Unravelling Emergence of Chirality in Click-Chemistry Polymers Down to the Single-Chain Level

doi:10.21203/rs.3.rs-5105035/v1

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Unravelling Emergence of Chirality in Click-Chemistry Polymers Down to the Single-Chain Level

https://doi.org/10.21203/rs.3.rs-5105035/v1

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Chirality plays a critical role in the structure and function of natural and synthetic polymers, impacting their mechanical, optical, and electronic properties. However, a comprehensive understanding of the hierarchical emergence of chirality from monomers to macromolecular assemblies remains elusive, largely because of current limitations in studying their chemical-structural properties at the nanoscale. Here, we unravel the emergence of different forms of chirality from small molecules to their resulting polymers and supramolecular assemblies. We leverage bulk spectroscopic methods combined with the development of mechanical-acoustical suppressed infrared nanospectroscopy, to empower chemical-structural analysis of single-polymer chains for the first time. This unprecedented sensitivity allows identifying key functional groups as a signature for different forms of chirality: CH groups for central chirality in small molecules; C = O groups for backbone and supramolecular chirality in heterogeneous polymers. This work opens a new single-molecule chemical angle of observation into chirality and polymers for the rational design in materials science, biotechnology, and medicine.

Physical sciences/Chemistry/Polymer chemistry/Polymer characterization

Physical sciences/Nanoscience and technology/Techniques and instrumentation/Imaging techniques

Physical sciences/Chemistry/Analytical chemistry/Infrared spectroscopy

Physical sciences/Chemistry/Analytical chemistry/Imaging studies

Physical sciences/Nanoscience and technology/Nanoscale materials/Structural properties

Nano-Analytical Chemistry

Chirality

SuFEx Chemistry

Polymers

Single Molecule Infrared Nanospectroscopy

Chirality plays crucial roles in the structure and functionalities of biomacromolecules and synthetic polymers.^1,2 Nucleic acids, proteins, natural carbohydrates, and several other naturally occurring polymers are composed of chiral building blocks contributing to their unique 3D structure and function, including biological recognition³, transcription⁴, and enzymatic processes.⁵ Taking inspiration from nature, polymer scientists have designed innovative materials where specific arrangement of chiral centres allows unique mechanical,⁶ optical,⁷ and electronic⁸ properties. However, understanding how to tune the molecular and chiral properties of polymers, to achieve new functionalities at the macromolecular level in advanced materials, remains a key challenge.⁹

A critical yet unresolved question is how chirality hierarchically emerges from the sub-molecular to the macromolecular level.¹ Previous work has studied the influence of molecular central chirality and polymerization conditions on polymer backbone chirality.^10–12 However, the mechanisms linking the stereochemistry of the small monomeric building blocks (central chirality), the arrangement and conformation of the polymer backbone (helical chirality), and the organization of single polymer chains in assemblies (supramolecular chirality) are still unclear.¹³ This knowledge gap arises from the limits of current state-of-the-art analytical methods to characterize the complex manifestation of chirality in polymer systems across multiple length scales.

Bulk circular dichroism (CD),¹⁴ infrared (IR) spectroscopy¹⁵, and vibrational circular dichroism¹⁶ can provide insights into chirality of molecular systems. However, bulk methods cannot obtain single-molecule level information on heterogeneous systems; thus, they are not able to differentiate the diverse manifestations of chirality at the (sub)-molecular (central), backbone (helical), and assemblies (supramolecular) level.^17–19 Nano-imaging methods allow visualizing polymeric structures down to the single-chain level and excel in detailing their molecular and chiral conformations. Cryogenic electron microscopy (cryo-EM)²⁰ collects 2D images of single helical molecules with < 0.5 nm resolution, but is limited by the need for extensive sampling and computational 3D model reconstruction,²¹ and has difficulties in reconstructing the structure of molecules with less-defined helices.²⁰ Atomic Force Microscopy (AFM)²² can image the 3-D morphology of single polymers chains and their assemblies with angstroms resolution in the Z direction, and < 1 nm resolution on the 2-D plane.²³ However, both cryo-EM and AFM fall short in discerning chemical information, thereby limiting the understanding of backbone and supramolecular chirality emergence as a function of the central chirality of monomers.¹⁸

To unravel polymer properties, the chemical power of bulk spectroscopy and single-molecule imaging of AFM have been combined in nano-chemical imaging and spectroscopic techniques. Tip-enhanced Raman spectroscopy²⁴ and scattering scanning near-field optical microscopy allowed chemical imaging of single small organic molecules,²⁵ 2D monolayers,²⁶ or larger polymer systems.²⁷ However, because of scattering, plasmonic, and tip-reproducibility effects causing spectral deformation or band suppression, they are limited in extracting structural data from (bio-)organic polymers.^28,29 As a key advantage compared to scattering methods, AFM-based photothermal infrared nanospectroscopy (AFM-IR) has proven to acquire robust and reproducible nano-chemical and structural information from (bio-)organic polymers.¹⁹ AFM-IR is free from scattering effects and independent of the surface/probe geometry.³⁰ While we have demonstrated to achieve AFM-IR single-molecule sensitivity, the method has up to now been limited to large biomolecules (> 400 kDa), while requiring the enhancement of its signal at a gold probe/surface nanogap.³¹ These technological and sensitivity constraints have precluded nanospectroscopic investigations of single polymer chains of smaller molecular weights and on different surfaces necessary to preserve their molecular conformations,³¹ thus restricting polymer analysis primarily to mesoscopic and sub-micrometre scales on relatively homogenous model systems, such as monolayers,³² block co-polymers³³ and polymer blends.^34,30

