Controlling Rheology of Fluid Interfaces through Microblock Length of Sequence-Controlled Amphiphilic Copolymers

Abstract

Previous studies have demonstrated that films of sequence-controlled amphiphilic copolymers display contact angles that depend on microblock size. This suggests that microblock length may provide a means of tuning surface and interfacial properties. In this work, the interfacial rheology of a series of sequence-controlled copolymers, prepared through the addition of bicyclo[4.2.0]oct-1(8)-ene-8-carboxamide (monomer A) and cyclohexene (monomer B) to generate sequences up to 24 monomeric units composed of (A_mB_n)_i microblocks, where m, n, and i range from 1 to 6. Interfacial rheometry is used to measure the mechanical properties of an air–water interface with these copolymers. As the microblock size increases, the interfacial storage modulus, G′, increases, which may be due to an increase in the size of interfacial hydrophobic domains. Small-angle X-ray scattering shows that the copolymers have a similar conformation in solution, suggesting that any variations in the mechanics of the interface are due to assembly at the interface, and not on solution association or bulk rheological properties. This is the first study demonstrating that microblock size can be used to control interfacial rheology of amphiphilic copolymers. Thus, the results provide a new strategy for controlling the dynamics of fluid interfaces through precision sequence-controlled polymers.

1. Introduction

Sequence-controlled polymers are macromolecules in which the sequence of monomer units of different chemical nature is controlled to some degree.^[1] Monomer sequence is a powerful tool to control the copolymer properties. Compared with homopolymers and block copolymers, sequence-controlled polymers display unique morphology and properties, such as folding and self-assembly,^[1] leading to possible implications for applications such as delivery. Previous reports have focused on the effect of sequence on the physical properties of copolymers with well-controlled microstructure. Meyer and co-workers performed comprehensive studies to show the importance of monomer sequence in copolymer properties.^[2] These studies focused on a family of complex isotactic, syndiotactic, and atactic repeating sequence poly(lactic-co-glycolic acid) copolymers, which show that monomer sequence can improve thermal properties and may have an impact on degradation and release behaviors.

Recently, a one-pot, non-templated synthetic approach for the preparation of copolymers (Figure 1) with well-controlled monomer sequences was developed in the Sampson group.^[3] In the same study, we examined the copolymers at varying microblock size.^[3] Contact angle measurements on spin-coated films in that study showed that increasing the microblock size from one to six can result in increasing water contact angles. This suggests that monomer sequence may be important in controlling the assembly of copolymers and in creating coatings with specific properties. In our earlier work, we focused on studying the properties of final dried films. However, to better understand how monomer sequence may affect the properties at air–water interfaces, we examined the interfacial rheology of this series of polymers prior to film formation in this work.

Figure 1.

a) Chemical structure of copolymers. b) Illustration of the copolymers. Modified from Li et al.^[3]

The study of rheology at the interface can have important implications for the properties and behavior of materials over a broad range of applications. Over the years, many studies have contributed to the understanding of interfacial rheology through dilational and shear rheology methods.^[4] While these methods both have their own advantages and disadvantages, advancements made in interfacial shear rheology measuring techniques have garnered much attention in recent years. Interfacial shear rheology data has often suffered from poor reproducibility due to the complexity of deformation in systems at the interface. This method of measurement is characterized by the functional relation between the deformation generated by the movement of the measuring probe (stress exerted in and on the interface), the deformation at the interface, and the resulting flow of the adjacent subphase. Therefore, the contact made by the measuring probe at the interface and the geometry of rheological flow field plays an important role and has to be carefully understood throughout the entire measuring process. Improvements in instrumentation have allowed researchers to more accurately probe the interfacial properties of a wide variety of complex materials. Recently, interfacial rheometers that utilize a double-wall ring (DWR) geometry that allows simultaneous and independent visualization of a planar interface have been developed and used to examine interfacial mechanics of particle-laden interfaces.^[5] Steady shear, linear viscoelasticity, nonlinear viscoelasticity, and extensional responses of carbon nanotube have been measured with this type of geometry, and the surface pressure and microstructure of carbon nanotube at an air–water interface were also studied.^[6] There are also reports of interfacial dynamics and rheology of polymer-grafted nanoparticles at air–water interface and xylene–water interface to characterize the impact of convection, step changes in the bulk concentration, and the interfacial mechanics.^[7] A rheological study on poly(lactic/glycolic) acid (PLGA) films at air–water interface explored detailed structural and dynamic properties of heterogeneous domains and their influences on the mechanical properties, and the interfacial rheology study results suggest that as-spread PLGA films are in a glassy state when subjected to deformation.^[8] Alternating copolymers with a substituted acetal-containing backbone exhibit viscoelastic behavior with a crossover frequency which decreases as the size of the R group on the acetal increases.^[9]

