thermo-draw history that undergo commonly different applications such as the food and
packing industry as well as cast film extrusion process. In situ WAXS and SAXS experiments using
small-scale industrial equipment such as the study of 2D stretching at different temperatures
and grades of low-density polyethylene (LDPE) has been conducted (Portale et al., 2013; van
Drongelen et al., 2014; Troisi et al., 2016) using a real film blowing machine connected to an
extruder. The dependence of the onset of crystallinity and the crystal orientation as a function
of different take-up and blow-up ratios of the polymer film has been studied. Flow-induced
crystallization has been largely investigated by in situ simultaneous SAXS-WAXS experiments in
combination with a slit-like rheometer (Balzano et al., 2010; Ma et al., 2012; Troisi et al., 2015)
to correlate the structure-property relationship of injection molding industrial process. The
formation under high shear rates of the so-called shish-kebab structures as a function of the
imposed stress in a wide range of polymers has been monitoring as well as the effect of
nucleating agents upon shear (Wilsens et al., 2018, 2019). Recently, the crystallization upon
precise and locally controlled uniaxial extensional flow of LDPE has been investigated by
dedicated rheo-SAXS experimental to decouple the effect of the extensional component from
the rotational component present in shear flow and thus, attain higher extensional forces.
(Pepe et al., 2020) Furthermore, polymer water/CO₂ emulsions at extreme conditions have
been monitoring (Alaimo et al., 2015) by the development of a versatile high-pressure X-ray
vessel (Hermida-Merino et al., 2014) fitted with a large exit opening angle that permits acquired
WAXS patterns. The internal structural development of block copolymer microparticules during
a dispersion polymerization has been probing by in situ simultaneous SAXS-WAXS experiments
(Alauhdin et al., 2019) under supercritical CO₂ to assess the degree of CO₂ affinity on the
different block on the final quenched structure. Moreover, Franeker et al. (Franeker et al.,
2017) reported recently the formation at different length-scales of sub-micrometer structures
during the thin film formation by spin coating of polymeric blends and for block-copolymers
(Fleury et al., 2019), studied by ultrafast grazing incidence small angle and wide X-ray scattering
(GISAXS/GIWAXS (Dong et al., 2020)) coupled with laser interferometry and laser scattering.

Experimental strategies to perform simultaneous SAXS-WAXS experiments in combination with non-X-ray based techniques with different sensitivities have been extensively discussed by Bras et al (Bras et al., 2014) as well as the design of appropriate complex sample environments. A pioneering combination of vibrational technics with synchrotron radiation light has been reported by Bras et al. (Bras, Derbyshire, Bogg et al., 1995) in which SAXS has been coupled to Fourier-transform infrared (FTIR) spectroscopy to study the driving force for structure development in polyurethanes. However, the technical requirements to conduct Fourier-transform infrared spectroscopy prevent to universalize the combination with other techniques and specific complex sample environments have to be integrated. In contrast, Raman spectroscopy has been widely exploited in combination with multiple techniques with the advent of the technological developments due to simple implementation, improvement of the time resolution and experimental advantages. (Gerrard & Maddam, 1986) Raman spectroscopy is a non-destructive technique that can be applied to different media (bulk, fibers, solution, gels and films), specifically aqueous solution without special experimental preparations, and requires small quantities of sample to be studied. Moreover, Raman spectroscopy is sensible
to molecular vibrations related to the bond polarization of the active moieties, (Ferraro *et al.*, 2003) which generally is attributed to apolar groups. Consequently, Raman spectroscopy conveys structural details of the chemical bonds, singularly for polymers, for both the monomer units to the macromolecular conformation and can be easily polarized to detect structural anisotropies. (2000) Furthermore, the Raman scattering response probes the covalent nature representative of polymeric backbone in particular the frequent unsaturated bonds as well as the chain stereoregularity. (Edwards *et al.*, 1993) In addition, the low frequencies spectral range is correlated to the collective motion associated with the transition time to identify and correlate the chemical conformation to the structure of nanomaterials. Quantitative analysis by Raman spectroscopy of the system composition requires the correction of the scattering bands intensity either by adding an internal standard or calibration with a different technique. (1999) The indirect relationship of Raman scattering intensity to the concentration arises from the uncertainty of the scatters volume illuminated by in the incident beam. However, the linear intensity dependence facilitates to distinguish different component concentrations. Particularly, the crystallinity content of semicrystalline polymers has been largely determined by Raman spectroscopy as well as the characterization of the related nanostructured is analyzed combined with SAXS/WAXS experiments due to the complementary sensitivities to different crystalline, interphase or amorphous phases. Moreover, the detailed description of the nanostructure composition by SAXS experiments has been greatly benefited by associated to Raman quantitative analysis to optimize the parameters of the correlation function, particularly at crystallinities around 50 % where electronic contrast hamper the phase assignment by SAXS.

Similarly to X-ray scattering techniques, polymerization under industrial has been largely employed by Raman spectroscopy to monitor both the kinetics and physical parameters of the reaction medium, especially in heterogeneous systems such as monomer droplets, polymer particles, micelles ( dropsit et al., 2019). Moreover, the optimization of the polymerization conditions in solution have been studied (Chevrel et al., 2012) by combined Raman spectroscopy to rheological measurements to relate the fluid mechanics with the synthetic mechanism of acrylic acid at different shear rates(Chevrel et al., 2014), enabling to be scaled to industrial environments by pilot plants studies (Kohlmann et al., 2016) (Chevrel et al., 2016). Likewise, the gel effect during the bulk polymerization of polystyrene has been probed (Brun, Youssef et al., 2013) by coupling Raman spectroscopy to rheology to correlate the viscosity development of the heterogeneous medium with the monomer conversion. Coupling Raman spectroscopy to drop tensiometry illustrated the versatility to monitor heterogenous polymerizations in a drop reaction vessel within an aqueous medium to account for the relation between monomer conversion with the interfacial tension and dilatation rheology(Dropsit et al., 2018). Remarkably, the scattering nature of the Raman response has been proved to discriminate simultaneously the kinetic reaction mechanism and phase separation or phase inversion physical phenomena during the polymerization by examining the reaction media with Raman probes with different analyzed volumes. (Brun et al., 2014) Moreover, the mechanism of deformation under uniaxial stretching of iPP has been monitored (Martin et al., 2013) by in situ polarized Raman spectroscopy, yielding detailed nanostructure evolution of the
amorphous and the crystalline domains. Beneficially, the orientation of the polymeric
amorphous domains and the stress distribution within the chemical bonds was directly probed
by Raman spectroscopy, in contrast to X-ray scattering techniques, as well as the degree of
crystallinity and orientation of the crystalline domains. Furthermore, ammonium nitrate phase
transitions upon thermal treatment have been monitored with in situ Raman spectroscopy and
DSC by relating with the group theory, the lattice vibrations to the required phase transition
energy. (Harju et al., 1992)

Particularly, Raman spectroscopy has been coupled to several X-ray based techniques across
the synchrotron community for both scattering such as micro-XRD (Davies et al., 2005), high
resolution (Boccaleri et al., 2007) XRD or spectroscopic techniques like UV-vis XAFS (Beale et al.,
2005) and sequential combinations of both as XANES-XRD. (Kongmark et al., 2009) Likewise,
Raman spectroscopy is beneficial to be associated with SAXS-WAXS techniques to analyze the
nanostructural evolution of polymeric systems upon external stimuli whilst obtaining
concurrently kinetic data as well as identifying potential X-ray beam damage. (McGeehan et al.,
2007) (Stanley et al., 2014) Simultaneous Raman spectroscopy with high time resolution (2s)
on back scattering mode using optical fiber probe with time-resolved SAXS-WAXS
measurements (Bryant et al., 1998) has been applied to follow in situ the phase transition of
high density polyethylene upon thermal treatment in combination with a Linkam calorimeter
stage. Moreover, the online combination of Raman spectroscopy with SAXS/WAXS has been
proved to advance the on the understanding of complex polymer processing such as reactive
solvent polymerizations where concomitant or sequential physicochemical phenomena occurs
that imply kinetic and morphological variations.

Firstly, the mechanism of formation and the polymer-solvent compound as well as the phase
transformation in syndiotactic polystyrene (sPS) gels was probed by harnessing the
simultaneous data at different lengthscales obtained by the combined Raman spectroscopy
SAXS WAXS setup (Rastogi et al., 1998; Goossens et al., 1999). The optical density differences
detected by Raman spectroscopy to identify the developed solid liquid phase separations upon
thermal treatment as well as the conformational changes during the phase transformations of
sPS with different solvents, (decalin, benzyl methacrylate, and cyclohexyl methacrylate).
Furthermore, the kinetics and polymerization mechanism of a reactive solvent as styrene with
PE has been (Goossens et al., 2001) monitored at different temperatures by the coupled
Raman spectroscopy with simultaneous SAXS/WAXS to follow the temperature dependence of
the sequential liquid-solid phase separation and subsequent PE crystallization that was
affected by solvent removal.

The structural development of both the crystalline and amorphous region under tensile
deformation of isotactic polypropylene (iPP) fiber (Ran et al., 2003) have been studied by in
situ simultaneous polarized Raman spectroscopy and synchrotron two-dimension (2D) WAXD
measurements. Raman spectroscopy has permitted to resolve the structural evolution during
the α to mesophase transition upon stretching to a packing disorder instead to the backbone
conformational conversion. Furthermore, Ponçot et al. (Ponçot et al., 2015) have conducted
simultaneous Raman spectroscopy and high energy wide angles X-rays scattering (HE-WAXS)
during uniaxial tensile in combination with a videotraction system of different polypropylene
blends to characterize the structural polymer micromechanism under the real mechanical
behavior. The multitechnique approach revealed a systematic correlation between the WAXS
and Raman spectroscopy mechanism of orientation of the polymer nanostructure was found.
However, WAXS orientation factors were found to be more sensible under low strain regimes
whereas chain orientation ratio detected by Raman spectroscopy were more accurate at high
strains values, highlighting the importance of the combined study on the same sample volume.

The coupling of SAXS with Raman spectroscopy (Haas et al., 2014) has been employed to study the dispersion properties of single-wall carbon nanotubes (SWNTs) in sodium dodecyl sulfate (SDS). The multiprobe environment (UV-Raman spectroscopies) has been controlled using an interface build with Python on SPEC (Certified Scientific Software) allowing the simultaneous triggering of both X-ray scattering and Raman spectrometer. Remarkably, the fitting of the SAXS profiles has been assisted by the Raman spectra by obtaining a proportional total number density of SWNTs as a function of the SWNT concentration from both SAXS and Raman data.

Summarizing, great efforts have been devoted to the combination of scattering techniques with spectroscopic methods, in particular, vibrational techniques, to cover the structural transformations at different length scales and to establish the driven force between the morphological transitions and the chemical conversions. Consequently, large amounts of data are generated that are tedious to analyze with slight temporal shifts if different techniques are not integrated under the same control unit. Moreover, acquisition time per frames will tend to differ among techniques sensitivities with the advent of new sources at synchrotron radiation facilities, highlighting the requirement of technical control on the triggering, acquisition, and files recording to avoid time discrepancies among variables under study.

Importantly, the correlation of the variables of a complex system that define the structural
transformations obtained from multiple techniques, particularly in the present study, WAXS,
SAXS and Raman spectroscopy as a function of temperature, is crucial to relate ambiguous
or/and unclear correlations among the crystallinity, long period, conformational changes and
phase transition. Generally, data analysis is focused on an area of interest of the spectra in
which evident changes are observed and compared to other techniques data, missing hidden
cross-correlations to the naked eye in the complete range of the recorded data. Particularly,
Raman spectra analysis typically consist in the deconvolution to a Voigt function (Sundius, 1973)
(Palmö *et al.*, 1988) of the Raman band (Lorentz-Gaussian outline) pre-processed by correcting
the raw data with a baseline, smoothing and normalizing the spectra to an independent band
of the process under study. However, the analysis subjectivity introduced on the normalization
and deconvolution inaccuracy of superimposed Raman bands entails generally tedious data
treatment and previous insightful knowledge about the Raman modes. A detailed assessment
of the quantification of polybutadine content in polystyrene by Raman spectroscopy using both
band fitting or chemometrics methods has shown the simplicity to obtain higher reproducible
and consistency results by chemometric procedures. (Brun, Bourson *et al.*, 2013)

Multivariate analysis (MVA) is a chemometric procedure defined as a statistical method used
to study complex sets of data that contain more than one variable, in particular the Principal
Component Analysis (PCA). Consequently, MVA across PCA assists to extract additional information from the data, especially in the case of very complicated phenomena and/or very weak changes. MVA have been largely employed in the spectroscopic field (Chapron et al., 2019). However, it has been scarcely applied to examine different parameters obtained from scattering techniques combined other different measurement as inverse gas chromatographic (Voelkel et al., 2006), SAXS/UV-vis (Haas et al., 2014; Caetano et al., 2017), SAXS and high performance anion exchange chromatography using pulsed amperometric detection (HPAEC-PAD) (Sanderson et al., 2006). Furthermore, chemometrics models has been rarely exploited to analyze the polymer transitions upon processing conditions. The importance of the spectral data pre-treatment by the adequate algorithm to correct the experimental artefacts has been (Witschnigg et al., 2010) evaluated with different material state (melt state to semi-crystalline solid state) and polymer postprocessing procedures (cooling down, heating up, ...) on the applied chemometrics model (PCA, partial least squares (PLS). The choice of a suitable spectral range of interest to apply the chemometric model has been found to be crucial to observe negligible variations not noticeable at first sight by the naked eye. Particularly, the crystallization mechanism at quiescent conditions of molten poly(3-hydroxybutyrate) (PHB) has been analyzed in detail by MCR and both homospectral and hetereospectral 2DCOS to reveal the sequential order of the multistep crystallization procedure. (Guo et al., 2012) Accordingly, simultaneous time-resolved SAXS/WAXS experiments have been conducted at DUBBLE, using Linkam calorimeter stage in combination with Raman spectroscopy with fine temperature control in a reproducible, versatile and simple setup. The attribution of phase transition and the understanding of the crystallization mechanism has been achieved by a multi-technique approach to correlate the organization of the polymer chains during the crystallization from a molecular level to its nanostructure. The polymer crystallization thermodynamics is generally studied by thermal techniques, and thus, the temperature control under the influence of the combined X-ray beam and laser probe on the sample is crucial to affect the sensitivity the Linkam device thermocouple. Consequently, it has been designed a combined Raman spectroscopy with time-resolved SAXS-WAXS setup coupled to the Linkam calorimetry stage in which the scattering detectors and Linkam stage are controlled under a permanent connection in parallel of all detectors and the Linkam simultaneously to avoid time uncertainties. Moreover, rigorous calibrations of the Linkam calorimeter stage with the Raman laser and the X-ray beam has been conducted to compare reliably the results to a standard DSC. Finally, the data generated by the different techniques (WAXS, SAXS and Raman spectroscopy as a function of temperature) will be correlated by applying the MVA data analysis. The multitechnique simultaneous analysis of the mechanism crystallization upon cooling from the melt will be applied to a high molecular weight PVDF blended with a fluoroelastomer. A multiprobe and detailed post-analysis approach is required to monitor the structural similarity as well as the physical impediments due to the high molecular weight interfere with the identification of the complex nanostructure under development with narrow temporal transitions.

Time-resolved simultaneous SAXS and WAXS experiments coupled to a Linkam calorimeter stage have become a routine technique at the synchrotron light community. The Raman probe was fitted with a focal lens and located in an angle to avoid disturbing the optical path of the
scattering X-rays as well as reduced the air gap to minimize the background during the X-ray
acquisition (Figure 3.1A shows a global picture of the multi-technique setup, a magnification of
the sample holder is seen in Figure 3.1B). The Linkam calorimeter-stage was placed vertically
to measure by simultaneous SAXS/WAXS in transmission mode whilst the Raman probe was
located in the exit window side, measuring in reflection mode.

<figure><img src="image_29.png" /><figcaption>Figure 3.1. Simultaneous SAXS/WAXS combined to both Linkam DSC and Raman spectroscopy setup at DUBBLE (BM26b, ESRF) (A) and a magnification of the sample environment (B).</figcaption></figure>

The experimental control of the combined DSC Linkam and X-ray scattering instrumentation
has to be harmonized to assign the thermal changes to the corresponding scattering patterns
due to the different technical sensibilities. A Permanent connection in parallel of all detectors
simultaneously has been established by the Multipurpose Unit for Synchronization, Sequencing
and Triggering (MUSST) electronic module as central unit (see Figure 3.2). The MUST unit is
fitted with an internal timer that it is independent and very flexible for triggering and
synchronizing different beamline components and user hardware such as the Linkam DSC. In
addition, the MUST built-in data storage capability can be used as data acquisition unit in
combination with other techniques with different experimental parameters. Moreover, MUST
unit offers the asynchronous periodic polling up to sub-millisecond resolution over extended
period. In a typical experiment, the use has to generate a configuration file with the thermal
protocol and launch the experiment with the selected acquisition times for the detectors. The
system will write into the header of each image the corresponding thermal data as well as a file
containing the DSC data with millisecond resolution time. Furthermore, new versions of the
acquisition systems will allow the motion integration and different devices (like Linkam) into a
Gui (GDA) in order to program several devices in sequential experimental order.
<figure><img src="image_30.png" /><figcaption>**Figure 3.2.** Schema of the MUSST module as central unit in parallel of all detectors simultaneously link to the GDA Gui.</figcaption></figure>

The versatile experimental procedure has been optimised to standardize efficiently the instrumental protocol and transfer it to DUBBLE community as well as couple it to another sample environment. The versatility to combine Raman spectroscopy with other techniques compare to FT-Infrared spectroscopy is mainly due to capacity to adapt the instrumentation to optimize the Raman spectra to the sample response. The components of the Raman spectrometer such as the laser wavelength (532 nm lasers for higher sensitivity or resonance enhancement, 633 nm lasers for cost effective solutions, 785 nm lasers for sample fluorescence suppression) and power (100-400 mW) can be tuning. Particularly, the flexibility of the probe that is fitted with a few meter fiber optics for both the incident laser beam and the back-scattering Raman radiation due to the wavelength employed falls in the visible light range. Moreover, the probe can feature focal lens that allows position distantly to the Raman spectrometer and the sample, enhancing the adaptability to combine Raman spectroscopy to a broad range of sample environments. Thanks to technological innovations of the Raman spectrometer components such as the laser, Notch and Edge filters, introduction of CCD detector, the acquisitions times are faster and compatible with synchrotron measurements, a few seconds are enough for a good quality spectrum. A probe lens with a focal length of 5 cm was arranged with angle of approximately 60 °, minimizing the sample to the flight tube distance and thus, the air scattering background (see Figure 3.1). The Raman spectra acquisition was synchronized remotely from the control cabin via the spectrometer computer. Moreover, the effect of the laser on the sample temperature and DSC profile has been assessed to avoid influencing the temperature reading, which is crucial to understand polymer crystallization and correlate it with offline measurements. Generally, spectrometers feature a shutter for the incident laser to obtain a better signal/noise ratio synchronized with the shutter aperture. An oscillating DSC signal with a superposed frequency associated to the opening shutter period was obtained (see Figure 3.3 a) when the laser shutter was on due to the intermittent laser
heat impinging on the sample whilst the deactivation of the laser shutter yields a constant DSC signal (see Figure 3.3, b).

<figure><img src="image_31.png" /><figcaption>Figure 3.3. Linkam DSC profile of the cooling and second heating ramps of the PVDF sample during the simultaneous SAXS/WAXS couple to the Linkam DSC combined with the Raman spectrometer with the activated laser shutter (left) and with the disactivated laser shutter (right).</figcaption></figure>

However, an output laser power of 150 mW has been selected to minimize the laser heat transfer to the sample and a calibration of the DSC temperature has to be conducted to account for both the X-rays and Raman laser on the sample temperature. The laser wavelength of the spectrometer is also a parameter that can be tailored. Commercial Raman spectrometers fitted with lasers of 532 nm and 785 nm are available and can be switched depending on the sample response. Laser wavelength of 785 nm are generally most suitable for obtaining good quality spectra to avoid the fluorescence effect (related to the aspect/color of the samples). The aspect of most of the polymers are normally white/yellowish, leading to potentially fluorescence effects on the spectra, and thus, a near infrared (IR) laser wavelength (785 nm) have been used.

Moreover, the focus of the laser beam into the sample permits to control the volume of the sample illuminated as well as the penetration by laser Gaussian-shaped beam. The beam size and penetration of the laser are parameters required to correlate the obtained data by diverse techniques with comparable sampling methodologies, consequently extinguishing the surface effects of the sample (skin effects). The waist dimension is a beam parameter that determine the amplitude of both the electric and magnetic field a long of a Gaussian beam at the focal point. The selected focal distance of 5 cm in the current experimental corresponds to a waist dimension (Long, 1977) of d = 20 µm and L = 2.1 mm obtained by the following equation:

$$d = \frac{4\lambda f}{\pi \phi} \qquad (1)$$

$$L = \frac{16\lambda n f^2}{\pi \phi^2} \qquad (2)$$

Being d, the diameter of the convergent lens, λ the wavelength of the laser (785 nm), f the focal distance (5 cm), φ the output laser beam diameter (2.6 mm, measured by the “razor blade”)(Chaudemanche, 2013; Kimura & Munakata, 1987), L the convergent lens distance and n the refractive index (1.42 for neat PVDF), the sample thickness was 0.8 mm. The selected
laser beam parameters permitted to probe throughout the sample to suppress the skin effect and surface orientation. In addition, a blank spectrum without the laser on is acquired to avoid instrumental noise and to improve the quality of the recorded spectra. However, the absorption depth and propagation of the laser beam through the material produces undesirable secondary effects related to the material response. Mainly, heat transfer to the samples occurs, the so-called photothermal effect is generated by laser beam when the energetic excitation rate of the material is slower to the thermalization time (Brown & Arnold, 2010). Typically, polymers are characterized in lower part of the thermalization times, in the order of 10⁻⁶ seconds. Moreover, photodecomposition can be produced if the excitation energy is analog to the bond energy. Similarly, X-ray interacts with the matter (Portale *et al.*, 2016a; Bras *et al.*, 2014), producing photoelectrons that can enhance the crystallization rate of materials, the growth of nanoparticules as well as material damaged at restrained X-ray fluxes as well as increasing the temperature by 3,4 Ks⁻¹ in adiabatic approximation. A compromise between the dose of both the laser and X-ray beam and the detrimental secondary effects need to be considered to design experiments with the correct time resolution of the structural and chemical process under study. Consequently, a calibration of the sample temperature in the coupled techniques experiment was conducted to correct both the X-ray and laser beam (785 nm) effects on the DSC thermocouple reading. The melting point of different well-known standards (Nylon, indium, benzoic acid, PVDF). have been used to cover the linearity response of the instrument in the temperature range under interest (see calibration curve in Figure 3.4).

