# Structural insights into nanofibers formed by templated polymerization-induced self-assembly through combined SANS and AFM characterizations

Gaëlle Mellot, Mathieu Chevereau, Laurent Bouteiller, François Stoffelbach, Marylene Vayer, Jutta Rieger, Christophe Sinturel

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## To cite this version:

Gaëlle Mellot, Mathieu Chevereau, Laurent Bouteiller, François Stoffelbach, Marylene Vayer, et al. Structural insights into nanofibers formed by templated polymerization-induced self-assembly through combined SANS and AFM characterizations. Journal of Molecular Liquids, 2025, 424, pp.127104. 10.1016/j.molliq.2025.127104. hal-05033481

HAL Id: hal-05033481

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Submitted on 3 Nov 2025

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# Structural Insights into Nanofibers formed by Templated Polymerization-Induced Self-assembly through Combined SANS and AFM Characterizations

Gaëlle MELLOT¹, Mathieu CHEVEREAU², Laurent BOUTEILLER¹, François STOFFELBACH¹, Marylene VAYER², Jutta RIEGER¹*, Christophe SINTUREL²*

1. Sorbonne Université, CNRS, Institut Parisien de Chimie Moléculaire, Equipe Chimie des Polymères, 4 Place Jussieu, 75005 Paris, France

2. Université d'Orléans, CNRS, Interfaces Confinement Matériaux et Nanostructures, 1 B Rue de la Férollerie, 45071 Orléans, France

Email: jutta.rieger@sorbonne-universite.fr, christophe.sinturel@univ-orleans.fr

## Abstract

This study combines Small-Angle Neutron Scattering (SANS) and Atomic Force Microscopy (AFM) to elucidate the formation and structural characteristics of poly(N,N-dimethyl acrylamide)-*block*-poly(2-methoxyethyl acrylate) (PDMAc-*b*-PMEA) block copolymer nanofibers prepared in water by the recently developed “Templated Polymerization-Induced Self-Assembly (PISA)” approach. The key feature of this technology is the use of a water-soluble macromolecular RAFT agent functionalized by a bisurea sticker (PDMAc-U₂). During the generation of the PMEA block, through reversible addition-fragmentation chain transfer (RAFT) radical dispersion polymerization, the bisurea sticker was believed to promote the 1D assembly of the block copolymers via directional H-bonding into colloidal nanofibers. In this work, Small Angle Neutron Scattering (SANS) experiments show that anisotropic objects already form at the very early stages of polymerization, then, in the course of polymerization, their diameter increases concomitantly with the increase of the PMEA chain length. Furthermore, the SANS data reveal that the nanofibers are organized in a PMEA-core/ PDMAc-shell, which was previously hypothesized but not demonstrated. Atom Force Microscopy (AFM) supported these findings through local differences in mechanical properties and detailed visualization of the nanofibers' internal structure. The combination of SANS and AFM results allowed us to corroborate the formation and internal structure of anisotropic particle obtained by the templated PISA approach. We believe that these nanostructured materials dispersed in aqueous solutions have potential for biomedical applications, while nanopatterning applications can be envisioned after deposition on surfaces.

## Keywords

PISA, supra molecular self-assembly, nanofibers, block-copolymer, AFM, SANS
## Highlights

* The internal structure of nanofibers formed via templated PISA is elucidated

* SANS and AFM data reveal a core/shell structure

* On surfaces, individual nanofibers flatten while maintaining an anisotropic shape

* In thin films, nanofibers self-assemble on surface in a regular pattern of lines

* Application of such nanofibers in surface nanopatterning can be envisioned

## Graphical abstract

<figure><img src="image_3.png" /></figure>
# 1. Introduction

Amphiphilic block copolymers are block copolymers that comprise polymer blocks of distinct polarity. The existence of at least one hydrophilic (polar) and one hydrophobic (nonpolar) block promotes their assembly into micellar aggregates in a selective solvent of one of the blocks. For application in the biomedical field, water is the most common solvent used for their preparation. In theory, the aggregates are then structured into a hydrophobic core surrounded by a hydrophilic shell. Depending on the respective block lengths, different morphologies, such as spheres, worms or vesicles, can be formed. This has been rationalized by the packing parameter¹ *P*, which takes into account the polymer block lengths, solvent and polymer interaction parameters, and which allows to predict the geometrical organisation of the chains in the assembly and thereby its morphology. In most cases, amphiphilic block copolymers assemblies cannot be simply obtained by direct dissolution in water. Different methods have therefore been devised to prepare them [1]; most of them require the use of an organic cosolvent [2], which has to be removed afterwards. Such extra steps are time-consuming and lead generally to low concentrated micellar solutions. In addition, for specific applications remaining traces of organic solvent can be detrimental.