Viscoelastic liquids and solids can generally be described in terms of their viscous response to stress described by the loss modulus, G″, and their elastic response to stress described by the storage modulus, G′. The storage modulus G′ represents the stored deformation energy, and the loss modulus G″ characterizes the deformation energy lost or dissipated through internal friction during flow or deformation. If G′ > G″ over the range of measurable frequencies, the solid-like or elastic character of the material dominates its response, the material can be termed a gel or viscoelastic solid. When G″ > G′, the liquid-like character of the material dominates its mechanical response. If G′ tends toward zero at low frequencies, the sample can be termed as viscoelastic liquid. The crossover point is defined as the frequency when G′ = G″, and this parameter is often related to a characteristic time for relaxation in the material.^[10]

To provide additional insight into the behavior of these polymers and understand whether any differences in the interfacial rheology were due to solution assembly or interfacial assembly, we also performed small-angle scattering on solutions. In small-angle X-ray scattering (SAXS) experiments, the scattering intensity I(q) is measured as a function of the magnitude of the scattering vector q(λ,θ) = (4π/λ)sin(θ/2). Here, λ is the wavelength of the incident radiation, θ is the scattering angle. The scattering intensity can be expressed as follows.

$$I(q)=NP(q)S(q)=\Phi \Delta p^{2}VP(q)S(q)$$ (1)

where N is the number of particles per unit volume and is equal to the product of volume fraction (Φ, unitless), the volume of the scattered objects (V, units of cm³), and the contrast term (Δp², units of cm⁻⁴). The form factor, P(q) contains information regarding particle size and shape. The structure factor, S(q), contains information concerning the relative position of scattering sources.

Here, we report results from our SAXS and interfacial rheology study on sequence-controlled copolymers. Synthesis of these polymers and contact angle measurements on dry films has been reported in a previous publication.^[3] The polymers are depicted in Figure 1.^[3]

2. Experimental Section

2.1. Preparation of Copolymer Solutions

Copolymers (Figure 1a) were prepared as described in Li et al.^[3] using an alternating ring-opening metathesis polymerization method for stepwise incorporation of one to six bicyclo[4.2.0]oct-1(8)-ene-8-carboxamide monomers into a growing, alternating oligomeric chain via sequential, stoichiometric monomer addition.^[3] Cyclohexene was used as the linker between carboxamides with hydrophobic or hydrophilic side chains. Sequence-controlled copolymers were synthesized with a constant ratio of hydrophilic and hydrophobic side chains presented in different alternating sequences (Figure 1b). NMR results were shown in their earlier paper,^[3] along with GPC results that indicate a M_n of 2.1–2.8 kDa for the series of P1 to P6, with disparities in the range of 1.1–1.3. Each 1.0% w/v copolymer solution was prepared by addition of anhydrous chloroform to a nitrogen-charged glass vial containing dried copolymer and equipped with a septum. The copolymers fully dissolved in a few seconds at 23 °C.