<figure><img src="image_32.png" /><figcaption>Figure 3.4. Calibration curve obtained under the simultaneous influence of both X-ray beam and Raman laser (785 nm) measured by the Linkam DSC-stage.</figcaption></figure>

Likewise, the temperature recorded in the header of every single SAXS and WAXS frame as well as Raman spectra is amended accordingly to correlate the thermodynamic transformation to the precise real temperature and to compare the measurements with other offline measurements. Moreover, the effect on the thermocouple reading and the material response of the laser beam impinging the sample has been evaluated individually to assess both the temperature fluctuation as well as the impact on the Raman spectra quality that would be related to damage effects by other techniques. The quality of the in-situ Raman spectra acquired in the experimental coupled with X-ray scattering and Linkam DSC techniques has been compared with the data obtained by combining a standard DSC and the same Raman spectrometer with a laser beam wavelength of 785 nm.
Validation of the Raman spectra quality by comparison to the DSC-Raman

A thermal protocol with identical ramp parameters has been carried out for PVDF Kynar 740 blended with a fluoroelastomer by using 3 different experimental setups to validate the temperature acquisition by either the in situ Linkam DSC combined with both Raman spectroscopy and X-ray scattering techniques or the combined DSC-Raman using a Q200 TA Instrument (Veitmann et al., 2015) and a single measurement in a Q200 TA Instrument DSC as a reference. A comparison of the amended thermograms with the corresponding calibration curve has been plotted as a function of the cooling temperature. A slight deviation has been found for (see Figure 3.5) the crystallization temperature (138.7 °C, 138.1 °C and 139.5 °C) even for the corrected temperatures of the respective different experiments. The minor temperature difference between the three crystallization temperatures can be explained by the difference of experimental error sources. The temperature discrepancy highlights the advantages of combined techniques to obtain equivalent thermodynamic conditions and compare the data to other independent measurements.

<figure><img src="image_33.png" /><figcaption>Figure 3.5. Comparison of the thermogram obtained from Linkam DSC, the Q200 TA DSC-Raman and a single Q200 TA DSC profile.</figcaption></figure>

Results

The correlation of the complementary data acquired by time resolved simultaneous SAXS-WAXS experiments coupled to the Linkam DSC stage in combination with Raman spectroscopy to study the crystallization from the melt of PVDF (Kynar 740) blended with a fluoroelastomer are briefly discussed. The nanostructure of PVDF has been largely investigated by X-ray scattering techniques as well as by Raman spectroscopy to tailor a broad range of applications in a variety of different industrial fields such as offshore pipes, construction materials, healthcare and pharmaceutics (Ameduri, 2009). PVDF is semi-crystalline polymer that crystallize in several crystalline phases defined by different chain configurations, namely α (TGTG'), β (TTTT), and γ (TTTGTTTG') phases. However, PVDF crystallize from the melt mainly in the α phase at moderate quenching temperatures and extremely high quenching temperatures or blending with other polymers such as PMMA, are required to obtain the desired electroactive β phase. PVDF has been blended with several polymers and inorganic fillers to enhance the its properties. However, the effect on the crystallization of PVDF by the addition of fluoroelastomer has slightly examined yet. The multi-technique approach to
monitoring the crystallization of the PVDF/fluoroelastomer blend benefits the interpretation of complex and faint structural transformations by the complementary simultaneous information related to different chemical sensibilities and length-scales. The β-phase quantification for instance, is a challenging task at low contents by WAXS measurements due to the broad profile of the reflections and the superposition with the amorphous phase. Similarly, Raman spectroscopy cannot rigorously quantify the amorphous phase due to overlapping with the α-phase, but is very sensitive to the β phase instead.

Moreover, a weak structural transition observed at 70 °C upon cooling has been controversially (Leonard *et al.*, 1984; Loufakis & Wunderlich, 1987; Nabata, 1990; Teyssedre *et al.*, 1993) assigned to the rigid amorphous phase densification. The mechanism of crystallization upon cooling from the melt (200 °C) to room temperature (25 °C) of the blend fluoroelastomer/PVDF the long period evolution and the crystallinity have been monitored simultaneously by SAXS (Figure 3.6A) and WAXS (Figure 3.6B) respectively, as a function of temperature. WAXS measurements have probed the transformation from the amorphous to the α-PVDF phase represented by the characteristics main reflections at the scattering values (q) of 12.6 nm⁻¹, 13.1 nm⁻¹, 14.2 nm⁻¹ and 18.8 nm⁻¹ that correspond to the (100), (020), (110) and (021) (hkl) planes of the monoclinic PVDF phase (pseudo-orthorhombic) (Lovinger, 1982), respectively. Note, WAXS requires a minimum of crystallinity (1% in polymer study (Wang *et al.*, 2000)) to produce at least one diffraction reflection. Moreover, the nanostructure parameters of the two phase lamellar system with diffuse boundaries has been The evolution of long period, crystallinity content, lamellar thickness and molecular conformation as a function of temperature during the crystallization from the melt of the blend PVDF/fluoroelastomer has been determined (see Figure 3.7).

<figure><img src="image_34.png" /><figcaption>Figure 3.6. 3D plots of the time-resolved SAXS (A), WAXS (B) patterns and Raman spectra (C) obtained in the in situ Raman spectroscopy combined with SAXS/WAXS coupled to the Linkam DSC setup, during the cooling step from the melt to RT of the blend PVDF/fluoroelastomer</figcaption></figure>
sample. Raman spectra from the Q200 TA DSC-Raman combined measurement (D) Validity of the in-situ Raman spectra quality of the blend PVDF/fluoroelastomer sample. The colour code in the 3D plot corresponds to a temperature range related to ΔT = 21.6 °C, in order to distinguish properly the evolution of the structural parameters upon temperature. The black arrows in the Raman spectra (C, D) correspond to the band position of the β-phase of the PVDF sample.

Raman spectroscopy is known to be sensible to phase variation in PVDF (α/β ratio) (Veitmann et al., 2015). Indeed, the vibrational bands at 796 cm⁻¹ and 840 cm⁻¹ of the α-phase (Kobayashi et al., 1975; Nallasamy & Mohan, 2005; Boerio & Koenig, 1971) corresponding to the rocking vibration mode r (CH₂) - vₛ (CF₂) and r (CF₃) (Yang et al., 2014), respectively, are monitored as a function of temperature (see Figure 6C). Initially, at the melt (200 °C) a broader band found slightly shifted to the α-phase, centered at 799 cm⁻¹, correspond to the rocking vibration mode of the amorphous phase of the PVDF. Note that the identification and quantification of the PVDF amorphous phase is complicated by Raman spectroscopy as the characteristic band of amorphous phase of PVDF is convoluted with the band of the α-phase at 796 cm⁻¹. However, as the temperature decreases the broad band become thinner due to standstill of the vibration group r (CH₂) - vₛ (CF₂), the transition between amorphous to crystalline phases and the phase change will be analyzed more in detail by MVA. According to the Raman spectra, the chain conformation (TGTḠ) of crystalline phase monoclinic (pseudo-orthorhombic) form polymer blend (PVDF/fluoroelastomer) remain unchanged during the crystallization confirming the β-phase absence revealed by WAXS. The integrated value of the characteristic α-vibration band (785 – 815 cm⁻¹) has been plotted (see Figure 3.7), to compare all in situ data obtained (Xc, Lp, Q and α-peak vs T). The slow increase of α-peak before Tc compare to x-ray data (Lp, Xc and Q) suggest that the crystalline phase at a crystallinity ≤ 1% (WAXS detection limits) can be detected by Raman spectroscopy.

The crystallization temperature observed to the Raman spectroscopy is about 139 °C due to a shift of the broad to thinner band center at 796 cm⁻¹ representative of the α crystalline phase, confirming the value from the Linkam DSC 138.7 °C (see Figure 3.5). Explication of the small difference will be explained in another part of the paper. In order to have the less possible difference between each machine a calibration curve have been elaborate according to literature melting point of standard materials.

Another point should be check, the quality of the in-situ Raman spectra. As explained previously, a combined DSC-Raman measurement has been conducted in order to verify this point. In Figure 3.6D, the 3D plot of the Raman spectra from this measurement done in the same condition than the in-situ Raman spectra is shown. The same trend and peaks are observed and present, respectively. The only difference is the smoothness of the spectra in the DSC-Raman spectra due to the background experiment. The data are comparable meaning that the in-situ Raman is of quality to follow the evolution of the characteristic peaks of the PVDF sample.

Further in detail SAXS analysis of the isotropic two-phase system (Ryan *et al.*, 1994), as the case for the blend PVDF/elastomer (amorphous and crystalline phase) was achieved by the invariant *Q*, also known as Porod's invariant (Glatter & Kratky, 1982). The invariant *Q* is related to the
electron density fluctuation between the two phases, during phase transformation and can be expressed by the given equation (3):

$$Q = \int_{0}^{\infty} q^{2} I(q) dq \propto 2\pi\Phi(1 - \Phi)(\rho_{c} - \rho_{a})^{2} \quad (3)$$

Where $I(q)$ is the scattering intensity, and $q$ is the scattering vector. In the second term of the equation (3), $\Phi$ and $\rho_a$ are the volume fraction and the electron density of the amorphous phase, and $(1 - \Phi)$ and $\rho_c$ are the volume fraction and electron density of the crystalline phase. The evolution of the invariant as a function of temperature (see in Figure 3.7) shows the phase transition at 140 °C (inflection point determined by derivative, available in SI) that corresponds to the crystallization of the PVDF sample, described by an increase of the electron density of the crystalline phase. Moreover, the drop of the invariant at 69–70 °C (inflection point determined by derivative, available in SI), suggest the compaction of the rigid amorphous part. The invariant is sensible in an evolving system to changes in the composition up to 50 vol. % (Xc PVDF sample max 51%, calculate from WAXS) when a maximum is reached, and inversion of the invariant will be observed.

<figure><img src="image_35.png" /></figure>
Figure 3.7. Porod's invariant, crystallinity (Xc) and long period evolution as a function of temperature during crystallization from the melt (200 °C) of PVDF sample. The dash line (≈ 145 °C) indicates the crystallization transition (Tc) and the dash line (≈ 70 °C) the transition of the rigid amorphous compaction (Trga).

Correlation data: chemometrics

The PCA is a reduction method of the original data using an orthogonal transformation to convert the data into linearly uncorrelated data, the so-called principal component (PC) to represent the largest variance possible. An advantage of the use of PCA analysis on low quality data is to be able to extract the data separating it from the noise. PCA was used to transform of the different measured variables into a data matrix that represent the relationship between crystallinity, long period and vibrational mode as a function of temperature. The data matrix is reduced by PCA to scores and loadings. The scores are the weights of the spectra/scattering profiles (here) in the different dimensions (PC), and the loadings are the representation of coefficients of the linear combination of the initial variables (here, the Raman band, or the scattering vector q, that indicate also the percentage of the variance attributed to each PC. The PCA has been applied to the crystallization from the melt to the Raman spectra and from the crystallization temperature to room temperature as the crystallization event cover the entire vector variance.
## 4-Controlling of fibrillization for drug delivery.

During my research, I have largely worked in the supramolecular field in collaboration with several groups across Europe in order to understand the structure-activity relationship and the gelation process that drives the self-assembly of potential gelators into gels. Gelation is not fully understood and the majority of gels are found serendipitously. However, it has been established that gelation is mainly controlled by the gelator design, gelator concentration, gelation conditions and solvent effects. Attempts to understand the molecular assembly of the gelators into molecular-scale fibrils and the subsequent hierarchal organization of the fibrils into spanning network that percolates the solution by fibril-fibril interaction has also been investigated. Particularly, I have investigated the important factors that control the fibrillization mechanism that in turn, govern the gelation process from a low molecular hydrogelator that forms monodispersed fibers and has been reported in the following manuscripts.

1. Francisco Rodríguez-Llansola, Beatriu Escuder, Juan F Miravet, Daniel Hermida-Merino, Ian W Hamley, Christine J Cardin, Wayne Hayes: Selective and highly efficient dye scavenging by a pH-responsive molecular hydrogelator. Chemical Communications 11/2010; 46(42):7960-2., DOI:10.1039/c0cc02338h

2. Hermida-Merino, D., Trebbin, M., Foerster, S., Rodriguez-Llansola, F. and Portale, G.; "Microfluidic Assisted Self-Assembly of pH-Sensitive Low-Molecular Weight Hydrogelators Close to the Minimum Gelation Concentration"; Macromol. Symp., (2015), 358: 59-66. doi: 10.1002/masy.

The physical characteristics of the low molecular hydrogel have triggered me to apply successfully for a visiting scientist grant in 2016 at the University of Castilla la Mancha, (1st classified) to the Pharmacy school in order to study the possibility to employ the nanofibers as nanocarriers for drug delivery applications.

Bioinspired supramolecular hydrogelators are great candidates to be employed as drug carriers as a result of their biocompatibility and capability to uptake and release specifically host molecules upon external stimuli. Drug uptake generally requires several steps and/or to covalently attach the nanocarrier resulting in inefficient industrial process to scale up. One pot passive loading of Doxorrubicin by an urea-derivative hydrogelator (HG1) was efficiently achieved by physical interactions attaining an impressive loading rate of DOX/HG1 molecules. The viability cell study confirmed the harmless impact of HG1 in healthy cells crucial to be applied as drug nanocarrier. Hydrogels were recently identified as ideal therapeutic local administration methodology in complex human anatomies after tumour resection due to their mechanical and adhesion properties. Particularly, glioblastoma multiforme is highly lethal tumour that is mainly treated by surgery resection and subsequently with local drug administration. However, recent studies showed limited therapeutic efficacy and local application of Doxorrubicine by injection of loaded hydrogel is a great candidate to enhance
the controlled drug release in the resection cavity. Molecular release profile studies monitored the Doxorrubicine regulated from the DOX discharge upon pH change from the internal fibre structure. Moreover, in vitro cells monitored the corroborated the therapeutic release efficiency whilst minimising the harmful side effects on healthy cells. Fluorescence lifetime studies on in vitro cells followed the release mechanism at the cellular scale of both DOX solution and DOX loaded in **HG1** formulations, confirming the molecular DOX release process described outside the cellular media. Fluorescence lifetime imagin probed the intracellular DOX release by both formulations (DOX solution and DOX loaded in **HG1**), highlighting the differences on the timescales and cellular pharmacological route traverse by DOX transported by **HG1**.

The design of ideal nanocarriers is a synthetic challenge that has to fulfil concomitantly ambivalent physical requirements. Bio-inspired smart materials that accomplish specific functions in complexes and dynamic environments represent a continuous synthetic challenge. The design framework required to conceive novel materials that feature targeted activities depends mainly on the structure-property relationship. Moreover, responsive smart materials to specific external stimuli involve subsequent structural developments in a dynamic media that introduce higher complexity.

Hydrogels are typically generated by the fine balance of self-assembled amphiphile monomers that can interact through covalent bonds, by noncovalent interactions or by a combination of both. The driven force that defines the internal structure of physical hydrogels together with the aqueous media equilibrium determines also the supramolecular assembled structures generated such as fibres, micelles, or ribbons. Among all, hydrogels based on small molecules or low-molecular-weight-hydrogelators (LMHG), are remarkable candidates to design responsive materials due to the dynamic nature of the non-covalent interactions that can be typically triggered by external stimuli (temperature, concentration, pH, light, salts concentration, etc).¹⁻⁷ Particularly, the structural complexity requisite of drug nanocarriers arise from the sequential multi functions achieved involved in different chemical media. The specifically drug upload must be efficient an, limiting the synthetic steps as well as avoiding the pharmaceutical compound damage. Moreover, the nanocarrier transport should enable the drug transference through the cellular membrane. Furthermore, the drug administration route define the initial nanocarrier mechanical properties. In addition, the therapeutic drug release must satisfy a regulated kinetic profile that minimise the pharmaceutical dosage and target the tumour cells. Importantly, pH responsive physical hydrogels combine the reversibility interaction capacity to specifically load and upload the therapuatic agents at human revelant pH and offers the mechanical characteristics to attain complex human anatomies to target explicit tumours.

Particularly, Glioblastoma multiforme (GBM) is the most common and aggressive of the primary brain tumours developed in adults. GBM appears firstly on the glial human cells and infiltrates diffusely into the unaffected surrounding parenchyma. Despite the relatively low incidence, 3.19 per 100,000 people per annum, the median survival of GBM patients is only approximately
12-18 months, and merely 10% of them currently survive for more than 5 years. The early and distant dissemination of malignant cells renders GBM a surgically incurable neoplasm.⁸ Aggressive resection surgery increases survival time in patients, however, recurrences occur 1-1.5 years after initial therapy.⁹,¹⁰ After resection, GBM patients generally follow the standard treatment: external beam radiation around the original tumour plus oral temozolomide (TMZ) daily.¹¹

Among the main challenges of GBM treatment, the difficulty of drug delivery to the tumour results in an insufficient amount of the most drugs at the site of action (f). The presence of the blood brain barrier (BBB) constitutes the main obstacle for the systemic treatment of brain tumours in general and GBM in particular. An increasing demand urges to develop novel strategies able to both enhance the drug efficiency and reduce the adverse side effects without affecting normal tissue(g).

Currently, novel therapeutic agents fail to show clinical relevance in the treatment of GBM, due to their toxicity or chemoresistance and therapeutic approaches focus on local delivery of therapeutic agents after surgery resection rejection emerges as an effective alternative.¹² Local therapeutic agent delivery provides tumour cells targeted exposure to high drug concentrations, minimising systemic healthy tissue damage as well as preventing the growth of cancer cells unattainable by surgery resection (j). Specially, FDA has approved the application of Gliadel® wafers as post-surgery medical care, supporting the therapeutic local intersticial treatment of antitumour drug to avoid the cancer recurrence.¹³ However, Gliadel® has shown poor and controversial results compared to the standard medical treatment from marketing approval.¹³,¹⁴ Particularly, the drug loading and the rigidity of the wafer structure limits the therapeutic efficiency into the anatomy of the resection cavity. Hydrogels have been largely proposed as candidates for drug delivery systems and potentially substitute films due to their mechanical properties. Indeed, hydrogels can be sufficiently loaded with the appropriate drug therapeutic dose and administered locally in the ¹⁶brain after a craniotomy via intracerebral implantation injection. Moreover, hydrogels have enough adherence capacity to stick to the cavity borders, covering an optimized area under treatment. Hydrogels were employed as drug delivery systems in the treatment of GBM, and in particular, a PLGA-based hydrogel, OncoGel™ was successfully tested as an adjuvant to radiation therapy.¹⁷,¹⁸ Recently, photopolymerizable hydrogels enable to TMZ delivery have shown a significantly decreased in the tumor growth in mice treated compared to controls.¹⁹,²⁰ In addition, a GBM xenograft model have shown the inhibition of tumour growth by the application of Doxorubicin (DOX) loaded in Mebiol™Gel, a thermoreversible polymeric gelator which is biocompatible and non-cytotoxic.²¹,²² Doxorubicin (DOX) is successfully employed in the treatment of several solid malignant tumours such as X.²³ However, high doses of DOX were required to be systemically administered to exert a therapeutic benefit.²⁴ However, the pharmacologic DOX systems currently under study feature drug carrier limitations such as low drug loading, short retention time, and thus, a rapid burst release was typically found.²⁴,²⁵ Moreover, DOX is unable to penetrate efficiently the central nervous system (CNS) due to the blood brain barrier restraining the application for GBM treatment (i).
Previously, a non-peptidic mono urea derivative hydrogelator (HG1) displaying the ability to intercalate selectively planar and positively charged guests into its negative charged structure was proven to load DOX (Figure 4.1a). Moreover, HG1 possesses a significant low minimum gelation concentration, indicating the percolation of the network formed by the self-assembled HG1 aggregates at low monomer ratio. The triggering of the highly monodisperse HG1 fibrilization is achieved by a fine pH change which maintains a low supersaturation point. The well-controlled uptake mechanism of guest molecules during fibrillation is an key HG1 feature to both diminish the uptake process time and enhancing the uptake capability which preserves the mechanical properties.³ In general, HG1 combines ideal physicochemical characteristics to be a promising drug delivery system for local therapeutic agent delivery in the GBM treatment. In the present work, the therapeutic transport performance of HG1 as a nanocarrier for the local DOX delivery in the treatment of GBM is assessed. Importantly the planar molecular structure of DOX facilitates the molecular encapsulation by the intercalation with HG1 molecules due to structural similarity and the possibility to interact through non-covalent bonds. Moreover, the passive loading of the hydrogel, the temperature dependence of the formulations as well as the interaction of DOX molecules with the hydrogel monomers within the network were evaluated. The DOX passive uptake mechanism and kinetics during the HG1 fibrillation process is controlled by the slow acidification produced by the hydrolysis of glucocono-δ-lactone (GDL). The controlled pH change by the GDL addition to the HG1 solution generates an unperturbed homogeneous nanofibrillation that promotes the uniform drug uptake/release throughout the entire volume at the nano length scale. Moreover, the biocompatibility of unloaded HG1 was tested in non-tumoral cells. Glioma cell lines C6 and U87 were selected as in vitro models for GBM and cytotoxicity of DOX-loaded HG1 was evaluated. Remarkably, the in vitro responsive behaviour to pH stimuli of DOX release profiles were found to be comparable to pharmacological standards at acidic and physiological values. Finally, DOX release profile in the glioma cell line C6 as well as the DOX distribution inside cells was evaluated by confocal and fluorescence lifetime microscopy. Notably, DOX-loaded HG1 showed a selective cytotoxic profile against glioblastoma cells without affecting the viability of non-tumoral cells.