In more recent years, polymerization-induced self-assembly (PISA) [3-6] has occurred as a straightforward alternative method to produce amphiphilic block copolymer assemblies in large amount, in a minimum of steps. In contrast to conventional methods, PISA combines the synthesis of block copolymers (in heterogeneous conditions) with their self-assembly in one single step. The number of synthesis steps is thereby reduced, and highly concentrated solutions (up to 30-40 wt%) can be recovered. In brief, in the presence of a first reactive solvophilic polymer block a second monomer - leading to an insoluble polymer - is polymerized in heterogeneous polymerization conditions. As the second block grows from the first block, amphiphilic diblock copolymers form that self-assemble *in-situ* into self-stabilized particles. Compared to conventional methods, the same diversity of morphologies can be prepared. TEM and SAXS analyses showed that higher order morphologies (such as worms/fibers or vesicles) are obtained from spherical precursor particles that are formed at the earliest stages of the polymerization [7-9]. This means that morphological transitions occur concomitantly to the growth of the length of the solvophobic block, as could be expected from the evolution of the packing parameter *P*. While experimental phase diagrams have been reported for a large variety of systems, their prediction remains difficult as out-of-equilibrium structures are often formed. In some cases, particularly in water, only spherical particles are observed despite systematic increase of the hydrophobic block [10,11]. Large particle diameters [12] and internal structuration [13] observed by cryoTEM
analyses suggest that the structure of such PISA particles can be complex and differ from the ideal core-shell structures with a complete segregation of the solvophilic and solvophobic block.

For various applications, anisotropic worm-like or fiber-like morphologies are of particular interest. Such morphologies can only be obtained in a very narrow window of polymerization conditions, e.g. for a distinct monomer concentration [14] at which PISA is performed and for precise lengths of the hydrophilic and the hydrophobic blocks [4,6,15]. This in turn means that the diameter of these assemblies, which is related to the block length, cannot be tuned because the morphology changes with the hydrophobic block length. Recently, we have therefore devised a novel PISA strategy termed “*templated PISA*”, which allowed us to synthesize block copolymer nanofibers in a large experimental window [16]. CryoTEM analyses suggested that long nanofibers form at early stages of the polymerization (i.e. for short hydrophobic blocks), and that during polymerization, the nanofiber diameter increases, rather than morphological transitions occurred. This *templated PISA* approach relies on the combination of conventional reversible addition-fragmentation chain transfer (RAFT) - PISA with directional supramolecular self-assembly. The key is the use of a macromolecular RAFT agent (macroRAFT agent) functionalized by a bisurea (U₂) sticker, which promotes via directional H-bonding the 1D assembly of the chains into nanofibers, and effectively locks the morphology during polymerization. CryoTEM analyses suggested that long nanofibers form during polymerization [16], although the macroRAFT agent is initially not fully organized in cylindrical micelles [17]. Systematic variation of the experimental conditions demonstrated the robustness of the *templated PISA* approach and revealed that poly(N,N-dimethylacrylamide)-*block*-poly(2-methoxyethyl acrylate) (PDMAc-*b*-PMEA) nanofibers can be produced at solids contents ranging from 7.5 to 28 wt% for number-average degrees of polymerization (DP<sub>ns</sub>) of the hydrophobic PMEA block varying from around 40 to more than 200. While cryoTEM analyses allowed us to identify the assemblies’ morphologies - letting us postulate a formation mechanism - this technique did not provide information on the internal organization of the polymer chains, nor the presence or not of a core-shell structure, as ideally expected in *templated PISA*.

In the present article, we studied the formation of PDMAc-*b*-PMEA nanofibers via *templated PISA* by Small-Angle Neutron Scattering (SANS) and combined this technique with Atomic Force Microscopy (AFM) to get a better understanding on the formation and the internal organisation of the nano-objects. SANS and AFM can reveal internal structuration in case the contrast in scattering length densities or mechanical properties is sufficient. In addition, SANS analyses provide average values on the entire sample in solution and do not require any sample preparation that could alter the sample.
## 2. Experimental section

### 2.1 Materials

2,2'-Azobis(2-amidinopropane) dihydrochloride (V-50, ≥ 97%, Aldrich) and ethylene glycol diacrylate (EGDA, 90%, Aldrich) were used as received. 2-Methoxyethyl acrylate (MEA, ≥ 98%, Aldrich) was distilled under vacuum prior to polymerization. Deuterated (D₂O, Erisotop) or deionized water was used for the polymerizations. All syntheses were conducted under argon atmosphere. The macroRAFT agents PDMAc₃₇-U₂ (*M*<sub>n,NMR</sub> = 4.3 kg mol⁻¹, *DP*<sub>n,NMR</sub> = 37, *M*<sub>n,SEC</sub> = 3.9 kg mol⁻¹, *D* = 1.17) and PDMAc₃₅-U₂ (*M*<sub>n,NMR</sub> = 4.1 kg mol⁻¹, *DP*<sub>n,NMR</sub> = 35, *M*<sub>n,SEC</sub> = 3.6 kg mol⁻¹, *D* = 1.27) were prepared as reported before [17,18].

### 2.2 Syntheses

In a typical experiment (PDMAc₃₇-*b*-PMEA₂₃₃-U₂, Table S1), 0.086 g of PDMAc₃₇-U₂ (*M*<sub>n</sub> = 4.3 kg mol⁻¹, 19 µmol) was placed in a 10 mL round bottomed flask. Then 1,3,5-trioxane (0.033 g, 0.18 mmol) was added. A solution of V-50 (6.3 mg, 0.023 mmol) in 3.88 g of deuterated water was prepared. 1.03 g of this solution (6 µmol of V-50) and 4.71 g of D₂O were added to the flask containing the macroRAFT agent and trioxane used as internal standard. After deoxygenation by bubbling with nitrogen for 20 min in an ice bath, distilled and deoxygenated MEA (0.64 mL, 4.98 mmol) was added. The septum-sealed flask containing the reaction mixture was heated at 60 °C in a thermostated oil bath and stirred at 500 rpm. Samples were periodically withdrawn to monitor the conversion by ¹H NMR in acetone-d₆. The polymerization was quenched at 90 min by immersion of the flask in iced water. The polymers were characterized by SEC in DMF (+ LiBr, 1 g L⁻¹).