2.2. Rheological Characterization

Interfacial rheological properties of copolymer solutions were measured using a DWR fixture on a TA Instruments DHR-III rheometer. The protocol for the interfacial rheology experiments was similar to that used by Won and co-workers^[8] for PLGA films. The DWR trough was loaded with 20 mL of water. The upper DWR fixture was lowered until it made initial contact with water; this point was determined by visual observation of the water surface. 60 μL of a 1% w/v solution of copolymer in chloroform was applied dropwise to the surface of the trough. A freshly prepared copolymer solution was used for each run. The solution was allowed to evaporate for 300 s. To ensure formation of a stable polymer interface, G′ and G″ were measured at a stress of 0.1% and a frequency of 1 rad s⁻¹ as a function of time. Samples typically took 20 min to reach equilibrium. After this time, stress sweeps and frequency sweeps were performed as described below. Before each copolymer system was tested, amplitude sweep tests were performed at a constant frequency but at variable amplitudes to ensure that the frequency sweep tests were performed within the linear viscoelastic (LVE) regime. All tests were performed at 25 °C.

2.3. Brightfield Microscopy Characterization

Brightfield microscopy of copolymer solutions on the air–water interface was performed using a Zeiss Axio Vert A.1 microscope equipped with a 20× NA0.5 air EC Plan-NeoFluar objective and Lumenera Infinity 3 monochrome CCD camera with a resolution of 1936 × 1456 and 4.54 × 4.54 μm pixels. The well was loaded with 1 mL of DI water. 50 μL of the solution was added to the surface of the copolymer films, followed by equilibrium in the dark for ≈1 day to ensure the homogeneous dispersion of the dye. Brightfield images were taken with 25 ms exposure time.

2.4. Confocal Microscopy Characterization

Confocal microscopy on the copolymer films was performed using a Zeiss Axio Examiner D1 modified with an Andor Differential Scanning Disk 2 confocal unit equipped with a 20× NA0.5 air EC Plan-NeoFluar objective and Piezo objective holder for acquiring Z-stacks. To prepare each sample film, a well was loaded with 50 μL of 1% w/v solution of copolymer in chloroform. The samples were put into vacuum drying oven for ≈1 day to form dry films. A 10⁻¹⁵ m solution of Rhodamine 110 (R110, Sigma-Aldrich) was formed by dissolving it in DI water. 50 μL of the solution was added to the surface of the copolymer films, followed by equilibrium in the dark for ≈1 day to ensure the homogeneous dispersion of the dye. Optimal z-stack step size was calculated using the Andor iQ3 software to provide Nyquist sampling. Excitation and emission filters used were: Green channel excitation 482/18, emission 525/45. Green fluorescent images were acquired with a 50 ms exposure time and analyzed using ImageJ.

2.5. SAXS Characterization

SAXS measurements were conducted on a Bruker Nanostar U in the high resolution configuration (Brookhaven National Lab, Brookhaven, NY). The wavelength of the beam was 0.15418 with Cu rotating anode source. The nominal distance from sample to detector (Vantec 2000 area detector) was 1.1 m and the actual distance was calibrated with silver behenate before the measurements. Copolymer solutions were prepared at concentrations of 1% w/v in THF. Sample was loaded and sealed into glass capillaries (diameter 1.0 mm) to avoid solvent evaporation. The capillary was fixed in the sample holder and scattering data for each sample was collected for 6 h. SAXS data analysis was performed using the SasView 4.2.2 small-angle scattering analysis software (http://www.sasview.org/).

3. Results and Discussion

3.1. Interfacial Rheology

Interfacial rheological characterization was performed at 25 °C on the air–water interface of copolymer samples prepared by spreading the copolymer on the water surface. The coexistence of hydrophobic and hydrophilic regions gives the copolymers an amphiphilic character, allowing copolymers to reside at water interfaces and to form stable viscoelastic films. Following the analysis of Wasan and co-workers, we evaluate the Boussinesq number, Bo, a dimensionless group that shows the ratio of the surface drag to the bulk:^[11]

$$Bo=\frac{\text{Surface drag}}{\text{Subphase drag}}=\frac{{\eta }_S\cdot \frac{V}{L_I}\cdot P_I}{\eta \cdot \frac{V}{L_S}\cdot A_S}=\frac{{\eta }_S}{\eta \cdot G}$$ (2)