Here, we gain molecular insights into the hierarchical emergence of chirality from single monomers (central) to polymers (backbone) and their assemblies (supramolecular). To reach this objective, we designed a system with monomers and polymers possessing “same connectivity, different chirality”. We use racemic and chiral monomers to produce both racemic and homochiral (> 99% ee) polymers of ~ 200 kDa, and their supramolecular assemblies, via enantiospecific click reactions.³⁵ We use a multiscale, correlative approach to identify and unravel the relationship among different forms of chirality (central, backbone, supramolecular), from bulk (CD,IR) to single-chain and supramolecular assemblies imaging (AFM). To uncover the correlation between our bulk and single-molecule imaging observations, we then push the limits of AFM-IR nano-chemical analysis. We reach unprecedented sensitivity down to the single-chain level on non-metallic surfaces, revealing the chemical-structural origin of the emergence of supramolecular chirality. Our work paves the way to generalise the characterisation of the mechanisms governing the emergence of chirality in polymers, which is key for the rational design of innovative functional chiral macromolecules in healthcare and materials science.

Multiscale Analysis of Origin and Emergence of Chirality

Figure 1 illustrates our framework to perform a multiscale analysis of the emergence of chirality. We designed a “same connectivity, different chirality” system to unravel the hierarchical emergence of chirality (Fig. 1a): i) from central chirality in small monomeric molecules: ii) to backbone chirality in single polymer chains; iii) and supramolecular order in their assemblies. We selected as model system chiral sulfur fluoride exchange (SuFEx) polymers, a novel class of environmentally friendly and sustainable polymers produced via click chemistry.^35,36

To investigate how chirality manifests from the central chirality of molecules to the backbone and supramolecular chirality of polymers, we implemented a multimodal and multiscale approach (Fig. 1b-e). We leveraged bulk Chiral High-Performance Liquid Chromatography (Chiral-HPLC), CD and IR spectroscopies (Fig. 1b-c) to evaluate and compare the average chemical-structural properties of racemic and chiral monomers and polymers. AFM and AFM-IR empowered us to gain a single-molecule understanding of polymer morphology and chemical properties. Phase-controlled AFM nano-imaging allowed visualizing directly backbone helical chirality at the single-chain level, and unravelling heterogeneity among individual polymer chains and their supramolecular self-assemblies (Fig. 1d).^37,38 For the first time, the development of AFM-IR enabled to correlate the morphological information with the corresponding molecular and chemical-structural properties of single-chains assemblies (Fig. 1e).

The Chiral Building Blocks, their Synthesis and Characterisation

Our framework with same connectivity but different chirality (Fig. 1a) included monomers, linkers, repeating units, and polymers (Fig. 2a-c, SI Fig. 1).^18,35

We synthetised two small molecules serving as building blocks of the polymers: a chiral molecule and a linker (Fig. 2a, SI Fig. 1, SI Methods 1–3). The chiral di(sulfonimidoyl fluoride) molecule (di-SF) had two chiral centres having R (red) or S (blue) configuration (Fig. 2a). Molecules with mixed configurations were termed rac-monomers. Enantiopure chiral di-SFs molecules were termed c-monomers. The linker is a symmetric bis(phenyl ether) (di-phenol) molecule, which is non-chiral and acted as linker for the polymerization (Fig. 2a, link-monomer). The monomers and link-monomer were repetitively bonded via the SuFEx reaction to form the polymers (Fig. 2b, SI Fig. 1, SI Methods 4–5). Rac-monomers may lead to non-helical polymers termed here rac-polymer, while the c-monomers may lead to polymers with helical backbone chirality, termed here c-polymers. Gel Permeation Chromatography (GPC) determined a molecular weight (MW) and polydispersity (PDI) of (SI Fig. 2): c-polymer, M_n~209 kDa with Ð = 1.79; rac-polymer, M_n~220 kDa with Ð = 1.42.

To mimic the smallest c- or rac-repeating unit of the polymers, we synthesised a molecule consisting of a c- or rac-monomer linked to a di-phenol link-monomer with an additional single methyl cap (Fig. 2c, SI Fig. 1, SI Methods 6–7). These repeating units allowed analysing the differences in chirality between monomers and the resulting polymers, focusing on structural differences rather than variations in the composition and presence of specific functional groups.

We finally synthetised both the RR and SS enantiomers of c-monomers, c-repeating units, and c-polymers; to compare their properties with the racemic species.

