0.5 h. HPLC analysis indicated that there is a greater conversion of indole in presence of RF compared to **FLPDA,** and similar products although in different yields, were observed as shown in **Figure 6** and **Table 1**. We next compared the yield of dyes after the reactions through the addition of 2 drops of 1M HCl to each reaction mixture to initiate dimerization of hydroxylated

Crocker and Fruk Flavin-Polydopamine Hybrid Photocatalyst

TABLE 1 | Amounts and conversions of major compounds identified by HPLC after acidification of reaction mixture using commercial standards as a calibration reference.

\*([*indigoid*] × *2*/[*consumed indole*]) *(Xu et al., 2012) with [indigoid] being determined by UV-Vis spectroscopy.*

indole species to form the indigoid dyes. As shown in **Figure 7B**, there was more indirubin obtained from the reaction conducted in the presence of **FLPDA** than **RF**, which was expected as 2.5 times more precursor **10a** can be detected in HPLC profile shown in **Figure 6**. Clearly, despite the decreased indole conversion, **FLPDA** appears to be more selective than its homogeneous counterpart toward dye conversion.

It is also worth noting that after the irradiation, **RF** could no longer be identified by TLC analysis within the reaction mixture and had degraded to lumichrome (LC), which was confirmed both by HPLC and LC-MS (9.22 min, m/z: 243) (**Figure 6**). Although there was some photodecomposition observed in **FLPDA**, the concentration of lumichrome was five times lower compared to **RF** (based on HPLC integration), indicating that the hybrid catalyst offers more protection toward photodecomposition of active flavin centers, and mitigates the loss of catalysis observed in homogenous systems.

### Further Control Experiments and Possible Mechanism of Action

As this is a novel hybrid system, we were also interested in unveiling the photocatalytic mechanism of **FLPDA.** To achieve this we employed a superoxide radical scavenger TEMPO and the singlet oxygen scavenger DABCO to quench these prominent reactive oxygen species (ROS) that may be liberated from reduced and/or photoexcited flavin species (Massey et al., 1969; Müller and Ahmad, 2011). As can be seen from the HPLC chromatograms of the reactions shown in **Figure 8**, the addition of these quenchers did not appear to inhibit the reaction completely. In fact, we observed that in the case of DABCO, the basic character of the species may have played a role in lowering the activity of **FLPDA** as flavins and PDA are less stable within a basic environment (Song et al., 1965; Yang et al., 2018). In terms of product selectivity, **10a** was not detected most likely due to direct oxidization to isatin **11** in basic condition (**Table 2**). Based on these observations it could be assumed that singlet oxygen plays a minimal role in **FLPDA** activity as both compounds **9** and **11** are obtained in the presence of singlet oxygen quencher.

The addition to the superoxide scavenger, TEMPO had no deleterious effect on **FLPDA** activity either and in fact, it enhanced the production of **9, 10a**, and **11**. In fact, TEMPO most probably acts as a redox mediator and co-catalyst in the reaction, as previously shown for the synthesis of isatin derivatives in the presence of hypervalent iodine (Sai Prathima et al., 2015). TEMPO's participation in the reactions was additionally proved by appearance of the by-product 2,2,6,6-tetramethylpiperidine (TMP) at a retention time of 2.49 min, m/z: 142 (LC-MS analysis). It should be noted that pigment melanin, composed of PDA units, is capable of superoxide quenching (Tada et al., 2010), and it is safe to assume that any superoxide radical generated by flavin would be quickly quenched by PDA in its immediate proximity.

We further investigated the mechanism of **FLPDA**'s photocatalytic activity by using a fluorescence-based assay to monitor the release of H2O2 liberated from the unstable 4ahydroperoxy-flavin species formed when reduced flavin moieties interact with oxygen (Usselman et al., 2014). We reasoned that the lower amount of H2O2 obtained in photo-reaction in presence of **FLPDA** compared to free flavin (riboflavin **RF)**, would indicate potential stabilization of 4a-hydroperoxy-flavin species, and its interaction with indole (**Scheme 5**) instead of direct transformation into reactive oxygen species. To minimize the loss of photons due to the scattering effect of a colloidal heterogeneous catalyst, low concentration of **FLPDA** (10µg/mL) was used, and the concentration of indole and effective concentration of free flavin were adjusted accordingly. The amount of H2O2 released after 1 h of irradiation was 3.62 ± 0.18 nmol for **FLPDA** and 5.13 ± 0.04 nmol for **RF** (see **Figure S4**). The lower value obtained for **FLPDA** suggest that the stabilization of 4a-hydroperoxy flavin species in **FLPDA** may occur, although this hypothesis can only be fully confirmed after the completion of already initiated EPR study.

Based on current data, we propose the mechanism shown in **Scheme 5**, which includes the formation of the 4a-hydroperoxy species as a result of indole and/or PDA photoreduction, and subsequent interaction with oxygen. This species could be stabilized by H-bonding with neighboring PDA moieties and be stable long enough to interact with the indole substrate. In addition, the appearance of compounds **12** and **13** indicates Type 1 photosensitization of indole through electron transfer to excited flavin moieties similar to that previously observed for riboflavin

TABLE 2 | Amounts and conversions of major compounds identified by HPLC after acidification of reaction mixture using commercial standards as a calibration reference.

and tryptophan (Silva et al., 2019). However, as stated above further studies using EPR spectroscopy are being undertaken to confirm our theory.

Although some "dark" reactions in the presence of flavin species have resulted in excellent yields of desired products (Arakawa et al., 2017; Chevalier et al., 2018), use of light to trigger the reaction cascade that leads to complex dye formation would have multiple advantages in an industrial setting, including temporal and spatial control, ultimately leading to the design of more efficient reactors.

### Recyclability

Finally, we were interested to estimate the reusability of **FLPDA** in this light-driven reaction as this is a desirable property for scaling up and industrial applications enhancing the green potential of the system. Previous work on a heterogeneous flavin-based photocatalytic system using mesoporous silica to immobilize flavin moieties suffered from a severe loss of activity after the first reuse, and a complete inactivation upon second and third attempt (Špacková et al., 2017 ˇ ). For our nanoparticle-based system, the catalyst was easily removed from the reaction mixture via centrifugation and maintained activity for 4 cycles. As a proof of concept study, activity was monitored by irradiation of **FLPDA** particles (50µg/ml) with indole (1.0 mM) for 1.5 h and measuring the changes in absorption of the reaction mixture between 300 and 600 nm as shown in **Figure 9A**. The area within this region was integrated and used as a measure for oxidized indole species production. The first run showed a decrease in relative activity compared to the initial run only by a factor of 1.36 (**Figure 9B**) and the subsequent three runs stayed within a similar range before minimum activity, although not a complete loss, was observed by run 6.

The enhanced stability of our system could be explained by the electron transfer from PDA to flavin moieties upon irradiation as proposed in our mechanism (**Scheme 5**). It has already been shown that an addition of an electron donor, such as 2- (N-morpholino)ethanesulfonic acid (MES) buffer, limits photoinduced decomposition reactions that curtail flavin activity upon irradiation (Alonso-De Castro et al., 2017). However, during the recycling steps, the number of retrieved particles decreased with each run. Loss of particles and the catalytic activity could point toward the photo-degradation of PDA moieties similar to what was previously observed in melanin samples (Ito et al., 2018). However, we did not notice any difference in the particle size and morphology after 2 h of irradiation (**Figures 9C,D**), which

disputes this argument, and indicates a mere physical loss of particles during recovery procedures.

Our ongoing work is focused on exploring the long-term stability and photo-degradation of **FLPDA**, as well as establishing a detailed mechanism using EPR studies. We believe this information will allow us to rationally design better catalytic systems with enantio-selective potential. We are also currently working on improving the reusability of **FLPDA** by addition of magnetic Fe3O4 nanoparticles during co-polymerization to enable magnetically aided retrieval, which has already proven effective in various application using PDA (Liu et al., 2013, 2017; Xie et al., 2014; Yu et al., 2016).

### CONCLUSION

We have developed a novel photocatalytic organic nanoparticle system made from flavin conjugated polydopamine. This was achieved through a convenient co-polymerization method involving dopamine and a flavin-dopamine analog, **FLDA**, to form **FLPDA**. The system effectively displays activity analogous to xenobiotic-degrading enzymes such as FMO, shown, using the oxidation reaction of indole, to form the indigo dye. In addition, it employs irradiation with visible light rather than an external cofactor for catalyst activation. The products observed in the reaction; indigo, indirubin, and related compounds, clearly indicated FMO-like activity of **FLPDA**. Additional ROS scavenging studies as well as the nature of some products, further confirmed that **FLPDA** displays mechanistic similarities to natural FMO, which is characterized by the formation of 4a-hydroperoxy flavin as a reactive intermediate. In contrast to our heterogeneous hybrid flavin, free flavin in a solution was characterized with higher photo-degradation and lower dye yields. In addition, **FLPDA** can easily be recovered from the reaction mixture and displays recyclable activity up to four cycles. We hope that this initial work sets the foundation to enable the design of more elaborate photocatalytic biopolymer hybrids comprising of flavin, PDA and additional functional moieties, to be applied in heterogeneous catalysis and artificial photosynthesis.

### REFERENCES

### AUTHOR CONTRIBUTIONS

LC carried out all the synthesis, characterization and analyses, and wrote the draft manuscript. LF was involved in the experimental planning and manuscript proofreading and corrections. LF was the academic lead of the project in charge of experimental planning and funding.

### FUNDING

The project was funded by the Department of Chemical Engineering and Biotechnology Start-up grant awarded to LF. LC would like to acknowledge support for DTP-EPSRC doctoral programme (EP/N509620/1).

#### ACKNOWLEDGMENTS

We would like to thank Dijana Matak, as well as Rhian Preece, Nitin Rustogi, and Sabine Bahn for their help with LC-MS analysis.

#### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2019.00278/full#supplementary-material

support and carbon adsorbent. ACS Appl. Mater. Interfaces 5, 9167–9171. doi: 10.1021/am402585y

intensity 365 nm light emitting diode. Free Radic. Biol. Med. 131, 133–143. doi: 10.1016/J.FREERADBIOMED.2018.11.026

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Crocker and Fruk. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

#### Chemical and Process Engineering (formerly Chemical Engineering)

Camille Petit

Dr Camille Petit is a Reader in the Chemical Engineering Department at Imperial College London. She leads the Multifunctional Materials Laboratory. She was a postdoctoral researcher at Columbia University and obtained her PhD in Chemistry from the City University of New York. Her research interests encompass the development of porous materials for energy and environmental applications. She focuses on metalorganic frameworks and nitride-based materials for molecular separations and solar energy conversion. She received the 2017 AIChE's 35 under 35 award, the 2017 ExxonMobil European Science & Engineering Program Award, the 2017 IOM3 Silver Medal, and the 2015 IChemE Sir Frederick Warner medal.

Kelly M. Schultz

Dr Kelly M. Schultz is an Associate Professor in the Department of Chemical and Biomolecular Engineering at Lehigh University. She obtained her PhD in Chemical Engineering advised by Eric Furst from the University of Delaware. She was a Howard Hughes Medical Institute postdoctoral research associate at the University of Colorado at Boulder working for Kristi Anseth. She began her position at Lehigh in 2013. She was named a TA Instruments Distinguished Young Rheologists in 2014, received an NSF CAREER Award in 2018, and a Libsch Early Career Research Award in 2019. Her group studies emerging hydrogel materials developed for consumer products and biological applications.

#### Florent Allais

Florent Allais is a Professor in Chemistry at AgroParisTech, Paris, France and the Director of the URD ABI in Pomacle, France. He obtained his PhD from the University of Florida in 2004 and his postdocs at ESPCI, Paris, France and ICSN, Gif-sur-Yvette, France. He has published approximately 80 papers and filed 16 patents. His multidisciplinary research team combine biotechnologies, green chemistry, and downstream processing approaches to develop durable processes to turn agro-resources into platform molecules as well as valuable chemicals that can be used in chemistry, in food/feed or cosmetics to name a few.

#### Tsuguyuki Saito

Dr Tsuguyuki Saito is an Associate Professor of the University of Tokyo, Japan. He obtained his PhD degree in 2008 from the University of Tokyo for his study on TEMPO-mediated oxidation of cellulosic fibers. He has further developed this oxidation as a tool to produce the finest type of cellulose nanofibers (CNF) with a width of approximately 3 nm from wood pulp in an energy-efficient way. For his ground-breaking discovery on CNF production, he was awarded the Marcus Wallenberg Prize by the Marcus Wallenberg Foundation in 2015, and the JSPS prize by the Japan Society for the Promotion of Science in 2019.

# Porous Boron Nitride Materials: Influence of Structure, Chemistry and Stability on the Adsorption of Organics

#### Sofia Marchesini, Xiyu Wang and Camille Petit\*

Department of Chemical Engineering, Barrer Centre, Imperial College London, London, United Kingdom

Porous boron nitride (BN) is structurally analogous to activated carbon. This material is gaining increasing attention for its potential in a range of adsorption and chemical separation applications, with a number of recent proof-of-concept studies on the removal of organics from water. Today though, the properties of porous BN—i.e., surface area, pore network, chemistry—that dictate adsorption of specific organics remain vastly unknown. Yet, they will need to be optimized to realize the full potential of the material in the envisioned applications. Here, a selection of porous BN materials with varied pore structures and chemistries were studied for the adsorption of different organic molecules, either directly, through vapor sorption analyses or as part of a water/organic mixture in the liquid phase. These separations are relevant to the industrial and environmental sectors and are envisioned to take advantage of the hydrophobic character of the BN sheets. The materials were tested and regenerated and their physical and chemical features were characterized before and after testing. This study allowed identifying the adsorption mechanisms, assessing the performance of porous BN compared to benchmarks in the field and outlining ways to improve the adsorption performance further.