Crosslinked nanofibers (PDMAc₃₅-*b*-PMEA₉₉-U₂-CL, Table S1) were produced following essentially the same procedure in deionized water. PDMAc₃₅-U₂ was used as macroRAFT agent and molar ratios identical to PDMAc₃₇-*b*-PMEA₉₇-U₂ were used ([MEA]₀/[CTA]₀ = 100, see Table S1). 1.5 mol% of EGDA relative to MEA were added after 1h45 of polymerization to crosslink the nanofibers while the polymerization was stopped at 4h45. In contrast to PDMAc₃₇-*b*-PMEA₉₇-U₂, after dissolution in DMF (+ LiBr, 1 g L⁻¹), the sample could not be filtered through a 0.2 µm membrane, as expected for crosslinked nanofibers. In addition, DLS analyses in DMF at 1 g L⁻¹ were also performed.
## 2.3 Characterization techniques

### NMR spectroscopy

¹H NMR spectra were recorded at 298 K, with a Bruker 400 MHz spectrometer in 5 mm diameter tubes. Monomer conversions were determined by ¹H NMR, spectroscopy in CDCl₃, DMSO-d₆ or acetone-d₆ by the relative integration of the internal reference (1,3,5-trioxane) peak at 5.1 ppm and the vinylic proton peaks of the monomers.

### Size Exclusion Chromatography (SEC)

The number-average molar mass ($M_n$), the weight-average molar mass ($M_w$) and the dispersity ($D = M_w/M_n$) were determined by size exclusion chromatography (SEC) in DMF (+ LiBr, 1 g L⁻¹) at 60 °C and at a flow rate of 0.8 mL min⁻¹, using two PSS GRAM 1000 Å columns (8 × 300 mm; separation limits: 1 to 1000 kg mol⁻¹) and one PSS GRAM 30 Å (8 × 300 mm; separation limits: 0.1 to 10 kg mol⁻¹) coupled to a differential refractive index (RI) detector and a UV detector. Toluene was used as flow rate marker. The OmniSEC 4.7 software was used for data acquisition and data analysis. Molar masses ($M_n$, and $M_w$) and dispersity ($D$) were calculated with a calibration curve based on PMMA standards, using the RI detector. The sample concentration was 5 mg mL⁻¹ in DMF (+ LiBr, 1 g L⁻¹) and all samples were filtered through 0.2 µm membrane before injection.

### Small-Angle Neutron Scattering (SANS)

The measurements were recorded at 20 °C at the LLB (Saclay, France) on the PA20 SANS instrument, at three sample to detector distances 1.1, 8 and 18 m and fixed wavelength equal to 6 Å to cover the 1.7 × 10⁻³ to 0.29 Å⁻¹ q-range, where the scattering vector $q$ is defined as usual, assuming elastic scattering ($q = (4π/λ)sin(θ/2)$, where $θ$ is the angle between incident and scattered beam). The SANS experiments were performed using hydrogenated solute in a deuterated solvent. This procedure ensures both a high contrast (between H and D) and a low incoherent scattering due to H thanks to the low concentration of the solute (less than 0.05 cm⁻¹ in all experiments). Data were corrected for the empty cell signal, the solute and solvent incoherent background. A light water standard was used to correct from the detector inhomogeneities and to normalize the scattered intensities to cm⁻¹ units. The data was fitted with the DANSE software SasView [19]. Details of the fits of the macroRAFT agent and the assembled block copolymers are given in SI and Figures S2-S7.

### Cryogenic Transmission Electron Microscopy (cryoTEM)

For cryoTEM analyses, the samples were diluted in water to 1 wt% before analysis. According to protocols reported elsewhere [20,21] thin liquid films of particle dispersions were prepared at room temperature by depositing 4 µL of the diluted sample and successive blotting. They were
flash frozen in liquid ethane and observed at -180 °C on a JEOL JEM-2100 LaB₆ microscope operating at 200 kV under low-dose conditions (10 electrons Å⁻² s⁻¹). Digital images were recorded on a Gatan Ultrascan 1000 CCD camera.

## Dynamic Light Scattering (DLS)

The z-average particle diameter (Dz) of the crosslinked nanofibers was measured by dynamic light scattering (DLS) at 25 °C, with a Zetasizer Nano S90 from Malvern (90° angle, 5 mW He-Ne laser at 633 nm). The analyses were performed in DMF at 0.1 wt% and treated by the cumulant method (see Figure S1).

## Atomic Force Microscopy (AFM)

For AFM imaging, all samples were dialyzed against water and diluted to a concentration of 1 g.L⁻¹ prior to deposition onto 10 × 10 mm² Si (100) substrates (generously supplied by STMicroelectronics Co). The substrates were pre-cleaned by sonication in tetrahydrofuran (THF), ethanol, and acetone for 10 minutes each. For individual nanofiber characterization, the solution was applied by spin-coating at 3000 rpm for 30 seconds. To form nanofiber thin films (assemblies), a drop of the solution was deposited onto the substrate and allowed to dry completely in a petri dish by natural evaporation.