Bulk Identification of Chirality

Our first aim was to prove the presence/absence of chirality in our building blocks. We applied bulk chiral HPLC (SI Fig. 3), CD (Fig. 2, SI Fig. 4) and Attenuated Total Reflection-Fourier Transform Infrared (ATR-FTIR) spectroscopy (Figs. 2–3, SI Fig. 5–7, SI Note 4).

Chiral-HPLC separated and identified different enantiomers in the monomers and repeating units (SI Fig. 3, SI Note 1). For both monomers and repeating units, we had: i) racemic (rac-) with mixed configurations (RR, RS/SR, SS); ii) enantiomer pure chiral (c-) species. Separation of chiral polymers was hampered in chiral-HPLC because of peak broadening due to polymers heterogeneity.³⁹

We then used CD to identify chirality in all c- versus rac-species (Fig. 2d-f, SI Fig. 4, SI Note 2).⁴⁰ Racemic molecules and polymers showed a flat spectrum, proving the absence of chiral asymmetry; similar behaviour was observed for isolated RS/SR monomers. All c-species (RR, SS) showed a CD spectrum with two peaks (238 nm, 290 nm), proving chiral asymmetry. However, CD does not provide direct information on the configuration of chiral centres, nor it differentiates between central chirality and new sources of chirality, such as backbone chirality emerging during polymer formation.⁴¹

To further identify chirality at the molecular level, we used ATR-FTIR spectroscopy to compare the c- versus rac-species (Figs. 2g-i, SI Figs. 5–7, SI Note 3). As symmetrical c-enantiomers (RR, SS) have the same chemical properties, we focused on the RR species. We acquired ATR-FTIR spectra in the mid-IR range (3200 − 1000 cm^− 1), and used chemometrics and Principal Component Analysis (PCA) to evaluate without bias the differences between the species in the regions of 1800 − 1200 cm^− 1 (Fig. 2g-i, SI Note 4), and 3200 − 2700 cm^− 1 (SI Fig. 5). The spectra and PCA score plots showed minor differences between rac- vs. χ-monomers (Fig. 2g-I, SI Fig. 6, SI Note 3). Instead, we observed statistically significant differences along the first principal component (PC1) for the repeating units (61%) and polymers (77%). Within 1800 − 1200 cm^− 1, these differences were associated with the: carbonyl region (C = O, 1780 − 1620 cm^− 1, ~ 50%); carbon-hydrogen region (C-H, 1550 − 1330 cm^− 1, ~ 20%); and sulphonyl region (S = O, 1300 − 1200 cm^− 1, ~ 30%). Within the 3200 − 2700 cm^− 1 (SI Fig. 6), the differences were associated to the methyl region (CH₃, 2925 − 2800 cm^− 1, ~ 60%). Overall, the FTIR analysis suggested that specific molecular bonds could be associated with chirality, with the largest differences observable in polymers.

Bulk Structural Differentiation of Forms of Chirality

We leveraged the molecular information contained in the IR spectra to unravel structural differences between c- versus rac-species (Fig. 3). PCA analysis (Fig. 2) focused our analysis on the most significant spectral regions: i) CH₃ region, 3020 − 2800 cm^− 1 (Fig. 3a); ii) the C = O region, 1780 − 1620 cm^− 1 (Fig. 3b). Our system with same connectivity but different chirality allowed then associating IR spectral differences purely to different structural features.

CH₃ groups are abundant and often located directly or closely to chiral centres or the backbone of polymers.^42,43 Within 3020 − 2800 cm^− 1 the spectra showed significant differences between rac- vs. c- species (Fig. 3a,c), which for monomers and repeating units could only be ascribed to central chirality.^8,42,44 Monomers showed only minor differences (SI Fig. 5–7). The repeating units showed a significantly higher IR absorption of the aromatic CH₃^as and CH₃^s groups for the rac- than the c-species (Fig. 3c); relative to the aliphatic CH₃^as (2970 cm^− 1) in the link-monomer. The weaker IR absorption of c-repeating units (RR) may be associated with two less co-planar aromatic CH₃ groups, with fixed angle respect to the backbone (SI Note 4). While rac-repeating units had a larger set of conformations (RR, SS, RS/SR) contributing to a higher IR absorption. Thus, central chirality conformation was associated to an increased planarity of the rac-molecules compared to the c-molecules. The polymers showed also differences in the C-H region, which however may also arise from polymerisation and phenomena such as backbone and supramolecular chirality.

The C = O stretching has been widely related to the structural conformation of the backbone of (bio)-polymers.^15,45 In the range 1700 − 1620 cm^− 1, monomers and repeating units did no show significant differences between rac- vs. c-species, while the polymers showed statistically significant differences (Fig. 3b, SI Fig. 7, SI Note 4). Thus, the spectral differences between rac- vs. c-polymers could be ascribed only to other forms chirality arising from polymerisation and supramolecular assembly. We observed that (Fig. 3b,d): i) the C = O peak shifted to lower wavenumbers from monomers (1663 cm^− 1), via the repeating units (1655 cm^− 1), to polymers (1648 cm^− 1); ii) c-polymers had higher absorption than rac-polymers at 1656 cm^− 1; iii) c-polymers exhibited a single C = O peak, while rac-polymers showed the emergence of a second C = O peak at 1723 cm^− 1.

