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**Conflict of Interest:** 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 © 2020 Zhang, Calderon, Ebbing, Elston, Byrd and Tsao. 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.
<figure><img src="image_50.png" /></figure>

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# Direct Graphene Synthesis on Lithium Niobate Substrate by Carbon Ion Implantation

Yuhang Xu¹, Fei Lu¹*, Kaijing Liu¹ and Changdong Ma²

¹School of Information Science and Engineering, Shandong University, Qingdao, China, ²Department of Radiation Oncology, Qilu Hospital of Shandong University, Jinan, China

For this study, a dozen-micron-scale four to seven layers of graphene were synthesized directly onto Lithium Niobate (LiNbO₃, LN) by carbon implantation in a nickel covered LiNbO₃ film. The 1.52E16 ions·cm⁻² carbon ions of 80 keV were implanted into the nickel film to a depth of 100 nm, which is close to the interface between the nickel film and the LiNbO₃. The implanted carbon atoms dissolved in nickel at an elevated temperature and diffused towards the interface between the nickel film and LiNbO₃ during annealing. After the nickel film fell off, the graphene material was found by Raman spectroscopy and confirmed by the Atomic Force Microscope topographic image. This implantation combined with patterned metal film suggests possibilities for engineering two-dimensional graphene with special confinement. It also provides a valuable and novel way for integrating graphene-wafer structure for other available substrate wafer in microelectronics and photonics devices.

**Keywords:** graphene, direct synthesis, lithium niobate substrate, carbon ion implantation, nickel film, annealing

OPEN ACCESS

**Edited by:**
Ming Xu,
Huazhong University of Science and
Technology, China

**Reviewed by:**
Jiangtao Di,
Suzhou Institute of Nano-Tech and
Nano-Bionics (CAS), China
Horacio Javier Salavagione,
Consejo Superior de Investigaciones
Científicas (CSIC), Spain

*Correspondence:
Fei Lu
lufei@sdu.edu.cn

**Specialty section:**
This article was submitted to Carbon-
Based Materials,
a section of the journal
Frontiers in Materials

**Received:** 13 June 2020
**Accepted:** 28 August 2020
**Published:** 10 November 2020

**Citation:**
Xu Y, Lu F, Liu K and Ma C (2020)
Direct Graphene Synthesis on Lithium
Niobate Substrate by Carbon
Ion Implantation.
Front. Mater. 7:572280.
doi: 10.3389/fmats.2020.572280

## INTRODUCTION

Graphene has received increasing attention since the successful preparation of it in 2004 (Novoselov et al., 2004). Single-layer graphene is a two-dimensional honeycomb lattice structure composed of a single layer of carbon atoms, and each carbon atom is bonded to the surrounding carbon atoms through SP² hybridization (Xu et al., 2013). This structure has excellent conductivity. The electron mobility of graphene is higher than that of carbon nanotubes or silicon crystals. At room temperature, it can reach 15000 cm² (V⁻¹·s⁻¹), while the resistivity is as low as about 10⁻⁶ Ω cm (Chen, 2015). Graphene also has high transmittance to light. The graphene monolayer absorbs only 2.3% of visible light, that is, it has a high transmittance (97.7%) in a wide spectral range from visible to infrared (Zhu et al., 2010). Graphene also has exceptionable mechanical properties, due to an elastic modulus of 1.0 TPa, its inherent tensile strength of 130 GPa and a two-dimensional ultimate plane strength of 42 N·m⁻¹ (Lee et al., 2008). As a material, graphene has excellent thermal conductivity (up to 5500 W (m⁻¹k⁻¹)) and high specific surface area (up to 2630 m²g⁻¹) (Wang et al., 2017). Due to this special structure, outstanding thermal conductivity and mechanical properties, graphene has aroused great interest from the scientific community in recent years (Geim and Novoselov, 2007; Geim, 2009; Ma et al., 2018; Martín-García et al., 2018), and is expected to trigger a revolution in the field of modern optoelectronic technology.

