping Wang and Prof. Biao Zuo for their assistance in morphology evolution measurements by AFM at Zhejiang Sci-Tech University.

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DOI: 10.1021/acs.macromol.8b01323 Macromolecules 2018, 51, 7368−7376

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DOI: 10.1021/acs.macromol.8b01323 Macromolecules 2018, 51, 7368−7376