We associated the shift to lower wavenumbers to electron delocalization caused by the backbone formation in polymers,⁴⁶ moreover the increased IR absorption at 1656 cm^− 1 for the c-polymers suggested the stabilisation of helices in the backbone⁴⁷ (Fig. 3d, SI Note 4). The emergence of the second C = O peak for the rac-polymer could be associated with (SI Note 4): i) intramolecular interactions in their RS/SR repeating units, via a p-p resonance between the carbonyl groups and phenyl groups, inducing a planar conformation; ii) enhanced intermolecular p-p stacking of the RS/SR phenyls leading to supramolecular assembly and possibly higher chiral order.

Overall, the IR structural analysis identified general key molecular vibrations useful to differentiate central chirality in small molecules (CH₃) and backbone/supramolecular chirality in polymers (C = O). However, bulk IR prevented analysing sample heterogeneity at the single-molecule level to prove: the helicity of c-polymers; and determine whether the splitting of the C = O absorption and the emergent peak at 1723 cm^− 1 are related to an intramolecular planar conformation or intermolecular assembly of the rac-polymers.

Single-molecule Imaging of Chirality with Angstrom Resolution

To unravel at the single-molecule level the origin of the emergence of chirality from small molecules to polymers, we further leveraged single-molecule 3D imaging. We investigated the presence of backbone helical chirality and the supramolecular assembly state of the rac- and c-polymers with angstrom resolution imaging via phase-controlled AFM (Fig. 4, SI Fig. 8–11).³⁷

We choose atomically-flat Highly Oriented Pyrolytic Graphite (HOPG) to preserve polymeric structure and mimic the hydrophobicity the diamond ATR-FTIR crystal (SI Fig. 8, Methods). AFM morphology maps showed significant differences in shape and size of c- versus rac-polymers (Fig. 4a, SI Fig. 9). A single-molecule statistical analysis of the volume of the polymers was conducted to estimate their molecular weight and polydispersity (SI Note 5). The c-polymers had AFM-measured molecular weight of 231.8 ± 138.9 kDa and PDI ~ 1.36; while the rac-polymers had higher molecular weight of 1219 ± 1003 kDa and PDI ~ 1.68. The comparison of single-molecule AFM and bulk GPC data (Fig. 4b, SI Fig. 2) demonstrated that c-polymers maintained a single-chain conformation, while the rac-polymers self-assembled on the surface.

Ultra-high resolution AFM 3D imaging was further performed to unravel the topological structure of single rac- and c-polymer chains. The AFM maps and the height cross-sectional profile of c-polymer chains showed a helical structure (Fig. 4c, SI Fig. 10). The rac-polymers did not show any periodic structure at the single-chain level (Fig. 4d). To quantitatively assess the presence of a helical structure in χ- vs. rac-polymers, we calculated the Fast Fourier Transform (FFT) of the cross-sectional profiles of the single-chains against any possible residual periodicity on the HOPG surface (SI Figs. 10–11). The FFT power spectra confirmed that χ-polymers had helical periodicity of ~ 40 nm, while rac-polymers single-chains had no backbone periodicity.

Although rac-polymers did not have helical conformation, the AFM analysis showed that rac-polymers further self-assembled into ordered and periodic structures (Fig. 4a,e). To unravel the supramolecular order of these rac-assemblies, we performed a single-molecule analysis of their shape as a function of their increasing size (Fig. 4e). The smallest and most abundant rac-assemblies showed: a constant height of ~ 0.4 nm (SI Fig. 10), which is the typical height of an organic chain on a surface; a transversal periodicity increasing in multiples of ~ 15 nm (Fig. 4f), corresponding to the convoluted width of a single chain. This analysis indicated that the rac-polymers initially self-assemble via lateral interactions of single-chains, which further assembly into larger supramolecular species with increasing height (Fig. 5).

Overall, the AFM single-molecule analysis was in excellent agreement with the IR structural results. AFM demonstrated that the c-polymers have a helical backbone structure, in agreement with the C = O shift at lower wavenumber and increased IR absorption at 1656 cm^− 1, which is a typical signature of helices in (bio)-polymers. It then showed that rac-polymers undergo self-assembly via lateral interactions; in agreement with FTIR conclusion that emergence of a new second C = O peak at 1723 cm^− 1 could be associated to intramolecular p-p resonances in the RS/SR polymeric-units with a co-planar conformation and/or their supramolecular assembly via intermolecular p-p stacking. However, structural analysis of rac-polymers by AFM nano-imaging and IR structural analysis separately could not unravel whether the emerging C = O at 1723 cm^− 1 had an intra- or inter-molecular origin. To unravel the origin of this structural signature, we further applied AFM-IR nano-chemical analysis.