The preparation and application of graphene have been explored in many studies over the past decade. Common preparation technologies include micro-mechanical peeling (Novoselov et al., 2004), pyrolysis of silicon carbide (Juang et al., 2009), Chemical Vapor Deposition (CVD) (Li et al.,
2009) and reduction of graphene oxide (Yin et al., 2011). The CVD method allows for the production of wafer-scale graphene layers, which has become the most popular technology. In recent years, a physical process has been proposed to synthesize graphene that involves the ion implantation of carbon into metal film and post-implantation annealing. Compared with a CVD method, ion implantation is a non-thermal equilibrium process, and has no requirements in terms of the carbon solubility of the material being implanted. The ability to control the number of layers of synthetic graphene may be possible by precisely adjusting the implantation dose (Garaj et al., 2010; Zhang et al., 2013). At present, transition metal nickel or copper are usually used to form graphene by carbon implantation (Garaj et al., 2010; Baraton et al., 2011a; Mun et al., 2012; Wang et al., 2013; Zhang et al., 2013; Lee et al., 2014; Kim et al., 2015; Wang et al., 2016). As the solubility of carbon in nickel decreases significantly with the decline of temperature, the carbon dissolved in nickel at high temperature will become saturated with the decrease of annealing temperature, and with the progress of thermal diffusion, they will eventually spread to the surface of the metal to form a single-layer or multi-layer graphene (Lander et al., 1952; Berry, 1973; Baraton et al., 2011b).

The application of graphene must be based on the graphene/
substrate structure, and a lot of research has been conducted in
recent years, but these studies are mainly focused on graphene/
Si(SiO₂) structure (Yoon et al., 2009; Garaj et al., 2010; Baraton
et al., 2011b; Wang et al., 2011; Kumari et al., 2014; Lee et al.,
2014; Dong et al., 2016; Onishi et al., 2017; Zuppella et al., 2017;
Faggio et al., 2019; Ma et al., 2020). Few studies have focused on
graphene/optical crystal, but among those that have , the
combination of graphene and Lithium Niobate (LiNbO₃, LN)
is one of the most promising structures for development. LN is a
nearly all-around artificial crystal, it has ferroelectric,
piezoelectric, pyroelectric, photovoltaic, photoelastic,
photorefractive and other properties, and has a wide range of
applications unmatched by other crystals, known as "silicon
materials in the optical field" (Graca et al., 2012; Boes et al.,
2018; He et al., 2019; Kong et al., 2019). Its combination with
graphene may have an impact on the quantum properties of two-
dimensional materials, resulting in new photoelectric properties,
which will inevitably have a wide range of potential applications
(Jin et al., 2013; Baeumer et al., 2015; Salas et al., 2017a; Salas et al.,
2018). For example, some studies have theoretically explored the
interfacial properties of the combination of graphene and LN
crystals (Wang et al., 2015; Salas et al., 2017b). One typical
example is the modulation of carrier doping of graphene via
LN polarization properties. The results show that by adjusting the
carrier doping in graphene in the graphene/LN structure, the
formation and propagation of surface plasmons polaritons can be
controlled.

Previously, graphene bonding to substrates was mostly based
on the graphene transfer method. However, the resulting
graphene/LN or Si (SiO₂) structures often suffer from
graphene defects caused by the transfer process which leads to
additional graphene degradation. Therefore, it's necessary to
study the direct synthesis of graphene/LN structure. The
formation of graphene on the LN surface using direct carbon

implantation of LN has been reported in our previous study (Liu
et al., 2019), which found that the carbon implanted into the LN
can accumulate on the crystal surface by out-diffusion at a
moderate temperature and eventually form graphene. The
carbon diffusion pathways, on the other hand, is related to the
distribution of damage of LN lattice induced by carbon
implantation. However, in this experiment only turbostratic
graphene was demonstrated on the LN surface. LN has a
complex structure, the solubility and diffusion behavior of
carbon in LN is still unclear.

In this article, we report the direct synthesis of graphene on the
surface of LN by annealing the post-carbon-implantation sample
consisting of surface transition metal film and LN substrate. In
the present experiment the carbon ions were implanted into the
metal film covering the surface of the crystal rather than into the
LN itself. This approach used the nickel as a catalyst, and graphene
growth was based on the large solubility change of carbon ions in
nickel films at different temperatures. Graphene was synthesized
at the interface of nickel/LN, and the realization of immediate
contact between graphene and LN omits the conventional
intermediate process of graphene transfer. Our research
provides a new method of preparation for applying graphene
on substrates without transfer. Importantly, it provides a new
approach that enables the direct synthesis of graphene that can
create specific shapes and sizes on any substrates, which is often
emphasized in photoelectric crystals such as LN.

MATERIALS AND METHODS

A schematic diagram of the experimental method used in this
study is shown in **Figure 1**. The substrate was Z-cut 10 × 10 ×
0.5 mm³ bulk LN with Chemical Mechanical Polishing on one
10 × 10 mm² surface for the following nickel coating. First, a layer
of nickel foil is deposited on this surface as an overlay, and then
the deposited sample is implanted with carbon ion. During the
annealing process, the implanted carbon moves to the interface
between the nickel foil and LN crystal. Ideally, a few layers of
graphene will be formed, and finally, the nickel foil is removed to
gain the graphene/LN structure.