Here, η_S is surface shear viscosity, η is the bulk viscosity, and V is a characteristic velocity. The parameters L_I, L_S, P_I, and A_S are the characteristic length scales for velocity decay at the interface and in the bulk, contact perimeter, and contact area; these are dependent upon the measurement geometry and can be combined into a single parameter G. For the DWR, G is 0.7 mm,^[12] which corresponds to the side length of the square ring. When ∣Bo∣ >> 1, interfacial stresses dominate and the surface rheological properties can be accurately obtained from the measurements, while for ∣Bo∣ ≤ 1, the subphase properties dominate the measurement.^[13] Since the smallest ∣η_S∣ we found was 0.84 Pa s m, Bo was estimated to be greater than or equal to 1.5 × 10⁶ for all samples. Therefore, it was unnecessary to make corrections for subphase flow effects, because these will have negligible effects on the data.

All polymers displayed a loss modulus G″ and storage modulus G′ that are dependent upon frequency, with G′ decaying at low frequencies (Figure 2a,b), indicating viscoelastic fluid behavior. Nearly all polymers showed a clear crossover of G′ and G″ in the measurable frequency range. P3 was selected as representative data in Figure 2c to illustrate the crossover behavior of the polymers. Data for all other polymers are provided in the Supporting Information. Most of the polymers display a crossover frequency, ω_x, in the interfacial rheology, which is defined as the frequency when G′ = G″. A characteristic relaxation time can then be calculated as τ ~ 1/ω_x. These are listed in Table 1. For the P5 polymer, no relaxation time is listed, as the G″ data for this polymer was somewhat noisy at lower frequencies (Figure S5, Supporting Information). As shown in Table 1, P1 to P3 have similar relaxation times, while H1 has a larger relaxation time than P1–P3. We expect that the relaxation time we observed for H1, P1, P2, and P3 to be some type of characteristic time for chain motion, similar to a Rouse time but modified for motion in two dimensions. As all copolymers P1–P6 are similar in molecular weight, in the absence of strong interchain interactions, we would expect this time scale to be similar for all copolymers. This is consistent with our SAXS results (Table 2), which show similar chain dimensions for copolymers P1 and P3, with a slightly larger size for H1, which would be consistent with a longer relaxation time for H1. Polymers P1 through P3 do contain very short hydrophobic segments which may lead to weak interchain interactions, but it is likely that any interchain associations driven by these very small hydrophobic segments are too short-lived to influence the dynamics.

Figure 2.

a) Interfacial storage modulus G′ and b) loss modulus G″ at variable frequencies for polymers with varying sequence and microblock lengths. c) An example of G′ and G″ for P3, showing viscoelastic fluid behavior and the crossover frequency.

Table 1.

Crossover points for each polymer. No value shown for P5 as G″ data were noisy and no crossover was observed in the experimentally accessible frequency range.

	Angular crossover frequency [rad s−1]	Angular crossover frequency [Hz]	Relaxation time [s]
H1	0.52	0.083	12
P1	1.4	0.22	4.5
P2	1.4	0.22	4.5
P3	1.4	0.22	4.5
P4	0.37	0.059	17
P5	—	—	—
P6	0.17	0.027	37
Pr	0.023	0.0037	270

Table 2.

Parameters from fitting SAXS data using a) a sphere model and b) polymer excluded volume model.