AFM imaging was performed in air at room temperature in peakForce QNM mode on an AFM Dimension Icon with ScanAsyst from Bruker, USA with ScanAsyst-air silicon-nitride cantilevers (resonance frequency 70 kHz, spring constant: 0.4 N m⁻¹) with a nominal tip radius of 2 nm. All the AFM images were 512*512 pixels.

# 3. Results

## 3.1 Synthesis and monitoring of nanofiber formation by SANS

To get more information about the mechanism of the nanofiber formation by *templated PISA*, we selected small angle neutron scattering (SANS). In contrast to cryoTEM experiments that only deliver information on the overall nanofiber diameter of imaged, selected parts of the sample, SANS provides values averaged on the entire sample directly in solution. In addition, it does not require any sample preparation that may alter the assembly. For core-shell structures, it can provide the dimensions of the shell and core, in case their solvation by D₂O is sufficiently different providing contrast for fitting separately the shell and the core (that would not have been possible using SAXS). A potential drawback of this technique is restricted access to SANS facilities and the need for deuterated solvents to provide contrast. In this study, we therefore performed the polymerization of MEA (a water soluble monomer) in deuterated water (D₂O) (synthesis PDMAc₃₇-
$b$-PMEA$_{233}$-U$_{2}$, Table S1) using the same bisurea end-capped trithiocarbonate (TTC) macroRAFT agent, PDMAc-U$_{2}$, as in our previous study performed in water². Samples were withdrawn at regular time intervals to determine the monomer conversion by ¹H-NMR and characterize the resulting PDMAc-$b$-PMEA-U$_{2}$ block copolymers by size exclusion chromatography (SEC) and the nano-objects by SANS.

Monitoring by ¹H-NMR and SEC (Figures 1a and b), showed that - during polymerization - the number-average molar mass, $M_n$, increased linearly with monomer conversion and dispersities were low at all conversions (Table S1). These results are very similar to the polymerizations performed in water [16], revealing that the change from water to D$_{2}$O does not have any significant impact on the polymerization kinetics and control, as expected.

For SANS analyses, the samples taken at regular time intervals were diluted with D$_{2}$O to 1 wt% before analysis. The evolution of the scattering profile recorded at 20 °C is displayed in Figure 1c. The figure also displays in light blue the SANS data of the pristine macro-RAFT agent, PDMAc-U$_{2}$, used for the polymerization in D$_{2}$O. Consistent with our previous study on PDMAc-U$_{2}$ of similar molar mass³, it could be fitted using a core-shell model with a core radius of 2.6 nm, matching quite well the theoretically predicted dimensions of the sticker region (as schematized in Figure S8) and a shell thickness of 2.7 nm (Figure S2 and Table S2). While the macroRAFT agent is assembled into mainly short cylinders (7 nm in lengths in average), upon polymerization, a significant increase in the scattered intensity as well as a clear $q^{-1}$ dependency over a large $q$ range is observed at low $q$ (Figure 1c). This means that long nanofibers form quickly at the earliest stages of polymerization (i.e. for short PMEA block length, with a number-average degree of polymerization, $DP_n$, of 30, PDMAc$_{37}$-b-PMEA$_{30}$-U$_{2}$).

Increasing the $DP_n$ of PMEA from 30 to 233, the $q^{-1}$ dependence of the scattered intensity persists, but the decline in scattered intensity observed at intermediate $q$ (10$^{-2}$ Å) shifts towards lower $q$ values suggesting an increase of the cylinder diameter (see coloured arrow in Figure 1c). To estimate the diameter of the nanofibers, we performed data fits using again the form factor of a core-shell cylinder with a circular cross-section, which fitted much better the data than a homogenous cylinder model (see the SI for more details and fits in Figures S3-7). The radius of the core, the shell thickness and the scattering length density of the shell were adjusted for each curve. As expected, the core dimensions and the total diameter of the nanofibers (2x shell thickness + 2x core radius) increased with increasing length of the PMEA block from 21 nm ($DP_n$ = 30) to 38 nm ($DP_n$ = 233) (Table 1). A similar shell thickness was calculated for all block copolymer assemblies (4.4 to 5.3 nm). However, it was almost twice as large as the shell of the macroRAFT
agent (2.7 nm). We believe that for the PDMAc-b-PMEA block copolymer assemblies, these “shell” values do not strictly relate to the PDMAc block length (which is the same in all cases), but rather correspond to a difference in hydration between a more hydrated shell and a less hydrated core. It is also reasonable to think that no abrupt change in hydration between these domains exists in these assemblies. Indeed, although PMEA block is not water-soluble, it can be hydrated by water [22-24] and it is possible that its hydration is increased by the vicinal PDMAc block. Such interfacial hydration of the hydrophobic block has already been proposed in the literature for amphiphilic block copolymer assemblies prepared by PISA [25]. Overall, these results are in good agreement with former cryoTEM analyses performed in water, which suggested that the nanofiber diameter increases with the *DP*<sub>n</sub> of the PMEA block [16]. This initial cryoTEM image analyses provided thus conclusions that are valid for the sample as a whole. Selected samples (*DP*<sub>n</sub> = 30 and 95) were also analysed by cryoTEM experiments (Figure S9), which match relatively well the diameters estimated by SANS.