Achieving Single Polymer Chain Nano-Chemical Analysis

Chemical-structural analysis of single polymer chains and assemblies paves the way for understanding the molecular origin of the emergence of supramolecular order in the rac-polymers. However, to date, AFM-IR single-molecule sensitivity has been limited to molecules > 400 kDa; and only if placed at a metallic nanogap between a gold surface and AFM probe. Indeed, the use of non-metallic substrates like HOPG would cause a further drop in AFM-IR sensitivity larger than 10-folds.

To allow for the first time the chemical analysis of single polymer chains of molecular weight down to ~ 200 kDa on non-metallic surfaces such as HOPG, we developed acoustical-mechanical suppressed (AMS) AFM-IR with sub-Angstrom noise level (Fig. 5, SI Fig. 12). We leveraged composite foams (Fig. 5a, Methods) to achieve RMS noise of 25 ± 5 pm in tapping-IR and 60 ± 5 pm in contact-IR mode; allowing chemical-imaging of single monoatomic steps of HOPG with ~ 3 Å height (Fig. 5b, SI Fig. 13). To reduce AFM thermal drift and allow single-molecule localisation with high-accuracy, we stabilised room temperature with variations < 0.1°C/h (Figure SI 13).⁴⁸

We then applied AMS AFM-IR to extract structural information of single polymer chains from nano-resolved IR maps and spectra as in bulk FTIR. We first focused on c-polymers (Fig. 5c-f). Tapping-IR allowed achieving multimodal imaging of morphology and IR absorption of a single polymer chain with a height down to ~ 5 Å (Fig. 5c); the phase locked loop (PLL) tracking of contact resonance frequency assured that chemical contrast was not affected by mechanical effects. Next, we pointed the AFM probe on a single polymer chain in the map (Fig. 5d) to acquire nano-localised IR spectra (Fig. 5e); for chains with a size down to ~ 0.8 nm height and ~ 5 nm width (Fig. 5f), corresponding to a molecular weight of ~ 200 kDa. While preserving better topography, tapping-IR detection was limited to the S = O region of a single c-polymer chain by low signal-to-noise ratio (SNR).

We thus employed contact-IR mode to detect the chemical properties of rac-polymers (Fig. 5g-k). Figure 5g shows a representative rac-polymer assembly. Instead of single wavenumber IR maps (Fig. 5c), we acquired 4D hyperspectral chemical maps of the polymers. Figure 5h shows four 3D-slices of the hyperspectral map, where each pixel represents an IR spectrum. The IR spectra showed the typical S = O absorption of the rac-polymer; while the HOPG surface showed significantly lower and uncorrelated IR absorption, which was subtracted to correct the spectra of the polymers (SI Fig. 13–14). The high SNR allowed acquiring nano-localised spectra of single rac-polymer assemblies in the full 1800 − 1200 cm^− 1 fingerprint region (Fig. 5i-j), for assemblies with a size down to ~ 4.5 nm in height and ~ 4 nm in width, corresponding to ~ 2–4 chains (Fig. 5k, SI Fig. 14).

The emerging supramolecular C = O stretching at 1723 cm^− 1 was observed in all acquired IR spectra of different rac-assemblies (SI Fig. 14), but with significant variations between different assemblies. We thus proceeded to a statistical analysis of their nano-chemical heterogeneity by AMS AFM-IR.

Unravelling the Origin of the Emergence of Supramolecular Chirality

We leveraged our AFM-IR unprecedented sensitivity to investigate how the chemical signature of the rac-polymers varied as a function of their size and supramolecular state, to unravel if the emerging C = O at 1723 cm^− 1 is related to intra- or inter-molecular interactions (Fig. 6).

We acquired AFM-IR spectra from assemblies with varying size, from the smallest one (Fig. 6a, ~ 4.5 nm height; ~4 nm deconvoluted radius) to the largest ones (Fig. 6b, ~ 10 nm height; ~6 nm deconvoluted radius). Although at the limit of AFM-IR sensitivity, smaller assemblies presented lower IR signal than larger ones in accordance with Beer-Lambert Law (Fig. 6c, SI Fig. 14). We then normalized the spectra to their maximum, to reduce variability due to signal intensity and ran a PCA analysis in the full fingerprint region (1750 − 1200 cm^− 1). The PCA analysis showed that smaller (3.0–6.4 nm) and larger assemblies (7.5–15 nm) belong to separate spectroscopic clusters (Fig. 6c, 95% confidence). Significant part of the variance related to the conformation of the C = O group (Fig. 6e), suggesting that rac-polymers of different sizes have different structural properties.

We consequently focused our analysis on the C = O groups (Fig. 6f, 1780 − 1600 cm^− 1), which were associated by IR to intra- and/or inter-molecular interactions in the rac-polymers. We normalised the spectra to the backbone C = O peak at 1646 cm^− 1, to investigate the structural differences related to the emergent C = O peak at 1723 cm^− 1, independently of the number of C = O bonds present in the assemblies. The shape of the emergent C = O peak was significantly different for smaller vs. larger assemblies; as proved by PCA analysis of the second derivatives of the spectra (Fig. 6g, 95% confidence). The PC1 loading plot of the second derivatives (Fig. 6h, 81% variance) proved that larger assemblies had increased absorption of the emergent C = O peak, which also shifted its maximum at higher wavenumber (1730 cm^− 1), compared to smaller assemblies (1710 cm^− 1). Since we normalised the spectra to the backbone C = O peak (1646 cm^− 1), these modifications could only be associated to an increasing number of intermolecular (p-p stacking) interactions between the rac-polymer chains in the larger assemblies; rather than the constant number of intramolecular interactions in the single chains.