The thin layer of nickel was deposited on the surface of LN by
resistance heating evaporation. As shown in **Figure 2A**, the
thickness of the nickel foil measured by Atomic Force
Microscope (AFM) is around 134 nm. The dose of carbon ion
implantation we used is 1.52E16 ions·cm⁻², which is selected to
match the carbon content of four-layer graphene, and the ion
energy of 80 keV ensures that the implanted carbons distribute
near the interface between the nickel foil and LN crystal.
**Figure 2B** shows the carbon ion distribution along the depth
profiles simulated by Stopping and Range of Ion in Matter
(SRIM) in the nickel/LN sample. According to SRIM
simulation, the peak of carbon concentration is at roughly
100 nm below the surface of nickel film and is close to the
interface between nickel and LN. The variation of the carbon
concentration with the depth of implantation approximately
follows the Gaussian distribution and the nickel film holds
more than 80% in all implanted carbon. Subsequently, the as-
<figure><img src="image_51.png" /><figcaption>**FIGURE 1** | A schematic diagram of the steps of the experiment.</figcaption></figure>

<figure><img src="image_52.png" /><figcaption>**FIGURE 2** | (A) AFM height profiles along the straight line from uncoated notch to the coated area (B) Structural schematic diagram of the coated bulk LN crystal with the distribution of carbon ions along with the depth profiles, simulated by SRIM.</figcaption></figure>

implanted sample was annealed in the furnace. It was first heated to 900 °C at a rate of 10 °C·min⁻¹, an kept at this temperature for 30 min, then allowed to naturally cooled to room temperature. The entire annealing process was carried out in N₂ ambient, with a gas flow rate of 100 sccm (standard-state cubic centimeter per minute) and pressure of 6E4Pa.

## RESULTS AND DISCUSSIONS

The sample was first observed with an optical microscope, the images of the sample before and after annealing were shown in **Figures 3A,B**, respectively. Wrinkles appeared on the surface of the annealed sample. This phenomenon occurs on the one hand, because of the poor adhesion between the deposited metal film and the LN substrate, and on the other, due to the different thermal expansion properties of the two materials. **Figure 3C** shows the enlarged two-dimensional AFM image of the wrinkles on the surface of the sample, and **Figures 3D, E** are the respective height profiles along the solid red line one and line two marked in **Figure 3C**, respectively. As revealed by these images, the average height of the folds is around 0.7–1.8 µm and the width is about 10 µm. To determine the cause of the wrinkles, another two kinds of samples (one without carbon ion implantation and the other

with carbon ion implantation) of the same coating treatment were annealed under the same conditions. It was found that both samples had similar wrinkles. This shows that the occurrence of wrinkles on the surface of this sample depends on the surface nickel film and the LN substrate themselves, and has nothing to do with carbon ion implantation.

No information of carbon material was detected by Raman spectroscopy on the upper surface of the nickel film, which excludes the case where carbon is precipitated from the nickel surface. The existence of wrinkles makes the nickel film easy to peel off. After the nickel film fell off, the LN surface was observed with a scanning electron microscope (SEM). A strong carbon signal was detected, indicating that the implanted carbon was mainly deposited at the interface between nickel and LN after annealing. Some carbon atoms converge into large clusters, which are distributed on the surface of LN. **Figures 4A, B** show the SEM image and its Energy Dispersive Spectrometer (EDS) mapping distribution of carbon from one of the clusters respectively. The highlights of the carbon distribution seen in **Figure 4B** fit well with the material distribution profile, and the Raman Spectra with the excitation wavelength at 532 nm from three, positions marked in **Figure 4A**, also confirmed that the aggregated carbon exhibits an amorphous structure. As can be seen in **Figure 4C**, the Raman curves at three locations show similar profiles. The D and G peaks
<figure><img src="image_53.png" /><figcaption>FIGURE 3 (A, B) The optical microscope images of the surface of the bulk LN crystal with nickel film before (A) and after annealing (B) (C) AFM topographic image of the wrinkles, and (D)–(E) AFM height profiles along the solid red line one and line two in (C).</figcaption></figure>

are merged into one peak, while the 2D peak shows a very gentle and wide span distribution. This characteristic indicates that the carbon here is amorphous carbon with poor crystallinity (Shimodaira and Masui, 2002; Kaniyoor and Ramaprabhu, 2012; Sathish-Kumar et al., 2012). **Figures 4D,E** indicate the AFM of this cluster and its height profiles along marked red solid lines respectively. The height of the carbon material is about 1 µm, which is consistent with the height of the bulge in the wrinkled nickel film in **Figure 3**, which is around 0.7–1.8 µm. It is reasonable to suppose that the amorphous carbon material gathered and developed in wrinkled bumps and that the bulged nickel film provides space for holding these carbon materials. Improving the deposit of nickel film and its adhesion to LN is one of the key factors for ensuring the uniformity of carbon distribution.