<figure><img src="image_4.png" /><figcaption>b</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>C</figcaption></figure>

<caption>d</caption>
<figure><img src="image_6.png" /><figcaption>**Figure 1.** Monitoring of the polymerization of PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub> (see Table S1 for synthesis conditions) with the simultaneous formation of nanofibers in D<sub>2</sub>O by a) <sup>1</sup>H-NMR, b) SEC in DMF + LiBr and c) by SANS (each curve is also plotted individually in the SI to better show the fine structure), d) Evolution of the nanofiber dimensions (SANS data fitted by a core-shell cylinder model).</figcaption></figure>

<table><caption>**Table 1.** Dimensions of assemblies of the PDMAc<sub>37</sub>-U<sub>2</sub> macroRAFT agent and PDMAc<sub>37</sub>-b-PMEA<sub>y</sub>-U<sub>2</sub> block copolymers cylindrical micelles in D<sub>2</sub>O, determined by fitting the SANS data.</caption><thead><tr><th>Expt.<sup>a</sup></th><th>DP<sub>n</sub>(PMEA)<sup>b</sup></th><th>Core radius (nm)</th><th>Shell thickness (nm)</th><th>Diameter of the cylinders (nm)<sup>c</sup></th></tr></thead><tbody><tr><td rowspan="2">PDMAc<sub>37</sub>-U<sub>2</sub></td><td>0</td><td>2.6</td><td>2.7</td><td>11</td></tr><tr><td>30</td><td>5.0</td><td>5.3</td><td>21</td></tr><tr><td rowspan="3">PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub></td><td>60</td><td>6.7</td><td>4.8</td><td>23</td></tr><tr><td>95</td><td>8.0</td><td>5.0</td><td>26</td></tr><tr><td>233</td><td>15</td><td>4.4</td><td>38</td></tr><tr><td>PDMAc<sub>37</sub>-b-PMEA<sub>97</sub>-U<sub>2</sub></td><td>97</td><td>9.2</td><td>4.3</td><td>27</td></tr></tbody></table>

A second polymerization in D<sub>2</sub>O was then performed using a 2.3 greater amount of PDMAc<sub>37</sub>-U<sub>2</sub> to produce nanofibers with a shorter PMEA block at full monomer conversion, i.e. presumably a smaller diameter than PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub>. A lower number-average degree of 100 was thereby targeted. Assuming that the self-assembly is strictly one-dimensional (1D), with one chain (unimer) per section, this should lead to nanofibers with a smaller diameter. If the overall diameter is controlled by the length of the PMEA block, it should be similar to that of the intermediate sample withdrawn at 38% monomer conversion (corresponding to DP<sub>n</sub>(PMEA) = 95)

<footnote><sup>a</sup> For synthesis conditions refer to Table S1</footnote>

<footnote><sup>b</sup> Number average degree of polymerization of the PMEA block. For PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub> samples were withdrawn during polymerization (see Table S1).</footnote>

<footnote><sup>c</sup> The final calculation of the diameter is rounded to values in nm to be in line with the resolution of the measurement by AFM.</footnote>
during the synthesis of PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub>. Again, high monomer conversion (96%) was reached, and PDMAc<sub>37</sub>-b-PMEA<sub>97</sub>-U<sub>2</sub> obtained within 2.5 h. Fitting of the SANS data revealed an overall diameter of 27 nm, which is very close to the diameter of the sample PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub> at 38% monomer conversion (DP<sub>n</sub>(PMEA) = 95; diameter = 26 nm). This reproducible correlation between DP<sub>n</sub>(PMEA) and diameter obtained via two different synthesis pathways corroborates our assumption [16] that the diameter of the nanofibers can be controlled via the DP<sub>n</sub> of the PMEA block because of a 1D organization of the polymer chains through directional hydrogen bonding with presumably one macromolecule per section.

## 3.2 Characterization of nanofibers by AFM

To get a deeper understanding of the internal organization of the nanofibers, they were analysed by AFM after deposition on a silicon wafer. Using the combination of different modes, AFM cannot only provide information on the morphology and characteristic dimensions, but it can also give information on internal structuration through differences in DMT modulus and adhesion. The ability to extract meaningful insights from this combination of AFM modes depends on the specific responses of each polymer to the individual modes. Considering the structural and compositional differences between the core (short segments tightly bound through supramolecular assembly via urea functions) and the arms (flexible, coiled polymer chains with amphiphilic properties), we hypothesize that this type of object could be ideally characterized using this approach.