Polymer	Ra [Å]	Rgb [Å]	m b
H1	34.6 ± 0.5	32.2 ± 1.1	3.0 ± 0.1
P1	24.9 ± 0.3	22.9 ± 0.6	2.4 ± 0.5
P2	—	—	—
P3	29.0 ± 0.4	24.2 ± 1.3	1.8 ± 0.6
P6	28.6 ± 0.3	26.4 ± 1.6	1.5 ± 0.3
Pr	25.1 ± 0.2	23.8 ± 1.0	1.2 ± 0.2

By contrast, for P4 and P6, this relaxation time increases beyond that of P3, and becomes larger as the microblock length increases. SAXS results (Table 2) show a similar size for P6 as for P1 and P3, so the increase cannot be due to a difference in chain size or dimensions. We suspect that for P4 and higher, the microblocks are long enough to form reversible associations between chains. If these reversible associations have a lifetime that are longer than the time for chain motion, they will hinder chain motion, akin to the “sticky Rouse time” described by Shen et al.,^[14] and increase the relaxation time. Pr displays an even longer relaxation time, consistent with the possible formation of larger microblocks due to the difference in reactivity rate of the hydrophilic versus hydrophobic monomer.^[3]

Homopolymer H1, which displays only the hydrophobic propyl side chain, forms interfaces with the largest value of G′ and G″ in this series. This suggests that the viscosity and elasticity of the interface are due, at least in part, to the hydrophobic character of the polymer. From P1 to P6, G′ increases systematically, except for P2. This trend can be seen more clearly in Figure 3, which shows G′ and G″ values for this series of samples at a single frequency of ω = 100 rad s⁻¹. Similar trends are observed at other frequencies and are reported in the Supporting Information. Figure 4 shows the complex viscosity as a function of frequency. All samples show shear-thinning behavior, and the complex viscosity follows the same trend as G′.

Figure 3.

Interfacial storage and loss modulus, G′ and G″, for ω = 100 rad s⁻¹.

Figure 4.

Interfacial complex viscosity at variable frequencies for polymers with varying sequence and microblock lengths.

From the interfacial rheology data, we observe that H1 has the highest G′ and G″ values while P2 has the lowest G′ and G″ values. We also see that G′ increases from P1 to P6 and also Pr. This may be due to an increase in the size of hydrophobic domains in the films as the length of the hydrophobic microblock increases. Previous contact angle measurements on this series of polymers^[3] showed that the H1 thin film displays a large contact angle (e.g., forms a hydrophobic film) and that the contact angle increases systematically as the microblock size ranges from P1 to P6, with the exception of P2, which does not follow this trend.^[3] This is consistent with our results from P1 to P6, where we expect the size of the hydrophobic regions to increase. This increase will likely lead to a more strongly associated copolymer film at the water surface, as stronger intermolecular associations often correspond to a stronger elastic response and a larger value of G′ in bulk rheology experiments of amphiphilic copolymers. By contrast, G″, which represents the viscous component of the mechanical response, does not show a clear trend with the microblock size in Figure 3. Again, in correspondence with bulk rheology, we may expect the viscous response to be dependent upon the overall molecular weight of the chains.

Regarding copolymer P2, which does not follow the trend of the other polymers in terms of G′ and G″ values, a similar trend was previously observed for this series of copolymers in the contact angle measurements on dried films. This behavior was attributed to the surface energy of the P2 thin film being dominated by chain orientation rather than microphase separation.^[3] Additional characterization of the polymer structure at the interface, using techniques such as neutron or X-ray reflectivity or grazing-incidence SAXS (GISAXS) could provide additional information on this aberration.

3.2. Microscopy Studies

Brightfield microscopy was performed to determine if any visible differences could be observed in the copolymer films at the air–water interface (Figure 5). Compared to P1 and P2, the copolymer solutions of H1, P3, P4, P5, P6, and Pr appear to form continuous films at the interface, as evidenced by the puckering that can be observed visually. By contrast, P1 and P2 do not appear to form the same types of films.

Figure 5.

Brightfield images of a) H1, b) P1, c) P2, d) P3, e) P4, f) P5, g) P6, and h) Pr on the air–water interface.