<figure><img src="image_36.png" /><figcaption>**Figure 4.1.** Molecular structures of a) HG1, b) Doxorubicin</figcaption></figure>

Typically, the poor mechanical properties of physical hydrogels in the swollen state limit the proliferation of drug release devices generated by low molecular weight hydrogelators. The drug uncontrolled release is generally arising from the weak interactions with the matrix scaffold generated by the physical network crosslinks of the hydrogelator assemblies.
Moreover, the drug vehicule system mechanical properties and thus, the efficiency of the drug release is intimately related to the drug uptake method. In an active loading drug uptake method, the drug is discharged from the hydrogel internal structure and diffuses on the damaged supramolecular network where remains entrapped. The main advantage of the active loading is the stability of the mechanical properties, avoiding the disruption of the previously formed network. However, a low drug loading capability is typically achieved along with large uptake process times due to the infiltration difficulty of the guest drug on the formed molecular aggregate. In contrast, the drug guest is entrapped during the network formation process in a passive loading, facilitating the drug uptake and increasing its loading efficiency. However, the passive loading is competitive assembly process that usually result in to detrimental poor mechanical properties of the loaded hydrogel²⁶ due to the acting as chain stoppers in physical hydrogels. The hydrogel **HG1** fibrillization mechanism was found to be controlled by the slow acidification that mantains constant a low supersaturation point. Similarly, a cooperative assembly between the drug and the hydrogelator is expected to yield also highly monodisperse nanofibers, however, affecting the kinetics of fibrillization. Importantly, the DOX addition to the **HG1** basic solution is expected to affect the GDL hydrolysis that affects the rate to attain the pKa of the monoprotonated **HG1**.

## 1. Sol-gel phase transition: fibrillation process mechanism.

The mechanism of drug uptake/release of **HG1** has been monitored at different length scales by SAXS, and FTIR to understand the key parameters that govern the passive/active loading of the DOX guest molecules by the nanofibers formed by **HG1**. The loading capacity and release/uptake profile will depend on the guest molecule incorporation method to the nanofiber carrier. The guest molecule structural features such as the molecular rigidity and planarity, as well as the interactions nature with the hydrogelator will determine the absorption method, either decorating the surface or being assimilated in the internal nanostructure. The hydrogel network formation upon pH change was monitored by SAXS from the self-assembled precursors originated from the dissolved molecules (DOX and **HG1**) at the supersaturation point to the complete percolation of the aggregates. Previously, the gelation mechanism of **HG1** was followed by multi-length scale approach, performing several techniques such as SAXS, rheology, FTIR and NMR. The initial **HG1** aggregates were found to arise from firstly the hydrophobic pockets resulted from the monoprotonation of the carboxylic acid group with the highest pKa, 6.4. The primary **HG1** concentration determined the formed aggregates cross-section and the subsequent enlogantion through cooperative hydrogen bonds between the nitro-urea groups (see scheme 4.2). Highly monodisperse nanofibers were formed by **HG1** under 15mg/ml characterised by a core-shell hollow cylinder shape ($r_{int}$1.4 nm and $r_{shell}$ca 1.6 nm). Importantly, the supramolecular polymerization of **HG1** was controlled maintaining a low supersaturation point by the pH change through the whole volume and reaction time.

SAXS experiments were performed to probe the gelation process in the presence of the DOX guest molecules during the passive loading at two distint uptaking concentrations, specifically 3 mg/ml and 12 mg/ml of DOX (see Figure X). The DOX loading concentration assessment aimed
to identify the role of DOX on the fiber structure as well as to perturbate the initial thermodynamic conditions. In principle, the monomer avability concentration at the onset of the nucleation at the supersaturation point would be affected by the different concentration of DOX. Furthermore, the DOX occurrence disturbed the pH change by altering the hydrolysis equilibrium of GDL and thus, the growth mechanism.

Similarly to the previous nanofibrillization of **HG1**, a rod-like structure with a core-shell cylinder shape described the self-assembled aggregates morphology of DOX in **HG1** hydrogel. The rapidly length growth of the rod-like self-assembly suggests a nucleation-elongation mechanism likewise the described by **HG1** fibers. Moreover, fibers with nearly structural parameters were obtained by both DOX loading concentrations in **HG1**, increasing only the polydispersity cross-section (see Figure 4.3). A simultaneous density core thickening with a polydispersity (see Figure 4.3 and table 1) enhancement indicates that the infiltration of DOX molecules in the internal fiber structure disrupted slightly the self-assembly motif whiles retaining the main physical interactions. Moreover, the high loading capacity of DOX by the **HG1** assembly manifested the cooperative involvement of DOX in the internal fiber structure, either likely due to hydrogen bonding or π-π interactions. However, the analogous cross-section for both DOX loading concentrations (3 mg/ml and 12 mg/ml)

and **HG1** nanofibers imply that the nucleation occurrence of **HG1** first determine the fibre cross-section. Subsequently, the DOX molecules compete with **HG1** in the fibre elongation process resulting in a slowdown of the self-assembly mechanism as well as the worsening of the interaction motif. However, the supramolecular interactions nature of the driven force during the elongation fibre development offered the flexibility to adapt the molecular assembly pattern. Moreover, the pH controlled change governed the monomer concentration availability in the fibre tip, avoiding the branching structure resulted from concentrations gradients.

Furthermore, the so-called Porod invariant (equation 1) ascribed to the total scattering intensity, Q, is related to volume fraction ($\phi$) transformation of the electron density ($n_e$) for bi-phasic systems and importantly, is independent of the scatter morphology. In the **HG1** system under supramolecular polymerization, the invariant is a powerful tool that affords to monitor the reaction kinetics and to assess the influence of DOX on the fibrillization mechanism.

$$Q = \int_{0}^{\infty} I(q)q^{2}dq \propto \langle n_{e} \rangle^{2} \phi_{1}\phi_{2} \quad \text{equation 1}$$

with $\langle n_e \rangle^2$ being $(\rho_{gel}-\rho_{sol})^2$ and $\phi_i$ the volume fraction of the **HG1** solution and **HG1** gel phases with time (see Figure 4.3) affords the assessment of the morphologic transformation.
<figure><img src="image_37.png" /><figcaption>Figure 4.2. Time resolved SAXS profiles a) DOX (3 mg/ml) loaded in HG1 b) DOX (12 mg/ml) loaded in HG1. 3) Comparison of SAXS profiles at 9000 s. d) SAXS invariant of</figcaption></figure>

DOX (3 mg/ml, 12 mg/ml) loaded in HG1 and HG1.

The SAXS invariant trend corresponding to either the passive loading of both DOX formulations (3 mg/ml, 12 mg/ml) in **HG1** compare to **HG1** gelation confirmed that DOX modified completely the kinetics of fibrillization. Recurrently, the linearization of the Avrami method to the **HG1** fibrillization confirmed the formation of aggregates in 1D (n~1)in agreement with the SAXS model with different velocity constants. However, the application of the Avrami model to the passive loading of both formulations proved the the polydisperisty increase (n~1,2, see Figure SX) in the uncomplete reaction under study. Remarkably, the mechanism of the nanofibre formation was found to depend simply on the initial **HG1** concentration despite the large presence of guest molecules that strongly interact through multiple weak interactions. Particularly, the nucleation step generated the fibre cross-section precursors that subsequently were elongated with the participation of the guest molecules, retarding the kinetics of fibrillization.

The nanofibers structure of the guest molecules uptake process during the active loading.

Moreover, frequency oscillation sweep of the hydrogels generated by the corresponding DOX
loading formulations (3 mg/ml and 12 mg/ml) under a non-destructive range (see Figure X)
were performed to probe the fibrillar network. The characteristic independent response for the
storage and loss moduli with frequency of the gel state was described for both DOX formulations (see Figure X) in **HG1**. Moreover, the both the storage and loss components featured similar modulus for both DOX formulations and **HG1** hydrogels related to fibrillar networks with similar strength. The virtual parallel rheological properties with absence of network relaxation in the frequency range under study is associated with the lack of fibre branching and with a density of crosslinks resulted from the long **HG1** fibres, in agreement with the morphology described by SAXS. Furthermore, the tan δ obtained (measured at 1,3 Hz of the dynamic stress) indicated a higher strength of the DOX loaded hydrogel compared with **HG1** (tan(δ) **HG1** = 3.6 tan(δ) DOX loaded in **HG1** = 8.7). However, the elastic component (G') was found continuously higher for **HG1** and the loss modulus component of DOX loaded in **HG1** slightly decreased, indicating larger resistance to flow.

<figure><img src="image_38.png" /><figcaption>**Figure 4.3.** a) frequency sweep of hydrogels formed by the both DOX passive loading formulations in **HG1** (3 mg/ml and 12 mg/ml) and **HG1** hydrogel passive loading DOX at pH 5 (black squares) and 7.4 (red circles).</figcaption></figure>

Moreover, the yield stress of the hydrogel network DOX (12 mg/ml) loaded in **HG1** and **HG1** was evaluated by a dynamic stress sweep to determine the point of the fibrillar network disruption. The yield stress is a typical mechanical parameter defined for nanostructurated liquids which is related to the processability, performance and long term storage material capability. The yield stress is largely employed in different fields such as food science, healthcare, paint applications and pharmaceutical as indicate the network stability to phase separation, sedimentation, aggregation, precipitation and resistance to vibration during transport. Particularly, the yield stress of the DOX loaded in **HG1** is a key parameter to assess the applicability of as injectable pharmacological hydrogel. The yield stress of the hydrogels under study featured a narrow transition (ca 10 Pa) with a decrease of both storage and loss modulus of 4 orders of magnitude. The yield stress was found around 22 Pa for both the DOX (12mg/ml) loaded in **HG1** and **HG1**, that confirmed the stability of the hydrogels network, similar for instance to the toothpast (10 Pa).

## 2. Molecular Release behaviour with pH.
Monitorization of the pH-dependent DOX release from the **HG1** nanofibers after the passive loading of DOX (12 mg/mL) were conducted by fluorescence spectroscopy (see Figure 4.4 a and b). The pH functional activity of an ideal nanocarrier must target a tight pH range to deliver the drug while retaining the interaction with the drug at acid pH, depending on the therapeutic application methodology. The DOX release from the **HG1** at both pH 5 and pH 7.4 was selected to mimic the media conditions at appropriate physiological conditions. The pH in the body controls important functional activities and ranges from the stomach with high acid media to the intestine featuring slight basic conditions. However, human blood is characterised by a pH 7.35/7.45 gradient and cancer cells in particular, is found to be in general slighter basic. The lack of fluorescence emission by **HG1** allows to follow the release mechanism of the guest molecules only. Only a 9% of DOX was released after 120 min at pH 5 (see Figure 3) while it a further increase in pH (pH = 7.4) produced a small further release of DOX (11%) after the releasing time (120 min). Moreover, the DOX release followed a slight different mechanism, either a diphasic or triphasic profile at pH 5 and pH 7.4, respectively. A slow diffusion from the hydrogel matrix with a linear zero-order release profile was found at pH 5 (see Figure 4.4). However, the DOX release at pH 7.4 suggested an initial rapid burst (phase I), diffusion from **HG1** matrix (phase II), and zero-order erosion stage (phase II) (see Figure 3). The release profile suggest that the DOX trapped in **HG1** matrix at pH 5 is more stable than at pH 7.4, minimizing the erosion stage and attaining easily a zero-order release. The percolated network is originated at acidic conditions due to the high hydrophobicity of the protonated **HG1** species that forms the assembly precusors at the pKa of the monoprotonated carboxylic species (pH=6.7). Moreover, the fibrillization driven is force is the hydrogen bonding interactions that are maintaining acid media. The profile release suggest a dilution effect of the gel network instead of disruption of the **DOX-HG1** interactions. Moreover, the lower percentage of DOX released by **HG1** at pH 5 and 7.4 compared to methylene blue indicates that DOX participate in the inner fiber structure interacting strongly with the likely hydrogen bonds.

<figure><img src="image_39.png" /><figcaption>Figure 4.4. a) % Release profile of passive loading DOX in HG1 at pH 5 (black squares) and 7.4 (red circles), b) DOX fluorescence spectroscopy during the release from HG1 at pH 7.4.</figcaption></figure>
Moreover, the centre of the DOX emission fluorescence band remained constant during both the uptake and release processes, (see Figure 3 b and SX in ESI), indicating the lack of π-π stacking interactions between the DOX and **HG1**. The DOX fluorescence band was simply shifted by the acid-base reaction due to likely dissociation of alcohol of DOX in basic media. However, the large quenching effect on the fluorescence intensity on both media, either in the DOX loaded in **HG1** (acid media) or the DOX disolved with **HG1** in basic media, revealed the strong interaction between the DOX and **HG1** molecules (see Figure 4.3) before and after the fibrillization process. In addition, the concurrence of the DOX fluorescence band for either the DOX in solution, or DOX dissolved together with **HG1** in basic media highlighted that DOX probed an analogous aqueous microenvironment (see Figure 4.3). Furthermore, the DOX fluorescence band monotonic increase together with the coincidence of fluorescence band position during the DOX release from the **HG1** fibrillar structure (see Figure 3b) confirmed the identical molecular DOX discharge mechanism to the upload process by **HG1**. Similarly to previous described by SAXS and rheology, the fluorescence spectra of DOX loaded in **HG1** indicated that DOX was placed inside the fibers with water molecules surrounding in a equivalent fashion than the DOX dissolved in water.

### 3. Biological studies

**Cell viability studies.** The effects of **HG1** on mitochondrial function and cell viability were investigated by determining the percentage of 3-(4,5-dimethylthiazol-2-yl)-2,5-diphenyltetrazolium bromide (MTT) transformed (% MTT transformed). MTT is metabolized at the mitochondrial level by viable and metabolically active cells to give an insoluble coloured formazan product. Insults that alter the metabolism of the cells affect the rate of MTT reduction into formazan. The reduction in the percentage of MTT has been largely employed as an index of cellular viability as the MTT transformation correlates with a reduction of the general metabolism (n).

The cell viability of the nanocarrier excipient **HG1** as a control test, has shown the innocuous effect on glioblastoma cells at the concentration range under study (% MTT transformed were 97.2 ± 2.7 and 96.3 ± 1.9, in C6 and U87 cells respectively) confirming the inert nature at typical therapeutic concentrations/doses and that acts as a mere vehicle for the therapeutic agents.
<figure><img src="image_40.png" /></figure>

Figure 4.5. MTT assay in Mouse fibroblast C3H10T1/2 cell were treated with different concentration of HG1 hydrogel 1 (0-1000 µM) for 72h.

Moreover, C6 and U87 cells were treated at increasing concentrations (0.01-1 µM) of free DOX and DOX-loaded **HG1** for 72 hours to determine the percentage of MTT transformed in culture media with a pH X. The treatment of C6 and U87 glioblastoma cells with free DOX or DOX-loaded in **HG1** reduced the glioblastoma viability in a concentration-dependent fashion (see Figure 4.6), producing higher potent effect in C6 cells compare to U87 at lower concentrations (see Figure 4a). Interestingly, in human glioblastoma cells, DOX-loaded in **HG1** displayed a potency very similar to free DOX, featuring nearly identical $IC_{50}$ values (see Figure 4.56.b), whilst the inhibition % in rat glioblastoma C6 cells treated with DOX-loaded in **HG1** resulted systematically lower (see Figure 4.6a).

a)

<figure><img src="image_41.png" /></figure>

b)

<figure><img src="image_42.png" /></figure>

Figure 4.6. Cytotoxic effect of vehicle(HG1), DOX, and DOX-loaded in HG1 on rat glioblastoma C6 (a) and U87 human cells (b). $IC_{50}$ values, expressed in micromolar, are included in the corresponding graph. Data are expressed as mean ± s.e.m. from at least three independent experiments.
Moreover, the side-effect of DOX-loaded in **HG1** on non-tumoral cells have been assessed by using a primary culture of rat astrocytes as the background of C6 rat glioblastoma cells. Rat astrocytes were incubated with free-DOX or DOX-loaded in **HG1** for different times (24-72 h) to determine the inhibition percentage of MTT. Remarkably, the rat astrocytes cellular viability was unaffected by the treatment with DOX-loaded in **HG1** at the studied incubation times (see Figure 4.7), whilst free-DOX induced a drastic reduction in cellular viability, attaining a reduction in the inhibition of MTT transformed below to 50% after 72 h of incubation (see Figure 4.7).

<figure><img src="image_43.png" /><figcaption>Figure 4,7. Cytotoxic effect of free DOX and DOX-loaded in HG1 on rat astrocytes. Data are expressed as mean ± s.e.m. from at least three independent experiments.</figcaption></figure>

In summary, **HG1** combines the therapeutic properties of an ideal excipient due to its inert nature together with an efficient upload and smart controlled release. **HG1** is a great candidate to develop novel drug formulation nanocarrier systems for sustained DOX delivery in the treatment of GBM, with an adequate safety profile against non-tumoral cells.

### In vitro release of DOX-loaded in HG1 hydrogel in culture media.

The time-course for the DOX level remaining in the culture media of the treated C6 glioblastoma cells either with free-DOX or DOX-loaded in **HG1** has been quantified by (m). Free-DOX was rapidly incorporated by the C6 cells and diminish monotonically with the incubation time, almost disappearing from the culture media within 24h (see Figure 4.8), in agreement with previously found in cells X(m). However, the amount of free-DOX in the culture media of the C6 cells treated with DOX-loaded in **HG1** in culture media slightly increased monotonically from a negligible DOX concentration to a plateau-like range at 9h of treatment (see Figure 4.8). Importantly, the free DOX levels in the culture media of both therapeutic administration methodologies are comparable from the 9 hours onwards of the treatment. The controlled DOX release generated by the **HG1** gel avoids the initial drug burst release stage. The lack of high DOX concentrations in contact with the cells across the delivery process by **HG1** gel minimize the toxic side-effects on healthy cells, as previously observed in the treatment of healthy rat astrocytes (see Figure 4.8).
<figure><img src="image_44.png" /></figure>

Figure 4.8. Time-course study of the percentage of DOX remaining in the culture medium of C6 cells treated with 0.5 µM DOX or 0.5 µM DOX-loaded in HG1. Data are expressed as mean ± s.e.m. from at least three independent experiments.

Release Dynamic of DOX-loaded in HG1 hydrogel by a passive loading formulation in culture cells.

<figure><img src="image_45.png" /></figure>

Figure 4.9. Dynamics of free DOX in buffer solution (red circles) and DOX-loaded in HG1 (black square) by a passive loading formulation.

The intracellular DOX release dynamics in culture C6 cells treated with either a PBS buffer solution of DOX and DOX-loaded in **HG1** at identical concentration (12 mg/mL) was monitored by following the DOX fluorescence lifetime. Remarkably, the fluorescence lifetime is a photophysical state function and thus, independent on the experimental conditions (excitation wavelength, light pulse duration, photobleaching, etc...) as well as the fluophore
concentration and fluorescence intensity. The fluorescence lifetime spectroscopy is a non
destructive technique applied to diverse media (solution, gel, films, bulk) related to nanosecond
and/or picosecond structural dynamics and interactions of the fluorophore that is ideal to
monitor the molecular dynamics in complex systems such as the cells. The time interval of a
fluorophore in the excited state is affected by quenching the fluorescence response due to
internal emission pathways such as conformation changes, side-chain rotation as well as
external irradiative relaxation routes such as interactions between adjacent groups, polarity,
solvent relaxation, temperature and media viscosity. Firstly, the fluorescence decay from the
excited state of the exclusive fluorophore (DOX) of the delivery system in both formulations is
studied by time-resolved fluorescence spectroscopy (TRF), to probe the DOX molecular
environment in either solution (PBS) or within the internal fiber structure. A mono-exponential
lifetime value of 1.02 ns for the DOX applied in solution was found, consistent with previously
reported lifetime values (see Figure 4.9)²⁷⁻³⁰ which DOX-loaded in **HG1** presented a
biexponential decay with an average lifetime of 0.47 ns. Specifically, the biexponential decay
showed a slow component of 0.93 ns (with an amplitude of 5%) related to DOX molecules at
the vicinity of the fiber, and a fast component of 0.20 ns (95% of amplitude)corresponded to
molecules likely trapped inside the fibres (see Figure 4.9). The molecular sensitivity of DOX to
changes in the local chemical environment affects its fluorescence response, decreasing
significantly the lifetime of the DOX trapped either in the neighbouring or the internal structure
of **HG1** fibre. Importantly, the amplitude/ratio of the DOX fluorescence decay which is
proportional to the corresponding DOX molecular state, confirmed a highly effective molecular
upload by the **HG1** fibers. The fluorophore probes two different molecular environments
featuring lower fluorescence lifetimes that the unassembled DOX in solution. Moreover, the
shorter fluorescence lifetime at the greater viscosity of the **HG1** hydrogel, where the
fluorophore is expected to possess a higher fluorescence lifetime due to a molecular
rigidification. The dramatic quenching effect of the DOX fluorescence lifetime and intensity in
the hydrogel is likely arising from the interaction with **HG1** or charge/proton transfer with the
donor or acceptor groups such as the nitro, amino, phenyl moieties.

Moreover, fluorescence lifetime was recorded by FLIM on C6 glioblastoma cells treated with the vehicle (HG1, control experiment), a DOX solution(PBS), or DOX-loaded in HG1 for different incubation times to probe the cellular DOX location. Particularly, the intracellular media at two different regions of interest, cytoplasm and nuclei, was studied by fitting the fluorescence lifetime to a biexponential decay analysis (see Figure 4.10 and Table 1). A similar decay patterns in both the cytoplasm and nuclei was found for both therapeutic formulations (DOX in solution applied to the C6 cells and DOX-loaded in HG1). A mono-exponential lifetime of 1.45 ns was obtained at nuclei during the incubation times under study, shifting to larger emission lifetimes the typical DOX value in solution (c.a 1.02 ns) due to higher molecular rigidity. Moreover, a slow component of 3 ns as well as a fast component of 0.46 were measured at the cytoplasm, suggesting that DOX molecules were located in areas with different local microviscosity and/or chemical environment (see table 1). The relation between amplitudes, A(%), of both lifetimes remains constant for the incubation times under study and for both formulation methods (DOX solution applied to the C6 cells and DOX-loaded in HG1). The absent
of new fluorescence lifetimes, at least above the FLIM resolution time 0,2 ns, in the DOX-loaded in **HG1** compare to DOX in solution, in the intracellular release dynamics suggested that DOX cross the cell without the vehicle **HG1X**.