#### Edited by:

Gil Garnier, Bioresource Processing Institute of Australia (BioPRIA), Australia

#### Reviewed by:

Petra Foerst, Technische Universität München, Germany Xiao-Yu Wu, Massachusetts Institute of Technology, United States

> \*Correspondence: Camille Petit<EMAIL_ADDRESS>
#### Specialty section:

This article was submitted to Chemical Engineering, a section of the journal Frontiers in Chemistry

Received: 28 November 2018 Accepted: 04 March 2019 Published: 26 March 2019

#### Citation:

Marchesini S, Wang X and Petit C (2019) Porous Boron Nitride Materials: Influence of Structure, Chemistry and Stability on the Adsorption of Organics. Front. Chem. 7:160. doi: 10.3389/fchem.2019.00160 Keywords: boron nitride, vapor sorption, adsorption, separations, water stability

# INTRODUCTION

Chemical separations account for more than half of the total energy consumed in industry owing to the many separation processes performed via distillation (Sholl and Lively, 2016). Adsorbent materials or membranes, which are able to separate molecules based on size exclusion principles and/or chemistries, offer a possibly less energy intensive and hence more sustainable route to separations (Sholl and Lively, 2016). Realizing the potential of adsorption and membrane separation requires the production of porous materials with tailorable properties and porosities. Today, research in this area primarily focuses on materials such as metal-organic frameworks (James, 2003), zeolites (Misaelides, 2011), carbon-based materials (Dias et al., 2007), covalent organic frameworks (Côté et al., 2005), and polymer of intrinsic microporosity (McKeown and Budd, 2006). A recent addition to this list is porous boron nitride (BN). Owing to its high thermal stability, "rich" chemistry, and high surface area [up to ∼2,000 m2 g −1 reported so far (Li et al., 2013b; Marchesini et al., 2017b)], porous BN has been tested for a range of different applications such as catalysis (Venegas et al., 2017), CO2 capture, oil spill clean-up (Lei et al., 2013), hydrogen storage (Portehault et al., 2010; Lei et al., 2014; Weng et al., 2014), and water cleaning (Zhang et al., 2012; Li et al., 2013a,b; Lei et al., 2015). Porous BN is somewhat analogous to activated carbon in terms of its turbostratic to amorphous structure with 6-member rings of alternating B and N atoms. Yet, its chemistry is obviously different from that of carbonaceous materials as the material is made of polar B-N bonds. BN materials are typically hydrophobic. Such a feature can be particularly advantageous for the adsorption of organics, either in the vapor phase or as part of a water/organic mixture in the liquid phase. In fact, a number of studies report the use of porous BN for the removal of organic molecules (e.g., dyes, pharmaceutical molecules, oils) from water Zhang et al., 2012; Li et al., 2013a,b; Lei et al., 2015. The sorption mechanisms are typically associated to: (i) π-π stacking interactions between the BN sheets and aromatic dye molecules (Zhang et al., 2012; Liu et al., 2014, 2015); (ii) electrostatic dipole-dipole interactions between the polar B-N bond and the polar dyes (Liu et al., 2014; Lin et al., 2016); (iii) physisorption in the micropores of BN (Liu et al., 2014; Lin et al., 2016), and (iv) hydrophobic interactions (Yu et al., 2018). The adsorption capacity can be highly dependent upon the pH of the solution in the case of dye sorption, due to electrostatic interactions with the polar dyes, which are weakened by competitive adsorption of protons or hydroxides (Singla et al., 2015; Li et al., 2016). For a similar reason, OH-functionalised BN adsorbs more cationic dyes (e.g., methylene blue) than its pristine counterpart (Li et al., 2015). Porous BN materials exhibit competitive sorption capacities of selected organics compared to other common adsorbents such as activated carbons (Yu et al., 2018). Interestingly, many studies on the removal of organics from water using porous BN, report multiple adsorption and desorption cycles with minimal (5–15%) decrease in performance (Li et al., 2013b, 2015; Lei et al., 2015; Zhao et al., 2015; Liu et al., 2016). This behavior is surprising given the demonstrated instability of some porous BN structures in the presence of water, which causes a decrease in porosity (Shankar et al., 2018). This raises a question around the effect of BN porosity in these applications. Overall, the studies above serve as proof-of-concept for the use of porous BN as adsorbent for a given organic. Today though, the materials parameters—i.e., surface area, pore network, chemistry—that dictate adsorption of specific organics remain vastly unknown.

Herein, we synthesized BN samples exhibiting a range of pore structures and crystallinity as well as distinct chemistries (Marchesini et al., 2017a; Shankar et al., 2018) The pore structure of the materials was controlled via manipulation of the precursor type and ratio used to synthesized BN (see our previous study Marchesini et al., 2017a). We tuned the crystallinity and chemistry—particularly the amount of O groups—using the synthesis temperature (Shankar et al., 2018). We tested all the porous BN samples for the adsorption of different organic molecules, either directly, through vapor sorption analyses or as part of a water/organic mixture in the liquid phase. The organics selected for this study cover different classes of chemicals in terms of polarity and molecular size: alkanes (linear and branched), cycloalkanes, aromatics alcohols, and dyes. We characterized the samples before and after adsorption and linked their materials features to the adsorption behaviors.

#### Materials Synthesis

We synthesized porous BN materials following a procedure previously developed by our group Marchesini et al., 2017b. Boron- and nitrogen- containing precursors were mixed and ground in different molar ratios. The precursors were placed in an alumina boat crucible and heated up to the 1050 or 1500◦C for 3.5 h (10◦C/min ramp rate) under nitrogen gas flow (50 cc/min during analysis, 3 h at 250 cc/min to purge). The furnace was cooled while flowing nitrogen until ambient temperature. We refer to porous BN samples synthesized at 1050◦C as "low temperature BN" samples, and to porous BN samples synthesized at 1500◦C as "high temperature BN" samples. We used three different sample formulations as follows:

#### Materials Characterization

Materials characterization techniques were used to: (i) verify that the synthesized samples cover a range of porosity, crystallinity and chemistry as envisioned initially and (ii) assess the impact of adsorption on the materials features and (iii) extract information of the adsorption mechanisms.

#### Structural Properties and Morphology

Nitrogen physisorption isotherms were measured using a porosity analyser (Micromeritics 3Flex) at −196 ◦C. Prior to nitrogen sorption analysis, the samples were degassed overnight at 120 ◦C at roughly 0.2 mbar pressure and degassed again in-situ on the porosity analyser for 4 h down to around 0.003 mbar at 120 ◦C. From the nitrogen isotherms, we derived the surface areas, total pore volumes, micro- and mesopore volumes of BN samples and pore size distributions (Rouquerol et al., 2013). The surface areas of the samples were calculated using the Brunauer-Emmett-Teller (BET) method (Brunauer et al., 1938). The total volume of pores was calculated from the volume of N2 adsorbed at P/P0 = 0.97. The micropore volume was determined using the Dubinin Radushkevich method (Chen and Yang, 1994). The pore size distributions were calculated using a DFT model for carbons with slit-shape pores using the Micromerics software.1 We assessed the crystallinity (or lack of) of our samples via powder X-ray diffraction (XRD) using an Xray diffractometer (PANalytical X'Pert PRO) in reflection mode. The operating conditions included an anode voltage of 40 kV and an emission current of 40 mA using monochromatic Cu Kα radiation (λ = 1.54178 Å).

#### Chemical Properties

The porous BN samples were characterized by Fourier Transform Infrared (FT-IR) spectroscopy. The samples were first ground in

1http://www.nldft.com

an agate mortar and spectra were collected in the range of 600– 4000 cm−1 using a Perkin-Elmer Spectrum 100 Spectrometer equipped with an attenuated total reflectance (ATR) accessory. X-ray Photoelectron Spectroscopy (XPS) was performed using a Thermo Scientific K-Alpha+ X-ray Photoelectron Spectrometer equipped with a MXR3 Al Kα monochromated X-ray source (hν = 1486.6 eV). X-ray gun power was set to 72 W (6 mA and 12 kV). All high resolution spectra (B 1s, N 1s, C 1s, and O 1s) were acquired using 20 eV pass energy and 0.1 eV step size. The samples were ground and mounted on the XPS sample holder using conductive carbon tape. Thermo Avantage software (ThermoFisher Scientific) was used to analyse the data. The XPS spectra were shifted to align the peak for adventitious carbon (C-C) at 285.0 eV.

#### Materials Testing Organic Vapor Sorption

We conducted these analyses to assess the "idealized" (i.e., single adsorbate, no competitive adsorption) adsorption of several molecules with different polarities and molecular sizes. We performed vapor sorption experiments on an IGA gravimetrical analyser (Hiden Isochema). The IGA analyser is equipped with a microbalance with ± 1 µg resolution. The samples (∼30 mg) were degassed in-situ at 120 ◦C for 4 h using a furnace attachment, prior to analysis. An ultra-high vacuum of up to 10−6 mbar could be created inside the chamber. Solvents were purified by cycles of evacuation followed by vapor equilibrium, repeated for at least 10 times. The solvents were vaporized and dosed up to 90 % of their saturation pressures at 25 ◦C: n-heptane (P0: 61.12 mbar), 2,3-dymethylpentane (P0: 91.82 mbar), toluene (P0: 37.93 mbar), methylcyclohexane (P0: 61.79 mbar), 1-butanol (P0: 9.22 mbar), methanol (P0: 169.24 mbar). The temperature was controlled using an external water bath equipped with a thermocouple which maintained a temperature stability of ± 0.3 ◦C. Equilibrium for each pressure point was predicted by the IGA software after a minimum wait of 20 min, by checking for weight changes using the least squares regression to extrapolate the asymptote (tolerance of 99 %). Pressure increase/decrease was performed with a rate of ∼0.25–1 of the step pressure change. All solvents were of analytical grade. We performed in situ-degas in between each experiment at 250 ◦C to regenerate porous BN. Solvent selectivity was calculated as the ratio of quantity adsorbed of two pure solvents measured under the same conditions, assuming equal partial pressures.

#### Water Vapor Sorption

Prior to testing the removal of organics from water, we assessed water sorption to evaluate the affinity of the samples with water and quantify their water adsorption capacity. Porous BN samples were degassed at 120 ◦C at roughly 0.2 mbar pressure and again in-situ on the porosity analyser (Micromeritics 3Flex) for 4 h down to around 0.003 mbar. We collected water vapor isotherms at 30 ◦C for porous BN samples at up to 90% relative humidity in a gas/vapor volumetric analyser.

#### Organics Removal From Water

We used a dye [Rhodamine B, RhB (HPLC, Sigma Aldrich)] as a representative organic molecule owing to the ease of monitoring dye sorption using UV-vis spectroscopy. In a typical test, 250 mL of stock solution (of 40 mg/L in deionized water) were vigorously stirred and 100 mg of porous BN powder were added to the solution while stirring. At 0, 5, 10, 30, 60, 120, 240 min, about 2 mL of solution were collected with a syringe and filtered with a syringe filter (0.45µm) to remove BN powder. We then analyzed these solutions using UV-vis spectrophotometry to determine the dye concentration. UV-vis spectra were collected between 400 and 650 nm (λmaxRhodamine = 554 nm) at a scan speed of 480 nm/min using a Perkin Elmer Lambda 35 spectrophotometer with quartz cuvettes. Calibration was performed by preparing solutions with decreasing dye concentrations (see calibration curve in **Supplementary Figure 1**, R2 > 0.99). After 240 min, BN samples were filtered using a Buchner filtration system, dried in the oven at 120 ◦C and then heated at 600 ◦C (10 ◦C/min) for 3 h in a tubular furnace under air gas flow (50 mL/min) before being reused for a second adsorption cycle. We selected this regeneration temperature based on the decomposition temperature of the dye, as determined by TGA analysis (Netzsch TG209 F1 Libra, 10 ◦C/min ramp rate, under air gas flow at 0.1 L/min; **Supplementary Figure 2A**).

### RESULTS AND DISCUSSION

First, we confirmed the "profiles" of the materials synthesized using a number of characterization techniques to assess the range of porosity, crystallinity, and functionalization. As expected, the porous BN materials were all micro/mesoporous but displayed different pore volumes and BET surface areas (**Table 1**; **Figure 1A**). The samples covered a range going from 734 to 1650 m2 g −1 in terms of surface area and 31 to 62% in terms of microporosity volume compared to total porosity. For all samples, the surface area and porosity decreased upon thermal treatment, as expected (Shankar et al., 2018). All samples contained oxygen atoms as part of their structure, as determined from XPS analyses (**Table 1**) and high temperature porous BN samples generally exhibited a lower oxygen content compared to low temperature porous BN samples. All porous BN samples displayed very similar XRD patterns with the same peak position for the (002) peak between the different samples, indicating similar d (002)-spacing (∼0.35 nm) (**Supplementary Figures 3A,B**). The absence of other reflections other than the (002) and (100) and the large width and low intensity of the peaks confirmed that all samples were turbostratic. Despite a similar lack of crystallinity in the bulk phase between the low and high temperature BN samples, our prior work identified localized regions of increased crystallinity and thereby hydrophobicity in the high temperature BN sample (Shankar et al., 2018). We assessed the effect of these crystalline and hydrophobic regions on the adsorption behavior below.

We then tested the low temperature porous BN materials (different pore structures and similar chemistries) for adsorption of C-7 hydrocarbons exhibiting different sizes/shapes (e.g., linear n-heptane or branched 2,3-dimethylpentane) and/or chemistries (e.g., aromatic toluene or non-aromatic methylcyclohexane) and two different alcohols (methanol and 1-butanol). Through this study, we intended to identify links between the materials physical properties, including porosity, and the adsorption TABLE 1 | Overview of the porosity and chemistry of the porous BN samples: textural parameters derived from nitrogen sorption isotherms at −196 ◦C and relative atomic percentages derived from XPS analysis.

performance. The results are presented in **Figure 1B** for methylcyclohexane and in **Supplementary Figure 4** for all other chemicals. The three samples exhibited different sorption isotherms. Given that their chemical composition was similar, the differences are attributed to differences in physisorption dictated by the pore structure of the materials. In fact, the organic vapor sorption isotherms followed the same patterns as the nitrogen sorption at −196 ◦C (compare **Figures 1A**,**B**). The total vapor sorption capacity increased with increasing surface areas and most of the vapors were adsorbed at low relative pressures, in the microporous regions. The presence of mesopores caused capillary condensation (e.g., see BN-MU0.25:5) as indicated by the presence of hystereses in the desorption curves. Most of the vapors desorbed (**Figure 1B**) as the pressure was lowered, confirming that little/no chemisorption occurred.