After dialysis against water and dilution to 1g.L<sup>-1</sup>, the samples were spin-coated onto a silicon wafer. Due to the low concentration of the solution and the deposition method, that forms a thin film of solution during the process, individual nanofibers could be recovered on the surface after drying. The nanofibers were characterized by AFM in Peak Force QNM mode, allowing for characterisation of their dimensions (height, length and width) as well as their nanomechanical properties (DMT modulus, adhesion, dissipation, and deformation). Figure 2 displays AFM images (height, DMT modulus and adhesion) of a representative single nanofiber of the sample PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub> (final sample), lying on the bare flat silicon substrate.
<figure><img src="image_7.png" /><figcaption>**Figure 2.** AFM height (a), DMT modulus (b), and adhesion (c) images of a single PDMAc₃₇-b-PMEA₂₃₃-U₂ nanofiber. The experimental nanofiber total width obtained from the measurements is scaled by the white rectangle. In (b) the red rectangle width scales the calculated sticker length (4.6 nm). In (c) the blue rectangle width scales the less adherent part of the nanofiber (20 nm – experimental measurement).</figcaption></figure>
From topographical AFM images (Figure 2a), the height and width of the nanofibers were extracted from the analysis of multiple line profiles (25 per image) taken perpendicularly to the long axis of the nanofibers (see Figure S10 for a typical example). All nanofibers display the similar flat shape with a width 51±3 nm and height 6.2±0.4 nm. Table 2 gives the dimensions extracted from these AFM measurements, the diameter determined by SANS in D₂O solution and the calculated width corresponding to a fully extended chain (all-trans planar conformation). From SANS measurements, we know that the nanofiber section is near-circular in solution with a diameter of 38 nm. The great difference in height and width measured by AFM suggests that upon drying, the nanofibers deform and spread on the surface as expected for a soft material and consistent with the low Tg of PMEA [23] (-35°C). While the nanofiber shrinks in the vertical direction due a collapse of the arms during the loss of solvation water, it stretches in the horizontal direction due to favourable interactions with the substrate. The experimental width is however smaller than the maximal theoretical width (141 nm) calculated for a fully extended conformation (see Figure S8 for calculation details) indicating that the arms keep a coiled structure around the central core. The illustration in Figure 3 proposes the structure of the nanofibers in solution (Figure 3a) and after deposition on the substrate that would result from a simple flattening process, while maintaining the core-shell structure (Figure 3b).

<figure><img src="image_8.png" /><figcaption>**Figure 3.** Proposed mechanism of nanofiber deposition: a) in water b) c) after deposition on silicon wafer. Experimental dimensions extracted from SANS (a) and AFM (b,c) measurements. PMMAc in blue, PMEA in red, sticker in yellow.</figcaption></figure>

On the DMT modulus image in Figure 2b, a contrast in the mechanical properties can be observed along the cross-sectional direction of the nanofibers. The nanofiber displays a bright (i.e. stiffer) central core and a darker (i.e. softer) shell. This contrast in modulus can clearly be observed on a cross-sectional line profile (Figure S11), which displays a marked increase in the modulus value in the centre on the nanofiber. Independently of the position on the nanofiber, the modulus is always lower on the nanofiber than on the bare silicon, which is consistent with the polymeric nature of the nanofiber [26]. The stiffer core zone in the middle of the nanofiber backbone can be related to organized domains formed by the 1D supramolecular assembly of the sticker. Due to
directional H-bonding between the urea functions, the bisurea stickers are expected to be
ordered along the nanofiber axis as schematized in Figure 4.
<table><caption>**Table 2.** Dimensions extracted from AFM measurements (¹), determined by SANS (²) and calculated considering a fully extended conformation (³).</caption><thead><tr><th rowspan="2">Name</th><th colspan="2">Topography</th><th colspan="2">DMT Modulus</th><th colspan="2">Adhesion</th><th rowspan="2">Diameter in D<sub>2</sub>O<sup>2</sup> (SANS)</th><th rowspan="2">Calc dimension (fully stretched)<sup>3</sup></th></tr><tr><th>DP<sub>n</sub><br>PDMAc</th><th>DP<sub>n,NMR</sub><br>PMEA*</th><th>Width<sup>1</sup></th><th>Height<sup>1</sup></th><th>Width<sup>1</sup></th><th>Width<sup>1</sup></th></tr></thead><tbody><tr><td>PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub></td><td>37</td><td>233</td><td>51±3 nm</td><td>6.2±0.4 nm</td><td>53±3 nm</td><td>51±3 nm</td><td>38 nm</td><td>141 nm</td></tr><tr><td>PDMAc<sub>37</sub>-b-PMEA<sub>97</sub>-U<sub>2</sub></td><td>37</td><td>97</td><td>35±2 nm</td><td>4.2±0.8 nm</td><td>37±3 nm</td><td>35±2 nm</td><td>27 nm</td><td>73 nm</td></tr><tr><td>PDMAc<sub>35</sub>-b-PMEA<sub>99</sub>-U<sub>2</sub>CL</td><td>35</td><td>99*</td><td>33±4 nm</td><td>5.8±0.6 nm</td><td>34±3 nm</td><td>33±3 nm</td><td></td><td>73 nm</td></tr></tbody></table>

<footnote>*theoretical DP<sub>n</sub> calculated by the initial [MEA]<sub>0</sub>/[macroRAFT]<sub>0</sub> ratio and the monomer conversion at the end of the polymerization without considering the crosslinker</footnote>
<figure><img src="image_9.png" /><figcaption>**Figure 4.** Schematic representation of the 1D assembly of the PDMAc-*b*-PMEA-U~2~ block copolymers in water driven by H-bonding.</figcaption></figure>