<table><caption>**Table 1.** Biexponential analysis of both formulations; DOX solution and DOX-loaded in **HG1** in the nuclei and cytoplasm of cells at different incubation times.</caption><thead><tr><th></th><th></th><th>time (h)</th><th>τ<sub>av</sub>(ns)</th><th>τ<sub>1</sub>(ns)</th><th>A<sub>1</sub>(%)</th><th>τ<sub>2</sub>(ns)</th><th>A<sub>2</sub>(%)</th></tr></thead><tbody><tr><td rowspan="6">Nuclei</td><td rowspan="3">DOX</td><td>3</td><td>1.48 ± 0.18</td><td>1.48 ± 0.18</td><td>100</td><td>-</td><td>-</td></tr><tr><td>6</td><td>1.49 ± 0.22</td><td>1.49 ± 0.22</td><td>100</td><td>-</td><td>-</td></tr><tr><td>24</td><td>1.44 ± 0.20</td><td>1.44 ± 0.20</td><td>100</td><td>-</td><td>-</td></tr><tr><td rowspan="3">DOX-<br>Loaded</td><td>3</td><td>1.45 ± 0.28</td><td>1.45 ± 0.28</td><td>100</td><td>-</td><td>-</td></tr><tr><td>6</td><td>1.44 ± 0.22</td><td>1.44 ± 0.22</td><td>100</td><td>-</td><td>-</td></tr><tr><td>24</td><td>1.51 ± 0.22</td><td>1.51 ± 0.22</td><td>100</td><td>-</td><td>-</td></tr><tr><td rowspan="6">Cytoplasm</td><td rowspan="3">DOX</td><td>3</td><td>1.95 ± 0.24</td><td>3.03 ± 0.40</td><td>18.2</td><td>0.49 ± 0.22</td><td>81.8</td></tr><tr><td>6</td><td>1.92 ± 0.24</td><td>2.91 ± 0.36</td><td>19.1</td><td>0.45 ± 0.24</td><td>80.9</td></tr><tr><td>24</td><td>1.86 ± 0.28</td><td>2.97 ± 0.44</td><td>24.6</td><td>0.44 ± 0.16</td><td>75.4</td></tr><tr><td rowspan="3">DOX-<br>Loaded</td><td>3</td><td>1.60 ± 0.16</td><td>2.80 ± 0.18</td><td>18.1</td><td>0.43 ± 0.24</td><td>81.9</td></tr><tr><td>6</td><td>1.88 ± 0.20</td><td>3.04 ± 0.40</td><td>18.6</td><td>0.47 ± 0.20</td><td>81.4</td></tr><tr><td>24</td><td>2.06 ± 0.20</td><td>3.36 ± 0.36</td><td>15.4</td><td>0.43 ± 0.14</td><td>84.6</td></tr></tbody></table>

<figure><img src="image_46.png" /></figure>
Figure 4.10. a)FLIM images of C6 cells treated with 0,5 µM of a DOX solution and DOX-loaded in HG1(a). Lifetime histograms for DOX applied in solution (b) and DOX-loaded in HG1 (c). Deconvoluted histogram at 6h in DOX-loaded HG1 treatment.(d)

Moreover, fluorescence lifetime imaging were reconstructed from the average lifetimes on the C6 glioblastoma cells treated with both therapeutic formulations (a DOX solution(PBS), or DOX-loaded in **HG1**) for different incubation times to highlight the DOX intracellular delivery. Firstly, FLIM image of control cells showed a very slight autofluorescence with an average lifetime value of 2.4 ns (see Figure S9 in ESI). The fluorescence lifetime distribution after 1h of treatment with the DOX solution showed that the fluorescence signal is mainly located in the cytoplasm as well as the nuclei membrane with an average lifetime value of 2.2 ns. Moreover, the DOX fluorescence signal tend to be displaced towards the nuclei along with DOX solution treatment time, and within 3h, DOX was found practically only in the nuclei. Importantly, the fluorescence lifetime decreased to 1.5 ns typically attributed to DOX-DNA intercalation,²⁷,²⁸,³⁰ remining at the nuclei for the incubation time under study (24 h). However, the fluorescence intensity of the C6 cells treated with DOX-loaded in **HG1** decreased due to the quenching effect of the interaction with **HG1**. As we commented above, it is only observed fluorescence lifetime of DOX when it is released. Initially, the fluorescence signal was concentrated in the cytoplasm with a lifetime of 2.5 ns at 1 h of DOX-loaded in **HG1** treatment. Similarly to the DOX solution formulation, the fluorescence signal shifted to nuclei membrane within 3h, decreasing the average lifetime decrease to 2 ns. Moreover, the DOX fluorescence signal appeared both at the nuclei membrane and nuclei at 6 h of DOX-loaded in **HG1** and, a further decreases of the lifetime fluorescence to 1.5 ns. However, the DOX fluorescence signal increase the average lifetime and shifted back to the nuclei membrane at 9 h of treatment and even to the cytoplasm after 24 h. The unusual DOX delivery behaviour is in agreement with the controlled release profile observed in "in vitro" and culture media for the DOX-loaded in **HG1** formulation (9 hours ~40% of release, see Figures 4.11b). Moreover, the lack of fluorescent activity avoided to monitor the **HG1** molecules during the DOX transport in the intracellular media. However, **HG1** detered the initial DOX burst release stage, reducing the high concentrations of DOX in contact with the cells compared to the DOX in solution formulation.

<figure><img src="image_47.png" /></figure>

<figure><img src="image_48.png" /></figure>
Figure 4.11. Lifetime fractions (%) from deconvoluted histograms extracted from FLIM images in Figure 8a, b, and c: a) buffer solution of DOX formulation, b) DOX-loaded in HG1 treatment. Errors bars represent the standard deviation as 2σ.

The delayed DOX cell penetration process transported by **HG1** was assessed by deconvolution of the Gaussian-like lifetime histogram, assigning the DOX average lifetimes values for cytoplasm, nuclei membrane, and nuclei to 2.5, 2.0 and 1.5 ns, respectively. The ratio of DOX molecules presents in the different areas of the cell (see Figure 4.11) was inferred from the relative areas of deconvoluted curves at each average lifetime (see Figures 4.10b, c, d). Importantly, DOX easily penetrated to the nuclei membrane and was then readily incorporated into the nuclei. However, the DOX infiltration was significantly slower for the DOX-loaded in the gel treatment and from 9 h onwards DOX mainly was found in the cytoplasm due to the critical release.

## CONCLUSION

The mono urea non-peptidic derivative hydrogelator (HG1) has been evaluated as a biocompatible and nontoxic carrier to local delivery of DOX in the treatment of GBM. Moreover, the remarkable controlled DOX uptake achieved by the cooperative nanofibrillization with HG1 highlighted the key role of the specific inter-supramolecular interactions between the DOX-hydrogelator to maximize the loading rate. In addition, the maintaining of the supersaturation point low, across the nanofibrillization and DOX uptake concomitant by the regulated pH change avoided the gradient molecular concentrations and consequently, the fibre branching. The inter molecular co-assembly mechanism resulted in long fibres, circumventing the detrimental chain stopper behaviour during the fibrillization generally found for guest molecules. The monitoring of fibrillization by SAXS showed the slight in the polydispersity increase of the fibre cross-section compared to HG1 fibres, however, retaining the mechanism of nucleation and elongation. Moreover, the hydrogel network formed with DOX featured analogous strength and yield stress to the HG1 gel, indicating a similar fibrillar percolation due to comparable fibre length and crosslinks. The release of DOX from HG1 is pH dependent due to the solubility of HG1 in basic media, yielding a controlled release over time. DOX-loaded in HG1 reduced the glioblastoma viability in a concentration-dependent manner, resulting in slightly IC50 lower values than direct treatment with a DOX solution. Moreover, the application of DOX-loaded in HG1 formulation avoided undesirable side effects over healthy cells at identical concentrations of DOX released after 24 hours of treatment with both formulations. Microscopy studies confirmed the DOX controlled release observed in the in vitro studies, with a slow DOX penetration in the DOX-loaded in HG1 treatment. Remarkably, HG1 hydrogel is a great candidate as local delivery super-nanocarrier. Particularly, DOX-loaded in HG1 hydrogel is an ideal alternative to the market therapeutic treatments to be applied into the tumour resection cavity. The HG1 hydrogel acts as DOX reservoir, releasing from the hydrogel matrix upon pH change and minimal systemic toxicity.
Moreover, I have been invited to give an oral contribution to the following conferences:

* Polymer Networks & Gels 2014 Conference, Correlating molecular structure to the mechanism of gelation of low molecular weight hydrogelators. Tokyo, Japan

* 2015 International Chemical Congress of Pacific Basin Societies Understanding the fibrillation of self-assembled low molecular weight hydrogelators. Honolulu, USA.

* XVII International Small Angle Scattering Conference - SAS2018XVII International Small Angle Scattering Conference - SAS2018, Investigation of the uptake/reléase during the fibrillization of self-assembled low molecular weight hydrogelators by SAXS/GISAXS. Traverse City, USA.

* 32nd Advanced Materials World Congress Conference organized by International Association of Advanced Materials entitled ‘Monitoring of the drug uptake/release by a self-assembled low molecular weight hydrogelator carrier’. Sydney, Australia.

Currently, we are working on several publications that will report the promising results as drug
nanocarriers as well as uptake/release mechanism.
# Selective and highly efficient dye scavenging by a pH-responsive molecular hydrogelator†

Francisco Rodríguez-Llansola,ª Beatriu Escuder,*,ª Juan F. Miravet,*,ª

Daniel Hermida-Merino,ᵇ Ian. W. Hamley,ᵇ Christine J. Cardinᵇ and Wayne Hayes*ᵇ

Received 2nd July 2010, Accepted 9th September 2010

DOI: 10.1039/c0cc02338h

**A structurally simple low molecular weight hydrogelator derived from isophthalic acid forms robust pH-responsive hydrogels capable of highly efficient and selective dye adsorption.**

Molecular gels represent an intriguing type of organic nano-
structured soft material. Molecular gelators form elongated
supramolecular aggregates that further evolve to nanofibrillar
networks which entrap the solvent to yield a gel.¹ This type of
soft matter is of interest in areas such as responsive materials,
catalysis or electronic and photonic materials.² Practical
applications of molecular hydrogelators are especially appealing
due to the biological relevance of water as a solvent. In recent
years several structurally diverse molecular hydrogelators have
been described.¹/*g A topical aspect within this field is the stimuli
responsiveness associated with molecular gels. In addition to
their characteristic temperature and concentration dependence,
the formation of molecular gels can be controlled by different
stimuli. In particular, the presence of pH sensitive groups in the
structure of the gelator can be exploited advantageously for the
controlled formation of molecular gels.¹g

Herein we report a structurally very simple, pH-responsive
bisaromatic hydrogelator (1) that is derived from isophthalic
acid and contains a urea functional group. The molecule
reported is a remarkable example of a molecular hydrogelator
capable of selective and extremely efficient incorporation of
dyes into the formed hydrogel network. In this paper we
additionally report on the fascinating layered solid state
structure of the gelator 1 which provides a framework to
account for the dramatic dye uptake capabilities of this urea
in its hydrogel form.

The interaction of molecular gel fibers with different guest species has been reported.³ The incorporation of dyes in hydrogels formed by a tripeptide derivative which forms gels at basic pH values has been described revealing that molecular hydrogels can potentially be used to remove industrial dyes from waste water.⁴ª The use of gels of ionic liquids for dye removal has also been reported recently.⁴ᵇ

The bisaromatic hydrogelator 1 was prepared in high yield
by direct addition of 5-aminoisophthalic acid to 4-nitrophenyl
isocyanate and exhaustively characterized (see SI). Hydrogels
were formed above a concentration of 0.9 mM (0.3 mg mL⁻¹)

by acidification with aqueous HCl of a solution of **1** in basic
water and sonication. Alternatively, for a slow and homo-
geneous acidification of the system, the hydrolysis of water
soluble glucono-δ-lactone can be used.⁵ Furthermore, the
gelation in neutral water was also possible in the presence of
small amounts (ca. 10% v/v) of polar organic solvents such as
methanol or DMSO when accompanied by gentle heating until
complete dissolution was attained and the solution was then
allowed to cool to 25 °C. The gels were found to be signifi-
cantly more stable upon increasing the concentration of
gelator (see SI, Fig. S1). Remarkably, the hydrogels were
stable at temperatures above the boiling point of water when
the gelator concentration was higher than 9 mM.

The gelation process by pH tuning provoked a colour
change from red-orange to yellow associated with the proto-
nation of the carboxylate groups (see Fig. 1). The observed
yellow colour in the gel samples was attributed to intermolecular
interactions. Solutions of compound **1** in MeOH are colourless
but upon increasing the proportion of water in the system,
aggregation takes place giving rise to the observed yellow
colour which correlates with the observed absorption above
400 nm (see SI, Fig. S3 and S4).

The morphology of the gel obtained by pH-tuning was
observed by Atomic Force Microscopy (AFM) to reveal the
presence of a network of nanofibers that were several microns
in length and less than 50 nm in width (Fig. 2).

Needle-like crystals of **1** suitable for single crystal X-ray
analysis were obtained from slow evaporation of a solution of

<figure><img src="image_49.png" /></figure>

<figure><img src="image_50.png" /><figcaption>**Fig. 1** (a) Pictures of vials containing a 3 mM aqueous solution of compound **1** (left) and the hydrogel formed upon acidification (right). (b) UV-vis spectra of compound **1** (1.5 mM) at different pH values.</figcaption></figure>

<footnote>ªDepartament de Química Inorgànica i Orgànica,
Universitat Jaume I, Avda. Sos Baynat s/n, 12071 Castelló, Spain.
E-mail: escuder@uji.es, miravet@uji.es</footnote>

<footnote>ᵇDepartment of Chemistry, University of Reading, Whiteknights,
Reading, UK RG6 6AD. E-mail: w.c.hayes@reading.ac.uk</footnote>

<footnote>† Electronic supplementary information (ESI) available: Experimental details and crystal structure description. See DOI: 10.1039/c0cc02338h</footnote>
<figure><img src="image_51.png" /><figcaption>**Fig. 2** Atomic force microscopy images of xerogels obtained from the hydrogel formed by **1** (0.9 mM). Left: phase contrast, right: topography.</figcaption></figure>

compound **1** in a 7:3 water:methanol mixture. The bisaromatic hydrogelator **1** crystallized in the monoclinic system with a dramatic layered structure. These layers are constituted by ribbons of molecules formed by intermolecular hydrogen bonding between the urea NH units and nitro moieties.⁶ The layer is constructed by connection of the ribbons by intermolecular hydrogen bonding among the carboxylic acid functional groups at the ribbon edges (see SI, Fig. S13–S15).

The X-ray powder diffraction pattern of a xerogel from a hydrogel of **1** revealed the crystalline nature of the fibrillar network (SI, Fig. S6) and showed that the arrangement found in the fibers is a polymorph of that described in the crystal structure mentioned above, which presents different diffraction peaks.

FT-IR studies unveiled that the xerogels formed by **1** presented a 2D arrangement of intermolecular hydrogen bonds related to that found in the crystal structure. The FT-IR and ATR grazing angle spectra showed that the carbonyl stretching band intensity was reduced significantly when the sample was irradiated in grazing angle mode (see Fig. 3a). According to previous studies,⁷ these results indicate a preferred orientation of the studied hydrogen bonding groups in a plane, suggesting that the hydrogen bonding array is oriented along the axis of a nanofiber.

Therefore, a model is proposed for **1** in which the π-π-stacking and hydrogen bonding interactions are orthogonally related with weaker aromatic interactions that lie perpendicular to the fiber axis (Fig. 3b). In the model, gel fibers can be seen to result from frustrated crystallization caused by the partial ionization of carboxylic acid units that would preclude the formation of extended 2D hydrogen arrays such as those

<figure><img src="image_52.png" /><figcaption>**Fig. 3** (a) Transmission vs. grazing angle ATR FT-IR spectra for xerogels of compound **1**. (b) Proposed orientation of the hydrogen bonding interactions within the fibers.</figcaption></figure>

<figure><img src="image_53.png" /><figcaption>**Fig. 4** Pictures of the methylene blue removal process. (1) Solution of methylene blue in water (500 mL, 8 mg mL⁻¹) and vial containing a hydrogel of **1** (2 mL, 20 mM). (2) Solution of methylene blue with the hydrogel fragments deposited at the bottom of the beaker. (3) Picture of the system after 72 h.</figcaption></figure>

observed in the crystal structure and favour the formation of elongated assemblies.

Interestingly, the crystalline layered structures have in common an inherent ability to mimic clays by intercalation of guest molecules.⁸ With this concept in mind, the potential for incorporation of guest molecules into the hydrogel fibers was tested. A preliminary study revealed that these hydrogels were very efficient in the adsorption of methylene blue. As shown in Fig. 4, a very small quantity of hydrogel (2 mL, 20 mM) is able to completely decolour a 500 mL solution of methylene blue with a dye concentration of 8 mg L⁻¹. It has to be noted that the adsorption of methylene blue into molecular hydrogels is not a general property. For example, adsorption of the dye was not observed when a hydrogelator that did not feature aromatic units that was available in our laboratories was studied (see SI, Fig. S8).

In a more detailed study, the adsorption of six different compounds into the gel was studied (see Table 1) including two planar, positively charged aromatic compounds that are used commonly as dyes, methylene blue and violet blue 2B. Additionally, two cationic aromatic compounds, 1-pyrene methylammonium hydrochloride and dopamine, and an anionic dye, indigo carmine, were used as possible guests. The last substrate studied was spermine, an aliphatic tetracationic substrate that is known to interact with negatively charged DNA strands. For the adsorption studies, solutions containing the different guests were deposited over a hydrogel obtained by pH-tuning. Since spermine does not exhibit a strong UV-vis chromophore, the adsorption was monitored by ¹H NMR spectroscopy for all of the above dyes.

Noticeably the hydrogel formed by compound **1** was extremely selective in the adsorption of the positively charged

<table><caption>**Table 1** Adsorption of different species (SI, Fig. S7) on the hydrogel formed by **1**ª</caption><thead><tr><td>Adsorbate</td><td>% Adsorbed</td></tr></thead><tbody><tr><td>Methylene blue</td><td>98</td></tr><tr><td>Methyl Violet 2B</td><td>97</td></tr><tr><td>1-Pyrenemethylamine</td><td>62</td></tr><tr><td>Spermine</td><td>&lt;1</td></tr><tr><td>Dopamine</td><td>&lt;1</td></tr><tr><td>Indigo carmine</td><td>&lt;1</td></tr></tbody></table>

<footnote>ª Determined by ¹H NMR. A solution in D₂O of these species (7.25 mM, 1 mL) with DMSO as internal standard was deposited for 48 h on 1 mL of the hydrogel formed with D₂O (14.5 mM). The structure of the adsorbates is depicted in the SI.</footnote>
dyes methylene blue and methyl violet 2B. When a solution of these dyes (7.25 mM, 1 mL) was deposited on 1 mL of the hydrogel (14.5 mM), the dyes could not be detected by NMR spectroscopy in the original solution after 48 h. The fact that the aliphatic tetracation spermine was not adsorbed at all under the same conditions supports the hypothesis that intercalation constitutes the main driving force for the aggregation rather than purely electrostatic interactions. Negligible adsorption into the gel was measured for positively charged dopamine or anionic indigo carmine. In contrast, 62% of the initial concentration of 1-pyrenemethylamine was adsorbed on the hydrogel fibers. The selectivity observed in the interaction with methylene blue and methyl violet 2B in comparison to the other aromatic substrates reveals that the adsorption is also selective to the nature of the aromatic moiety.

For a solution of the dye deposited over a hydrogel formed in a vial the maximum amount of dye adsorption and the kinetics of the process are very dependent on the vial dimensions, namely, on the contact area between gel and solution (see SI, Fig. S9). These results indicate that for this experimental setup a diffusion controlled adsorption is taking place. For example, for vials with a diameter of 2.7 cm quantitative dye removal was observed in less than 6 h. Importantly, the removal of dyes from solutions was tested taking advantage of the pH-responsive nature of hydrogel formed by **1**. For example, when aqueous HCl is added to a solution of methylene blue and the gelator a gel is formed. Filtration of the system afforded a colourless solution and, according to UV-vis spectroscopy, the dye was removed very effectively (>>99%) (see SI, Fig. S11). Similar results were obtained for methyl violet 2B but for indigo carmine this procedure was significantly less efficient. Noticeably, the dye removal efficiency using this procedure is unprecedented—the dye-hydrogelator ratio was almost equimolar in the formed gels (see SI, Table S1). The measured removal efficiency parameter determined by UV-vis spectroscopy for methylene blue and methyl violet 2B was *ca.* 800 mg of dye per gram of gelator, a significant enhancement when compared to the values described for related systems. For example, the tripeptide based hydrogel aforementioned exhibited values of 7.36 mg g⁻¹ for Rhodamine B and 10.56 mg g⁻¹ for the dye direct red 80.⁴ª The dye removal efficiency is also notable when compared to the adsorption of methylene blue to a common adsorbent such as active charcoal which exhibits an efficiency of up to *ca.* 400 mg g⁻¹.⁹

Regarding the interaction of the gelator and the dyes, a direct proof of the intimate interaction taking place in this system is the fact that the thermal stability of the gels was significantly affected by the presence of the dyes (see SI, Fig. S10).

Finally, gel regeneration can be achieved easily. For example, after repetitive extractions with chloroform to remove the dye, the gel can be regenerated by dissolution in basic water followed by acidification.

In summary, a new simple urea-based compound capable of forming robust, pH-responsive hydrogels with outstanding dye scavenging properties has been described. It is to be noted that the hydrogels are comprised of a self-assembled nano-fibrillar network as revealed by microscopy studies. Interestingly, in this case a crystal structure of the gelator can be obtained

which serves as a reference for a possible structural arrangement in the gel nanofibers which, as shown by X-ray diffraction studies, possess a microcrystalline structure. A related hydrogen bonding array was inferred for the xerogel by grazing angle IR spectroscopy. A remarkable fact is that the nanofibers interact selectively with positively charged dyes such as methylene blue and methyl violet 2B. The nanofibrillar structure provides a stimuli responsive functional material which presents a very high aspect ratio. The pH responsiveness of the system gives an outstanding added value for potential practical applications. Dye solutions at basic pH containing the ionized, soluble, hydrogelator can be transformed easily into hydrogels upon acidification. Filtration of these systems affords colorless dye-free solution. With this procedure the hydrogels present an unprecedented capacity of dye removal with *ca.* 0.8 g of dye being captured by 1 g of the gelating component of the hydrogel. The extraordinary dye removal capabilities were found to be specific for this particular gelator and were attributed to its acidic character and its aromatic units. In addition to employment in industrial scale tasks such as waste water treatment, applications for hydrogels of this type in controlled drug release (especially gene therapy agents such as DNA intercalating drugs) are also envisaged.

Work at Reading was supported by EPSRC grant EP/G026203/1 to WCH and IWH. Work at Castelló was supported by MCI, grant CTQ2009-13961 and Universitat Jaume I, grants P1-1B2009-42 and P1-1B2007-11. F.R-L. thanks Generalitat Valenciana for a FPI fellowship.