We also evaluated water sorption for different low temperature porous BN samples at 30 ◦C (**Figure 1C**). Water uptake increased with the surface areas. Importantly, the water did not completely desorb upon reducing the pressure as observed from the "open" hystereses. As highlighted in a previous study (Shankar et al., 2018), this behavior originates from the hydrolysis of the BN materials through the following reaction: 2BN + 3H2O → B2O3 + 2NH3. The hydrolysis of porous BN takes place due to the presence of defects in the structure (Motojima et al., 1982; Alkoy et al., 1997; Streletskii et al., 2009; Cao et al., 2012; Shankar et al., 2018). Upon thermal treatment, as mentioned above, we form localized regions of increased crystallinity and hydrophobicity. Hence, the high temperature porous BN showed a reduced water uptake capacity, especially at low relative pressures (**Figure 1C**). At higher pressures, the remaining defects sites react with water leading to decomposition of the material (Shankar et al., 2018).

From the organic vapor sorption isotherms, we calculated the sorption capacity at 0.6 P/P0 (**Figure 1D**). At this relative pressure, organic vapor uptake had ceased and BN materials were saturated. Sorption capacities of low temperature porous BN samples increased with increasing micropore volumes for all organics considered (**Figure 1D**). We found a linear trend between BN micropore volume and sorption capacity for porous BN materials with similar chemistries. This aligns with the behavior observed in activated carbons (Lillo-Ródenas et al., 2005). All porous BN materials adsorbed more toluene than any of the other tested adsorbates. We hypothesized that this is due to the presence of π-π interactions between the h-BN rings in porous BN and the aromatic rings in toluene, which contribute to improved physisorption. To put the results in perspective, porous BN exhibited toluene and methanol sorption capacities of the same order of magnitude to activated carbons with similar surface areas (toluene: 700 vs. 950 mg g−1 ; methanol: 600 vs. 800 mg g−1 ) (Xu et al., 2015). While the materials absorbed large amounts of organics, they were not selective (**Supplementary Figure 5**; linear n-heptane vs. branched 2,3-dimethylpentane; aromatic toluene vs. non-aromatic methylcyclohexane). This is related to the multimodal pore size distribution of the material i.e., no narrow pore size distribution—that prevent size exclusion (**Supplementary Figure 6**). Similarly, we tested a high temperature porous BN sample for organic vapor sorption to evaluate the effect of the localized hydrophobic regions of the material (**Figure 1D**). Despite its lower porosity compared to the low temperature BN samples, the high temperature BN sample exhibited increased uptake of all organics. Further support for this observation is provided below when investigating dye sorption.

The adsorption potential of all porous BN materials was finally evaluated for a practical application. We employed

the porous BN samples (low and high temperatures) for the removal of organics (here Rhodamine B, RhB) from water. The results are presented in **Figure 2**. All samples achieved complete removal of the dye within 4 h or less (**Figures 2A,B**). The time needed to reach complete removal reduced or remained unchanged (for BN-U5) when using the high temperature samples despite their initially lower surface area. Again, high temperature porous BN materials displayed a higher adsorption capacity for organics. We attribute this observation to the increased hydrophobicity and water stability of the high temperature BN samples, as well as to the presence of a larger quantity of π electrons, which allowed greater affinity for the adsorbate while preventing in-situ collapse of the material.

We regenerated the better performing low and high temperature samples (BN-MU1:5 and BN-MU0.25:5, respectively) at 600 ◦C in air and tested them for a second adsorption cycle. The results are reported in **Figure 2C**. While the low temperature BN sample exhibited a decreased capacity after regeneration (10 % decrease at 4 h), the high temperature BN sample maintained its original capacity. This is related to the greater stability of the high temperature sample in the aqueous environment compared to the low temperature sample. Analyses are presented below to confirm this aspect. Interestingly, the high temperature porous BN sample adsorbed more dye in the 2nd cycle than in the 1st cycle for the same exposure time. We explain this by the greater dispersibility of regenerated sample. Indeed, after regeneration and prior to the second sorption run, the sample exhibited finer particles.

We characterized the BN samples after the 1st cycle of dye removal from water and subsequent regeneration in air at 600 ◦C. This was done to confirm whether a collapse (partial or complete) of the porosity occurred and influenced the adsorption in the 2nd cycle as observed above. To quantify the impact of this potential collapse on the adsorption behavior, both the low and high temperature BN samples were analyzed—the high temperature BN samples being a priori less prone to collapse. We note that TGA and XPS analyses confirmed that the dye was completely removed upon regeneration in air at 600 ◦C (**Supplementary Figure 2**): no carbon residues were observed after regeneration.

First, we analyzed the changes in porosities of the BN samples before testing and after regeneration (**Figure 3**). All low temperature porous BN samples showed a significant decrease in surface area and micropore volume after regeneration pointing to a collapse of the structure via hydrolysis (Shankar et al., 2018). Interestingly, the percentage of micropores decreased significantly in all porous BN samples, indicating that the

micropores were initially most susceptible to hydrolysis. This hypothesis is supported by the fact that porous BN samples with higher micropore volumes (BN-MU1:5 > BN-MU0.25:5 > BN-U5), were more unstable, as indicated by a larger decrease in surface area and pore volume. For instance, highly microporous BN-MU1:5 exhibited a significant decrease in surface area from 1650 to 29 m2 g −1 , while the surface area of less microporous BN-U5 decreased from 1092 to 486 m2 g −1 . This is not surprising as the more porous samples are probably more defective and therefore more prone to hydrolysis. Despite the large decrease in surface area after regeneration, the RhB adsorption capacity only decreased by about 10 %. We hypothesized that the RhB concentration used was not high enough to saturate the adsorbent. Therefore, despite a reduced surface area after regeneration, the adsorption capacity was not significantly affected. The high temperature BN samples exhibited higher stability than the low temperature BN samples after testing and

regeneration. Indeed, the samples maintained a high porosity after testing and regeneration. This aligns with the better performance of the high temperature BN samples observed before (**Figure 2**).

Continuing on the characterization of the materials post testing and regeneration, all samples showed a higher relative intensity of the (002)/(100) peaks, possibly indicating a preferred orientation of the crystals along the (002) plane (**Supplementary Figure 3**). In addition, the peak position for the (002) peak (∼25◦ ) shifted to higher angles, indicative of a decrease of the d(002)-spacing and the full width at half maximum (FWHM) decreased, demonstrating an increase in crystallite size (larger number of 2D layers). These three factors together showed that the samples exhibited higher crystallinity after exposure to water, probably as amorphous BN was more readily decomposed in water, resulting in an apparent increase in crystallinity.

After demonstrating that the porosity of the high temperature BN samples was maintained upon adsorption testing and regeneration, we investigated any potential changes in the chemistry of the materials using XPS analyses. The relative atomic percentages derived from XPS confirmed that the oxygen content increased in all low temperature porous BN samples after adsorption testing and regeneration (**Figures 4A,B**). This further supports the hydrolysis of the low temperature BN samples, causing the oxidation of BN and formation of a larger quantity of boronoxynitride and B-OH groups (Lin et al., 2011). Samples with higher micropore volumes and surface areas (BN-MU1:5 > BN-MU0.25:5 > BN-U5) showed higher increases in oxygen content, indicative of more extensive hydrolysis, in line with the previously described structural changes. The relative nitrogen content decreased in all samples after exposure to water, confirming that a decomposition reaction is occurring, in which water reacts with BN to release ammonia: 3H2O + 2BN → B2O3 + 2NH3 (Lin et al., 2011). On the other hand, the high temperature porous BN samples did not show any major variations in oxygen content from XPS analysis for two of the samples (**Figures 4C,D**) (BN-U5-1500C and BN-MU0.25:5-1500C).

#### CONCLUSIONS

We presented here a study on the adsorption behavior of porous boron nitride. For this, we synthesized porous BN samples with varied porosity, crystallinity and chemistry in order to identify the effect of these materials features on the adsorption of different organics. The main conclusions of this work are summarized below:

### REFERENCES

#### DATA AVAILABILITY

All datasets generated for this study are included in the manuscript and the **Supplementary Files**.

#### AUTHOR CONTRIBUTIONS

SM and CP designed the experiments. SM and XW carried out the experiments. SM, XW, and CP analyzed the data. SM and CP wrote the manuscript. All authors have read and agreed the submission of the manuscript.

#### ACKNOWLEDGMENTS

The authors would like to thank Dr. Les Bolton and Prof. Glenn Sunley from BP for their technical input. The authors would also like to acknowledge the funding and technical support from BP through the BP International Center for Advanced Materials (BP-ICAM), which made this research possible, as well as EPSRC for the funding through the CDT in Advanced Characterization of Materials (CDT-ACM) (EP/L015277/1). We also acknowledge support from the Department of Chemical Engineering at Imperial College London.

#### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2019.00160/full#supplementary-material

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Marchesini, Wang and Petit. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Structural Changes in Polymeric Gel Scaffolds Around the Overlap Concentration

#### Han Zhang, Matthew D. Wehrman and Kelly M. Schultz\*

*Department of Chemical and Biomolecular Engineering, Lehigh University, Bethlehem, PA, United States*

Cross-linked polymeric gels are an important class of materials with applications that broadly range from synthetic wound healing scaffolds to materials used in enhanced oil recovery. To effectively design these materials for each unique applications a deeper understanding of the structure and rheological properties as a function of polymeric interactions is required. Increasing the concentration of polymer in each scaffold increases physical interactions between the molecules that can be reflected in the material structure. To characterize the structure and material properties, we use multiple particle tracking microrheology (MPT) to measure scaffolds during gelation. In MPT, fluorescently labeled probe particles are embedded in the material and the Brownian motion of these particles is captured using video microscopy. Particle motion is related to rheological properties using the Generalized Stokes-Einstein Relation. In this work, we characterize gelation of a photopolymerized scaffold composed of a poly(ethylene glycol) (PEG)-acrylate backbone and a PEG-dithiol cross-linker. Scaffolds with backbone concentrations below and above the overlap concentration, concentration where polymer pervaded volume begins to overlap, are characterized. Using time-cure superposition (TCS) we determine the critical relaxation exponent, *n*, of each scaffold. The critical relaxation exponent is a quantitative measure of the scaffold structure and is similar to a complex modulus, *G*∗ , which is a measure of energy storage and dissipation. Our results show that below the overlap concentration the scaffold is a tightly cross-linked network, *navg* = 0.40 ± 0.03, which stores energy but can also dissipate energy. As polymeric interactions increase, we measure a step change in the critical relaxation exponent above the overlap concentration to *navg* = 0.20 ± 0.03. After the overlap concentration the scaffold has transitioned to a more tightly cross-linked network that primarily stores energy. Additionally, continuing to increase concentration results in no change in the scaffold structure. Therefore, we determined that the properties of this scaffold can be tuned above and below the overlap concentration by changing the polymer concentration but the structure will remain the same in each concentration regime. This is advantageous for a wide range of applications that require scaffolds with varying stiffness and the same scaffold architecture.

Keywords: multiple particle tracking microrheology, hydrogel scaffolds, poly(ethylene glycol), photopolymerization, time-cure superposition

#### Edited by:

*Gil Garnier, Bioresource Processing Institute of Australia (BioPRIA), Australia*

#### Reviewed by:

*Erica Wanless, University of Newcastle, Australia Eric Furst, University of Delaware, United States*

> \*Correspondence: *Kelly M. Schultz<EMAIL_ADDRESS>
#### Specialty section:

*This article was submitted to Chemical Engineering, a section of the journal Frontiers in Chemistry*

Received: *05 December 2018* Accepted: *23 April 2019* Published: *08 May 2019*

#### Citation:

*Zhang H, Wehrman MD and Schultz KM (2019) Structural Changes in Polymeric Gel Scaffolds Around the Overlap Concentration. Front.  This complexity grows out of the vast array of applications and high demand on gel versatility. Gels are part of everyday life from commonly used personal, fabric and home care products to exotic biomaterials designed to mimic the extracellular matrix (ECM) (Winter and Chambon, 1986; Scanlan and Winter, 1991; West and Hubbell, 1999; Raeber et al., 2005, 2007; Yamaguchi et al., 2007; Benton et al., 2009; Fairbanks et al., 2009b; Zustiak and Leach, 2010; Wylie et al., 2011; Wehrman et al., 2018). Cross-linked gels have also played a significant role in enhanced oil recovery (Moradi-Araghi et al., 1988; Zolfaghari et al., 2006; He et al., 2009; Jung et al., 2013; Tongwa et al., 2013). These materials are used to decrease permeability in high permeability zones near naturally fractured carbonates that require water shutoff but cannot be permanently plugged. Rheological measurement is a critical means for characterizing and validating gelation strategies and gaining insight into their structure and properties. Quantitatively identifying dynamic scaffold structure and properties and the relation to material function is crucial in advancing the design of these materials. These history-dependent systems are characterized during gelation to establish a quantitative framework to understand how polymeric interactions, i.e. overlap and entanglement, within macromer solutions change the gelation reaction and influence final material properties in a chain-growth system. We characterize the scaffold structure and rheological properties of a well-defined photopolymerized hydrogel scaffold with increasing polymeric interactions due to an increase in the concentration of the polymer backbone. This work can be leveraged to design gels with highly-engineered microstructures and properties that can be tailored throughout the phase transition.