Overall, AMS AFM-IR proved the ability of studying the heterogeneous chemical-structural properties of single polymeric assemblies, composed of only few chains. The analysis enabled identifying the splitting of the C = O group absorption and the second peak at 1725 cm^− 1 as key signature of intermolecular p-p stacking and the emergence of supramolecular order and chirality in the rac-polymers (Fig. 6i).

Despite over a century of investigations, the hierarchical emergence of chirality over multiple chemical and topological scales has remained elusive due to technological limitations in characterizing different forms of chirality at the molecular nanoscale.⁹

To bridge this gap, we have elucidated here for the first time the complex interplay between central chirality in small molecules and the emergence of backbone and supramolecular chirality in polymers from sub-molecular scales to macro-molecular scales. To reach this objective, we have developed acoustical-mechanical suppressed AFM-IR to achieve for the first time chemical-structural analysis of the heterogeneity and assembly of single-polymer chain(s). This unprecedented sensitivity, combined with nano-imaging and bulk spectroscopy, enabled distinguishing between different forms of chirality and their structural manifestation at the molecular level via bonds generally present and abundant in small molecules, polymers and their supramolecular assemblies (Fig. 6i). We could first identify CH₃ groups conformation as a key signature of central chirality in small molecules, which is in large part insensitive to backbone conformation. Second, we generalised the idea that C = O carbonyl is sensitive to backbone structure in protein molecules and assemblies^14,15 to identify the conformation of C = O groups in polymers, as key sensitive signature^49,50 of helical chirality via twisting of their backbone and supramolecular order chirality via p-p intermolecular stacking of their chains.

Overall, these results contribute to understanding the hierarchical emergence of chirality in polymers and pose a new framework for offering key insights into the structural-chemical determinants of chirality across multiple scales. This understanding will be in turn fundamental for the rational design of advanced chiral materials with tailored properties and may help to accelerate the development of novel materials in electronics, chemical industry, biotechnology, and healthcare.

Finally, this study pushes the limits of nano-analytical chemistry down to the identification of the molecular fingerprint and the structural analysis of a single polymer chain on any surface, without requiring the enhancement of metallic surfaces. These capabilities pave the way to have a new single-molecule and angstrom resolved window of observation into the chemical-structural properties of virtually any (bio-)organic material in nature.

Preparation of SUFEX Monomers, Repeating Units and Polymers

Monomers were prepared following a previously established protocol (See SI Methods 2). The monomers were characterized by NMR and HRMS (See SI Methods 3). Polymers were prepared following a previously established protocol (See SI Methods 4). The polymers were characterized by NMR (See SI Methods 5). A new procedure for the preparation of the repeating units was developed here, where first monomers where modified and then reacted with link-monomers, to obtain the wanted rac- and c-repeating units (See SI Methods 6).

We verified the molecular weight, chemical identity, and purity of our synthesized compounds using state-of-the-art Nuclear Magnetic Resonance (NMR) and High-Resolution Mass Spectrometry (HRMS) (SI Methods 1–7, SI Data Set).

NMR Measurements

The Bruker Avance III 400 MHz spectrometer was used to record ¹H NMR, ¹³C NMR, and ¹⁹F NMR spectra at a temperature of 298 K. The chemical shifts are listed in ppm on the δ scale and coupling constants were recorded in Hertz (Hz). The 19F NMR was not referenced. All other chemical shifts are calibrated relative to the signals corresponding of the non-deuterated solvents (CHCl₃: δ = 7.26 ppm for ¹H and 77.16 ppm for ¹³C, DMSO: δ = 2.50 ppm for ¹H and 39.52 ppm for ¹³C, CH3CN: δ = 1.94 ppm for ¹H and 1.32 and 118.26 ppm for ¹³C). Abbreviations are used in the description of NMR data as follows: chemical shift (δ = ppm), multiplicity (s = singlet, d = doublet, t = triplet, q = quartet, dt = doublet of triplets, m = multiplet, br = broadened), coupling constant (J, Hz).

HR-MS measurements

The Quadrupole Time-of-Flight (QTOF) micro-spectrometer (Thermo Fisher Scientific Inc., United States) was used to obtain high-resolution mass spectra (HR-MS) through electrospray ionization (ESI) in either positive mode (ESI+) or negative mode (ESI-).

HPLC measurements

Enantioselectivity was monitored using high-performance liquid chromatography (HPLC, Agilent 1100 bearing a UV-Vis detector) with a CHIRALPAK® IA column (5 µm, 4.6 mm x 250 mm). As references, racemic products or nonchiral mixtures, including diastereoisomers, were synthesized to determine the column conditions for baseline separation. Unless specifically mentioned, these were found to be n-hexane/dichloromethane/iso-propanol (80/20/2.5) as the mobile phase, a flow rate of 0.5 mL/min, room temperature, and a detector wavelength of 240 nm.