Confocal microscopy was performed to see if any large-scale structures could be observed in the copolymer films. Fluorescent images are shown in Figure 6, and the corresponding brightfield images in Figure 7. Data on H1, P1, P2, P3, P5, P6, and Pr are shown; limitations on sample availability prevented us from performing experiments on P4. The light regions in Figure 6 indicate regions of the copolymer films with higher concentrations of the hydrophobic fluorescent dye. P1 and P2 did not show much large hydrophobic aggregation. From P3 to P6 to Pr, we could see spherical hydrophobic aggregates with sizes of about 1.5 μm. From P5 to P6 to Pr, the density and the size of the spherical aggregates appear to increase. We believe this is due to the presence of larger hydrophobic microblocks in P6 as compared to P5, and even longer runs of hydrophobic segments present in Pr.

Figure 6.

Confocal fluorescent images on the polymer films of a) H1, b) P1, c) P2, d) P3, e) P5, f) P6, and g) Pr. Shown are single representative images from z-stack acquisitions. All images were acquired with identical microscope exposure settings and displayed with identical brightness/contrast settings.

Figure 7.

Brightfield images of the polymer films of a) H1, b) P1, c) P2, d) P3, e) P5, (f) P6, and g) Pr corresponding to the confocal fluorescent images displayed in Figure 6.

3.3. SAXS Studies

Solution behavior of sequence-controlled copolymers P1, P2, P3, P6, random copolymer Pr, and homopolymers H1 was investigated by measuring SAXS on 1% w/v polymer in THF. Due to sample limitations, we were unable to obtain SAXS data on copolymers P4 and P5. Preliminary solution SAXS data on these polymers in THF had been reported by Li and Sampson^[3] and qualitatively showed similar solution structure for all copolymers; however, these data were taken over a limited range of q and thus were not amenable to quantitative data fitting to obtain chain dimensions. Here, we present data over a wider range of q and have fit the data to provide more information about chain size and conformation.

As discussed above, because the copolymers are too hydrophobic to dissolve directly in water, for the interfacial rheology experiments, we used the protocol developed by Won and co-workers^[8] for interfacial rheology measurements of PLGA copolymers, whereby the polymers are first dissolved in a solvent such as chloroform, which is then cast onto the air–water interface and allowed to evaporate before interfacial rheology experiments are performed. Thus, we felt that the conformation of the polymers in a solvent capable of dissolving the polymer would be appropriate to probe, as any aggregates formed in the solution state would likely persist in at the air–water interface once the solvent evaporates. SAXS experiments in chloroform were attempted; however, contrast in this solvent was very poor due to the presence of Cl in the solvent, which scatters strongly and results in a high background in SAXS experiments. Thus, THF, which also was capable of dissolving the polymer was chosen as the solvent for our SAXS studies.

To further understand the solution behavior, the polymers were fit to a sphere model (Figure 8a) and the polymer excluded volume model (Figure 8b), with the exception of P2. As can be seen from Figure 8, excess scattering was observed a low q for this copolymer, which prevented us from obtaining accurate values for the dimensions of this copolymer from the SAXS fits. This type of behavior is typically attributed to the presence of clustering or larger-scale structures in the sample. It is not clear what might be causing the low q scattering in this sample. No evidence of high molecular weight contaminants were found in this sample, and the rheology and microscopy studies are not consistent with cluster formation or aggregation; thus, we expect that this may be an artifact due to the presence of air bubbles or dust being trapped in the SAXS sample capillary.

Figure 8.

SAXS curve plotted for varied 1% w/v copolymers in THF. a) The SAXS pattern fits of the sphere model (except for P2). b) The SAXS pattern fits of the polymer excluded volume model (except for P2). Data are vertically shifted for clarity.

The scattered intensity was fit using a spherical form factor (Equation (3)):^[15]

$$I(q)=\frac{Scale}{V}{\left[3V(\Delta \rho )\frac{\sin (qr)-qrcos(qr)}{(qr{)}^3}\right]}^2+\text{Background}$$ (3)

where r is the radius of the sphere, V is the volume of the sphere given by V = 4πr³/3, and Δρ is the scattering length density difference between the scatterer and the solvent.