In this work, polymeric interactions are varied to determine the change in the gelation reaction and final material properties. We are using a neutral polymer, poly(ethylene glycol) (PEG). Due to this, polymeric interactions are defined as only physical interactions between the macromolecules in solution. Polymeric solutions have three regimes in the concentrationviscosity curve: dilute, semi-dilute and entangled (Graessley, 1980; Doi and Edwards, 1988; Pavlov et al., 2003; Rubinstein and Colby, 2003; Wehrman et al., 2018). In the dilute regime, polymers do not interact. In the semi-dilute regime, the pervaded volume of the polymers begin to overlap. In the entangled regime, polymer molecules physically interact and entangle (Graessley, 1980; Doi and Edwards, 1988; Pavlov et al., 2003; Rubinstein and Colby, 2003; Wehrman et al., 2018). The overlap concentration, c ∗ , is defined in a good solvent as

$$c^\* = \frac{M}{\frac{4}{3}\pi R\_\mathcal{g}^3 N\_A} \tag{1}$$

where M is the molecular weight, NA is Avogadro's number and Rg is the radius of gyration defined for a star polymer as R star g = R arm g h 3f −2 f i 1 2 where R arm g is the radius of gyration of a single arm and f is the number of functional groups on the polymer. The entanglement concentration, c ∗∗, is defined as

$$c^{\*\*} = \frac{\nu \left(M\_s / N\_A\right)}{b^6} \tag{2}$$

where v is the excluded volume parameter, Ms is the monomer molecular weight and b is the Kuhn length (Graessley, 1980; Doi and Edwards, 1988; Larson, 1999; Pavlov et al., 2003; Rubinstein and Colby, 2003). The gel characterized in this work is a fourarm star PEG-acrylate cross-linked with a linear PEG-dithiol. This is a photopolymerization reaction that uses lithium phenyl-2,4,6-trimethylbenzoylphosphinate (LAP) as a photoinitiator. PEG is chosen due to the versatility of this molecule. PEG is a hydrophilic molecule that can be functionalized with many different chemistries (Bhat and Timasheff, 1992; Iza et al., 1998; Kienberger et al., 2000; Bryant and Anseth, 2002; Hansen et al., 2003; Rubinstein and Colby, 2003; Lee et al., 2007; Aimetti et al., 2009; Fairbanks et al., 2009a). Due to this versatility, PEG has been widely used as the basis for polymeric gel scaffolds that are used in applications from consumer care products to implantable biomaterials. With the development of each new chemistry, the change in the rheology and structure must be understood to design these materials for each application.

Bulk rheological measurements have been the classic way to characterize gelling materials and polymeric solutions (Muthukumar and Winter, 1986; Winter and Chambon, 1986; Chambon and Winter, 1987; Winter, 1987; Adolf and Martin, 1990; Scanlan and Winter, 1991; Izuka et al., 1992; Winter and Mours, 1997; Larson, 1999; Rubinstein and Colby, 2003; Larsen, 2008; Larsen and Furst, 2008; Larsen et al., 2008). In these measurements the viscous, G ′′, and elastic component, G ′ , of the complex modulus, G ∗ , is measured. When a material gels the viscous component decreases as the elastic component simultaneously increases. At the gel point, G ′ and G ′′ are parallel over all frequencies (Muthukumar and Winter, 1986; Winter and Chambon, 1986; Chambon and Winter, 1987; Winter, 1987; Adolf and Martin, 1990; Scanlan and Winter, 1991; Winter and Mours, 1997; Larson, 1999). To determine the point at which a material gels, which is defined as the first sample-spanning network cluster, time-cure superposition (TCS) is used. Timecure superposition is the superposition of viscoelastic functions at different extents of reaction (Muthukumar and Winter, 1986; Winter and Chambon, 1986; Chambon and Winter, 1987; Winter, 1987; Adolf and Martin, 1990; Scanlan and Winter, 1991; Winter and Mours, 1997; Larson, 1999). TCS is used to determine the critical gel time, tc , and critical relaxation exponent, n. The critical relaxation exponent indicates the structure of the scaffold and is a measure similar to a complex modulus, indicating how much energy the scaffold can store and dissipate. Although bulk rheology has been the classic way to collect data of gelling systems, microrheological characterization can also be used.

Microrheological characterization has been used extensively to characterize gelling materials (Freundlich and Seifriz, 1922; Heilbronn, 1922; Seifriz, 1924; Valentine et al., 1996, 2001; Crocker et al., 2000; Gardel et al., 2005; Panorchan et al., 2006; Slopek et al., 2006; Veerman et al., 2006; Caggioni et al., 2007; Wong Po Foo et al., 2009; Corrigan and Donald, 2009a; Mulyasasmita et al., 2011; Schultz and Furst, 2012; Schultz and Anseth, 2013; Furst and Squires, 2017). Our work focuses on passive microrheological characterization of gelling scaffolds. Multiple particle tracking microrheology (MPT), a passive microrheological technique, measures the Brownian motion of probe particles embedded in a material which is related to rheological properties using the Generalized Stokes-Einstein Relation (GSER) (Mason and Weitz, 1995; Crocker and Grier, 1996; Mason, 2000; Valentine et al., 2004; Savin and Doyle, 2005; Squires and Mason, 2010; Furst and Squires, 2017). Advantages of MPT that make it ideal for characterization of gelling scaffolds are: small sample size (1 − 50 µL), measurement of a large frequency range (up to MHz), short acquisition time enabling measurement of steady state properties of an evolving material and sensitivity in the low moduli range enabling characterization of the fragile microstructure of a gel at the sol-gel transition (Mason et al., 1997; Waigh, 2005, 2016; Squires and Mason, 2010; Schultz and Furst, 2012; Wehrman et al., 2016; Furst and Squires, 2017; Daviran et al., 2018). Additionally, TCS has been adapted to determine the critical gelation time and critical relaxation exponent using MPT measurements (Larsen and Furst, 2008; Larsen et al., 2008; Corrigan and Donald, 2009a,b; Schultz et al., 2009a). Due to these advantages, we use MPT to characterize the critical transitions of PEG-acrylate gels as polymeric interactions are increased.

Previous work determined the overlap concentration, c ∗ = 13 ± 4 wt%, of the PEG-acrylate backbone using bulk rheology (Wehrman et al., 2018). This previous investigation also characterized the change in material properties and structure when the PEG-acrylate backbone is below (3 wt%) and above (10 and 18 wt%) the overlap concentration. This work characterized a change in structure to a more tightly cross-linked scaffold as PEG-acrylate concentration is increased (Wehrman et al., 2018). The present work expands upon this work to determine if there is a gradual or step change in the structure of the material below and above c ∗ . Scaffolds are characterized changing the PEG-acrylate concentration below c ∗ from 4 − 9 wt% and above c ∗ at 13 and 15 wt%. Data from the previous study is also included for comparison. First, the kinetics of gelation are measured with MPT. The logarithmic slope of the mean-squared displacement, α, quantifies the change in material properties during scaffold gelation. For this chain-growth gel, there is an initial step when polymers are adding to growing chains and there is no measurable change in the rheology. Once chains begin to cross-link and form the gel network, rapid gelation is measured. Normalization of the UV exposure time with the final time of gelation, defined as the time when α ≈ 0.02, results in all curves collapsing onto a master curve indicating that polymeric interactions do not change the kinetics of gelation. TCS is used to analyze MPT data at PEG-acrylate backbone concentrations above and below c ∗ . We calculate a critical relaxation exponent, n, and critical gelation time, tc , for each scaffold. The critical relaxation exponent is constant above and below c ∗ with a step change at c ∗ . The normalized critical gelation time has no change as a function of PEG-acrylate concentration. Together, these results characterize a change in scaffold structure that is only dependent on the overlap concentration. This will enable the scaffold rheological properties to be tailored without changing the structure. This work provides additional information about scaffold properties and structure which can be used to tailor these materials for applications that include decreasing permeability during enhanced oil recovery, rheological modification of fabric and home care products and as synthetic implantable materials for wound healing.

### MATERIALS AND METHODS

#### Hydrogel Scaffold

The hydrogel scaffold is composed of a four-arm star PEGacrylate which is cross-linked with PEG-dithiol. PEG-acrylate has a molecular weight of Mn = 20, 000 g mol−1 (JenKem Technology) and is end-functionalized with acrylates (f = 4 where f is functionality). The cross-linker is a PEG end-functionalized with thiols (Mn = 1, 500 g mol−1 , f = 2, Aldrich). These two molecules undergo a chaingrowth polymerization upon exposure to ultraviolet (UV) light (PhotoFluor LM-75, output range 340 − 800 nm, 89 North, Inc.). The reaction is initiated by lithium phenyl-2,4,6 trimethylbenzoylphosphinate (LAP). LAP is synthesized using previously published protocols (Fairbanks et al., 2009a). Carboxylated polystyrene probe particles (2a = 0.97 ± 0.01 µm where a is the particle radius, Polysciences, Inc.) are added to the polymer precursor solution to enable multiple particle tracking microrheological measurements.

Hydrogel samples maintain an acrylate:thiol ratio of 1.4 : 1. The concentration of PEG-acrylate is changed by 1 wt% from 4− 9 wt%. Higher PEG-acrylate concentrated samples are measured and are 13 and 15 wt%. Data from 3, 10 and 18 wt% are included from a previous study for comparison (Wehrman et al., 2018). During sample preparation, all concentrated stock solutions are kept on ice for 30 min prior to addition to the polymer precursor solution. This ensures that there are no reactions in the solutions, especially thiol-thiol reactions in the PEG-dithiol stock solution. The polymer precursor solutions consist of PEG-acrylate, PEGdithiol, 1.5 mM LAP, 0.052% solids per volume probe particles and water. Solutions are well-mixed prior to injection into the sample chamber. Hydrogel samples are gelled by exposure to UV light and MPT data are collected.

### Multiple Particle Tracking Microrheology

Multiple particle tracking microrheology (MPT) measures the change in rheological properties during scaffold gelation. In MPT, the Brownian motion of 1 µm fluorescently labeled probe particles (carboxylated polystyrene particles, 2a = 0.97 ± 0.01 µm, Polysciences, Inc.) is measured. Data are collected using video microscopy on an inverted fluorescent microscope (Zeiss Observer Z1, Carl Zeiss AG). A high numerical aperture waterimmersion objective is used to capture data, which maximizes the pixels per probe particle (63× water-immersion objective, N.A. 1.3, 1× optovar, Carl Zeiss AG). Our equipment is calibrated to minimize static and dynamic particle tracking errors (Savin and Doyle, 2005). Data are collected at 30 frames per second and an exposure time of 1000 µs (Phantom Miro M120, 1024 × 1024 pixels, Vision Research Inc.).

After data acquisition, the brightness-weighted centroid of each probe particle is determined using classical tracking algorithms (Crocker and Grier, 1996; Crocker and Weeks, 2011; Furst and Squires, 2017). These algorithms determine the center of each particle in each frame. Those particle centers are linked together into trajectories using a probability distribution function that accounts for Brownian motion (Crocker and Grier, 1996; Crocker and Weeks, 2011; Furst and Squires, 2017). From the particle positions, the ensemble-averaged mean-squared displacement (MSD, h1r 2 (τ )i) is calculated from the twodimensional data using h1r 2 (τ )i = h1x 2 (τ )i + h1y 2 (τ )i. From the MSD, rheological properties can be calculated using the Generalized Stokes-Einstein Relation

$$
\langle \Delta r^2(\mathbf{r}) \rangle = \frac{k\_B T}{\pi a} J(\mathbf{r}) \tag{3}
$$

where kB is the Boltzmann constant, T is the temperature, a is the particle radius and J (τ ) is the creep compliance.

To measure gelation using MPT we first add probe particles into the polymer precursor solution. Briefly, 1 µm probe particles are washed 3× by dilution and centrifugation at 5, 000 RPM for 5 min (Centrifuge 5425, Eppendrof) to remove any excess dye. Probe particles are then sonicated (40 kHz, Emerson Industrial Automation) for 15 min to ensure that there are no particle aggregates. We add probes at a final concentration of 0.052% solids per volume to the polymer precursor solution. We then inject our polymer precursor solution into a sample chamber. Sample chambers are made on a standard glass slide (75 × 25 × 1 mm, Thermo Fisher Scientific), with 0.16 mm thick spacers and a coverglass (22 × 22 × 0.16 mm) as the top of the chamber. After the polymer precursor solution has filled the sample chamber it is sealed with an air-curing epoxy (Gorilla Glue Company). During curing, sample chambers are kept in the dark and allowed to cure for 15 min.

Each sample is exposed to UV light and MPT data are collected. This is repeated until ≈30 min after complete gelation is measured. Neutral density filters (Chroma Technology) are used to lower the intensity of the UV light. This is done to slow the gelation reaction enabling acquisition of MPT data during the sol-gel transition. The neutral density filters used are 32 and 10% transmission (Chroma Technologies) for 4−15 and 18 wt% PEGacrylate, respectively. Due to the arbitrary time of UV exposure, data are reported as a normalized UV exposure time which is UV exposure time divided by final time of gelation. The final time of gelation is the time when the logarithmic slope of the MSD, α, is ≤0.02. All MPT experiments are repeated at least 3× to ensure repeatability and critical values are reported as the average ± the standard deviation.

#### RESULTS AND DISCUSSION

This work characterizes the change in rheological properties and scaffold structure as polymeric interactions are increased. Previous work determined the overlap concentration, c ∗ , of this PEG-acrylate backbone using bulk rheology and MPT. The value of c ∗ = 13 ± 4 wt% (Wehrman et al., 2018). We consider scaffolds with PEG-acrylate concentrations c ≤ 9 wt% to be below c ∗ and c > 9 wt% to be above c ∗ . This uses the lower limit of c ∗ as the transition point. This value is chosen not only from the measurements of c ∗ but also from the characterization of scaffold properties presented here. We measure the same trend in gelation kinetics as a function of UV exposure, regardless of PEG-acrylate concentration. Previous work found that below the overlap concentration the material had a more open structure than samples at and above the overlap concentration. This work expands this initial study determining that there is a step change in the scaffold structure when the overlap concentration is reached. In the previous study, we hypothesized that there would be a gradual decrease in the value of n below the overlap concentration and a constant value above c ∗ . The gradual decrease in n below c ∗ would have indicated a more densely cross-linked structure as the concentration is increased. Instead, we measure no change in the dilute, c < c ∗ , and semi-dilute, c > c ∗ , concentration regimes, but there is a step change in the structure of the scaffold at the transition. This result indicates that the structure remains constant above and below the overlap concentration regardless of PEG-acrylate concentration. This information can lead to the design of new materials, where the moduli of the material can be tailored for applications, such as the mimic of a tissue or a material to reduce permeability in enhanced oil recovery without changing the scaffold structure.