We also found other forms of carbon structures on the surface of the LN, as shown in **Figure 5A** is the SEM image from part of the sample surface, and **Figures 5(B)–(D)** are the corresponding EDS mapping diagrams of Nb, C and O elements at the same area, respectively. It can be seen that the elements Nb and O, which come from the LN substrate, are distributed uniformly. While the distribution of element C seen in **Figure 5C** is consistent with the morphology of **Figure 5A**, indicating that the spots on the surface are carbon materials. Subsequently, Raman spectra with 532 nm excitation wavelength for the five marked positions one to five in

**Figures 5A,C**, which correspond to exactly the same positions, were obtained, as shown in **Figure 6**.

The Raman spectra at positions one and two are similar, revealing spectral features and uniquely characteristic of graphene films: a G peak (1,580 cm⁻¹) associated with sp² carbon bond stretching, a D peak (1,345 cm⁻¹) arising from symmetry-breaking features, such as graphene defects and domain boundaries, and a 2D peak (2,690 cm⁻¹) sensitive to the graphene inter layer interactions (Garaj et al., 2010; Han et al., 2011; Mun et al., 2012; Lee et al., 2014). The ratios of Iₔ to Iₐ (a measure of disorder) and I₂ₔ to Iₐ (determined by the number of stacked graphene layers) are very similar, and are often used as the main parameters to judge the structural defects and the number of layers of graphene (Garaj et al., 2010; Lee et al., 2014). The corresponding ratio values of 1.1 and ~0.5 indicate that the number of graphene layers is four to seven layers (Yu et al., 2008; Wang et al., 2013; Zhang et al., 2013). At the same time, the ratios of Iₔ to Iₐ are also higher at positions one and two, indicating that the graphene is defective. The graphene we obtained had many defects, including the mixing of multiple structures, graphene fragmentation, folding and structural distortion. This is the main reason for the larger Iₔ/Iₐ ratio. The spectra at positions three to five are shown in **Figure 6B**, where three distinct Raman peaks of D, G, and 2D are also visible. However, the ratio of I₂ₔ to Iₐ is much lower than those of
<figure><img src="image_54.png" /><figcaption>FIGURE 4 | (A) SEM image of the amorphous carbon material on LN (B) EDS mapping for C element in the region shown in (A) (C) Raman Spectra of three positions shown in (A) (D) AFM topographic image of the amorphous carbon material, and (E) AFM height profiles along the solid red line in (D).</figcaption></figure>

<figure><img src="image_55.png" /><figcaption>FIGURE 5 | (A) SEM image of graphene on LN (B)–(D) EDS mapping for Nb, C, and O elements in the region shown in (A).</figcaption></figure>
<figure><img src="image_56.png" /><figcaption>**FIGURE 6** | Raman Spectra of (A) positions one and two, (B) positions three, four, and five in Figure 5.</figcaption></figure>

positions one and two, which is only around 0.3, suggesting that the graphene here could be a stacked multilayer structure. This result indicates that the carbon materials at these positions have the structural characteristics of turbostratic graphene (Garlow et al., 2016; Liu et al., 2019). Turbostratic graphene is a multilayer structure with a disordered arrangement, and its properties are significantly different from the structured two-dimensional multilayer graphene. Since our research focus is to obtain a two-dimensional multilayer graphene structure on LN, the results of positions one and two are of particular note.