GPC measurements

Gel Permeation Chromatography (GPC) was performed using an Agilent Technologies 1200 system, equipped with a diode-array detector (DAD, G1315D), a refractive index detector (RID, G1362A), autosampler (ALS, G1329A), bin pump (G1312A), solvent degasser (G1322A), and a VARIAN GPC/SEC column from PLgel 5 µm MIXED-D (300 × 7.5 mm). The system was calibrated with EasyVial PS-M and EasyVial PS-L polystyrene standards (Agilent Technologies), GPC polystyrene standards with a combined range from Mp 400 to 2,000,000 Da (Fluka). HPLC-grade tetrahydrofuran (THF, Biosolve®, unstabilized, HPLC grade) was used as a mobile phase for low molecular weight polymers and high molecular weight polymers were measured using HPLC grade N,N-Dimethylformamide (DMF, Biosolve®, unstabilized, HPLC grade) with 0.1% LiBr (ReagentPlus®, ≥ 99% from Sigma-Aldrich).. All GPC measurements were performed at room temperature or at 25°C using an injection volume of 10 µL (1.0 mg/mL). The polymer number average molecular weight (M_n) and polydispersity (Đ) relative to a set of polystyrenes were determined using DAD (at 270 nm or 290 nm) or RI detectors.

FTIR measurements, Data Treatment and Analysis

Attenuated total reflection infrared spectroscopy (ATR-FTIR) was performed using a Bruker Tensor 27 IR equipped with Platinum ATR accessory with a diamond ATR element. Spectra were collected by the built-in software OPUS (version 7.8, Bruker, USA). The samples were deposited on the ATR crystal by 2 times of 0.5 µL droplet at a concentration of 2 mg⋅mL^− 1 in THF. Spectra were acquired at the range of 3500 cm^− 1 − 500 cm^− 1 with a resolution of 4 cm^− 1. All spectra were processed using OriginPro (version 2022b, OriginLab, USA). For each sample, 5 spectra of 512 co-averages were collected. Since the spectral background line shape slightly depends on wavenumber regions and at lower wavenumber range (2500 − 500 cm^− 1 range) and higher wavenumber range (3500 − 2500 cm^− 1 range) does not share the same baseline. The spectra were first cut into two ranges that contains relevant information: lower wavenumber range (1800 − 1000 cm^− 1 range) and higher wavenumber range (3150 − 2700 cm^− 1 range). The final spectra for each range were obtained by averaging 5 individual spectra. Each spectrum was smoothed using a Savitzky–Golay filter (second order, 7pts), baseline-corrected (zero-order for 1800 − 1000 cm^− 1 range and first-order for 3150 − 2700 cm^− 1 range) and normalizing to the range [0,1].

Circular dichroism measurements

Circular Dichroism (CD) measurements were done on a Jasco J-715 spectropolarimeter equipped with a Peltier holder. CD spectra were measured at a concentration of 20 µg⋅mL^− 1 in a Hellma® Quartz Glass high performance macro absorption cell with a 0.1 cm pathlength. Measurements were performed with a scanning speed of 50 nm ⋅min^− 1 with a data pitch of 0.5 nm at 20°C and were averaged from 16 scans. All spectra were processed using OriginPro (version 2022b, OriginLab, USA). The final spectra were obtained by first subtraction of THF background, secondly, applying a Savitzky–Golay filter (second order, 21 pts) and finally converting to molar ellipticity by the following formular: \(\:=\frac{M\times\:{\theta\:}_{obs}}{10cl}\), where \(\:M\) is averaged molecular weight in g⋅mol^− 1, \(\:l\) is the path length (0.1 cm) and \(\:c\) is the concentration [mg⋅mL^− 1].

AFM Sample Preparation

Samples deposition for atomic force microscopy (AFM) analysis was performed on two differently charged substrates: hydrophobic highly ordered pyrolytic graphite (HOPG) and negatively charged Mica. Sample preparation was realized by two different methods: manual deposition and vacuum drying. For manual deposition, a 10 µl droplet of 2 µg/ml polymers solution in THF was applied at room temperature for 40 seconds on substrate. A similar droplet was used for HOPG. in the case of HOPG. After the incubation time on the surface. Then the sample was rinsed with 1 ml Milli-Q water and then dried with a gentle flow of nitrogen gas. For Vacuum drying, a 10 µl droplet of 2 µg/ml polymers solution in THF was applied at room temperature for 3 seconds on substrate under vacuum.

AFM Measurements

Atomic force microscopy (AFM) measurements were performed on a NX10 AFM (Park systems, South Korea) operating in True Non-Contact™ Mode with a silicon tip PPP-NCSTR (Park systems, nominal spring constant 7.4 Nm^− 1, nominal tip radii ~ 7nm). The image was acquired with a tip velocity of 0.8 µm per second. Three map sizes and two pixelization setting were employed: 4×4 µm² with 4 nm/px, 1×1 µm² and 0.5×0.5 µm² with 0.5 nm/px. Image flattening and cross-sectional dimension analysis of single polymers chains were performed by SPIP (Image Metrology, Denmark) software.