### Scaffold Gelation

Multiple particle tracking microrheology measures the scaffold properties during gelation. After UV exposure, MPT data are collected. From the MPT data, the ensemble-averaged meansquared displacement is calculated for all the measured probe particles. Previous work that used MPT to measure this hydrogel scaffold determined that the scaffold evolves homogeneous during gelation with each particle probing the same material properties (Wehrman et al., 2018). The logarithmic slope of the MSD, α = d logh1r 2 (τ )i d log τ , is a measure of the state of the material (Stauffer et al., 1982; Winter and Chambon, 1986; Adolf and Martin, 1990; Larsen and Furst, 2008; Corrigan and Donald, 2009a,b; Schultz et al., 2009a, 2012; Schultz and Furst, 2012; Wehrman et al., 2016; Daviran et al., 2018). When α = 1 probe particles are freely diffusing and the material is a liquid. When α → 0 probe particles are completely arrested in the gel scaffold. This occurs at a low moduli value, G ′ ≈ 4 Pa (Waigh, 2005; Schultz and Furst, 2011, 2012; Furst and Squires, 2017). Due to the low maximum moduli value, bulk rheology should be used to supplement MPT to measure the equilibrated material properties. Finally, when 0 < α < 1 the material is a viscoelastic sol or gel and probe particle movement is restricted.

To quantitatively determine when the sol-gel transition occurs, time-cure superposition (TCS) is used to analyze the critical transition and will be discussed in detail below. From TCS, the critical relaxation exponent is determined, n. α = n is the sol-gel transition (Larsen and Furst, 2008; Corrigan and Donald, 2009a,b; Schultz et al., 2009a,b; Schultz and Anseth, 2013; Adibnia and Hill, 2016; Wehrman et al., 2016; Escobar et al., 2017; Daviran et al., 2018). This is when the first sample-spanning network cluster has formed and is the definition of a gel (Stauffer et al., 1982; Muthukumar and Winter, 1986; Winter and Chambon, 1986; Chambon and Winter, 1987; Winter, 1987; Adolf and Martin, 1990; Larsen and Furst, 2008; Corrigan and Donald, 2009a,b; Schultz et al., 2009a,b; Schultz and Anseth, 2013; Adibnia and Hill, 2016; Wehrman et al., 2016; Escobar et al., 2017; Daviran et al., 2018). To determine the state of the material α is compared to n. When α < n the material is a gel and when α > n the material is a sol (Larsen and Furst, 2008; Corrigan and Donald, 2009a,b; Schultz et al., 2009a,b; Schultz and Anseth, 2013; Adibnia and Hill, 2016; Wehrman et al., 2016; Escobar et al., 2017) ; (Daviran et al., 2018).

**Figures 1A,B** show the change in α as the polymer precursor solution is exposed to UV light. Upon initial exposure to UV light there is lag prior to scaffold gelation. This lag is due to radicals being formed in the polymer precursor solution and the growth of polymeric chains that are not cross-linking into a network structure. During chain-growth, cross-linking occurs by the addition of polymers to growing network chains. Those chains then cross-link together to form the scaffold network structure (Rubinstein and Colby, 2003; Tibbitt et al., 2013; Payamyar et al., 2016). This cross-linking of chains leads to the steep slope in α vs. normalized UV exposure. Prior to the decrease in α, probe particles are freely diffusing. During this time polymer chains are forming but are not large enough to restrict probe particle motion. When these chains start to cross-link, the network structure is forming and probe particle movement becomes restricted and, eventually, arrested in the gel scaffold. At the point of probe particle arrest, the network structure will continue to grow, but MPT can no longer measure the change in material properties.

For all PEG-acrylate samples the kinetics of gelation follow the same trend that is described above. **Figure 1A** shows PEGacrylate samples with concentrations below c ∗ and **Figure 1B** are concentrations above c ∗ . This is further illustrated in **Figure S1**, where all concentrations are plotted together. In **Figures 1A,B**, UV exposure time is normalized by the UV exposure when complete gelation is measured, defined as α ≤ 0.02. This value is used because it is a slope where no probe particle movement is measured. Upon normalization, all data sets collapse and have a lag in gelation and then a steep slope, where the value of α decreases rapidly with time. Therefore, from these experiments, we can conclude that there is no change in the mechanism of

gelation when polymeric interactions are added into the system by concentrating the PEG-acrylate backbone over c ∗ .

#### Time-Cure Superposition

Time-cure superposition (TCS) is used to analyze each scaffold gelation reaction and determine the critical relaxation exponent, n, and the critical gelation time, tc . TCS is the superposition of viscoelastic functions at different extents of reaction (Muthukumar and Winter, 1986; Winter and Chambon, 1986; Chambon and Winter, 1987; Winter, 1987; Adolf and Martin, 1990; Larsen and Furst, 2008; Corrigan and Donald, 2009a,b; Schultz et al., 2009a,b; Schultz and Anseth, 2013; Adibnia and Hill, 2016; Wehrman et al., 2016; Escobar et al., 2017; Daviran et al., 2018). This analysis exploits the self-similarity of measurements of the scaffold material properties prior to and after gelation to shift data onto master curves. In MPT, the shortest lag times, 0.03 ≤ τ ≤ 1s, measure the longest relaxation times, τR, of the material. This is because the short lag times are equivalent to high frequency measurements in bulk rheology. The longest relaxation times in the pre- and post-gel give distinct curvature to MSDs measured with MPT. In the pre-gel, the longest relaxation time of the polymers are measured. In the postgel, the longest relaxation time of the gel network is being probed (Muthukumar and Winter, 1986; Winter and Chambon, 1986; Chambon and Winter, 1987; Winter, 1987; Adolf and Martin, 1990; Larsen and Furst, 2008; Corrigan and Donald, 2009a,b; Schultz et al., 2009a,b; Schultz and Anseth, 2013; Adibnia and Hill, 2016; Wehrman et al., 2016; Escobar et al., 2017; Daviran et al., 2018). Since these measurements are probing the same relaxation time, over a large time scale, the data can be shifted into pre- and post-gel master curves by shifting along the lag time and MSD axes.

The shift factors determine the critical values, namely the critical relaxation exponent and critical gelation time (Muthukumar and Winter, 1986; Winter and Chambon, 1986; Chambon and Winter, 1987; Winter, 1987; Adolf and Martin, 1990; Larsen and Furst, 2008; Corrigan and Donald, 2009a,b; Schultz et al., 2009a,b; Schultz and Anseth, 2013; Adibnia and Hill, 2016; Wehrman et al., 2016; Escobar et al., 2017; Daviran et al., 2018). The lag time shift factor, a, is related to the inverse of the longest relaxation time, τR, and the distance away from the critical gelation time, |t−tc| tc , by a scaling factor y by

$$a \sim \tau\_R^{-1} \sim \left(\frac{|t - t\_c|}{t\_c}\right)^{\mathcal{V}}.\tag{4}$$

Similarly the MSD shift factor, b, is related to the inverse of the steady state creep compliance, Je, and the distance away from the critical gelation time by a scaling factor z by

$$b \sim J\_c^{-1} \sim \left(\frac{|t - t\_c|}{t\_c}\right)^z. \tag{5}$$

The critical relaxation exponent is the ratio of the two scaling exponent

$$m = \frac{z}{\nu}.\tag{6}$$

The critical relaxation exponent determines the state of the material as described in detail above. This is the value of α where the first sample-spanning network cluster forms during gelation. Additionally, n can be thought of as a complex modulus, G ∗ , with both a viscous and elastic component. When n > 0.5 the material is an open porous network which dissipates more energy than it stores. When n < 0.5 the scaffold is a tightly cross-linked network that stores more energy. When n = 0.5 the scaffold is a percolated network which stores and dissipates equal amounts of energy (Stauffer et al., 1982; Schultz et al., 2009a; Wehrman et al., 2018).

All data taken of PEG-acrylate gelation is analyzed with TCS to determine the value of the critical relaxation exponent, n, and the critical gelation time, tc . TCS for a 4 wt% PEG-acrylate gelation is shown in **Figure 2**. An example of TCS for all other PEG-acrylate concentrations are in the Supplementary Material, **Figures S2–S8**. The ensemble-averaged MSD is calculated from MPT data, **Figure 2A**. The MSDs decrease in both magnitude and α as UV exposure time is increased. Initially, probe particles are freely diffusing in the polymer solution. As UV exposure increases the scaffold network grows and probe particle movement decreases. As mentioned above, at the sol-gel transition there is a change in the shape of MSD curves due to the change in the relaxation times measured from the relaxation of polymers to the relaxation of a network. This change in shape enables the MSD curves to be shifted into two master curves, a pre- and post-gel master curve, **Figure 2B**. The start of the pregel master curve is the polymeric solution with α = 1. As the value of a · τ decreases the value of α also decreases which is due to the growing polymer chains restricting the movement of the probe particles. Additionally, the curvature of the pre-gel curve shows that at short lag times the MSD curves up slightly, which is the measure of the polymeric longest relaxation time, τR (Furst and Squires, 2017; Wehrman et al., 2018). At α = n the preand post-gel master curves meet. In the pre-gel curve this is the last measurement before the sample-spanning network cluster forms and in the post-gel curve this is the first measurement after the sample-spanning network cluster has formed. In the post-gel master curve as a·τ increases, there is a decrease in α until α → 0 and probe particles are completely arrested in the polymerized network scaffold. In the post-gel master curve, the curvature is reversed and at short lag times, the MSDs curve down, indicative of measurements of the longest relaxation time of a network.

The shift factors are used to determine the critical relaxation exponent and critical gelation time. In **Figure 2C** both the lag time and MSD shift factors diverge at the sol-gel transition. This is the divergence of the viscosity upon network formation and the emergence of the elastic modulus, G ′ , as the network continues to grow. This divergence happens at the critical gelation time, which is tc = 11.6 min for this 4 wt% PEG-acrylate network. The critical relaxation exponent is the ratio of scaling exponents. Scaling exponents y and z are calculated by fitting the log a and log b vs. the log |t−tc| tc , **Figure 2D**. The ratio of the scaling exponents determine the critical relaxation exponent, which is n = 0.26 ± 0.08 for this hydrogel scaffold.

This analysis is done for all MSD data collected both above and below c ∗ . Since the absolute value of UV exposure can be dependent on the size of the sample, we normalize the critical gelation time by the final time of gelation tc,norm = tc tfinal . The critical gelation time, tc , is defined as the time when the first sample-spanning network cluster is formed in the material. The final time of gelation, tfinal, is again defined as the time when α ≤ 0.02. The normalized critical gelation time is tc,norm = 0.86±0.17 and tc,norm = 0.96 ± 0.06 for PEG-acrylate concentrations below and above c ∗ , respectively. A plot of these values is provided in the Supplementary Material, **Figure S9**. These values are within error of each other, indicating that the critical gelation time

increased. (B) MSD curves shifted into pre- and post-gel master curves using shifting factors *a* and *b*, which (C) diverge at the critical gelation time, *tc*, at the sol-gel transition. (D) The scaling exponents *y* and *z* are determined from the slope of log *a* and log *b* vs. the logarithm of the distance away from the critical gelation time, log |*t*−*tc*| *tc* . The critical relaxation exponent, *n* is calculated from the scaling factors.

occurs at the same point within the scaffold gelation reaction regardless of the backbone concentration. The first samplespanning cluster will require a similar amount of cross-links regardless of polymeric interactions and, therefore, occurs at a similar point in the gelation process. This is also supported by the collapse of MPT measurements throughout gelation in **Figures 1A,B**. These results show that scaffold gelation proceeds through the same reaction mechanism regardless of polymeric interactions in the precursor solution.

Finally, the critical relaxation exponent is calculated for scaffolds above and below the overlap concentration of PEGacrylate. **Figure 3** is a plot of the average value of n vs. PEGacrylate concentration. As described previously, scaffold with c ≤ 9 wt% are considered below c ∗ and c > 9 wt% are considered above c ∗ . The critical relaxation exponent indicates the scaffold structure and is a measure similar to a complex modulus, indicating how much energy the scaffold can store and dissipate. The value of n has a step change at the overlap concentration. Below c ∗ navg = 0.40±0.03 and above c ∗ navg = 0.20±0.03. This step change in the value of n indicates that there is a change in structure below and above the overlap concentration, but within the dilute and semi-dilute concentration regimes there is no structural change. For all concentrations, a tightly cross-linked network that stores more energy than it dissipates is measured.

When c > c ∗ the network is more tightly cross-linked with a smaller pore structure than when c < c ∗ . This is due to polymeric interactions. When c > c ∗ , polymers are interacting and are able to enter into the pervaded volume of other chains. Due to the polymers being physically closer during network formation there is a high likelihood of arm interpenetration during network formation. This leads to a scaffold with a smaller porous structure that can store energy. This result indicates that the scaffold structure is unchanged by the change in backbone concentration. The only factor that changes scaffold structure is whether the backbone concentration is above or below c ∗ . Therefore, this scaffold can be tailored to change the elastic moduli of the material without changing the structure of the scaffold enabling the material to be tailored for desired applications.