Figure 8a shows SAXS data for copolymers and homopolymers fit to the sphere model, with the exception of P2. The sphere model includes the radius R. The detailed parameters are summarized in Table 2. For all copolymers (except for P2, whose SAXS data could not be fit with sphere model), R was 25–30 Å. The uncertainties provided for the values of the fit parameters are based on goodness-of-fit of the model. The data themselves are quite scattered, so other methods of estimating errors in the fit parameters may lead to higher uncertainties. However, the large number of data points in a typical SAXS experiment typically provides good confidence in the values of fitting parameters, even in the case where the data themselves are scattered, assuming the model used provides a good fit to the data. Additionally, the agreement between the radius derived from the sphere model and the radius of gyration of the polymer excluded volume model, described below, provides further confidence in the values of the fit parameters.

The polymer excluded volume model describes the scattering from polymer chains with excluded volume effects. It was originally presented by Benoit^[16] and was later put into analytical form by Hammouda.^[17] This model includes two parameters: the polymer radius of gyration R_g, and a parameter m which is related to the excluded volume parameter, ν, as m = 1/ν. In dilute solution, the conformation of a polymer chain depends on the interaction between chain segments and solvent molecules. In a good solvent, a chain expands from its unperturbed dimensions to maximize the number of segment–solvent contacts, and the coil is expanded compared to a Gaussian chain, and the exponent is ν = 3/5. In a poor solvent, the chains will contract to minimize interactions with the solvent, and the exponent is ν = 1/3. However, competing with this effect is the tendency for chains to expand to reduce unfavorable segment–segment interactions, which is the excluded volume effect. If these two effects are in perfect balance, the polymer molecule will adopt unperturbed dimensions. The solvent is then said to be a theta solvent and ν = 1/2.

Figure 8b shows SAXS data for copolymers and homopolymers fit to the polymer excluded volume model. This model includes the polymer radius of gyration, R_g, and a parameter m that is related to the excluded volume parameter, ν, as m = 1/ν. The detailed parameters are also summarized in Table 2. For all copolymers (except for P2, whose SAXS data could not be fit with polymer excluded volume model), R_g was 23–26 Å. This range is similar to the sphere model, which gives confidence in our fit values. We would expect values for the parameter m of 1.67 for flexible polymer chains swollen by a good solvent, 2 for chains in a theta solvent, and close to 3 for chains that are collapsed in a poor solvent.

Both models show that the homopolymer H1 has a larger, more expanded structure. The value of the parameter m is ≈3 for H1 which indicates that the H1 polymer behaves as if in poor solvent. This result corresponds well with the fact that the H1 is hydrophobic. However, the uncertainties in the values of m are large. Similarly, for the copolymers, the fitting results suggest that the five polymers are behaving as if in theta to good solvent from P1 to Pr. However, most of the changes observed are within the uncertainty of the value of m. So, there do not appear to be significant variations in the solution structure of the copolymers. This suggests that the variations observed in the interfacial rheology are due to changes in the interfacial structure and assembly of these copolymers, rather than any type of association in solution or changes in the bulk rheology.

4. Conclusions

Interfacial rheology was performed on the air–water interfaces containing sequence-controlled copolymer samples with varying microblock size, along with the corresponding homopolymers. SAXS was also conducted on these polymers in THF, demonstrating that samples were found to have similar spherical conformations in solution. Although the monomer sequence seemed to have little impact on the solution morphology, the air–water interfacial properties were found to be strongly dependent on microblock size. As microblock size increases from 1 to 6, the storage modulus was found to increase. We expect this arises from the formation of larger hydrophobic domains at the interface. This trend is similar to previous contact angle results on dried films. An exception was observed when two carboxamides comprise the microblock; this sample displayed the lowest storage modulus. The relaxation time also increased with increasing microblock size. This work provides a new strategy for controlling the viscoelasticity and stress relaxation dynamics of fluid interfaces through the use of precision sequence-controlled polymers.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.