It is therefore reasonable to postulate that this backbone exhibits a higher mechanical stiffness than the surrounding amorphous and soft PMEA [23], and will thus influence the mechanical response of the nanofiber, even if this rigid core is buried inside the nanofibers (as depicted in Figure 3b). From the cross-sectional line profile, a total width (core and shell) of 53±3 nm is measured, consistent with the result in height mode. However, it is difficult to estimate the width of the stiffer central part since the DMT modulus does not show any plateau in this region (see the line profile in Figure S11). Instead of showing a clear distinction between a softer (shell) and a stiffer (core) domain, a gradual variation in DMT modulus is observed. This phenomenon can be related to the convolution of the mechanical properties of each domain, also modulated by the thickness variation along this direction. Consequently, instead of extracting dimensions from the experimental data (image or profile), we have positioned marks on the AFM image (Figure 2b) using the width calculated for the stickers in Figure S8. The distance between two urea groups (blue zone) is calculated as 1.4 nm, whereas 4.6 nm was calculated when the lateral moieties of the sticker (i.e. the alkyl chain and the TTC function) are included. As seen in Figure 2b, marks having a width of 4.6 nm fit well with the image, indicating that AFM in the DMT modulus mode can identify the position of the bisurea rigid backbone in the horizontal direction.
On the adhesion image, the nanofibers can be distinguished from the silicon surface due to
difference in adhesion properties, because polymers usually exhibit a higher level of adhesion
compared to silicon [27]. From the cross-sectional line profile (see Figure S12), a total width (core
and shell) of 51±3 nm is measured, consistent with the result in height and DMT modulus mode.
Very interestingly, “core” and “shell” regions could be also observed in adhesion mode (the core
exhibiting lower adhesion) indicating that the external surface of the deposited nanofibers is
heterogeneous (whereas it is homogenously composed of a PDMAc shell in solution). This result
suggests that the nanofiber structure is probably more complex than the simple picture of a
flattened core-shell structure (cf. Figure 3b) but involves additional reorganization of the
PMEA/PDMAc arms around the sticker. Particularly, it can be anticipated that due to the
hydrophilic nature of the PDMAc extremities, the arms extremities will migrate towards the
surface, leaving a PMEA surface (Figure 3c) on the top of the nanofibers. Compared to the DMT
modulus image (Figure 2b), the core region is much larger (~20 nm) confirming that it cannot be
attributed to the urea backbone alone. In this case, the core-shell structure observed in adhesion
mode might rather be attributed to the difference in adhesion between the PMEA core and the
PDMAc shell. In view of the progressive change in adhesion over the nanofiber width, it is however
not possible to extract distinctive domains and correlate them to the degrees of polymerization of
the outer PDMAc (DPn = 37) and inner PMEA (DPn = 233) blocks.

The influence of the degree of polymerization and crosslinking of the PMEA arms on the nanofibers
dimensions was also investigated by AFM. Figure 5a displays the AFM DMT modulus images for
PDMAc<sub>37</sub>-*b*-PMEA<sub>97</sub>-U<sub>2</sub> (height and adhesion images are shown in Figure S13-14). Compared to
PDMAc<sub>37</sub>-*b*-PMEA<sub>233</sub>-U<sub>2</sub>, decreasing the size of the PMEA block reduces the size of the nanofibers
(width reduced from 51±3 nm to 35±2 nm and height reduced from 6.2±0.4 nm to 4.2±0.8 nm).
Again, the width is laying between the diameter of the nanofiber determined by SANS in solution
(27 nm) and the dimension of the fully extended conformation (73 nm). A contrast in the
mechanical properties is again observed along the cross-section of the nanofibers, with a brighter
(stiffer) central zone and a darker (softer) outer zone (Figure S15). Although the dimensions of the
stiff core cannot be extracted with good accuracy from the profile image, the width of the central
bright domains fits again quite well the calculated width of the bis-urea sticker (4.6 nm, scaled by
the red rectangle on Figure 5a).

Furthermore, nanofibers with a crosslinked PMEA block were synthesized using the same
conditions as for PDMAc<sub>37</sub>-*b*-PMEA<sub>97</sub>-U<sub>2</sub> except that 1.5 mol% of difunctional ethylene glycol
diacrylate (EGDA, with respect to MEA) was added as a crosslinker to the polymerization medium.
After drying, the solid was dissolved in DMF which is a good solvent of PDMAc-*b*-PMEA-U<sub>2</sub>
unimers. Dynamic light scattering (Figure S1) analysis at 0.1 wt% in DMF revealed the presence of light scattering nano-objects which are too large to correspond to unimers confirming the successful crosslinking of the nanoobjects. The AFM images of crosslinked PDMAc<sub>35</sub>-b-PMEA<sub>99</sub>-U<sub>2</sub>-CL and non-crosslinked PDMAc<sub>37</sub>-b-PMEA<sub>97</sub>-U<sub>2</sub> were compared (Figure S13 for height, S14 for adhesion, Figure 5 for DMT modulus, Figure S15 and S16 for typical DMT modulus line profiles for these samples). A significant increase in height (5.8±0.6 nm instead of 4.2±0.8 nm) and a slight decrease in width (33±4 nm instead of 35±2 nm) was systematically observed (Table 2). This indicates a reduced collapse of the nanofibers on the surface, consistent with the crosslinking of the PMEA block. The modulus image of the crosslinked sample (Figure 5b) reveals a hard central zone which is larger than the initial 4.6 nm estimated for the non-crosslinked sample (Figure 5a) (scaled by the red rectangle). The presence of a larger stiff central zone can be related to the increase of the T<sub>g</sub> of the PMEA blocks upon crosslinking. In this case, the stiff zone (which width was estimated to ca. 10 nm) includes the sticker and the crosslinked PMEA core, in contrast to the non-crosslinked nanofibers (PDMAc<sub>37</sub>-b-PMEA<sub>97</sub>-U<sub>2</sub> and PDMAc<sub>37</sub>-b-PMEA<sub>233</sub>-U<sub>2</sub>) for which this zone only corresponds to the stiff backbone made by the sticker (4.6 nm).