### CONCLUSIONS

This work characterizes the change in rheological properties and scaffold structure during the photopolymerization of a four-arm star PEG-acrylate:PEG-dithiol hydrogel scaffold as polymeric interactions are added to the system. MPT is used to measure the change in scaffold properties as UV exposure time is increased. The gelation mechanism is the same for all scaffolds, regardless of polymeric interactions. The rheology of all scaffolds indicate that they follow a typical chain-growth polymerization. Upon UV exposure, chains of polymers begin to grow in the system and gelation is measured when these chains begin to crosslink into a network. Measured mean-squared displacements are further analyzed using time-cure superposition. The critical relaxation exponent and critical gelation time are determined. The normalized critical gelation time is independent of polymeric interactions. This result supports measurements that all scaffolds follow the same gelation mechanism. The critical relaxation exponent is sensitive to polymeric interactions. A step change in the value of the critical relaxation exponent is measured when polymeric interactions increase above the overlap concentration. This decrease in the value of n indicates that scaffolds made with backbone concentrations above c ∗ have a smaller porous network which stores more energy than scaffold made with PEG-acrylate concentrations below c ∗ . This is due to the increased physical interactions, which likely lead to an interpenetrated network structure when c > c ∗ .

The results of this work provide a predictability of material properties, structure and kinetics during gelation.

#### This information can be used to determine the feasibility of these scaffolds for desired applications, particularly when these materials are made in a new environment and require a specific modulus and structure. A stable structure below and above c ∗ enables properties of the material to be tailored while the structure remains constant. The wide applicability of hydrogel scaffolds has lead to the need to more precisely design these materials. This work informs this design by building a knowledge base which can be exploited to minimize trial-and-error when a scaffold is used in a new application.

### AUTHOR CONTRIBUTIONS

HZ was responsible for design of experiments and collection and analysis of data. MW was responsible for design of experiments. KS was responsible for design of experiments, data analysis and writing the manuscript.

### FUNDING

Funding for this work was provided by the American Chemical Society Petroleum Research Fund (54462-DNI7). Acknowledgment is made to the Donors of the American Chemical Society Petroleum Research Fund for support of this research.

### ACKNOWLEDGMENTS

The authors would like to acknowledge Stephanie Cook, Melissa J. Milstrey, and Danielle Rafanelli for their help with experiments. The authors would also like to acknowledge Nan Wu for carefully reading the manuscript.

### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2019.00317/full#supplementary-material

Additional graphs for α vs. normalized UV exposure for all PEG-acrylate concentrations, TCS of each PEG-acrylate concentration and tc,norm for all PEG-acrylate concentrations are provided. Additionally, a table with the values of n and tc,norm are provided for each PEG-acrylate concentration and can be found in the Supplemental data.

### REFERENCES

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Zhang, Wehrman and Schultz. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Preparation of Renewable Epoxy-Amine Resins With Tunable Thermo-Mechanical Properties, Wettability and Degradation Abilities From Lignocellulose- and Plant Oils-Derived Components

#### Louis Hollande1,2, Izia Do Marcolino1 , Patrick Balaguer 3 , Sandra Domenek 2 , Richard A. Gross 4 and Florent Allais 1 \*

1 URD ABI, CEBB, AgroParisTech, Pomacle, France, 2 UMR GENIAL, AgroParisTech, INRA, Université Paris-Saclay, Massy, France, 3 Institut de Recherche en Cancérologie de Montpellier, Val d'Aurelle, Montpellier, France, 4 Department of Chemistry and Chemical Biology, Rensselaer Polytechnic Institute, Troy, NY, United States

#### Edited by:

Gil Garnier, Bioresource Processing Institute of Australia (BioPRIA), Australia

#### Reviewed by:

Guillermo Javier Copello, Consejo Nacional de Investigaciones Científicas y Técnicas (CONICET), Argentina Juan Carlos Serrano-Ruiz, Universidad Loyola Andalucía, Spain

\*Correspondence:

Florent Allais<EMAIL_ADDRESS>
#### Specialty section:

This article was submitted to Chemical Engineering, a section of the journal Frontiers in Chemistry

Received: 28 November 2018 Accepted: 04 March 2019 Published: 27 March 2019

#### Citation:

Hollande L, Do Marcolino I, Balaguer P, Domenek S, Gross RA and Allais F (2019) Preparation of Renewable Epoxy-Amine Resins With Tunable Thermo-Mechanical Properties, Wettability and Degradation Abilities From Lignocellulose- and Plant Oils-Derived Components. Front. Chem. 7:159. doi: 10.3389/fchem.2019.00159 One-hundred percent renewable triphenol—GTF—(glycerol trihydroferulate) and novel bisphenols—GDFx–(glycerol dihydroferulate) were prepared from lignocellulose-derived ferulic acid and vegetal oil components (fatty acids and glycerol) using highly selective lipase-catalyzed transesterifications. Estrogenic activity tests revealed no endocrine disruption for GDFx bisphenols. Triethyl-benzyl-ammonium chloride (TEBAC) mediated glycidylation of all bis/triphenols, afforded innocuous bio-based epoxy precursors GDFxEPO and GTF-EPO. GDFxEPO were then cured with conventional and renewable diamines, and some of them in presence of GTF-EPO. Thermo-mechanical analyses (TGA, DSC, and DMA) and degradation studies in acidic aqueous solutions of the resulting epoxy-amine resins showed excellent thermal stabilities (Td5% = 282–310◦C), glass transition temperatures (Tg) ranging from 3 to 62◦C, tunable tan α, and tunable degradability, respectively. It has been shown that the thermo-mechanical properties, wettability, and degradability of these epoxy-amine resins, can be finely tailored by judiciously selecting the diamine nature, the GTF-EPO content, and the fatty acid chain length.

Keywords: ferulic acid, glycerol, lipase, bio-based thermosets, wettability, degradability

### INTRODUCTION

Thermoset polymers are widely used in industrial applications thanks to their versatile performance, good durability, and excellent chemical resistance provided by their highly cross-linked structure (Ellis, 1993; Auvergne et al., 2014; Ramon et al., 2018). It is therefore very common to find these polymers in a broad range of applications such as maintenance coating, adhesives for aerospace (Prolongo et al., 2009) and automobile (Holbery and Houston, 2006) industries, or binders in composites (Gojny et al., 2006). Epoxy resins are one of the most important thermoset materials and are usually synthesized by reacting (poly)phenolic compounds with epichlorohydrin under basic conditions (Bruins, 1968).

The thermoset polymers sector, and more generally the polymer industry, has adopted the current trend which consists in developing greener and more sustainable chemicals and technologies. For instance, Solvay or Dow Chemical have already initiated the transition with the commercialization of bio-based epichlorohydrin prepared from the chlorination of bio-based glycerol, a by-product of biodiesel production (Strebelle et al., 2001; Freddy, 2006). Furthermore, a number of renewable resources—including fatty acids (Biermann et al., 2000; Maisonneuve et al., 2013; Laurichesse et al., 2014), lignocellulosic biomass (Isikgor and Becer, 2015), and enzymatic products—have been reported as alternative feedstocks in polymers synthesis.

Bisphenol A (BPA) is commonly used to prepare commercial high-performance epoxy resins (aka DGEBA, DiGlycidylEther of Bisphenol A). Although DGEBA, upon curing with diamines, leads to cross-linked materials with exceptional properties such as strong adhesion, mechanical integrity and chemical resistance, BPA is recognized as an endocrine disruptive chemical and many studies have been conducted to replace it by innocuous bio-based chemicals. Lignans and lignin-derived chemicals that can be obtained through pyrolysis (Celikbag et al., 2017; Barde et al., 2018) or controlled depolymerization (Pandey and Kim, 2011; Shuai et al., 2016) are the most used biomass-derived chemical platforms for the design of bio-based BPA-substitutes, for the preparation of sustainable thermosets with high thermomechanical properties. Indeed, the aromatic moieties present in these bio-based chemicals confer rigidity to the resulting polymers (Wang et al., 2017; Feghali et al., 2018). For example, there have been many reports on epoxy resins from eugenol (Wan et al., 2016), vanillin (Fache et al., 2015a,b; Hernandez et al., 2016; Mauck et al., 2017; Nicastro et al., 2018; Savonnet et al., 2018; Zhao et al., 2018), ferulic and sinapic acids (Maiorana et al., 2016; Janvier et al., 2017; Ménard et al., 2017), guaiacols (Maiorana et al., 2015), and creosol (Meylemans et al., 2011) with promising performances capable of competing with current DGEBA-based materials. However, few researchers have considered the potential toxicity of these BPA substitutes, especially endocrine disruption, which remains a crucial health issue for both plastic sector workers and consumers (Jiang et al., 2018). The degradability of the resulting materials and their behavior regarding water are two other aspects that are also under-investigated. In fact, in most of the studies described above, the thermo-mechanical (e.g., Tg ) and physico-chemical (e.g., wettability) properties of the materials cannot be finely tuned because of the lack of apolar moieties in the biobased aromatic-epoxies.

Plant oil-derived fatty acids are highly valuable building blocks and are frequently used as renewable resources for the synthesis of thermoplastics/thermosets (Meier, 2018). Their long aliphatic chain and their double bond(s) allow a wide number of chemical modifications and provide materials with tunable properties, such as glass transition temperature (Tg ), chemical resistance, (bio)degradability, and adhesion strength. Combining the rigidity of bio-based aromatic building blocks and the apolar moiety of the alkyl chain of a fatty acid in a single BPA-substitute thus appears as an interesting approach to finely tune the Tg , the wettability and the degradability of the resulting epoxy-amine resins, after curing with diamines.

In this present work, we aim at developing a series of biobased epoxy resins, from naturally occurring ferulic acid and plant oils-derived components (i.e., glycerol and fatty acids), to access epoxy-amine thermoset materials with tunable thermomechanical properties, wettability and chemical degradation ability. Through highly selective chemo-enzymatic pathways, several polyphenolic compounds, named GDFx and GTF, have been prepared with tunable flexibility and degree of functionality (i.e., 2 or 3). The determination of the toxicity of these compounds has been carried out by determining their binding to human hormone receptors (ERα, PXR, and AR) and was benchmarked against that of commercial and controversial BPA. GDFx and GTF were then converted to their corresponding glycidyl ether by TEBAC-mediated glycidylation in presence of epichlorohydrin. Characterization studies were then performed on the corresponding epoxy resins formulations (i.e., molecular weight distribution and EEW). The latter were finally cured with three crosslinkers (i.e., IPDA, DA10, and DIFFA) and the thermo-mechanical properties as well as wettability and chemical degradation abilities of the resulting epoxy-amine resins were investigated.

### EXPERIMENTAL

### Material

Ferulic acid, lauric acid, palmitic acid, stearic acid, benzyl bromide, N,N-Dimethylpyridin-4-amine (DMAP), N,N-Diisopropylcarbodiimide (DIC), benzyltriethyl ammonium chloride (TEBAC), and palladium supported on carbon, were supplied by Sigma-Aldrich. Glycerol was purchased from Alfa Aesar. Epichlorohydrin was purchased from Acros Organics. Isophorone diamine (IPDA) and Decane diamine (DA10) were purchased from TCI. Candida antarctica Lipase B (CAL-B) immobilized on resin (LC200291, 10,000 propyl laurate units.g−1 ) was obtained from Novozyme. All reactants were used as received. All solvents were bought either from ThermoFisher Scientific or VWR. Deuterated chloroform (CDCl3) was purchased from Euriso-top.

### Methods

#### Purification

Column chromatographies were carried out with an automated flash chromatography (PuriFlash 4100, Interchim) and prepacked INTERCHIM PF-30SI-HP (30µm silica gel) columns using a gradient of cyclohexane and ethyl acetate for the elution.

#### Characterization

FT-IR analyses were performed on Cary 630 FT-IR with ATR. NMR analyses were recorded on a Bruker Fourier 300. 1H NMR spectra of samples were recorded in CDCl3 at 300 MHz, chemicals shifts were reported in parts per million (CDCl3 residual signal at δ = 7.26 ppm). 13C NMR spectra of samples were recorded at 75 MHz (CDCl3 signal at δ = 77.16 ppm). HRMS were recorded by the PLANET platform at URCA on a Micromass GC-TOF.

#### Gel Permeation Chromatography (GPC)

Gel permeation chromatography (GPC) was performed at 40◦C on an Infinity 1260 system from Agilent Technologies with PLgel 5µM MIXED-D column in THF (flow rate 1 mL.min−1 ) using conventional PS calibration and UV detection at 280 nm.

#### Thermo-Gravimetric Analyses (TGA)

Thermo-gravimetric analyses (TGA) were recorded on a Q500, from TA. About 10 mg of each sample was heated at 10◦C.min−1 from 30 to 500◦C under nitrogen or oxygen flow (60 mL.min−1 ).

#### Differential Scanning Calorimetry (DSC)

Differential scanning calorimetry (DSC) thermograms were obtained using a DSC TA Q20, under inert atmosphere (N2), with a calibration using indium, n-octadecane and n-octane standards. For each sample, about 10 mg was weighed in a pan which was then sealed and submitted to three heat/cool/heat cycles: heating from 30 to 200◦C at 10◦C.min−1 , cooling from 200◦C to −50◦C at 20◦C.min−1 . Glass transition temperatures (Tg ) were determined at the inflection value in the heat capacity jump.

#### Dynamical Mechanical Analyses (DMA)

Dynamical Mechanical Analyses (DMA) were performed on a DMA Q800 from the TA Instrument. The DMA samples had a rectangular geometry (length: 40 mm, width: 8 mm, thickness: 1.5 mm). Uniaxial stretching of samples was performed while heating at a rate of 3◦C.min−1 from 30 to 200◦C, keeping frequency at 1 Hz. Deformation was kept at 0.001% (amplitude of 7µm) to stay in the linear viscoelastic region. The storage modulus (E') and tan δ curves as a function of the temperature, were recorded and analyzed using the TA Instruments Universal Analysis 2000 software. The temperatures Tα ( ◦C) were determined as the temperatures at the peak maximum of the tan δ curves.