## Notes and references

1. (a) *Molecular Gels: Materials with Self-Assembled Fibrillar Networks*, ed. R. G. Weiss and P. Terech, Springer, 2005; (b) P. Terech and R. G. Weiss, *Chem. Rev.*, 1997, **97**, 3133; (c) D. J. Abdallah and R. G. Weiss, *Adv. Mater.*, 2000, **12**, 1237; (d) J. van Esch and B. L. Feringa, *Angew. Chem., Int. Ed.*, 2000, **39**, 2263; (e) O. Gronwald and S. Shinkai, *Chem.-Eur. J.*, 2001, **7**, 4328; (f) L. A. Estroff and A. D. Hamilton, *Chem. Rev.*, 2004, **104**, 1201; (g) M. De Loos, B. L. Feringa and J. H. van Esch, *Eur. J. Org. Chem.*, 2005, 3615; (h) A. R. Hirst and D. K. Smith, *Chem.-Eur. J.*, 2005, **11**, 5496.

2. (a) A. R. Hirst, B. Escuder, J. F. Miravet and D. K. Smith, *Angew. Chem., Int. Ed.*, 2008, **47**, 8002; (b) S. Banerjee, R. K. Das and U. Maitra, *J. Mater. Chem.*, 2009, **19**, 6649; (c) B. Escuder, F. Rodríguez-Llansola and J. F. Miravet, *New J. Chem.*, 2010, **34**, 1044.

3. (a) M.-O. M. Piepenbrock, G. O. Lloyd, N. Clarke and J. W. Steed, *Chem. Rev.*, 2010, **110**, 1960; (b) P. Mukhopadhyay, Y. Iwashita, M. Shirakawa, S.-I. Kawano, N. Fujita and S. Shinkai, *Angew. Chem., Int. Ed.*, 2006, **45**, 1592; (c) B. Escuder, J. F. Miravet and J. A. Sáez, *Org. Biomol. Chem.*, 2008, **6**, 4378; (d) U. Maitra, S. Mukhopadhyay, A. Sarkar, P. Rao and S. S. Indi, *Angew. Chem., Int. Ed.*, 2001, **40**, 2281.

4. (a) B. Adhikari, G. Palui and A. Banerjee, *Soft Matter*, 2009, **5**, 3452; (b) S. Dutta, D. Das, A. Dasgupta and P. K. Das, *Chem.-Eur. J.*, 2010, **16**, 1493.

5. D. J. Adams, M. F. Butler, W. J. Frith, M. Kirkland, L. Mullen and P. Sanderson, *Soft Matter*, 2009, **5**, 1856.

6. S. George, A. Nangia, C.-K. Lam, T. C. W. Mak and J.-F. Nicoud, *Chem. Commun.*, 2004, 1202.

7. S. E. Paramonov, H. Jun and J. D. Hartgerink, *J. Am. Chem. Soc.*, 2006, **128**, 7291.

8. M. J. Zaworotko, *Chem. Commun.*, 2001, 1.

9. S. Mukherjee and S. Bhattacharya, *J. Am. Chem. Soc.*, 1949, **71**, 1725.
5-Caroline Grant (from Oct 1st 2017 to Nov. 30th 2020) in the project "NODENS" which is a Caroline Fellowship program co-founded by European Union Horizon 2020 Marie Skłodowska-Curie actions and Irish Research Council.

In order to continue the investigation on the gelation mechanism, I successfully obtained a joint application for Caroline Grant with Bing Wu, to spend 36 months (from Dec. 1st 2017 to Nov. 30th 2020) working as a guest researcher in the project "NODENS" which is a Caroline Fellowship program cofounded by European Union Horizon 2020 Marie Skłodowska-Curie actions and Irish Research Council.

The post-doctoral fellow, Bing Wu, has generated novel materials in the first part of the project that has been investigated by SAXS measurements. I have successfully applied for beamtime to both ESRF and ALBA in order to understand the structure development in situ. I have estimated the value on beamtime by the number of awarded shifts (5500€ per shift).

* 2018093178 - Study of the nanostructure double network formation of a modified polypeptide derivativ Alba sincrotrón 2019 - 49.500,0 €

* SC-4908- In-Situ SAXS-WAXS Analyses of 'Microgel' Gel Structure in UV-cured Polyacrylate Systems, ESRF, 2019 - 49.500,0 €

* 2018022840 - Study of the structure of a microgel formation mechanism during the UV-curing of PEG-based polyacrylate. Alba sincrotrón 2018 - 49.500,0 €

* 26-02-900-Structural investigation of the double network formation of a modified polypeptide derivatives. ESRF 2018 -33.000,0 €

* 26-02-899-Investigation of the nanostructure development of a PEG-polyacrylate during UV curing, ESRF, 2018 - 49.500,0 €

We are working on the data analysis and I am focussing on the post-doctoral formation to analyze the SAXS data. I have promoted the post-doctoral fellow to apply for a world-recognized course on synchrotron techniques 'Hercules'. Bing Wu will attend it in March 2020 and we are working on 2 manuscripts on his work that are under preparation.

We have also successfully applied for beamtime for a project on (MA-4398-)In-situ SAXS-WAXS Analyses of 3D-Graphene-Nanocomposite Supercapacitor in the charging-discharging cycle at ESRF (49.500,0 €). We have recently published the obtained results:

* Cuili Xiang, Yin Liu, Ying Yin, Pengru Huang, Yongjin Zou*, Marcus Fehse, Zhe She, Fen Xu, Dipanjan Banerjee, Daniel Hermida Merino, Alessandro Longo, Heinz-Bernhard Kraatz, Dermot F. Brougham, Bing Wu*, Lixian Sun, ACS Appl. Energy Mater. 2019, 2, 5, 3389-3399.
6-Scientific granted project by Fapesp (Sao Paulo Federal Research Council) for the investigation of 'In situ studies of nanoparticles nucleation and growth assisted by simultaneous SAXS and WAXS and XAS in combination with UV-visible spectroscopy', 2017,

I have applied together with José Fernando Queiruga Rey from the Federal University ABC to the Fundação de Amparo à Pesquisa do Estado de São Paulo (FAPESP), for a project entitled 'Estudo de nucleação e crescimento de nanopartículas assistido simultaneamente por técnicas de Espalhamento de Raios X a baixos e médios ângulos (SAXS/ WAXS) e espectroscopia de absorção de Raios X (XAS) em combinação com Espectroscopia UV-Vis'. The project was awarded with 30.000,0 € in order to cover the subsistence as well as the traveling and accommodation of the associate Professor, José Fernando Queiruga Rey. The project has been developed at DUBBLE at ESRF from the 21/08/2017 to the 20/08/2018 under my supervision. Moreover, the project has been actively continued as a result of the promising results obtained by several accepted proposals. A Master thesis has been focused on the understanding in detail of thermal activated nucleation and growth process of CsHSO₄ in Nafion by in situ combined SAXS-WAXS/EXAFS experiments. Moreover, a Master thesis has been established to study the nucleation and growth of gold nanoparticules in mild conditions to incorportate them in a future step in hydrogels for developing drug delivery systems with sensing properties.

SAXS-WAXS/ EXAFS in situ studies of thermal activated nucleation and growth process of CsHSO₄ in Nafion

The study of metal-organic hybrid systems has been initially selected as it provides a novel approach to combine the mechanical support that offers the polymeric material with the improvement of targeted properties. Nafion has been largely employed as solid electrolytes in Polymer Electrolyte Membranes (PEM) fuel cell applications in the recent years. However, intermediates temperatures (<90 °C), relative high humidity and pure hydrogen as fuel are required to yield high conductivity performances. Hybrid materials have been suggested to be ideal candidates to improve the conductivity performance of Nafion membranes at higher temperatures and lower humidity conditions whilst maintaining the chemical stability and mechanical strength that is particularly significant to operate above 100 °C. Hybrid materials are generally generated by the incorporation of hydrophilic inorganic nanoparticles in the polymer structure. The precursor solution impregnates the Nafion membrane followed by a thermal treatment that induces the nucleation and growth of the nanoparticules inside the membrane. The nanostructure of Nafion-Cs nanoparticles (N-CSP) composites has been studied to enhance the proton conductivity at working conditions. The main advantage of hybrid membranes is the maintenance of high protonic conductivity at low humidity conditions. The objective of the synthesis of Nafion-Cs nanoparticles (N-CSP) composites is to increase protonic conductivity in relation to Nafion under anhydrous conditions. CsHSO₄ is an excellent protonic
conductor at high temperatures and low relative humidity when compared with pure Nafion. The addition of Cs nanoparticles to the Nafion above the percolation fraction may promote an increase in the conductivity of such composite films and allow the operation of direct ethanol fuel cells (DEFC) without gas humidification. The presence of large amounts of water in the DEFCs reduces the gases diffusion in the electrodes of the cell and limit the operating temperature of the DEFC in T <100 °C. The nanomorphology of Nafion membranes has been studied in details in the last years, explaining the relationship between morphology and the high-water diffusion coefficient and protonic charge carriers. A detailed study of nanomorphology of the hybrid membranes can play a fundamental role in the design of novel materials.

The thermal activated nucleation and growth process of CsHSO₄ nanoparticles have been studied by in situ combined X-ray absorption fine structure (XAFS) with Small Angle X Ray Scattering (SAXS) measurements at Cs edge. The combined study approach permitted to determine both the chemical nature of the main species in the solution as well as to monitor their morphological evolution in situ.

A detailed study of nanomorphology of the hybrid membranes can play a fundamental role in the design of novel hybrid materials. Preliminary results (see Figure 6.1) shown an improvement in the conductivity of the hybrid membranes followed by changes at the Nafion morphology, as the ionomer peak is more resolved than the pure Nafion counterpart. Previously, it has been hypothesized the internal distribution of the Cs nanoparticles within the Nafion channels.

<figure><img src="image_54.png" /><figcaption>Figure 6.11. a) SAXS heating profiles of Nafion membranes impregnated with Cs precursor solution; b) SAXS isothermal profile showing the time nanoparticles grown process; c) Conductivity of Hybrid nafion nanoparticles at different temperatures.</figcaption></figure>

Moreover, simultaneous XAFS/SAXS measurements performed at bm26a/DUBBLE (see Figure 6.2a) during the heating step have shown that the ionomer peak is modified during the growth of (Figure 6.2b) the Cs nanoparticles (Figure 6.2c), showing that the thermal process play an important role in the nucleation and growth of the nanoparticles.
<figure><img src="image_55.png" /></figure>

<figure><img src="image_56.png" /><figcaption>Figure 6.2. a) SAXS-XAS setup at bm26a in the ESRF; b)SAXS heating profiles of Nafion membranes impregnated with Cs precursor solution; c) EXAFS measurement of the element Cs analyzed on the K-shell</figcaption></figure>

During the project, I have successfully applied for beamtime to ESRF in order to study the
structure of the N-CSP composites.

* 26-01-1159-study of in situ thermal activated nucleation and growth process of CsHSO4 nanoparticles by simultaneous XAFS/ SAXS measurements, 2018. ESRF, 49.500,0 €

The preliminary measurements performed have lead to the following publication:

1. Bruno R. Matos, Rodolfo Politano, José Fernando Q. Rey, Daniel Hermida-Merino, Ulrich Schade, Ljiljana Puskar, Fabio C. Fonseca: *Interplay of α/β-Relaxation Dynamics and the Shape of Ionomer Building Blocks*. Scientific Reports 12/2018; 8(1)., DOI:10.1038/s41598-018-31368-8

We are currently working on the in-depth structural characterization and more manuscripts are
expected to be published in the next future as well as the Master thesis of Victor Cruz de Souza.

Furthermore, I have continued the collaboration with José Fernando Queiruga. Rey at Universidade Federal do ABC by co-directing the Master thesis of Murillo Donizeti dos Santos Rosa entitled ‘Nucleação e crescimento de nanopartículas de ouro incorporado em hidrogel’.

Nanostructured materials with metallic nanoparticles, particularly gold nanoparticles, have been probed their applicability in a wide range of fields such as biomedical materials. Herein, the nucleating and grown mechanism of plasmonic nanoparticles produced by an environmental friendly chemical route have been resolved by in-situ SAXS/ UV-vis spectroscopy.

Likewise, I have firstly successfully applied to Alba sincrotrón for the investigation of gold nanoparticles in mild conditions by simultaneous SAXS/UV spectroscopy experiments.

* 2018093066 - Investigation of the nucleation and growth mechanism of gold nanoparticules by in situ UV-vis spectroscopy and SAXS. date:14/07/2019. 49.500,0 €, Alba sincrotrón
Moreover, I have developed a novel cell that allows monitoring the growth of the gold nanoparticles in flow to minimize the X-ray influence by simultaneous SAXS- UV-vis spectroscopy in order to correlate the changes on the plasmon with the morphological evolution (see Figure 6.3).

<figure><img src="image_57.png" /><figcaption>Figure 6.3. a)SAXS/UV spectroscopy setup at bm 11 in Alba. b) A magnification of the flow closed-loop system circulated with the help of a peristaltic pump.</figcaption></figure>

Applications of nanostructured materials in advanced materials and biomedicine have been extremely important at the development of nanoscience and nanotechnology. Metallic nanoparticles, like gold and silver nanoparticles, are very stable and shows to be very interesting due to their electronic, magnetic and optical properties are related with the size (quantum effects). Metallic nanoparticles like gold and silver nanoparticles are interesting since they catalytic and biological properties can change the materials applications. Nowadays, synthetic routes achieve with high efficiency nanoparticles with tuning sizes and shapes at high production rates. However, the development of biocompatible chemical routes that retained the synthetic efficiency and nanoparticles stability are required to meet environmental policies. The use of biomolecules as reducing and/or stabilizing agents in a chemical or photochemical synthesis are ideal candidates to follow environmental friendly approaches. The effect of tetrachloroauric acid speciation during the nanoparticles synthesis in the replacement of Cl by H₂O and D₂O at the gold in aqueous medium has been studied. Thus, the process of gold speciation was subjected to the differential capacity of exchange of the deuterium dioxide associated with the gold ion by water molecules of the reaction medium. The gold speciation will be promoted by the pH variation at the gold salt precursor solution. Likewise, the gold nanoparticles could be chemically obtained by the mixture of 0.45 mM gold salt solution (AuP) at a pH 10 buffer HEPES 2.5mM /sodium phosphate (B1) or HEPES 30 mM/sodium phosphate (B2). The HEPES buffer features a piperazine molecule which acts either as reductor and stabilizing agent since it can produce free radicals at the central nitrogen able to reduce the gold ion.

The in-situ measurements were carried out during the nucleation and grown of gold nanoparticles by the addition of 0.45 mM of gold salt solution (H₂O or D₂O) in a reductor medium (HEPES (0.25mM or 30mM) + sodium phosphate) with pH 10. According to the following scheme: AuC₄⁻+AuCl₃(OH)⁺+AuCl₃(H₂O) and AuC₄⁻+AuCl₃(OD)⁺+AuCl₃(D₂O) both
solutions with proportion 5:2:3. The precursor solution and buffer were mixed in a controlled way by the stopping flow device and injected in a capillary to follow the reaction from the beginning.

<figure><img src="image_58.png" /></figure>

Figure 6.4 a) Uv-Vis absorption spectra of a solution mixture of HAuCl₄ and HEPES with different times ([AuP = 0.45 mM, H₂O]/([B1 HEPES = 2.5 mM, D₂O]). b) UV-vis absorption spectra of a solution mixture of HAuCl₄ and HEPES with different times ([AuP = 0.45 mM, D₂O])/([B2 [HEPES = 2.5 mM, D₂O)]). c) SAXS experimental data of the nucleation and growth mechanisms of a solution mixture of HAuCl₄ and HEPES ([AuP = 0.45 mM, D₂O])/([B1 HEPES = 2.5 mM, H₂O])

The nanoparticule growth has been followed by SAXS measurements as well as the simultaneous monitoring of the absorption gold band at different conditions will allow us to understand the effect of the chemical environment can affect the process of nucleation and growth process to obtain a particular morphology. A detailed analysis of both the maximum absorption band of gold nanoparticles at different times of D₂O/AuP suspensions in B1 Buffer (see Figure 6.4a) as well as in H₂O/AuP suspensions in B2 Buffer (see Figure 6.4b) with the SAXS experients (see Figure 6.4c) during the nucleation and grown process will permit us to correlate the plasmonic properties with the nanoparticules dimensions. Moreover, the fitting of the SAXS profiles will be used to estimate the size and size distribution of the gold nanoparticules.

References

1. Colvin, V.L., *The potential environmental impact of engineered nanomaterials*. Nature Biotechnology, 2003. **21**(10): p. 1166-1170.

2. Daniel, M.C. and D. Astruc, *Gold nanoparticles: Assembly, supramolecular chemistry, quantum-size-related properties, and applications toward biology, catalysis, and nanotechnology*. Chemical Reviews, 2004. **104**(1): p. 293-346.

3. Sarikaya, M., et al., *Molecular biomimetics: nanotechnology through biology*. Nature Materials, 2003. **2**(9): p. 577-585.

4. Gurunathan, S., et al., *A green chemistry approach for synthesizing biocompatible gold nanoparticles*. Nanoscale Research Letters, 2014. **9**(1): p. 9-248.
SCIENTIFIC REPORTS

<figure><img src="image_59.png" /></figure>

OPEN

# Interplay of α/β-Relaxation Dynamics and the Shape of Ionomer Building Blocks

Received: 18 May 2018

Accepted: 13 August 2018

Published online: 07 September 2018

Bruno R. Matos<sup>1</sup>, Rodolfo Politano<sup>1</sup>, José Fernando Q. Rey<sup>3,4</sup>, Daniel Hermida-Merino<sup>4</sup>, Ulrich Schade<sup>2</sup>, Ljiljana Puskar<sup>2</sup> & Fabio C. Fonseca<sup>1</sup>

The relation between the α/β relaxations and the shape of the building blocks of ionomer materials is a key factor for programming an important temperature-dependent property: the memory of shape. However, the morphology of ionomers is indirectly obtained via modeling of small angle X-ray scattering (SAXS) data owing to the hardly accessible image characterization of the nanometric building blocks – micelle-like cylindrical polymeric aggregates (radius ~2–6 nm and length >100 nm). Herein, broadband dielectric spectroscopy (BDS) measurements, free from electrode polarization effects, allowed identifying the time and temperature dependence of the polarization of different length scales of the ionomer matrix, and more importantly, by directly providing the aspect ratio of the radius and length of the polymeric aggregates for each desired temperature. This finding is essential for controlling the shape of ionomer based functional products under several stimuli conditions, thereby advancing remarkable applications, such as four dimensional (4D) printing.

Ionomers are considered shape memory polymers due to the ability of fixing, or programming, two or multi macroscopic shapes on different types of stimuli, such as, temperature and moisture content<fnref n="1" />. Such adaptive materials are envisioned as a revolutionary concept and a massive world trend of production of functional parts<fnref n="1" />,². Associated with additive manufacturing, the shape memory polymers have the ability of reconfiguring the matrix structure to a temporary desired shape according to the ambient stimulus, adding a new dimension to the final product: time. The advance of shape memory ionomers is valuable for several high-performance applications such as artificial muscle design, electrocatalyst binders, polymer electrolyte fuel cells, and four dimensional (4D) printing<fnref n="1" />⁻⁵.

Perfluorinated ionomers (PIs) are a specific family of ionomeric materials that displays outstanding physical properties, such as, thin-film nanoconfinement, exceedingly high ion transport, elevated mechanical resistance, and memory of shape³,⁶,⁷. All of these properties are known to be linked to the PIs morphology. The PIs morphology has been extensively modeled by SAXS in which the building blocks of their microstructure are assigned to long cylinder like polymeric aggregates (r ~ 2.5 nm and L > 100 nm), adjacently packed together mainly by temporary ionic crosslinks and a low degree of crystallinity⁵,⁶. Many efforts have been made to characterize the shape of such polymeric aggregates by image characterization techniques⁸,⁹. The most promising results have been obtained by transmission electron microscopy (TEM) and atomic force microscopy (AFM)⁸,⁹. However, due to instability of the sample during TEM measurements and the interaction of the sample with the tip in AFM, most of the available data display artifacts and are not representative of the PI morphology⁸,⁹. Moreover, by means of such techniques, the in situ temperature characterization of the polymeric aggregates is very problematic. This issue is aggravated if the morphology information as a function of relative humidity is required. Such drawbacks have been overcome with the use of small angle X-ray scattering (SAXS) in which the ionomer morphology at different sets of temperature and relative humidity can be indirectly obtained via modeling⁹,¹⁰. In addition, usually at the α-transition, the peak associated with the radial correlation of the polymeric aggregates (ionomer peak) disappears, and as such, the mechanism for this thermal-triggered disordering, or loss of correlation of the polymeric aggregates, has not been fully understood or modeled⁹,¹¹. The absence of a solid empirical finding relating

<footnote n="1">¹Instituto de Pesquisas Energéticas e Nucleares, IPEN-CNEN/SP, São Paulo, 05508000, Brazil. ²Methoden der Materialentwicklung, Helmholtz-Zentrum für Materialien und Energie GmbH, Berlin, 12489, Germany. ³Universidade Federal do ABC, UFABC, Santo André, 09219170, Brazil. ⁴Netherlands Organisation for Scientific Research (NWO), DUBBLE@ESRF, Grenoble, 38000, France. Correspondence and requests for materials should be addressed to B.R.M. (email: brmatos@usp.br)</footnote>
the temperature and shape of the polymeric aggregates makes the SAXS modeling arbitrary and speculative. Therefore, this issue requires a modeling-independent characterization of the shape and size of the polymeric aggregates as a function of *T* for understanding the shape memory property and advancing 4D printing, for example.