<figure><img src="image_10.png" /><figcaption>b</figcaption></figure>

<figure><img src="image_11.png" /></figure>

**Figure 5.** AFM DMT modulus images of single PDMAc<sub>37</sub>-b-PMEA<sub>97</sub>-U<sub>2</sub> (a) and PDMAc<sub>35</sub>-b-PMEA<sub>99</sub>-U<sub>2</sub>-CL (b) nanofibers (see Figure S13 and S14 for height and adhesion). The red rectangles width scales the calculated sticker length (4.6 nm), fitting quite well the core width in (a) but not in (b), indicating an increased core width upon crosslinking.

**3.3 Characterization of thin films of nanofibers by AFM**
Assemblies of nanofibers were also characterized after slow evaporation of a drop of the nanofibers colloidal suspension of samples PDMAc₃₇-b-PMEA97-U₂ and PDMAc₃₅-b-PMEA99-U₂-CL (1g.L⁻¹). After complete drying, a ca 500 nm thick film was obtained, resulting from the superposition of multiple layers of nanofibers. The AFM of the surface reveals a 2D pattern of lines for both samples (Figure 6). Bright lines are observed in height and DMT modulus for both samples (see also Figure S17) indicating that the higher domains are also stiffer. In addition, in both samples, the distance between the bright lines corresponds to approximately 20 nm (average distance taken between the maxima on multiple line profiles as shown on Figure S17). These dimensions and the results obtained for the individual nanofibers let us assume that the brighter lines correspond to the position of the stiff backbone made of the stickers, whereas the darker lines correspond to the PMEA-b-PDMAc domains. Considering a total width of ca 30 nm for individually deposited nanofibers (see Table 2), the observed pattern (20 nm distance) suggests that the nanofibers densely pack, with low stretching and/or interdigitation of the arms. The modulus AFM images (Figure S17d) show a higher level of stiffness for the crosslinked sample than for the non-crosslinked sample (Figure S17c), in good agreement with the results obtained for the individual nanofibers reported above.

<figure><img src="image_12.png" /><figcaption>**Figure 6.** AFM height images of the surface of the dry film left after drop casting solution (1g.L⁻¹) of PDMAc₃₇-b-PMEA97-U₂ (a) and PDMAc₃₅-b-PMEA99-U₂-CL (b)</figcaption></figure>
4. Conclusions

This integrated SANS and AFM study enabled us to confirm the structural properties of PDMAc-*b*-PMEA nanofibers formed via *templated PISA*, which were proposed in previous work¹⁶ but not fully demonstrated at that time. Compared to the earlier study, we were able to demonstrate in this work that the internal structure of the nanofibers consists of a PMEA core and a PDMAc shell, as confirmed by both SANS and AFM investigations. Furthermore, this study offers unprecedented insights into the morphology of the nanofibers in their dried state, whereas the previous research was limited to their characterization in solution.

SANS analysis revealed the early formation of anisotropic objects in significant quantities, with a gradual increase in diameter as polymerization progresses. Unlike cryoTEM, SANS provides a comprehensive view of the entire sample. The data also support the presence of a core-shell structure in the polymer particles, consistent with our earlier hypothesis, though not previously proven. However, it is important to highlight that core-shell structuring is not a universal feature of all PISA-formed assemblies, as outlined in the introduction.

AFM analysis further corroborated the core-shell structure of individual particles through observed differences in mechanical modulus and adhesion properties. This nanoscale characterization provides direct and detailed visualization of nanofiber structure after deposition, revealing features such as width and internal organization. Despite noticeable flattening in the vertical dimension and stretching in the horizontal, the anisotropic structure remains intact, suggesting potential for surface templating. This was further confirmed in thin-film configurations, where the nanofibers exhibited dense and regular alignment.

These findings enhance our understanding of the formation and structural evolution of anisotropic particles during polymerization, providing deeper insights into their behavior in both solution and on surfaces. Furthermore, they open the door towards future applications of such anisotropic structures. In aqueous solutions, their core-shell structure makes them promising materials for biomedical application. Interactions of the shell with the physiological environment could be tuned through rational chemical modification of the shell, while the core could be used as a reservoir to encapsulate drugs by specific or non-specific interactions. On surfaces, the periodic patterns formed by the nanofibers packing offers very exciting prospects for lithographic applications. The possibility to decrease the size of the shell (while maintaining a core-shell structure) would enable the formation of periodic structure with ultra-small pitch. The challenge would be then to produce a monolayer in order to facilitate the pattern transfer towards the underlying substrate. Langmuir-Blodgett deposition could be used for that purpose.
Acknowledgements

The authors thank the French Agence Nationale de la Recherche for funding (“PISAForFilms” ANR-17-CE09-0031-01, “JASUR” ANR-22-CE06-0016-01 and “Fitt-In” ANR-23-CE06-0019-01. Jacques Jestin (LLB, Saclay) is acknowledged for assistance with SANS experiments and Jean-Michel Guigner (IMPMC, Sorbonne Université) for performing cryoTEM experiments.
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