#### Synthesis of Resin Epoxy Precursors

**Scheme 1** presents the complete synthetic pathway using the following molecules: ferulic acid (**a**), ethyl dihydroferulate (**b**), benzylated ethyl ferulate (**c**), benzylated glycerol diferulate (GDFoBn) (**d**), glycerol tri-dihydroferulate (GTF) (**e**) and lipophilic glycerol dihydroferulates (GDFx) (**f**, **g** and **h**). Compound syntheses were carried out following procedures previously described in the literature (Ménard et al., 2017; Hollande and Domenek, 2018). Full characterizations and detailed procedures for new compounds are given in the **Electronic Supplementary Information** (ESI). 1H & 13C NMR spectra, FTIR spectrum as well as TGA trace of GTF-EPO can be found in a previous published work (Ménard et al., 2017).

#### Glycidylation

Phenolic precursors (GDFx or GTF; 1 eq) were dissolved in epichlorohydrin (10 eq/OH). Triethyl benzyl chloride (TEBAC) (0.2 eq) was added and the suspension was stirred for 2 h at 80◦C. The reaction medium was cooled down to room temperature and NaOH (5M, 2 eq/OH) was added. The biphasic solution was vigorously stirred for 4 h at room temperature then extracted with ethyl acetate (3 × 100 mL). The organic layers were washed with brine (80 mL), dried over anhydrous MgSO4, filtered and concentrated in vacuo.

#### Epoxide Equivalent Weight (EEW)

Experimental values of EEW for the epoxide resins synthesized herein were determined by ASTM D1652. Samples were dissolved in dichloromethane (DCM) with tetraethylammonium bromide and titrated in triplicate, using perchloric acid solution in acetic acid (1 N). Crystal violet was used as an indicator and the end point was determined when the solution turned from blue to green for longer than 30 s. Detailed results are provided in the ESI.

#### Preparation of Cured Epoxy Materials

This method considers that the epoxy resins prepared herein contained small amounts of higher molecular weights fractions. The weight of the resins and cross-linkers were determined by using the EEW values and amine hydrogen equivalent weight (AHEW). Weights of reaction components for the formulated resin systems were determined using equation (1). Equations (2) and (3) were used to calculate the AHEW and to obtain the parts by weight of diamine per hundred parts resin (phr).

$$EEW\ of\ mix = \begin{array}{c} \text{Total Weight} \\ \begin{array}{c} \frac{Weight\ of\ resist\ A}{EEW\ of\ residue\ A} \\ \end{array} + \begin{array}{c} \text{Weight of\ residue\ B} \\ \end{array} \end{array} \tag{1}$$

$$AHEW = \begin{array}{c} \text{molecular weight of amino} \\ \hline \text{number of active hydrogen} \end{array} \tag{2}$$

$$phr = \frac{AHEW \times 100}{EEW} \tag{3}$$

For instance, if phr = 20, then 20 g of diamine would be needed for 100 g of epoxy resin. Epoxy resins were vigorously mixed with the appropriate amount of crosslinker at room temperature in a disposable aluminum pan. In the case of the solid diamine DA10, an additional stirring time at 120◦C was necessary to obtain a complete dissolution of the amine in the epoxy precursor. Mixtures were then transferred into rectangular stainless steel molds (length: 40 mm, width: 8 mm, thickness: 1.5 mm). The method of curing was adapted from a literature procedure (Janvier et al., 2017). In summary, the resins were maintained under ambient conditions in the molds for 2 h and then for 2 h at 60◦C, for 20 h at 110◦C and for 1 h at 150◦C.

#### Hydrolytic Degradation Assays

Cured resin samples with dimensions about 4 × 8 × 1.5 mm, with weights ranging from 75 to 95 mg, were placed in a 10 mL sealed tube. An aqueous solution of HCl (3M) was added. The vials were heated at 60◦C for 40 min for equilibration. Incubations were then continued for 5, 20, 26, 43, and 50 h. The residual solid was then washed with deionized water, dried then weighed to determine the mass loss.

#### Measurement of Contact Angle

The contact angle (θ) of liquid water against the cured epoxy resins was measured by the sessile drop method at room

temperature with a drop shape analysis system (DIGIDROP and Visiodrop software). The contact angle was measured at three different locations for each thermoset and the average value was reported.

#### Estrogen Binding Affinity Testing

The agonistic and antagonistic potentials of bisphenols were analyzed following a literature method (Delfosse et al., 2012) in which ERα, PXR, and AR transcriptional activities were monitored by using corresponding reporter cells HELN ERα, U2OS AR, and HG5LN PXR cells, respectively. Activities were measured in relative light units (RLU) and 100% activities were assigned to the RLU value obtained with 10 nM agonist control (estradiol (E2), R 1881, SR 12813). The vehicle (DMSO) was tested as a control without any compound.

### RESULTS AND DISCUSSION

#### Synthesis of the Bio-Based Epoxy Resin Chemo-Enzymatic Synthesis

Chemical esterification of p-hydroxycinnamic acids with alcohol is unselective and side reactions generally lead to an unwanted product that needs to be removed by purification steps, generating waste to dispose of. In opposition, the lipase-mediated enzymatic synthesis offers some advantages for synthesizing esters, such as milder reaction conditions and selectivity. Hence, the selective synthesis of the bis/tri-ferulate resin epoxy precursors **e**, **f**, **g** and **h**, involving glycerol linker group, was based on chemo-enzymatic methods recently published (**Scheme 1**) (Pion et al., 2013; Hollande and Domenek, 2018).

Ferulic acid was first transformed either in ethyl dihydroferulate (**Scheme 1**, b) or in benzylated ethyl ferulate (**Scheme 1**, c), through a two-step one-pot process. Afterwards, a specific affinity of Candida antarctica Lipase B (CAL-B) toward glycerol allowed controlling the degree of functionalization of the latter, giving access to the symmetric bisphenol (**Scheme 1**, d) or to the fully functionalized glycerol triferulate (**Scheme 1**, e) in excellent yields and purities. Fatty acids were then grafted onto the available secondary alcohol of the intermediary compound d, immediately followed by a palladium-catalyzed hydrogenation to simultaneously reduce unsaturation and cleave the benzyl group, thus providing the GDFx compounds (**Scheme 1**, f, g, h. It is noteworthy to mention that all the syntheses described in **Scheme 1** have been performed at the multi-gram scale. Moreover, in a recent work, we have not only demonstrated the feasibility of the lipase-mediated transesterification of dihydroferulate ethyl ester at the kilo-scale (Teixeira et al., 2017), but also devised a sustainable and industrially relevant membrane-based technique, allowing both the purification of the targets and the recycling of solvents and unreacted reagents.

#### Endocrine Receptor Activities of GDFx Bisphenols

Nowadays, the development of new platform chemicals is deeply regulated, especially for bisphenolic compounds—such as BPA and other controversial endocrine disrupting chemicals (EDCs). Herein, endocrine activities of newly created GDFx bisphenols were investigated by evaluating their ability to interact with three types of receptors ERα, PXR, and AR. ERα is a member of the nuclear hormone receptors family, and its activity is regulated by the steroid estrogen sex hormone 17β-estradiol (E2). PXR is a member of the steroid and xenobiotic sensing nuclear receptors family and AR is a nuclear type androgen receptor that is activated by binding with any of the androgenic hormones. In **Figure 1A**, sex hormone E2 induces a 100% ERα activity at 5.10−10 M, which plateaus at higher concentrations. For BPA, estrogenic activity was observed at molarities of 10−7 and slightly above 60% at concentrations of 10−5 . In contrast, all newly created GDFx bisphenols, GDF10, GDF14, and GDF16, exhibited an activity similar to that of the control, even with increasing concentrations. In addition, **Figure 1B** reveals that the highest concentration tested for GDFx family (i.e., 10−5 M) did not induce any abnormal receptor activities on both PXR and AR. Finally, an endocrine disruption assay (ESI p22) indicated that phenolic structures were neither agonist nor antagonist ligands. Based on these results, that demonstrated the innocuousness of these newly created GDFx bisphenols, we were encouraged to further explore their use for the production of epoxies and epoxy-amine resins.

#### Glycidylation

Epoxy resins synthesis consists in the glycidylation of the different GDFx bisphenols, using the classical procedure, i.e., large excess of epichlorohydrin under alkaline conditions. However, to prevent the partial hydrolysis of the internal ester moieties observed by Maiorana et al. (2016) during classical glycidylation of bisferulate esters, milder temperatures, and shorter reaction times were applied. Thus, to offset this restriction, a catalytic amount of triethylbenzyl ammonium chloride (TEBAC) as phase transfer catalyst—well known for allowing higher epoxy functionality—was used (Fache et al., 2015a). Under such conditions, the phenolate formed is able to attack each one of the carbons of epichlorohydrin (**Scheme 2**), two of them, (b) and (c), leading to the epoxy ring opening. Thus, a basic treatment is necessary to close the open forms. Yet, not all the resulting species are diepoxies resins. Side reactions can occur, the most important being (b) the oxetane formation or (d) the formation of branched molecules due to the oligomerization of chlorinated intermediate (Ellis, 1993).

Resulting crude di/triglycidyl ether bis/triferulate epoxy resins, named GDFxEPO, and GTF-EPO, respectively, were obtained in high isolated yields (>85%). All the resins prepared herein (see ESI, p18) showed the appearance of an epoxy vibrational band at 912 cm−1 characteristic of oxirane C-O group on their FT-IR spectra. 1H NMR analyses showed the expected chemical shifts for the proton on the tertiary carbon of the oxirane ring at 3.40 ppm (see ESI, p17). However, signals were also observed between 3.75 and 3.95 ppm, which could be attributed to the oxetane or the oligomeric production epoxies pictured in **Scheme 2** (b) and (d). In order to investigate the potential side formation of such oligomeric epoxies, molecular weight distributions of GDFxEPO and GTF-EPO resins were measured by gel permeation chromatography. The results are shown in **Figure 2** and the ESI (p20-21).

The distribution of molecular species (Mn) in epoxy resins was analogous for each sample. Experimentally determined values (GPC) described the normalized weight distribution of epoxy systems in two monodisperse fractions according to the Mw/Mn values closed to 1.0. The first corresponds to the monomeric glycidyl ether of GDF10, GDF14, GDF16, and GTF, with Mn of 840, 970, 1,070, and 860 g/mol, respectively, whereas the second corresponds to Mn above 2,000 g/mol, suggesting the formation of oligomeric by-products through the dimerization of chlorinated intermediate depicted in **Scheme 2**. Nevertheless, the contribution of the high molecular weight tail to the cumulative molecular weight distribution is <5%. Considering the relatively low concentration of by-products estimated through 1H NRM and GPC analyses, we decided to use these epoxy resins directly without further purification. **Table 1** displays values of experimental and theoretical epoxide

FIGURE 1 | (A) Estrogenic activity (%) as function of concentration of E2, BPA and synthesized bisphenols. (B) Activities (%) of Erα, PXR and AR induced by synthesized bisphenol at 10−5 molarity.

equivalent weights (EEW) where the latter assumes 100% conversion to the corresponding monomeric glycidyl ferulate structures. Experimental EEWs values are 11 to 15% greater than that of the corresponding theoretical values. This result is attributed to the formation of dimers which results in the increase of the EEW of the synthesized resins above the theoretical value. Finally, **Table 1** also reports on the thermal behaviors of uncured resins. TGA analysis revealed thermostability (Td 5%) in the range of 311–340◦C under an inert atmosphere, and 304–321◦C under an oxidative atmosphere. Furthermore, the alkyl chain length does not significantly impact the degradation temperature.

#### Thermosets Synthesis

The chemical structures of all amine hardeners or curing agents are reported in **Figure 3**. Two bio-based (DA10 and DIFFA) and one extensively used fossil-based (IPDA), presenting different rigidities, were selected.

To ensure that epoxy resin has a good reactivity toward diamines, a DSC analysis of an equimolar mixture of resins and diamines was carried out (**Figure 4**). An exothermic peak was observed around 100◦C corresponding to the epoxy ringopening reaction to form the crosslinked polymer. In order to ensure chain mobility during gelation and to obtain optimal crosslinking content, a multistep temperature program was

TABLE 1 | Epoxy resins characterization.

\*Determined by TGA (10◦C.min−1 ).

aUnder N2 flow.

bUnder O2 flow.

elaborated. Mixtures were first left for 2 h at room temperature followed by two additional hours at 60◦C. Formulations were then cured at T ≈ TDSCcrosslinkingpeak (20 h at 110◦C), then at T > TDSCcrosslinkingpeak (1 h at 150◦C). Each difunctionalized GDFxEPO epoxy precursor was cured in presence of each curing agent leading to 9 formulations (**Table 2**, thermosets 1–9). In order to tune the properties of these thermosets, and as the formulations with diamines and difunctionalized GDFxEPO (i.e., thermosets 1–9) did not provide materials rigid enough to perform DMA analyses, 1 formulation with GTF-EPO (**Table 2**, thermoset 15) as well as 5 co-formulations with

IPDA, in variable ratios of difunctionalized (GDF14EPO) and tri-functionalized (GTF-EPO) epoxy resin (**Table 2**, thermosets 10–14) were also prepared.

#### Characterization of the Thermosets Thermogravimetric Analyses (TGA)

The thermal stability of cured thermosets was studied by thermal gravimetric analysis (TGA) under nitrogen flow. Td5% was defined as the temperature at which the thermosets lost 5% of its initial mass, Tdmax as the temperature at which the kinetic degradation of the thermoset occurs at the maximum rate and W%char corresponds to the relative amount of stable residue at a high temperature. **Table 2** sums up the values for the thermosets prepared and all thermograms are displayed in ESI p23-25. The lowest Td5% was obtained with DIFFA (**Table 2**; thermosets 4– 6), which might be explained by the initiation of the degradation by rupture of the curing agent segment, as reported by Ménard et al. (2017). Overall, higher Td5% and Tdmaxwere measured with IPDA- and DA10-containing resins. Concerning the high temperature stable char content (W%char), values were around 10% for GDFxEPO resins; moreover, the higher the GTF-EPO ratio in epoxy resin system, the higher the char content. This parameter would be particularly interesting for the development of flame-retardant materials.