Broadband dielectric spectroscopy (BDS) is a powerful tool to study the dynamics of the polarization due to ion-hopping within ionomeric matrices in a wide range of length scales⁶,¹². However, the dielectric spectrum usually displays electrode polarization contributions that mask the characteristic relaxation frequencies within the ionomer matrix¹³. Specifically, this feature has been a main issue for the characterization of PIs, due to the high ion conductivity, and for more than 40 years the assignment of the dielectric spectrum of PIs is an open problem⁶,¹⁴,¹⁵. Even more recently, from 2014 up to the present day, the old interpretations of PI dielectric spectra have not thrived over the years⁶,¹⁵. From 2014 onwards, more recent findings showed a careful characterization of the dielectric spectrum of solutions and hydrated solid films of a one of the mostly studied PI, Nafion®¹⁶. Dilute Nafion solutions represented a simple system for the characterization of the mechanisms of ion-hopping polarization within the polymeric aggregates¹⁶. The striking result was that PIs exhibit polyelectrolyte-like polarizations instead of the conventional segmental motion relaxations observed in hydrocarbon ionomers, for example. The α and β relaxations observed in the dielectric spectrum were assigned to the ion-hopping polarization along the longitudinal and radial directions of the polymeric aggregates, respectively¹⁶,¹⁷. By adding salt to Nafion solution a rod-to-coil transition occurred and the size and shape of the polymeric aggregates were estimated considering the position of α and β relaxations¹⁶. An excellent match was obtained for the dimensions of the polymeric aggregates as estimated by SAXS and BDS¹⁶,¹⁷. This finding opens the investigation of the shape/relaxation relation for solid PI films.

Here, the aspect ratio of the polymeric aggregates was directly obtained at a broad range of temperature for Nafion by BDS. The electrode polarizations at very low frequencies ($f < 10^{-2}$ Hz) were eliminated in this work by the use of a 4-probe setup or a two-probe having high specific surface area electrodes. Initially we performed variable temperature SAXS, infrared spectroscopy (FTIR), and dynamic mechanical analysis (DMA) on Nafion in order to determine the morphology of the ionic domains, state of the electrostatic interactions among ionic groups, and the thermal transition within the ionic phase, respectively, at rigorously the same experimental conditions. Posteriorly, the SAXS, FTIR and DMA data are confronted to BDS measurements of Nafion.

Mid and far infrared (MIR and FIR) spectroscopy and SAXS measurements were performed on Nafion during increasing temperature in the 30–200 °C *T*-range (RH ~ 0%) under N₂. Figure 1a and b shows the MIR data in the functional groups region (4000–1250 cm⁻¹) collected during the first and second heating, respectively. Figure 1c shows the 2D plot of FIR (400–20 cm⁻¹) data for Nafion measured in the first heating. Figure 1d shows the SAXS patterns for Nafion during the first heating, and the inset shows *T*⁻¹/² vs *q*² plot in the low *q*-range.

Relevant features in the Nafion MIR spectra are observed on the first and second heating: (i) at 200 °C, the OH stretching at 2980 cm⁻¹ displays a blueshift (~220 cm⁻¹); followed by the development of a high frequency (3500 cm⁻¹) band assigned to dangling OH groups in SO₃H; (ii) the water bands at 1660, 1710 and 2200 cm⁻¹ disappeared revealing the underlying CF₂ overtone at 2370 cm⁻¹; and (iii) the bands associated with RSO₃⁻ (1060 cm⁻¹, not shown¹⁸) gradually disappeared leading to the gradual appearance of SO₂ vibration in RSO₃H (1410 cm⁻¹)¹⁸,¹⁹. Notably, the initial coordination water of Nafion samples results is monohydrated sulfonic groups (RSO₃H.H₂O)²⁰. Previous thermogravimetric analysis on vacuum dried Nafion samples evidenced that heating from 30–200 °C removed the residual water molecules in the samples²⁰. The water removal is in agreement with absence of MIR water bands and the association of the protonic charges with the sulfonic groups at *T* > 100 °C (hereafter mentioned as “dry samples”) indicating that the proton diffusion takes place mostly via ion-hopping. Further minimization of the water content is attained in the second heating (Fig. 1b) in which protons are associated with the sulfonic groups from 40 to 200 °C, due to the presence of the 1410 cm⁻¹ band in the entire *T*-range; whilst the broad OH stretching MIR band (3000 cm⁻¹) is mostly due to clustering RSO₃H in the sample¹⁹. Importantly, the ionic crosslinks persisted in the entire temperature range investigated. The blueshift of the OH stretching in RSO₃H (2980 cm⁻¹) takes place in the α-transition *T*-range (*T*~α~ ~ 120 °C), having the onset and endset at 100 and 140 °C, respectively. A blueshift of the OH stretching frequency is reported for solid-liquid-vapor transition of H₂O and corresponds to the weakening of the hydrogen bonding²¹. The weakening of the OH bonding in Nafion with increasing *T* can be understood considering the counterelastic forces imposed by the main and side chains against the ionic aggregation as the temperature increases. This feature suggests that the enough *KT* energy is being provided allowing segments of the polymer chains to escape from the temporary ionic crosslinks.

Cation exchange experiments in the FIR region reported the cation motion frequency of the H⁺-SO₃⁻ form in Nafion appearing as broad band centered at 240 cm⁻¹, which can also be seen in Fig. 1c¹⁹,²². The most prominent features reported in the spectral region of 50–400 cm⁻¹ are the CF₂ rocking band at 204 cm⁻¹, and three bands present in the proton form of Nafion~240, 335, and 375 cm⁻¹, which are assigned to H-bond stretching involving the hydronium ion¹⁹,²². It was previously observed that decrease of hydration promoted a decrease of the intensity of 335 and 375 cm⁻¹ bands¹⁹, following the behavior of the water bands in the MIR range (Fig. 1a and b). However, these bands persisted for *T* > 120 °C, suggesting that these bands are also probably related to OH stretching in RSO₃H.

It is difficult to infer feature changes of the cation motion band (240 cm⁻¹) upon heating due to the partial overlapping with 335 cm⁻¹ band (Fig. 1c). Nonetheless, it is possible to notice a broadening and intensity reduction for *T* > 140 °C of the cation motion band. At high *T* (>180 °C), the absorbances of both 240 and 335 cm⁻¹ bands are diminished and only a broad coalesced absorbance is observed. These findings reveal important aspects of the cation-motion band of perfluorinated ionomers. As the cation motion band of ionomers is related to the proton transport across the length scales of the ionic heterogeneities of the sample, the broadening of such band for *T* > 140 °C suggests that the ion-hopping occurs over more disordered ionic aggregates.
<figure><img src="image_60.png" /><figcaption>**Figure 1.** *In situ* MIR temperature measurements of Nafion during the first and second heating in the transmission mode in the functional groups region, 4000–1250 cm⁻¹ (**a** and **b**). *In situ* FIR temperature measurements of Nafion during the first heating in the transmission mode using the synchrotron source from the IRIS beamline (**c**). *In situ* SAXS temperature measurements of Nafion during the first heating using the synchrotron source from the ESRF beamline (**d**). The inset shows *I*⁻¹/² vs *q*² plot in the low *q*-range.</figcaption></figure>

Usually, SAXS patterns of Nafion membranes at low RH (RH ~ 10%) display a scattering maxima centered at *q* ~1.9 nm⁻¹ due to the ionomer peak²³, which can be clearly observed in Fig. 1d for low temperatures. Previously, ultra small angle X-ray scattering (USAXS) measurements of Nafion in the hydrated form displayed low-*q* scattering upturn (*q* < 0.03 nm⁻¹), being attributed to large-scale density fluctuations associated with the length of the polymeric aggregates (*l* ≈ 350 nm)²³,²⁴. Similarly, the Zimm approximation (I⁻¹/² vs *q*²) is showed in the inset of Fig. 1d to estimate the length of the polymeric aggregates of Nafion in the dry form²³,²⁴. In the first and second heatings, the length exhibits a weak dependence on temperature. For the first heating, the length decreases from 182 to 174 nm in the 20–200 °C *T*-range, whereas in the second heating the length decreases from 154 to 150 nm at the same *T*-interval. The reduction of the size of the polymeric aggregates, in the sequence 350 nm (hydrated Nafion)²³ > 182 nm (1st heating) > 150 nm (2nd heating), is possibly due to the less expanded polymeric aggregates due to the water elimination and coordination of protons with the sulfonic acid groups, minimizing the electrostatic repulsions among bare sulfonic acid groups in the sample containing water molecules.

The ionomer peak corresponds to the radial correlation of the cylindrical polymeric aggregates, and at 120 °C the radius is 2.74 nm (Fig. 1d). The SAXS patterns of Nafion measured in the first and second heatings are similar in the 30–200 °C *T*-range. With increasing *T*, in the 30–140 °C range, the ionomer peak decreases in intensity, broadens, and displaces to higher *q*-values (from 1.72 to 2.48 nm⁻¹). In the second heating, the reduction in SAXS peak intensity in the 30–140 °C range is minimized. Such a less temperature-dependent peak intensity is supported by FTIR (Fig. 1b) and can be related to the loss of residual water molecules and irreversible modification of the ionic network arrangement during the first heating. Most pronounced changes in the SAXS patterns are observed at *T* ~ 140 °C, for which the ionomer peak is totally suppressed. The reduction of the ionomer peak intensity could be linked to the water evaporation. However, MIR results evidenced that the effect of the removal of coordinated water from sulfonic acid groups for *T* > 100 °C are minimal. Therefore, the changes of the ionomer peak features for *T* > 140 °C can be attributed to a loss of ordering of the polymeric aggregates. Such finding is in good agreement with the FIR data, confirming that a structuring of the ionic phase takes place due to the weakening of the electrostatic interactions at the α-transition. Previously, the disappearance of the ionomer peak has been observed for drawn Nafion films in the direction perpendicular to the stretching axis, which was attributed to a misalignment of the polymeric aggregates²⁴. In this scenario, a tentative explanation for the suppression of the ionomer peak at the α-transition would be the misalignment of the polymeric aggregates, possibly due to the minimization of the electrostatic repulsion among sulfonic acid groups due to water removal and the weakening of the electrostatic interactions among SO₃H dipoles, which are the “pillars” for the cylindrical shape of the polymeric aggregates. However, the water removal effect and the weakening of the electrostatic interactions do not explain the reason why the suppression of the ionomer peak takes place at 140 °C. It is also worth noting that the only information that can be extracted from both FIR and SAXS measurements is a disordering of the ionic phase
<figure><img src="image_61.png" /><figcaption>**Figure 2.** Electric loss modulus of Nafion using three different electrodes setup: 2-probe SS; 4-probe SS; and 2-probe CC (a). Electric loss modulus of Nafion with decreasing relative humidity using 2-probe CC (b).</figcaption></figure>

<figure><img src="image_62.png" /><figcaption>**Figure 3.** Imaginary component of the proton conductivity as a function of frequency in three temperature ranges: (a) from 40 to 120 °C; (b) from 130 to 160 °C; and (c) from 170 to 200 °C.</figcaption></figure>

with increasing *T*. Both techniques remain silent with respect to the mechanism of this disordering transition, mechanism of which can be assessed by BDS²⁵.

In BDS, the characteristic frequencies of both ion-hopping and α-relaxation are commonly overlapped with electrode polarization at low frequency (typically in the 10⁻²–10² Hz range), hindering the investigation of the dynamics of the dielectric dispersions¹⁷,²⁵. BDS measurements of Nafion using 4-probe and 2-probe (using high surface specific area electrodes) in the through-plane setup are free from electrode contributions. Such measurements allowed identifying the ion-hopping (proton hopping) as well as the α and β relaxations.

Figure 2 shows the dielectric spectroscopy measurements (in the electric modulus formalism)/*M*\* = *M*′ + *iM*\") of Nafion performed with decreasing relative humidity from 100% to 2%. Such measurement allows determining the position of the relaxations at high *RH* in which the ion-hopping dispersion (Fig. 2a) and α-relaxation (Fig. 2b) are more easily identified¹⁷, and monitor its displacement as the sample dries, thereby determining the position of the relaxations in the dry film. The identification of ion-hopping dispersion (σ-dispersion) of Nafion, as showed in Fig. 2a, was performed using three different setups: (i) 2-probe using flat stainless steel (SS) electrodes; (ii) 4-probe SS electrodes; and (iii) 2-probe carbon cloth (CC) electrodes. Both the *M*''(*f*) representation and the 4-probe setup eliminates electrode polarization effects revealing the σ-dispersion at *f* ~ 10⁻¹ Hz for Nafion at *RH* = 100%¹³. Such a feature cannot be detected using the 2-probe SS setup. However, highly reproducible DS measurements with a well-defined σ-dispersion are obtained using high specific surface area 2-probe CC electrodes in a broad range of frequency and temperature. The elimination of the electrode polarization in such frequency range allows determining the position of the σ-dispersion and α-relaxation as the relative humidity decreases, as shown in Fig. 2b.

At 130 °C (*RH* = 60%) the relaxations are displaced to a higher frequency range in which the σ-dispersion and the α-relaxation are identified at *f*<sub>σ</sub> ~ 10<sup>-1</sup> and *f*<sub>α</sub> ~ 10<sup>2</sup> Hz, respectively (Fig. 2b). As the *RH* decreases from 60 to 2%, the α-relaxation decreases in frequency from 10<sup>2</sup> to 10<sup>-1</sup> Hz, superposing with σ-dispersion (~ 10<sup>-2</sup> Hz), indicating a lower *RH*-dependence of *f*<sub>σ</sub> compared to *f*<sub>α</sub>. The σ-dispersion frequency of Nafion at dry conditions is in good agreement with previous investigations of *ac* conductive dispersion (*f* ~ 10<sup>-2</sup>–10<sup>-3</sup> Hz)¹⁷. Therefore, the results of Fig. 2b indicate that at dry conditions (*T* = 30 °C and *RH* ~ 0%) initial position of the α and β relaxations with respect to the σ-dispersion are: *f*<sub>σ</sub> ~ *f*<sub>α</sub> and *f*<sub>σ</sub> < *f*<sub>β</sub>. Such relations are helpful to understand the dynamics role of the ion-hopping on the *T*-dependence of the relaxation dynamics of the ionomer, as shown in Fig. 3.

Figure 3 shows the imaginary conductivity spectra (σ\* = σ' + iσ') of Nafion at three *T*-ranges at *RH* ~ 0% during the second heating. In Fig. 3a, only the β-relaxation can be observed in the 40–120 °C *T*-range, indicating that at *RH* ~ 0%, the alpha-relaxation is displaced even further to lower frequencies.

The ion-hopping characteristic frequency, *f*<sub>σ</sub>, refers to the dynamics of the temporary crosslinks²⁵, while α- and β-relaxations represent the ion-hopping polarization due to the higher activation energy across the length scales
of the polymeric aggregates¹³. The protonic charges diffuse via ion-hopping in the polymeric matrix and polarize at the polymeric aggregates. In this context, two limiting cases are relevant for understanding of Nafion spectra: (i) for $f_{β}>f_{σ}>f_{α}$, the proton diffusion is not limited by the longitudinal polarization of the polymeric aggregates; and (ii) for $f_{β}>f_{α}>f_{σ}$, the proton diffusion is reduced by the polarization of the protonic charges across the radius and length of the polymeric aggregates. The change from behavior (i) to (ii) takes place at 120 °C, as shown in Fig. 3c and d, indicating that the α-transition of Nafion is associated with the presence of an additional polarization of charges existing at $T > 120$ °C.

Figure 3b shows that as the temperature increases from 40 to 120 °C, the β-relaxation displaces to higher f and a shoulder is seen at 120 °C due to α-relaxation high-f displacement. In the 140–170 °C T-interval (Fig. 3c), the β-relaxation displaces to low f until it superseded with the α-relaxation at $f = 10^{-1}$ Hz. Both the high-f and low-f displacements of α and β-relaxations, respectively, represent a change of conformation of the polymeric aggregates. Therefore, the blueshift of α-relaxation indicates a reduction of the length of the polymeric aggregates, whereas the redshift of β-relaxation indicates a thickening of the polymeric aggregates. The coincidence of both relaxations for $T > 170$ °C suggests a reduction of the aspect ratio of the polymeric aggregates to a nearly spherical shape.

The length and radius of the polymeric aggregates (L) are associated with the radial ($f_{β}$) and longitudinal ($f_{α}$) polarization frequencies by the following equation¹²,²⁵:

$$f_{\alpha,\beta}^{-1} \approx \frac{L^2}{D}, \qquad (1)$$

where D is the diffusion coefficient of protons counterions obtained by Nernst-Einstein relation: $D = \sigma k_{b} T / ne^{2}$); where $k_{b}$ is the Boltzmann constant, T is the absolute temperature, e is the elementary charge, and n is the charge concentration. In the $f_{β}>f_{σ}>f_{α}$ regime, at 120 °C, the radius of the polymeric aggregates is~6.47 nm. Due to the difficulty in assessing the α-relaxation from 30 to 120 °C it is not possible to precisely obtain the length of the polymeric aggregates. However, the high-f shift of the α-relaxation in Fig. 2b ($f \sim 10^{-2}$ to $10^{-1}$ Hz) suggests a reduction of the length from 287 to 90 nm, in which the latter is the lowest length achieved by the aggregates at $T = 160$ °C (the endset of α-transition). In the 170–200 °C range, the α-relaxation frequency is nearly constant resulting in an increase of the length of the polymeric aggregates ranges from~141 to 210 nm, which can be an outcome of the thermal expansion of the aggregates. It is worth emphasizing that the radius and length obtained by BDS have similar magnitude as the ones obtained by SAXS. The FIR, SAXS and BDS are in very good agreement indicating that the properties of Nafion membranes are profoundly modified above the α-transition due to the crossover from the two regimes and the suppression of the radial correlation of the polymeric aggregates.

Figure 4a shows the radial correlation of the polymeric aggregates (r) estimated by SAXS and DS. Figure 4b combines FIR, SAXS and BDS data evidencing the suppression of the radial correlation length at the α-transition, as determined by dynamic mechanical analysis (DMA). Figure 4c shows a schematic representation of the proposed conformation transition of Nafion due to the crossover from regime i to ii.

In Fig. 4a, as the temperature increases from 30 to 120 °C, the radius obtained by SAXS and BDS of the polymeric aggregates decreases, following a similar power law with increasing T, -0.02 and -0.05, respectively.

The α-transition as determined by DMA, the elastic modulus drops to values close to zero confirming that the long-range motion of the main chains is no longer restricted by the ionic crosslinks. Since the ion-hopping is correlated with the motion of the main and side chains, the α- and β-relaxations are associated with the long and short range motions of the polymer backbone, respectively. Similarly, considering the segmental motion of the main chains, the two limiting cases can be devised as: (i) For $f_{β}>f_{σ}>f_{α}$, the long range motion of the main chains is restricted by the ionic crosslinks; and (ii) for $f_{β}>f_{α}>f_{σ}$, only the short range motions are restricted by the ionic clustering. This conformation transition is only possible due to the $f_{β}>f_{α}>f_{σ}$ relation at $T > 120$ °C, which allowed the long range motion to escape from the restrictions imposed by the ionic crosslinks.

The combined techniques showed that the length scales of the polarizations probed by BDS are compatible with the ones observed by SAXS. Such findings allow correlating conformation changes of the polymeric aggregates with the ion-hopping. Therefore, the understanding of the evolution of the size and ordering of the polymeric aggregates and the ion-hopping with increasing T would help establish the relationship between the ionic network and microstructure of PIs, which is crucial for tailoring new high-performance ionomers²⁶. The reversibility of the shape of the polymeric aggregates due to annealing and water swelling of Nafion is the next step of this investigation.

## Conclusions

Broadband dielectric spectroscopy (BDS) with free electrode polarization effects is shown to be a unique tool capable of revealing the origin of the atypical dynamics of α-transition in perfluorinated ionomers. The combination of variable temperature IR, SAXS, and DS allowed for monitoring of the electrostatic interactions and morphology from different length scales in Nafion membranes. These characterizations provide further insights into the origin of α-transition. Namely, the α-transition is a critical temperature ($T = 120$ °C) separating two regimes. Below α-transition ($T < 120$ °C) the shape of the polymeric aggregates are locked into nanocylinders due to the fast ion-hopping, which restricts the translational motion of the main and side chains. Above α-transition ($T > 120$ °C), the relaxation frequency of α surpasses the ion-hopping permitting the long-range motion of the main chains, which relaxes in a more spherical conformation. Such finding allows mastering the temperature-shape relationship in ion-containing polymers as well as tuning the desired shape of the building blocks of the ionomer morphology.
<figure><img src="image_63.png" /><figcaption>**Figure 4.** The α-transition as measured by DMA and the disordering of ionic domains of the polymeric aggregates as determined by different techniques, FIR (band position), SAXS (ionomer peak position) and BDS (β-relaxation position) (a); the radius of the polymeric aggregates as estimated by SAXS and BDS (b); and schematic representation of the proposed conformation transition of Nafion due to the crossover from regime i to ii (c).</figcaption></figure>

## Methods

Commercial Nafion membranes (EW = 1,100 g eq⁻¹) were obtained from DuPont with different thicknesses (N211, N115). The membranes were pre-treated by standard cleaning and activation protocols¹⁷. The obtained film was then post-treated in 3% (w/w) H₂O₂ and 0.5 M H₂SO₄, with intermediate steps in H₂O to remove excess chemicals. Nafion membranes in the non-ionic form (RSO₂F) were used to help identifying the IR bands. In all measurements, the water content in commercial Nafion samples (EW = 1100 meqg⁻¹) was minimized by keeping them in N₂ for 24 h, which also allowed us to better estimate the initial hydration condition of the membrane²⁰.

Variable temperature mid infrared (MIR FT) and far infrared (FIR FT) spectra were measured at the IRIS beamline at the electron storage ring BESSY II of Helmholtz Zentrum Berlin. MIR FT was measured in transmission mode in the Bruker Vertex 70/v using a KBr beamsplitter, a DLaTGS detector fitted with an internal global source. For these measurements a modified Harrick demountable FT IR transmission cell was employed which allowed in-vacuum heating. The cell was equipped with ZnSe windows of 2 mm thickness and a 630 µm PTFE spacer. The spectra were acquired with 4 cm⁻¹ resolution and 32 scans were averaged. The references were recorded through the empty channel inside the spectrometer sample compartment under vacuum. FIR FT measurements in spectral region between 600 and 30 cm⁻¹ were performed using the Harrick cell equipped with 2 mm thick Ge windows, with 2 cm⁻¹ resolution using infrared synchrotron radiation as a source together with a liquid helium cooled silicon bolometer and a silicon beamsplitter. The reference was obtained by measuring the empty Harrick cell with identical conditions in the entire range of temperature investigated. Previous to the measurements, the membranes were preconditioned in the Harrick cell by purging N₂ (RH ~ 0.1%) for 24 h at room T. The relative humidity was recorded using a humidity sensor (Sensirion EK-H4) at the outlet of the cell. The data collection was performed on N211 and N115 samples in the 30-200 °C T-range with a step of 10 degrees in two successive heatings. The FTIR bands in the 1100-1200 cm⁻¹ range are saturating (not shown). Measurements
were performed after ~10 min at each temperature for system stabilization, when no noticeable spectral changes were observed.