We initially optimised the deposition of the polymers on three substrates with different charges: positive gold, negative mica, hydrophobic Highly Oriented Pyrolytic Graphite (HOPG) (SI Fig. 8). HOPG was chosen as a substrate for high-resolution imaging because of its atomically flat surface, favourable interactions with hydrophobic polymers, and similar hydrophobicity as the diamond ATR-FTIR crystal, thus allowing to compare the molecules’ behaviour on the surface. Each image was first flattened with a first order polynomial fit for the global correction, and then, a zero order least mean square (LMS) fit line-wise correction. Afterwards, the image was further flattened by first and second order LMS fit line-wise correction with area of interest masked.

Acoustical-Mechanical Suppressed AFM-IR

AFM-IR is based on the mechanical detection of the photothermal expansion of a sample when absorbing IR light. To achieve AMS AFM-IR, we leveraged composite foams to suppress any source of mechanical and acoustic noise transmission caused by cables and vibrating parts inside and outside our AFM-IR system (Fig. 5a). The noise suppr ession was quantified on the acquired AFM-IR maps via their RMS noise and the integration of their FFT spectral noise within 0-500 kHz (Figure SI 12). In tapping mode, AMS AFM-IR achieved RMS noise level of 25 ± 5 pm to allow imaging single monoatomic steps of HOPG with a height of ~ 3 Angstroms (Fig. 5b). In contact mode, the AMS suppression led to RMS down to 60 ± 5 pm, against a standard RMS noise of 215 ± 15 pm. To further allow single-molecule localisation with high-accuracy, we stabilised room temperature with variations below 0.1°C per hour to reduce AFM thermal drift (Figure SI 13).

AFM-IR measurements, maps treatment and analysis

Analysis by AFM-IR was performed with a nanoIR3 (Bruker, USA) on atomically flat HOPG substrates. The substrate roughness and chemical response were also characterized by AFM-IR in the SI Fig. 13. The root mean square roughness of the AFM maps was measured by SPIP (Image metrology, Denmark). The morphology of the polymer samples was scanned by the nanoIR microscopy system, with a line scan rate within 0.1–0.5 Hz and in contact mode. All AFM maps were acquired with a resolution between 1–30 nm/pixel. A silicon gold coated PR-EX-nIR2 (Bruker, USA) cantilever with a nominal radius of ~ 30 nm and an elastic constant of about 0.2 N/m was used. The AFM images were treated and analysed using SPIP software. The height images were corrected using a first order polynomial flattening.

All measurements were performed at room temperature under controlled nitrogen atmosphere with relative humidity below 3%. Both spectra and images were acquired by using phase loop (PLL) tracking of contact resonance, the phase was set to zero to the desired resonant frequency around the IR amplitude maximum, and tracked with an integral gain I = 0.2 and proportional gain P = 1.

FUNDING

The research presented in this manuscript has received funding from the Dutch Sector Plan beta.

AUTHORS Contributions

F.S.R. and X.L. conceived the project. X.L. performed the CD, IR, AFM experiments and analysed the relative data. F.S.R. and X.L. performed the AFM-IR experiments and analysed the relative data. S.P. synthetised the molecules and the polymers, performed the HPLC, NMR, GPC, MS experiments and analysed the relative data. X.L. and F.S.R. wrote the original draft of the manuscript. F.S.R., J.V.G., H.Z. supervised X.L. H.Z. supervised S.P. All the authors revised and commented the manuscript.

Competing interests

The authors declare no competing interests.

Data

All data needed to evaluate the conclusions in the paper are present in the paper and/or the Supplementary Materials. The source data underlying Figure 1-6 and SI Figure 1-14 are provided as a Source Data file. Additional data are available from authors upon request.

Corresponding Author

Correspondence should be addressed to: F.S. Ruggeri ([email protected]) , H. Zuilhof ([email protected]).

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Unravelling Emergence of Chirality in Click-Chemistry Polymers Down to the Single-Chain Level

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Abstract

Figures

INTRODUCTION

RESULTS AND DISCUSSION

Multiscale Analysis of Origin and Emergence of Chirality

The Chiral Building Blocks, their Synthesis and Characterisation

Bulk Identification of Chirality

Bulk Structural Differentiation of Forms of Chirality

Single-molecule Imaging of Chirality with Angstrom Resolution

Achieving Single Polymer Chain Nano-Chemical Analysis

Unravelling the Origin of the Emergence of Supramolecular Chirality

CONCLUSIONS

MATERIALS AND METHODS

Preparation of SUFEX Monomers, Repeating Units and Polymers

NMR Measurements

HR-MS measurements

HPLC measurements

GPC measurements

FTIR measurements, Data Treatment and Analysis

Circular dichroism measurements

AFM Sample Preparation

AFM Measurements

Acoustical-Mechanical Suppressed AFM-IR

AFM-IR measurements, maps treatment and analysis

Declarations

References

Additional Declarations

Supplementary Files

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