#### Differential Scanning Calorimetric Analyses (DSC)

The degree of cure of the resins and glass transition temperatures were both determined through DSC. Analyses of post-cured resins showed no residual exothermic peak above or below the glass transition corresponding to further curing reactions (ESI, p26-28). These results indicate that if additional cure ever occurred during DSC heating, it is below the instrument level of detection. Therefore, consumption of epoxy moieties and concurrent events of cross-linking during the curing cycle

aDetermined by TGA (10◦C.min-1, N2 flow).

bDetermined by DSC on the second heat-cool-heat cycle (10◦C.min-1, N2 flow).

cDetermined by DMA (frequency 1 Hz, amplitude 7 µm, 3◦C.min-1). E'glassy at 30◦C, E'elastic at 150◦C.

adopted herein results in a degree of cure that closely approaches its upper limit. Furthermore, the curing of difunctionalized resins (**Table 2**, entries 1–9 and 15) with hardeners presenting various rigidities (DA10 < DIFFA < IPDA due to their aliphatic, aromatic, and cyclic configuration), respectively, lead to a restricted range of Tg values (from 3 to 23◦C). Hence, the most flexible amine (DA10) provided the lowest Tg values, while the IPDA-containing thermosets led to the highest Tg . Nevertheless, for a given diamine, thermosets exhibit quite similar Tg , meaning that the impact of the diamine on the rigidity of the network outweighs that of the alkyl chain length (**Figure 5A)**, Entries 1–3; 4–6; 7–9).

When IPDA reacted with tri-functionalized resins, i.e., GTF-EPO, the Tg value increased up to 60◦C due to the higher epoxy functionally (2 vs. 3) and higher cross-link density (**Figure 5A**, 15). Moreover, as shown in **Figure 5B**, it is noteworthy to

FIGURE 6 | (A) Tan δ curves of (blue) 12, (green) 13, (red) 14, and (yellow) 15 thermosets; (B) Storage modulus (E'glassy) at 30◦C of (blue) 12, (green) 13, (red) 14, and (yellow) 15 thermosets.

mention that the Tg increased linearly with the tri-functionalized resin GTF-EPO content (R² = 0.9898).

#### Dynamic Mechanical Analyses (DMA)

Mechanical properties of the cured resins were evaluated by DMA on each sample with a glass transition upwards of 30◦C (thermoset 12–15). **Figure 6** and **Table 2** provide (a) tan δ values as a function of temperature and (b) storage modulus values at the glassy state at 30◦C. Similar to the previous linear relationship observed for Tg and GTF-EPOwt%, the same trend was observed for tan δ that increased with the increasing incorporation of the GTF-EPO to the GDF14EPO system. As discussed previously, the thermal behaviors observed with the addition of GTF-EPO was a consequence of the decrease of EEWs of the co-formulated systems relative to the GDFxEPO resin. In other words, co-formulated resins will lead to thermosets with relatively higher crosslink densities. On the other hand, when the ratio of GDF14EPO in GTF-EPO does not exceed 25 wt%, thermosets exhibit a higher storage modulus than that of the GTF-EPO resin (**Figure 6B**; entries 13 and 14 vs. 15). It is worth mentioning that the storage modulus at 30◦C of epoxy systems which contain 10 wt% of GDF14EPO is 50% higher than that of the GTF-EPO thermoset (2,523 vs. 1,685 MPa). This rather unexpected outcome may be explained by the more efficient interaction/distribution of lipophilic GTFxEPO epoxies during the gelation/curing induced by their "surfactant" properties.

#### Degradation and Wettability Behavior

To gain further insight on how surface hydrophobicity varies as a function of the alkyl chain length, contact angles were measured (**Table 3**). Whatever the curing agent, increasing the alkyl chain length from 12 to 18 carbons (i.e., lauric acid -> palmitic acid -> stearic acid) resulted in an increase of the contact angle (**Table 3**; thermosets 1 vs. 3, 4 vs. 6, and 7 vs. 9). However, it is worth mentioning that the nature of the curing agent does not significantly impact the hydrophilicity of thermosets as for a same TABLE 3 | Characterization of weight loss plots and contact angle.

resin GDFxEPO cured with the three different curing agents (i.e., DA10, DIFFA, and IPDA) contact angles differ by < ±7 ◦ .

In addition, thermosets weight loss, determined as a function of the incubation time in acidic aqueous solutions, was plotted (**Figure 7**). Data showed a decrease in mass upon incubation with linear behavior, according to the R² values for weight loss plots (**Table 3**). Finally, it is noteworthy to mention that the length of the fatty acid chain grafted onto GDFx epoxies had a significant effect on the hydrolytic degradation rate of cured thermosets. Indeed, an increase in hydrophobicity of the thermoset, by increasing the length of the fatty acids graft, was consistent with a corresponding decrease of the degradation rate measured in mg.h−1 and the susceptibility of that surface to hydrolysis.

### CONCLUSION

Herein, the chemo-enzymatic synthesis, thermo-mechanical properties, wettability and acidic hydrolysis of 100% renewable

ferulic-, glycerol- and fatty acids-based bis- and triphenols - GDFx and GTF - are described. The estrogenic activity of GDFx with fatty acid chain lengths of 12, 16, and 18-carbons was quantified and compared to bisphenol A and 17β-estradiol and showed no significant activity for ERα, PXR, and AR receptors. Bio-based GDFx and GTF bis/triphenols were then successfully converted to their corresponding di- and triglycidyl ether epoxy resins (i.e., GDFxEPO and GTF-EPO, respectively) through a TEBAC-mediated glycidylation. GDFxEPO was then cured with DIFFA, DA10 and IPDA diamines; GDF14EPO was also coformulated with tri-functionalized GTF-EPO and IPDA. All the resulting epoxy-amines exhibited relatively high thermostability with Td5% ranging from 282 to 310◦C and Tg between 3 and 62◦C. Surprisingly, in the case of GDFxEPO-diamine resins, for a given diamine, the chain length of the fatty

#### REFERENCES

acid moiety did not significantly impact the Tg . However, one can tailor the Tg by playing with the diamine nature and achieve Tg up to 23◦C with IPDA and GDF14EPO. To further tailor the Tg and to reach higher values, GTF-EPO must be added to the formulation. In such formulations, DSC and DMA analyses showed that the Tg and the storage modulus can also be modulated by finely adjusting the GTF-EPO content. Finally, with regards to wettability and degradability, the chain length of the fatty acid was found to provide a simple but powerful approach to tailor the wettability and the susceptibility to hydrolysis of the GDFxEPO-based epoxyamine resins. Indeed, the shortest fatty acid provided the highest wetting and hydrolysis rate, and vice-versa. This study therefore demonstrates the great potential of combining ferulic acid, glycerol and fatty acids using chemo-enzymatic processes for the preparation of epoxies and epoxy-amine networks with tunable properties.

#### AUTHOR CONTRIBUTIONS

FA and RG conceived the research. FA and SD managed the research. LH and ID performed the chemo-enzymatic reactions and the characterizations. PB performed the endocrine disruption assays. FA and LH provided the technical guidelines, reviewed the results, wrote, and drafted the article. FA, RG, and SD reviewed and approved the article.

#### ACKNOWLEDGMENTS

The authors are grateful to Région Grand Est, Conseil Départemental de la Marne and Grand Reims for financial support.

#### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2019.00159/full#supplementary-material

ACS Sustain. Chem. Eng. 6, 14812–14819. doi: 10.1021/acssuschemeng. 8b03340

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Hollande, Do Marcolino, Balaguer, Domenek, Gross and Allais. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Nanocellulose Xerogels With High Porosities and Large Specific Surface Areas

Shunsuke Yamasaki 1 , Wataru Sakuma1 , Hiroaki Yasui 1 , Kazuho Daicho1 , Tsuguyuki Saito1 \*, Shuji Fujisawa1 , Akira Isogai 1 and Kazuyoshi Kanamori 2

*1 Department of Biomaterial Sciences, Graduate School of Agricultural and Life Sciences, The University of Tokyo, Tokyo, Japan, 2 Department of Chemistry, Graduate School of Science, Kyoto University, Kyoto, Japan*

Xerogels are defined as porous structures that are obtained by evaporative drying of wet gels. One challenge is producing xerogels with high porosity and large specific surface areas, which are structurally comparable to supercritical-dried aerogels. Herein, we report on cellulose xerogels with a truly aerogel-like porous structure. These xerogels have a monolithic form with porosities and specific surface areas in the ranges of 71–76% and 340–411 m2 /g, respectively. Our strategy is based on combining three concepts: (1) the use of a very fine type of cellulose nanofibers (CNFs) with a width of ∼3 nm as the skeletal component of the xerogel; (2) increasing the stiffness of wet CNF gels by reinforcing the inter-CNF interactions to sustain their dry shrinkage; and (3) solvent-exchange of wet gels with low-polarity solvents, such as hexane and pentane, to reduce the capillary force on drying. The synergistic effects of combining these approaches lead to improvements in the porous structure in the CNF xerogels.

Keywords: cellulose nanofiber, xerogel, aerogel, porous material, ambient pressure drying

## INTRODUCTION

Aerogels are defined from a processing perspective as porous structures that are obtained by supercritical drying of wet gels (Hüsing and Schubert, 1998). In general, aerogels are mesoporous and combine a high porosity and large specific surface area (SSA). Owing to their specific porous structure, aerogels can excellently function as thermal, acoustic, and electrical insulators as well as catalyst supports, separators, and adsorbents. Some aerogels can be optically transparent yet thermally insulating, and have drawn interest as transparent insulators that might reduce thermal energy loses from windows in offices and automobiles (Hayase et al., 2014; Zu et al., 2018). Aerogels were first reported in 1931 (Kistler, 1931); however, it remains difficult to put aerogels to practical use because of two problems: one is the mechanical brittleness of aerogels; and the other is the process of supercritical drying, which limits productivity and scalability. Therefore, much effort has recently been made to overcome these problems.

In terms of the brittleness, one approach is to use cellulose nanofibers (CNFs) as the skeletal units of aerogels. CNFs are wood-derived emerging materials, which have excellent mechanical strength (∼3 GPa) (Saito et al., 2013) and solvent tolerance. CNFs are thus suitable building blocks for aerogels. Pääkkö et al. (2008) demonstrated the flexibility of structurally aerogel-like CNF cryogels, and Kobayashi et al. (2014) have reported on mechanically tough aerogels with good optical transparency and thermal insulating properties, based on a very fine type of CNFs with a width of ∼3 nm.

#### Edited by:

*Gil Garnier, Bioresource Processing Institute of Australia (BioPRIA), Australia*

#### Reviewed by:

*Dong Gu, Wuhan University, China Makoto Tabata, Chiba University, Japan Lorenz Ratke, Institute of Materials Research, German Aerospace Center, Germany*

#### \*Correspondence:

*Tsuguyuki Saito<EMAIL_ADDRESS>
#### Specialty section:

*This article was submitted to Chemical Engineering, a section of the journal Frontiers in Chemistry*

Received: *22 February 2019* Accepted: *23 April 2019* Published: *07 May 2019*

#### Citation:

*Yamasaki S, Sakuma W, Yasui H, Daicho K, Saito T, Fujisawa S, Isogai A and Kanamori K (2019) Nanocellulose Xerogels With High Porosities and Large Specific Surface Areas. Front. Chem. 7:316. doi: 10.3389/fchem.2019.00316*

**169**

One approach to addressing the processing problem is to switch focus from aerogels to xerogels. Xerogels are defined as porous structures that are obtained by evaporative drying of wet gels. However, producing structurally aerogel-like xerogels under ambient conditions remains particularly challenging. Even thick block-like hydrogels shrink considerably during evaporative drying owing to capillary forces; hence, the resulting xerogels inevitably become thin films with a very low porosity. One approach is to exchange the solvent of the wet gels, which is typically water, to a solvent of lower polarity, such as hexane and pentane, to reduce the capillary force. This concept was reported in a paper by Prakash et al. (1995), in which aerogel-like silica xerogels were formed as a thin film with a high porosity (∼99%) through solvent exchange after chemically modifying the silica skeleton to be hydrophobic. Later, Kanamori et al. (2007) further developed the idea and reported that monolithic xerogels can be optically transparent and even exhibit rubbery compression with the use of organosilicons as the skeletal precursor.

In the field of cellulose science, Toivonen et al. (2015) reported a landmark study of CNF xerogels, which were similarly produced through a solvent exchange process. Here CNFs were formed into a wet cake by vacuum filtration of their water dispersion, followed by drying under ambient conditions after solvent exchange with octane. The resulting xerogels were mesoporous and in a film form with a thickness of ∼25µm. Their porosity and SSA were ∼60% and 200 m2 /g, respectively, which are high characteristics among xerogels but still not within the structural range of supercritical-dried CNF aerogels (porosity 60–99%, SSA 300–600 m2 /g) (Sehaqui et al., 2011; Kobayashi et al., 2014; Sakai et al., 2016). In another route to xerogel formation, freeze-thawing of aqueous CNF dispersions has been used to form block-like hydrogels, followed by solvent exchange. The resulting xerogels were monolithic with a high porosity of over 98% but were macroporous. Their SSA values were ∼30 m2 /g (Li et al., 2017; Erlandsson et al., 2018). Furthermore, a non-CNF skeletal, different type of cellulosic xerogels with a high porosity (∼90%) has been produced by a molecular dissolutionregeneration process; however, the SSA values of these materials were ∼10–30 m2 /g (Pour et al., 2016).

Here we report on CNF xerogels with a truly aerogel-like porous structure. The xerogels are in a monolithic form with high porosities and SSA in the ranges of 71–76% and 340–411 m2 /g, respectively. Our strategy is based on three ideas: (1) the use of a very fine type of CNFs with a width of ∼3 nm; (2) increasing the stiffness of precursor hydrogels by reinforcing inter-CNF interactions; and (3) solvent exchange as described above. To our knowledge, these ideas have not yet been combined, and the synergistic effects of these factors lead to the development of porous structure in CNF xerogels.

### MATERIALS AND METHODS

#### Materials