The samples were examined by SAXS at the beam line BM26-B27,28 at the European Synchrotron Radiation Facility (ESRF) in Grenoble, France. The energy of X-ray source was 12 keV (λ eV0.1033 nm), and sample to detector distance was 3005 mm. Data was recorded using the Pilatus 1 M detector, with 172 × 172 µm pixel size. The samples were heated at 2 °C/min from 20 °C to 250 °C, cooled at 40 °C/min to 20 °C, heated again at 2 °C/min to 300 °C and finally cooled at 10 °C/min to 20 °C, in a Linkam DSC600 hot stage for the simulatenous SAXS temperature-dependent study. The SAXS patterns were reduced by BUBBLE and by the homemade XRTools software. The scattering Intensity (I) vs scattering vector, q sc4π sin θλ⁻¹, where λ is the X-ray wavelength and 2θ is the scattering angle. The scattering vector was calibrated using silver behenate. The scattering patterns were also corrected for transmission, normalized upon primary beam fluctuations and background scattering before data integration.

Dynamic mechanical analysis (DMA) was carried out under nitrogen flow in a Netzsch DMA 242 E in tensile mode. Rectangular Nafion samples (15 × 6 mm²) were cut and tested, respecting the machine direction. The measurements were performed from 30–200 °C with 10 °C steps under N2 flowand amplitude of 4 µm and oscillation frequency of 1 Hz.

Broadband dielectric spectroscopy (BDS) data were collected in a specially designed air-tight sample holder able to measure the proton conductivity in the range T = 30–200 °C with RH = 0% (dry N₂ purge)¹⁷. Temperature controllers connected to band heaters placed externally around the cylindrical chambers are monitored by thermocouples (type K) inserted inside the metallic walls. The sample holder is capable of controlling both the temperature (from room temperature up to ~200 °C) and the relative humidity (RH, from ~3 to 100%). Nafion samples were sandwiched between stainless steel spring-load contact terminals (electrically insulated from the chamber walls) with carbon cloth to facilitate water equilibration. In this experimental apparatus, the RH of the sample chamber can be calculated by, RH = ρ(Tᵢ)/P(Tᵢ) × 100, where ρ is the vapor partial pressure, P is the saturated vapor partial pressure, and Tᵢ and Tᵣ are, respectively, the water reservoir and sample chamber temperatures. A Solartron 1260 frequency response analyzer was used in the frequency (f) range of 4 mHz to 3 MHz applying an ac amplitude of 100 mV. The complex conductivity (σ* = 2πfε₀ε*) and electric modulus (M* = M + iM*) = 1/ε* representations were used throughout this study, in which the dielectric loss (ε'(f)) was obtained from:

$$ε'(f) = - \frac{dcos(\theta(f))}{2\pi f \varepsilon_0 S |Z(f)|} \quad (2)$$

where ε' and ε' are the real and imaginary parts of the dielectric permittivity; ε₀ is the vacuum permittivity (~8.854 × 10⁻¹⁴ Fcm⁻¹); S is the electrode active area, d is the thickness of the membrane; |Z| and θ are the modulus and phase angle of impedance. The frequency dependent conductivity was obtained using the relation: σ' = 2πfε₀ε''¹³.

## References

1. Tibbits, S. 4D Printing: Multi-Material Shape Change. *Architect. Design.* **84**, 116–121 (2014).

2. Ge, Q., Qi, H. J. & Dunn, M. L. Active materials by four-dimension printing. *Appl. Phys. Lett.* **103**, 131901 (2013).

3. Xie, T. Tunable polymer multi-shape memory effect. *Nature* **464**, 267–270 (2010).

4. Zhao, Z. *et al*. Three-Dimensional Printed Shape Memory Objects Based on an Olefin Ionomer of Zinc-Neutralized Poly(ethylene-co-methacrylic acid). *ACS Appl. Mater. Interfaces* **9**, 27239–27249 (2017).

5. Page, K. A., Park, J. K., Moore, R. B. & Sakai, V. G. Direct Analysis of the Ion-Hopping Process Associated with the α-Relaxation in Perfluorosulfonate Ionomers Using Quasielastic Neutron Scattering. *Macromolecules* **42**, 2729–2736 (2009).

6. Osborn, S. J. *et al*. Glass Transition Temperature of Perfluorosulfonic Acid Ionomers. *Macromolecules* **40**, 3886–3890 (2007).

7. Matos, B. R., Goulart, G. A., Santiago, E. I., Muccillo, R. & Fonseca, F. C. Proton conductivity of perfluorosulfonate ionomers at high temperature and high relative humidity. *Appl. Phys. Lett.* **104**, 091904 (2014).

8. Lopez-Haro, M. *et al*. Three-dimensional analysis of Nafion layers in fuel cell electrodes. *Nature Comm.* **5**, 5229 (2014).

9. Kusoglu, A. & Weber, A. Z. New Insights into Perfluorinated Sulfonic-Acid Ionomers. *Chem. Rev.* **117**, 987–1104 (2017).

10. Schmidt-Rohr, K. & Chen, Q. Parallel cylindrical water nanochannels in Nafion fuel-cell membranes. *Nature Mater.* **7**, 75–83 (2008).

11. Jackson, D. A., Koberstein, J. T. & Weiss, R. A. Small-Angle X-Ray Scattering Studies of Zinc Stearate-Filled SulfonatedPoly(ethylene-co-propylene-co-ethylidenepropylene-co-ethylidene norbornene) Ionomers. *J. Polym. Sci. Part B: Polym. Phys.* **37**, 3141–3150 (1999).

12. Colby, R. H. *et al*. Dynamics of Lightly Sulfonated Polystyrene Ionomers. *Phys. Rev. Lett.* **81**, 3976 (1998).

13. Schonhals, A. & Kremer, F., In Broadband Dielectric Spectroscopy, Schonhals, A. & Kremer, F., Editors, p. 59 (Springer Verlag, Berlin, 2003).

14. Yeo, S. C. & Eisenberg, A. Physical properties and supermolecular structure of perfluorinated ion-containing (nafion) polymers. *J. Appl. Polym. Sci.* **21**, 875 (1977).

15. Di Noto, V., Negro, E., Sanchez, J.-Y. & Iojoiu, C. Structure-relaxation interplay of a new nanostructured membrane based on tetraethylammonium trifluoromethanesulfonate ionic liquid and neutralized nafion 117 for high-temperature fuel cells. *J. Am. Chem. Soc.* **132**, 2183 (2010).

16. Matos, B. R., Santiago, E. I., Rey, J. F. Q. & Fonseca, F. C. Origin of α and β relaxations of Nafion. *Phys. Rev. E* **89**, 052601 (2014).

17. Matos, B. R. *et al*. dc Proton conductivity at low-frequency in Nafion conductivity spectrum probed by time-resolved SAXS measurements and impedance spectroscopy. *J. Polym. Sci. Part B: Polym. Phys.* **53**, 822 (2015).

18. Puskar, L. *et al*. Infrared dynamics study of thermally treated perfluoroimide acid proton exchange membranes. *Phys. Chem. Chem. Phys.* **19**, 626–635 (2017).

19. Bernardina, S. D. *et al*. Mechanism of Ionization, Hydration, and Intermolecular H-Bonding in Proton Conducting Nanostructured Ionomers. *J. Phys. Chem. C* **118**, 25468–25479 (2014).

20. Matos, B. R. *et al*. Proton and cesium conductivity in perfluorosulfonate ionomers at low and high relative humidity. *Solid State Ionics* **301**, 86–94 (2017).

21. Eisenberg, D. & Kauzmann, W. *The Structure and Properties of Water*. (Oxford University Press, New York, 1969; pp. 300).

22. Bawagana, A. D. O. *et al*. Far-infrared studies on Nafion and perfluoroimide acid (PFIA) and their alkali salts. *Vibr. Spectr.* **75**, 213–217 (2014).
23. Gebel, G. & Lambard, J. Small-Angle Scattering Study of Water-Swollen Perfluorinated Ionomer Membranes. *Macromolecules* **30**, 7914–7920 (1997).

24. van der Heijden, P. C., Rubatat, L. & Diat, O. Orientation of Drawn Nafion at Molecular and Mesoscopic Scales. *Macromolecules* **37**, 5327–5336 (2004).

25. Tierney, N. K. & Register, R. A. Ion Hopping in Ethylene-Methacrylic Acid Ionomer Melts As Probed by Rheometry and Cation Diffusion Measurements. *Macromolecules* **35**, 2358–2364 (2002).

26. Matos, B. R. In Situ Fabrication of Nafion–Titanate Hybrid Electrolytes for High-Temperature Direct Ethanol Fuel Cell, https://doi.org/10.1021/jp405754v.

27. Bras, W. et al. Recent experiments on a combined small-angle/wide-angle X-ray scattering beam line at the ESRF. *J. Appl. Cryst.* **36**, 791–794 (2013).

28. Portale, G., Cavallo, D., Alfonso, G. C., Hermida-Merino, D. & Van Drongelen, M. Polymer crystallization studies under processing-relevant conditions at the SAXS/WAXS DUBBLE beamline at the ESRF. *J. Appl. Cryst.* **46**, 1681–1689 (2013).

## Acknowledgements

We thank the Brazilian funding agencies CAPES, CNPq, and FAPESP (grant number 2013/50151-5, 2014/09087-4, 2014/50279-4 and 2016/14785-8). Thanks are also due to LNLS (18801 and 20160239) and CNEN. FCF is a CNPq fellow. We thank the ESRF and the HZB for the allocation of synchrotron radiation beamtime.

## Author Contributions

R.P. and B.R.M. performed the DMA experiments and analyzed the data. B.R.M., L.P. and U.S. performed the FTIR experiments and discussed the data. J.F.Q.R. and D.H.M. performed the SAXS experiments and analyzed the data. B.R.M. and F.C.F performed the BDS experiments and discussed the data. B.R.M. wrote the manuscript. All the authors revised the manuscript.

## Additional Information

**Competing Interests:** The authors declare no competing interests.

**Publisher's note:** Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.

<figure><img src="image_64.png" /></figure>

**Open Access** This article is licensed under a Creative Commons Attribution 4.0 International License, which permits use, sharing, adaptation, distribution and reproduction in any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons license, and indicate if changes were made. The images or other third party material in this article are included in the article's Creative Commons license, unless indicated otherwise in a credit line to the material. If material is not included in the article's Creative Commons license and your intended use is not permitted by statutory regulation or exceeds the permitted use, you will need to obtain permission directly from the copyright holder. To view a copy of this license, visit http://creativecommons.org/licenses/by/4.0/.

© The Author(s) 2018
## 7- Co-proposer PhD Thesis project awarded ; 'Nouvelle génération de biopolymères nanostructurés pour la fabrication additive et l'électrofilage, Université de Lorraine.

I have been awarded with 92.500,0 € as co-applicant with prof. Patrice Bourson by the Region of Grand Est to finance the development of a thesis to investigate novel nanostructured biopolymers to generate fibers by electrospinning. The thesis has recently started, 01/10/2019 and the founds will cover the project until the 30/09/2022. The project involves several scientific partners to integrate different expertise to cover from the tailored synthesis of novel biodegradable polyesters to the characterization of the structure-properties relationship and their applications in the biomedical field (see Figure 7.1). I have led the organization of the scientific cooperation using my previous connections across Europe to identify the task required as well as the timelines to achieve the project. Firstly, I have conceived the polymer derivatives to proceed with the nanostructural investigation to initiate the formulation of drug delivery systems. We have implemented an Erasmus project between the Université de Lorraine and Universidad de Castilla la Mancha at a first stage to synthesise the aimed polymer derivatives and the nanostructure characterization will be continued in DUBBLE under my supervision.

<figure><img src="image_65.png" /><figcaption>Figure 7.1. Planning of the objectives for the generation of novel stereo-PLA-block-copolymers from the synthesis to the applications.</figcaption></figure>

Plastics revolution started in 1869 when John Wesley Hyatt found that a practical plastic material could be made by mixing nitrocellulose, camphor and alcohol: celluloid(1). Since then, plastics have replaced a large number of materials on the market due to the synthesis of a wide variety of new polymers with a great variability of applications and properties (i.e., polyvinyl chloride, polystyrene). In 2018, the production of plastics worldwide was 359 million metric tons (2) and it is expected to double within the next 20 years. According to the United Nations, five billion plastic bags are used each year and one million plastic bottles are bought every
minute. Plastic production projections based on current trends estimate that will be nearly 12 billion tons of plastic waste in landfills and oceans and consequences could be devastating by 2050 (3,4). Plastic debris in oceans was first observed in 1960s and since then, plastics reputation has worsened due to the lack of biodegradability as well as the potential threat to the human health. Over time, plastics in the sea break down into smaller fragments called micro-plastics that when consumed by marine animals could enter the human food chain.

Among the actions that could reduce the magnitude of the problem (public education, regulatory and economic policy instruments), the substitution of polyolefin for bio-based and biodegradable plastics is one of the most important. However, the reduction of plastic production is limited due to the lack of substitute materials that can satisfy the perceived advantages of plastic and consumer practices. Bio-based plastics such as starch blends, bio-PET and bio-PE are already been employed in several industrial applications and its production is expected to increase to 6.1 million tons by 2021.

During my second granted scientific visit to the Universidad de Castilla la Mancha, at the Pharmacy School, I have established a collaboration to synthesize novel block copolymers stereocomplexes of PLLA-PLDA at high M<sub>w</sub> in order to generate bio-based plastics with properties similar to polyolefins. The grant as a senior scientific visitor has been awarded 2500 € to cover my subsistence, traveling and accommodation for a month.

PLA is a well-studied bio-based polymer that is usually synthesized by ROP generating two enantiomers, poly (L-lactic acid) (PLLA) and poly (D-lactic acid) (PDLA) due to the presence of a carbon chiral centre in its backbone. PLA is the bio-based polymer most widely employed in diverse fields such as food handling, fibre manufacturing, textile industry or biomedical. Particularly, biomedical field applications include tissue engineering (tissue regeneration, internal fixation, wound treatment, suture thread), stents, drug delivery systems.

Polylactic acid (PLA) possesses particular interest to manufacturers due to its biocompatibility, processability (easy to be handled in fused deposition modeling machines -3D printing-) and its eco-friendliness (derived from sugar resources by well-known chemical processes, biodegradable, recyclable and compostable).

One of the main advantages of PLA is to be bioassimilable, i.e. it is completely metabolized to water and carbon dioxide in the human body and as such totally eliminated (Food and drugs administration (FDA) approved). Moreover, toxicity studies have shown the inert effect as a solid form in humans. The degradation kinetics depend on several physicochemical parameters such as pH, molecular weight or thickness of the sample, and thus, can be modulated by playing with the PLA chirality to continue or reduce the degradation rate, respectively. In addition, the degradation of a highly crystalline form to lactic acid will last for months whilst an amorphous derivative will degrade in weeks.

However, the PLA applicability is slightly inferior to conventional petroleum-based polymers in terms of thermal resistance, mechanical properties and crystallinity.

Isotactic PLA, composed of identical repeat units of either L-LA or D-LA, is a semicrystalline polymer that generates either PLLA or PDLA respectively, that crystallize in two enantiomeric
homochiral strands. Moreover, PLA can cocrystallize if both enantiomers are present as a stereocomplex (SC) phase, featuring properties superior to the homochiral (HC) crystal phase.

Different strategies have been evaluated to improve the physicochemical properties such as the use of plasticizers or nucleating agents. Moreover, the equimolar blend of PLLA and PDLA enantiomers is a strategy largely used to generate SC crystallites, which present a melting point 50°C higher than its homocrystals (HC) counterparts due to the strong interactions between L- and D-lactyl unit sequences (Tsuji, 2016). However, the SC crystallization of the blended enantiomers (PLLA and PDLA) diminishes for high molecular weight (HM~w~) PLA and enantiomeric HC are obtained instead (Competitive Stereocomplexation 2015). Block copolymers formed by enantiomeric PLA strands are required to retain the SC crystallization in HM~w~ polymers.

Previous synthetic strategies for obtaining stereo-block-PLA have been focused on the lengthening of building blocks from pre-polymers by a stepwise fashion or the attachment of pre-synthesized PDLA and PLLA strands by reactive end-group linkers. However, limitations on the preciseness and complexity of the desired polymer architecture are the main drawbacks of the prepolymer extension synthetic approaches. A straightforward strategy to achieve stereo-regularity in the polymeric backbone is the one-pot sequential monomer addition to a truly living polymerization catalyst, requiring catalysts that are extremely rare. Recently, highly active Zn (II) and cloro-magnesium complexes in the ROP catalysis of homo-chiral lactides have been introduced to synthetize HM~w~ stereo-n-block co-polymers. Stereo-di-block copolymers characterized with a high T~m~ and ΔH~m~ have been successfully synthetized, suggesting the formation of SC crystals (Rosen, Goldberg, Venditto, & Kol, 2016). Nevertheless, an exhaustive characterization of the stereo di-block-copolymers is needed to understand the mechanism of crystallization of the SC phase as well as to correlate the crystalline structure to its physicochemical properties. The crystallization mechanism of enantiomeric PLA (PLLA and PDLA) has been largely investigated as well as the equimolar PLLA/PDLA physical blend. The latter, showed the exclusive formation of SC crystals only at low M~w~ samples, and a progressive transformation of the HC/SC ratio as M~w~ increased. Importantly, the crystallization mechanism of stereo di-block co-polymers PLA has never been investigated.

Chiral N,N,O-scorpionate zinc alkyls previously synthetized by the group at Universidad de Castilla la Mancha (Otero 2012) showed to be effective and stereoselective initiators for ROP of lactides, acting as single-site initiators at mild conditions and without the need for a cocatalyst or activator. Herein, a series of PLA stereo-diblock-copolymers attaining different M~w~ were synthetized, through a scorpionate zinc alkyl catalyst, as a proof of concept, to achieve highly stereo-regular derivatives.
Summarising, the aim of thesis during the first year was the optimization of the synthetic route and its characterization (see Table 1). Polymerizations of homochiral L- and D-LA, and stereo di-block co-polymers were conducted in toluene at 90°C.

<table><caption>Table 1. Summary of the PLLA polymers synthesis optimization.</caption><thead><tr><th>Entry</th><th>Di-block copolymer (theoretical block length)</th><th>Temperature (°C)</th><th>Time (min)</th><th>% conversion</th><th>M<sub>w</sub> (theor.) (Da)</th><th>M<sub>w</sub> (exper.) (Da)</th><th>M<sub>w</sub>/M<sub>n</sub></th><th>Di-block copolymer (block length)</th></tr></thead><tbody><tr><td>1</td><td>PLLA-b-PLLA (50L-b-50D)</td><td>90</td><td>50</td><td>95</td><td>14400</td><td>32823</td><td>1.93</td><td>PLLA-b-PLLA (115L-b-115D)</td></tr><tr><td>2</td><td>PLLA-b-PLLA (100D-b-100L)</td><td>90</td><td>60</td><td>94</td><td>28800</td><td>79107</td><td>2.38</td><td>PLLA-b-PLLA (275D-b-275L)</td></tr><tr><td>3</td><td>PLLA-b-PLLA (100L-b-100D)</td><td>90</td><td>60</td><td>95</td><td>28800</td><td>79320</td><td>2.25</td><td>PLLA-b-PLLA (275L-b-275D)</td></tr><tr><td>4</td><td>PLLA-b-PLLA (200L-b-200D)</td><td>75</td><td>120</td><td>81</td><td>57600</td><td>72508</td><td>2.20</td><td>PLLA-b-PLLA (250L-b-250D)</td></tr><tr><td>5</td><td>PLLA-b-PLLA (200L-b-200D)</td><td>90</td><td>90</td><td>90</td><td>57600</td><td>63313</td><td>2.20</td><td>PLLA-b-PLLA (220L-b-220D)</td></tr><tr><td>6</td><td>PLLA-b-PLLA (300L-b-300D)</td><td>90</td><td>100</td><td>97</td><td>86400</td><td>107912</td><td>1.69</td><td>PLLA-b-PLLA (375L-b-375D)</td></tr><tr><td>7</td><td>PLLA-b-PLLA (500L-b-500D)</td><td>90</td><td>120</td><td>84</td><td>144000</td><td>94181</td><td>2.11</td><td>PLLA-b-PLLA (330L-b-330D)</td></tr><tr><td>8</td><td>PLLA-b-PLLA (500L-b-500D)</td><td>90</td><td>7h</td><td>85</td><td>144000</td><td>150094</td><td>2.31</td><td>PLLA-b-PLLA (520L-b-520D)</td></tr></tbody></table>

<figure><img src="image_66.png" /><figcaption>Figure 7.2. ¹H NMR.</figcaption></figure>

Different M<sub>w</sub> were obtained by varying the ratio [LA]/[catalyst] and M<sub>w</sub> distributions obtained were fairly narrow (~2). The polymerizations kinetics shown, including high molecular weight polymerizations, high conversion rates, above 80%. Polymerization of L-lactide and D-Lactide yielded stereochemically pure PLA (¹H NMR and homonuclear decoupling ¹H NMR, see Figure 7.2), indicating that epimerization of the monomer and polymer was absent under the polymerization conditions. Stereo di-block co-polymers were readily prepared by the one-pot sequential addition method.

<table><thead><tr><th>tetrad</th><th>Probability of ESC (non-Bernoulli)</th><th>Probability of CEC (Bernoulli)</th></tr></thead><tbody><tr><td>mmm</td><td>[P<sub>m</sub><sup>2</sup>+(1-P<sub>m</sub>)<sup>2</sup>+P<sub>m</sub><sup>3</sup>+(1-P<sub>m</sub>)<sup>3</sup>]/2</td><td>P<sub>m</sub><sup>2</sup>+0.5 P<sub>m</sub>P<sub>r</sub></td></tr><tr><td>mmr</td><td>[P<sub>m</sub><sup>2</sup>(1-P<sub>m</sub>)+P<sub>m</sub>(1-P<sub>m</sub>)<sup>2</sup>]/2</td><td>0.5 P<sub>m</sub>P<sub>r</sub></td></tr><tr><td>rmm</td><td>[P<sub>m</sub><sup>2</sup>(1-P<sub>m</sub>)+P<sub>m</sub>(1-P<sub>m</sub>)<sup>2</sup>]/2</td><td>0.5 P<sub>m</sub>P<sub>r</sub></td></tr><tr><td>rmr</td><td>[P<sub>m</sub><sup>2</sup>(1-P<sub>m</sub>)+P<sub>m</sub>(1-P<sub>m</sub>)<sup>2</sup>]/2</td><td>0.5 P<sub>m</sub><sup>2</sup></td></tr><tr><td>rrm</td><td>[P<sub>m</sub>(1-P<sub>m</sub>)]</td><td>0.5 (P<sub>m</sub><sup>2</sup>+P<sub>m</sub>P<sub>r</sub>)</td></tr></tbody></table>

<figure><img src="image_67.png" /><figcaption>Figure 7.3. Homonuclear decoupling ¹H NMR</figcaption></figure>

Monomer consumption was found higher than 85% for both block sequences followed by ¹H NMR. Moreover, stereo di-block co-polymers tacticity was calculated from deconvoluted
homonuclear decoupling ¹H NMR spectra by using both Bernoullan and non-Bernoullan statistics (see Figure 7.3). The stereocontrol achieved by a particular initiator is expressed in terms of probability to generate certain stereosequences. Pm: probability of isotactic enchainment and Pr: probability of sindio/heterotactic enchainment (platel2008). Besides, Pm + Pr = 1 and while Pm=0.5 idicates a random insertion, Pm=1.0 indicates isotactic PLA (orhan2018). The deconvoluted methine region of the homonuclear decoupled ¹H NMR spectrum of 300L-b-300D stereo di-block co-polymer is shown in Figure 3, as an example. The peaks were assigned to the appropriate tetrads in accordance with literature (Ovitt2002).

Stereocontrolled polymerization can be mediated by two distinct mechanisms, namely, chain end control (CEC) or enantiomorphic site control (ESC). In the former case, control of the chirality is associated with the propagating chain end that in the transition state of the next monomer insertion defines the chirality of the next monomer unit to be inserted. In contrast, in polymerizations mediated through an ESC mechanism, the chirality of the catalyst determines the chirality of the next monomer unit.

<table><caption>Table 2. Summary of the tacticity of PLA derivatives for both methodologies.</caption><thead><tr><th></th><th>Pm (ESC)</th><th>Pm (CEC)</th></tr></thead><tbody><tr><td>PLLA-b-PLLA (50L-b-50D)</td><td>0.96</td><td>0.88</td></tr><tr><td>PLLA-b-PLLA (100D-b-100L)</td><td>0.93</td><td>0.83</td></tr><tr><td>PLLA-b-PLLA (100L-b-100D)</td><td>0.90</td><td>0.96</td></tr><tr><td>PLLA-b-PLLA (200L-b-200D)</td><td></td><td></td></tr><tr><td>PLLA-b-PLLA (200L-b-200D)</td><td></td><td></td></tr><tr><td>PLLA-b-PLLA (300L-b-300D)</td><td>0.99</td><td>0.96</td></tr><tr><td>PLLA-b-PLLA (500L-b-500D)_20</td><td>0.82</td><td>0.69</td></tr><tr><td>PLLA-b-PLLA (500L-b-500D)_19</td><td>0.98</td><td>0.92</td></tr></tbody></table>

The isomer monomer insertion creates a stereoerror from which « standard » growth would form a new stereoblock of opposite configuration to the previous formed (-LLLLLDDDDDLLLLL-). Assuming isotactic-PLA would be formed by a CEC mechanism exclusively when polymerization is run from rac-LA, the following relative tetrad intensities would be expected: [mrm]=[rmm]≠ [rmr]. Conversely, the ESC mechanism should generate single insertion stereoerrors of the type -LLLLLLELDLLLL-. In ESC mechanism, the tetrad ratio should then follow the sequent probabilites; [rmr]=[mmr]=[rmm]=2/[mrm] (Orhan2018_homedecoupling). Generally, stereocontrol in the ROP of PLA is more commonly achieved by CEC, even if the catalyst contains a chiral component. However, any conclusion related to the mechanism follow by our catalyst can not be drawn from the tetrad probabilities, due to the assumptions of the models applied do not correspond to those of our synthesis, i.e., the model is usually applied for polymerizations run from rac-or meso-LA, while our synthesis is based on the one-pot sequential addition of enantiopure LA monomers. However, the polymerization mechanism followed by our catalyst can be elucidated from a racemic PLA polymerization and thus, the appropriate Pm values (those obtained from ESC or CEC mechanism) will be considered.

Moreover, it has been investigated the polymer crystallization mechanism of those derivatives in collaboration with the group of prof. Sanjay Rastogi at Maastricht University.
<figure><img src="image_68.png" /><figcaption>Figure 4. a) PLLA and PDLA structure; b) SC crystalline lattice; c) Raman studies; d) WAXS studies; e) HC- cooling curve ; f) SC- cooling curve ; g) HC- heating curve ; h) SC- heating curve .</figcaption></figure>

<table><thead><tr><th rowspan="2">Cooling rate</th><th colspan="3">HC-PLLA</th><th colspan="3">SC-PLLA</th></tr><tr><th>T<sub>c</sub> (°C)</th><th>T<sub>m</sub> (°C)</th><th>ΔH<sub>m</sub> (J.g<sup>-1</sup>)</th><th>T<sub>c</sub> (°C)</th><th>T<sub>m</sub> (°C)</th><th>ΔH<sub>m</sub> (J.g<sup>-1</sup>)</th></tr></thead><tbody><tr><td>10°C/min</td><td>100.48</td><td>175.50</td><td>57.04</td><td>108.70*</td><td>183.55</td><td>30.82</td></tr><tr><td>8°C/min</td><td>105.51</td><td>176.90</td><td>55.03</td><td>132.40</td><td>207.93</td><td>47.91</td></tr><tr><td>6°C/min</td><td>107.73</td><td>175.70</td><td>55.48</td><td>134.69</td><td>205.32</td><td>49.34</td></tr><tr><td>5°C/min</td><td>107.81</td><td>176.12</td><td>55.67</td><td>137.94</td><td>205.41</td><td>54.03</td></tr><tr><td>3°C/min</td><td>111.48</td><td>178.28</td><td>55.34</td><td>112.35</td><td>188.82</td><td>30.06</td></tr><tr><td>1°C/min</td><td>119.25</td><td>174.29</td><td>59.14</td><td>103.13</td><td>174.10</td><td>26.88</td></tr></tbody></table>

*: cold crystallization

<caption>**Table 3.** DSC results of HC-PLLA and SC-PLLA upon non-isothermal melt crystallization and subsequent melting.</caption>

The crystallization mechanism of enantiomeric PLA (PLLA and PDLA) has been largely investigated as well as the equimolar PLLA/PDLA physical blend. The latter, showed the exclusive formation of SC crystals only at low M<sub>w</sub> samples, and a progressive transformation of the HC/SC ratio as M<sub>w</sub> increased. Importantly, the crystallization mechanism of stereo di-block co-polymers PLA has never been investigated as previously mentioned. Preliminary independent Raman and SAXS/WAXS measurements at room temperature confirmed that stereo-diblock-copolymers crystallize exclusively as SC (see Figure 7.4 c and d), compared to physically blended equimolar enantiomers, in which a competition between stereocomplexation and homoenantiomeric crystallization occurs. Moreover, crystallization studies on PLA stereo-diblock-copolymers by DSC (Image 4e-h) showed dramatic differences on the crystalline populations depending on cooling rate (see ΔH<sub>m</sub> and T<sub>m</sub> in table 3). Importantly, PLA stereo-diblock-copolymer crystalline populations could be associated to either SC (8, 6 and 5°C/min) or HC (1°C/min) despite been covalently linked. Moreover, a non-reported crystalline phase is obtained when cooling at 10 and 3°C/min that could be related to the number of monomeric units participating in the SC. A fully characterization of the crystallization mechanism from the melt will be performed by simultaneous SAXS-WAXS in combination with Raman spectroscopy in the next future and compare them to the corresponding enantiomeric and physically blended counterparts. Raman spectroscopy will be used to correlate the chain conformation evolution of either the SC phase (3<sub>1</sub> helix 908 cm<sup>-1</sup>) or the α phase (10<sub>3</sub> helix 921 cm<sup>-1</sup>) to the crystallinity triggering resolved by WAXS as well as possible phase transformations, as suggested by DSC depending on the thermal history. The long period will be monitored by SAXS and detailed multivariate 2D analysis will be carried out to correlate Raman spectra to scattering profiles to understand complex multistep crystallization mechanism. Moreover, crystallization studies
have been conducted by polarized optical microscopy (POM) of both PLLA and PDLA homopolymers as well as HMw stereo di-block co-polymer at different isothermal crystallization temperatures (Tic) after 5 and 15 min, respectively, from the start of the isothermal step. It has been found that the higher the selected Tic the lower the density of spherulites for all PLA derivatives. Moreover, higher density of spherulites in PLLA samples were observed compared to that of PDLA, indicating that crystallization in PLLA is faster than in PDLA. Photomicrographs of HMw co-polymer were acquired after 15 minutes from the start of the isothermal step, rather than 5, because the latency period until spherulites start to grow up is higher for the co-polymer than for homopolymers (see Figure 7.5).

<figure><img src="image_69.png" /><figcaption>Figure 5. Polarization photomicrographs of PLLA (A-E) and PDLA (F-J) homopolymers and (K-O) HMw stereo di-block co-polymer at different isothermal crystallization temperatures (Tic) after 5 and 15 min, respectively, from the start of the isothermal step.</figcaption></figure>

Moreover, the thesis student has initiated to investigate the possibility to use the novel generated derivatives as drug carriers at Universidad de Castilla la Mancha following our recent work distributed as research highlight by Alba to the National press:

Enrique Niza, Cristina Nieto-Jiménez, María del Mar Noblejas-López, Iván Bravo, José Antonio Castro-Osma, Felipe de la Cruz-Martínez, Marc Martínez de Sarasa Buchaca, Daniel Hermida-Merino, Eduardo Solano, Alberto Ocaña, Carlos Alonso-Moreno. Poly(Cyclohexene Phthalate) Nanoparticles for Controlled Dasatinib Delivery in Breast Cancer Therapy. Nanomaterials 2019, 9, 1208. DOI: 10.3390/nano9091208 en la prensa por ser elegido por su importancia en la investigación fundamental y en la aplicación de los rayos X por el sincrotrón ALBA. https://www.sincrotronalba.es/es/actualidad/noticias/nuevas-nanoparticulas-biocompatibles-contra-el-cancer-de-mama

Novel drug delivery systems (DDS) are a crucial strategy to target specifically a smart therapeutic dose and local administration at a suitable rate, for an established period and at a set concentration. Generally, the drug can diffuse through paths or pores in polymer matrices or alternatively, the drug is physiosorbed on the nanocarrier surface and undergoes a very rapid release ("burst release").

As previously described, PLA is a great candidate for the design of drug delivery systems due to its degradable nature in humid environments and the nontoxic resorbed products or excreted by the human body. Drug release rates from PLA depend on a range of variables such as the physicochemical properties of the polymer, the type of formulation and the percentage
of drug loading. DDS generated from PLA block co-polymers feature progressive drug release profiles for long term treatments, due to the strong interactions with the therapeutic drug. A series of PLA derivatives were synthesized by ROP to formulate nanoparticles drug delivery systems of Doxorrubicin (DOX) and Dasatinib (DAS) to elucidate the relationship between the crystalline phase of the internal nanoparticles structure and the drug release. DOX has demonstrated very limited efficacy in the treatment of gliomas due to the short half-life of DOX in biological experiments. DAS is an FDA-approved multi-target compound for the treatment of chronic myeloid leukemia. However, the extremely low solubility of DAS diminishes its efficacy. The development of efficient delivery systems of DOX and DAS could potentially improve the therapeutic efficacy. Three different formulations (racemic-PLA, homopolymer-PLA and diblock-copolymer-PLA) for both drugs were recently synthesized by double emulsion and solvent evaporation method with a final hydrodynamic radius smaller than 200 nm, and loading efficacy (LE) results show important differences among them presenting higher loading efficacy for DOX than for DAS (see Table 4).

<table><caption>Table 4. Hydrodynamic radius (nm) and loading efficacy (LE) of three different formulations. SC: copolymer; RAC: racemic; HC: homopolymer.</caption><thead><tr><th rowspan="2">Formulation</th><th rowspan="2">Drug load (mg)</th><th colspan="2">DASATINIB</th><th colspan="2">DOXORUBICIN</th></tr><tr><th>Average size (nm)</th><th>LE (%)</th><th>Average size (nm)</th><th>LE (%)</th></tr></thead><tbody><tr><td>SC-PLA</td><td>0,5</td><td>147,4</td><td>0,01</td><td>130,9</td><td>0,11</td></tr><tr><td>RAC-PLA</td><td>0,5</td><td>143,1</td><td>0,03</td><td>140,4</td><td>0,10</td></tr><tr><td>HC-PLA</td><td>0,5</td><td>143,9</td><td>0,09</td><td>171,2</td><td>0,18</td></tr><tr><td>SC-PLA</td><td>1</td><td>133,0</td><td>0,13</td><td>144,1</td><td>0,25</td></tr><tr><td>RAC-PLA</td><td>1</td><td>191,4</td><td>0,13</td><td>125,1</td><td>0,10</td></tr><tr><td>HC-PLA</td><td>1</td><td>129,7</td><td>0,11</td><td>119,3</td><td>0,22</td></tr><tr><td>SC-PLA</td><td>1,5</td><td>133,2</td><td>0,03</td><td>133,4</td><td>0,16</td></tr><tr><td>RAC-PLA</td><td>1,5</td><td>127,3</td><td>0,01</td><td>160,7</td><td>0,16</td></tr><tr><td>HC-PLA</td><td>1,5</td><td>161,6</td><td>0,03</td><td>106,7</td><td>0,04</td></tr></tbody></table>

Different release profiles among formulations are expected from the kinetically trapped structure generated during the nanoparticles formation from the double emulsion system. Different crystalline domains were observed by thermal analysis during the 1st heating cycle of the formed nanoparticles after precipitation compared to the raw material. SAXS/WAXS previewed measurements will elucidate the effect of the nanostructure on the interaction between the DOX and DAS and the PLA derivatives and particularly, the effect of the drugs on the crystallinity of the PLA derivative that will determine the drugs release profile.

Finally, I am currently working in the first manuscript about the synthesis of the stereo-block PLA copolymers of the project together with the student Carmen Moya.
## 8- Study of the nanostructure of the fish gelatin hydrogels for drug delivery applications

Biomaterials are ideal candidates to supply sustainedly functionalised nanomaterials that in turn, could potentially substitute synthetic materials in dayly use applications. However, the structure-properties relantionship of biomaterials have been scarcely investigated. Particularly, marine bioresources are uncommonly employed for designing novel soft matter systems with revelant industrial applications. Likewise, I have been stimulated to contact the Institute of marine research of the CSIC organization (Spanish council research) to obtain fish gelatin to characterise the physcochemical properties and its structural correlation in collaboration with the University of Vigo. Gelatin is a biopolymer, widely used in the food, pharmaceutical, and and photographic industries, due to its excellent biocompatibility, biodegradability properties. Gel colloids offer functionality and versatility in a wide variety of applications such as hygienic products, agriculture, drug delivery systems, food additives, pharmaceuticals, biomedical. Recently, gelatin was proved to promote joint and skin health. The skin and bones of cattle and pigshave often been used commercially for the production of gelatin. However, in recent years, fish gelatin has become increasingly important as a natural alternative source due to the livestock crisis and the increased demand for non-bovine and non-swine gelatin for religious and social reasons. In addition, fish skin is an important by-product of the fish processing industry considered for a long-time waste and pollutant. The valorization of value-added products derived from fisheries waste sources is a key step towards the improvement of fishing industry environmental sustainability and as a contribution to the circular economy.

In particular, shark and tuna have been identified as source species to obtain gelatin due to the abundance as a by-product of the fisheries industries, and their properties have been subject to study.

However, Shark/Tuna fish gelatin possess structural inherent limitations to replace mammalian sources despite a similar composition. The main limitations of Shark and Tuna fish gelatin are the low melting point and the gel strength whilst possessing a higher viscosity than mammalian gelatin, although the relatively lower gel strength and melting temperature than mammalian gelatin, may provide a better alternative in the food industry. Food systems require low melting points and moderate storage modulus to simplify the flavor release and thus, provide better sensory characteristics.

Particularly, the shark fish gel formed is often unstable and has poor rheological properties, limiting its field of application. However, the physical properties of tuna skin gelatin are comparable to mammalian gels. Thermal analysis of Shark/Tuna fish gels has determined various phases according to water content and the temperature of formation. Importantly, glass transition temperature has been determined to correlate the structural to the physical properties such as crystallization temperature, adhesion, collapse, and molecular mobility of biological materials.

Moreover, gelatine-based fish and Bovine gels have been extensively studied in the field of drug delivery systems as a versatile biomaterial and have numerous advantageous capabilities such as biocompatibility, biodegradability, low antigenicity, and multifunctionality.

The general objective of this study is to characterize the nanostructure of Tuna/Shark gelatin based hydrogels to exploit them as drug nanocarrier vehicles.
Hydrogels are polymer networks extensively expanded by the incorporation of water. Hydrogels that are usually referred to as hydrophilic gels are covalent or physical networks of polymeric chains eventually found as colloidal gels [1], in which water is the dispersion medium.

The ability of hydrogels to absorb water arises from hydrophilic functional groups attached either to the main chain or as pendant groups to the polymeric backbone, while their resistance to dissolution arises from cross-links between network chains.

Natural polymers are popular for current biomedical applications as are typically derived from living organisms. Moreover, natural polymers are generally non-toxic, biocompatible, and produce negligible inflammatory response in the host organism due to their biological origin [2], [3].

The desired mechanical property of the hydrogel can be achieved, by changing the degree of crosslinking, Hence, there is an optimum degree of crosslinking to achieve a relatively strong and yet elastic hydrogel.

Current physicochemical studies extensively focus on the understanding of the structure-
properties of natural polymers to design bioinspired functional nanomaterials. In particular,
gelatin has been an attractive candidate for preparing hydrogels with long-term biomedical
applications due to the large number of functional groups in the backbone and the facility to
be crosslinked to enhance their mechanical properties[2].

As previously mentioned, the vast majority of commercial gelatin is obtained from mammalian species in particular from pig, cattle and bovine gelatins [4]. However, the increased demand for an alternative to gelatin sources is increasingly required due to socio-cultural and religious reasons related to mammalian species. By-products from the fish processing industry represent a potential and promising source for gelatin production [5], [6].

Moreover, novel environmental and economic policy guidelines promote the rational use of marine resources, and in particular, the revaluation of fishing industry waste. Fish is an excellent nutritional source due to the high content of proteins in the muscles as well as high levels of essential amino acids. Collagen is the main structural protein in the animal kingdom that is estimated to represent the 75% of the total weight in fish processing industry residue [7], and thus further processing to yield gelatin can help to offset harmful environmental effects.

Moreover, a large part of the by-products generated during the fish filleting consists of skin and bones with large amounts of collagen that possess a wide range of applications[8].

Furthermore, obtaining valuable by-products from fisheries and reducing industrial waste is an attractive research topic due to sustainability issues. The effective use of fishery resources turns increasingly important and thus the demand for more efficient use of fish waste will only increase. Moreover, fish gelatin materials comply with the recommendations of the European Union to develop the Circular Economy [9], specifically, in the field of fishing, aquaculture, and exploitation of marine resources.

Sustainable development is an overarching objective that requires an interdisciplinary strategy to address the societal challenge concerning climate action, environment, resource efficiency, and raw materials. In a green environmental society valorization of abundant and available biowastes with high potential to manufacture value-added products is the first step to close the loop between waste and consumption in line with the main goal of the circular economy.

Collagen is mainly composed by glycine, proline, and hydroxyproline, which denature in the presence of dilute acid standards (% H2SO4, citric acid), turning in a soluble protein such as
gelatin in dissolved solutions [10]. Several methods are used by the industry to manufacture gelatin from collagen. In general, gelatin is obtained using a sequence of three types of processing steps: pretreatments to remove noncollagen impurities and prepare the collagen for extraction, one or more water extraction steps to convert collagen into gelatin, and finally, a series of refinement and recovery processes to get a highly purified dried gelatin. The main step of the gelatin production process is firstly to remove undesirable side-products that will interfere with the gelatin extraction. Subsequently, the water insoluble collagen thus obtained, is converted into the water soluble gelatin while maintaining maximum yield and functional properties [11].

For gelatin production the raw material may be any collagen-containing tissue. Hides, skins and bones from mammalian sources such as porcine and bovine are preferred, but gelatins are also produced from the skins of cold and warm water fish species as well as minor quantities from avian sources.

Different mammalian and fish gelatin sources. Mammalian gelatins have significantly superior physical properties such as higher viscosity, melting and setting temperature, and faster setting time compared to fish gelatin counterparts. Porcine skin-derived gelatin is the most popular source (46%), followed by bovine skin (29.4%), bovine bone (23.1%) and other sources (1.5%) [12].

The lower content of amino acids in fish gelatin is the main structural difference with the mammalian counterparts. The higher amino acid sequence enhances the ordered interchain interactions that stabilises the network formed in a hydrogel. Fish gelatin gel contains fewer physical cross-links as a result, between clusters compared to mammalian gelatin, or alternatively exhibits weaker interaction between clusters [13].

The mechanical properties of most commercial gelatins, derived from mammalian sources, has been conducted. However very little work has been performed on gelatins from fish origin.

Importantly, the unhydrolyzed collagen coil-helix or renaturation temperature that will define the mechanical properties, depends upon the proportion of the iminoacid in both proline and hydroxyproline.

The iminoacid proportion is generally, 24% for mammals and 16%–18% for most fish species. The collagen from cold-water fish, has a very low hydroxyproline content. Consequently, the gelatin extracted from cold water fish collagen has a very low gelling and melting temperature [14]. Mammalian gelatin possesses a gelation temperatures values of 25°C, in contrast at warm-water fish gelatin that features a gelation temperature of 29°C and cold-water fish gelatin that gelates at 10°C.

Mammalian gelatin is characterised by the highest concentration of proline and hydroxy proline that help to stabilize the triple helical structures. The triple helical structures act as junctions in a network that eventually forms the gel[15].

The applicability, functionality, and commercial value of fish gelatins depend essentially on physicochemical properties such as viscosity, strength and gel rigidity, intumescent capacity, thermal stability, as well as pH. The physicochemical properties are influenced in particular by amino acids composition, molecular weight, and proportion of α-chains.

The melting and gelling temperatures (the temperature at which a gelatin solution changes from solid to liquid and vice versa) of fish gelatin hydrogel are relatively low compared to cattle and pig gelatins due to their smaller amounts of proline and hydroxyproline [16]. Gelatin fish is