The SEM image around position one in **Figure 5** is shown in **Figure 7A**, some layered structures were observed. Subsequently, we conducted AFM scanning around positions one and two in **Figure 5**, and the results are shown in **Figures 7B-E**. A depression in the dark spot can be found at the test center point due to the loss of graphene caused by multiple Raman spectroscopy measurements. The height difference between the depression and the surrounding area represents the thickness of the multilayer graphene here. **Figures 7B,C** are the AFM topographic image around locations one and two, and **Figures 7D,E** are the profiles along the three different solid yellow lines in **Figures 7B,C**, respectively. It can be seen that there is a downward hollow in **Figures 7D,E**, which is the edge depression of a few layers of graphene. Although the three height change curves of each set in the figure are slightly different, on average, the depth of the depression at position one and two are about 46 and 70 nm respectively. These results are consistent with the estimations from our Raman results, an $I_{2D}/I_G$ ratio of 0.58 in position one representing four to six layers of graphene, while 0.51 in position two represents five to seven layers of graphene. They also all correspond to the thickness

values of four to six layers and five to seven layers multilayer graphene measured with AFM in references (Wu et al., 2009; Lee et al., 2014).

The direct two-dimensional synthesis of graphene on the surface of LN by carbon implantation has been proved to be feasible in our experimental scheme, although there are still many areas that need improvement. The current experimental results show that as a host for carbon ions, the quality of the nickel film and the uniform adhesion between the nickel and the LN substrate are important factors that affect the formation and structure of graphene. One of the improvements to the current experimental method is to use magnetron sputtering instead of vacuum evaporation to grow nickel thin films, which will increase the uniformity of nickel and the adhesion to the substrate.

The control of the annealing process is also an important factor, because proper annealing may reduce the thermal mismatch between the two materials during expansion and contraction. Thermal annealing is the only way to form graphene through carbon precipitation. The sample is firstly annealed at high temperature to accommodate more carbon and then cooled down so that the carbon becomes saturated in nickel as the temperature decreases, and the graphene is formed when the carbon overflows to the interface between nickel and LN with thermal diffusion. Since the diffusion property of carbon depends on several variables, such as the concentration gradient of carbon, the distance from carbon to the nickel/LN interface, the temperature range and temperature rise and fall speed, and ambient gas, etc., make the determination of annealing conditions particularly difficult. Therefore, the establishment of appropriate annealing conditions requires
<figure><img src="image_57.png" /><figcaption>FIGURE 7 | (A) SEM image around position one in Figure 5, (B,C) are the AFM topographic image around locations one and two in Figure 5, (D,E) are the profiles along the three different solid yellow lines in figure (B) and figure (C).</figcaption></figure>

repeated adjustments, improvements, and measurements, which are currently being carried out in our research.

CONCLUSION

In our experiment, the 1.52E16 ions·cm⁻² carbon ions of 80 keV were implanted to the nickel/LN substrate, which consists of an LN crystal covered with a layer of 134 nm thick nickel metal film. The implantation depth was 100 nm according to simulation by SRIM, which is close to the interface between the nickel film and LN. During post-implanted annealing, the sample was heated to 900°C and held for 30 min, then cooled down naturally *in situ*. After annealing, the nickel film on the surface of LN was wrinkled and easy to fall off. This can be attributed to the weak bonding between the nickel film and LN substrate and their mismatch thermal expansion coefficients. No signal of carbon material was detected by Raman spectroscopy from the upper surface of the nickel film, which

excludes a result of carbon precipitation from the upper surface of the nickel film. After the nickel film fell off, two kinds of carbon structures were found on the surface of the LN substrate by SEM, EDs, and Raman spectroscopy. First, amorphous carbon with poor crystallinity gathered and developed under the bulge of the wrinkled nickel film. Second, the graphene structure with an uneven number of layers mixed with turbostratic graphene and four to seven layers of two-dimensional graphene, as confirmed by AFM topographic image. The direct synthesis of two-dimensional graphene on the surface of LN was proved in our experimental scheme, which draws on the experiences of predecessors in the synthesis of graphene by carbon ion implantation of nickel, but our experiment omitted the process of film transfer. In particular, implantation combined with patterned metal film suggests the possibility of engineering two-dimensional graphene with special confinement, which is often emphasized in photoelectric crystals such as LN. The flexibility of photolithography technology in the metal pattern may makes it achievable to potentially form an air-graphene-wafer system with
special nanostructure, such as graphene nano-disk. Direct synthesis of graphene with patterned and even distribution on the LN is our next research goal, and further experiments are in progress.

## DATA AVAILABILITY STATEMENT

The raw data supporting the conclusions of this article will be made available by the authors, without undue reservation.

## AUTHOR CONTRIBUTIONS

YX, and FL: conceptualization, data curation, formal analysis, methodology, and writing - review and editing. FL: funding

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acquisition, supervision and validation. YX, FL, and KL: investigation. YX, KL, and CM: project administration. YX: visualization and writing - original draft.

## FUNDING

This study was funded by the National Natural Science Foundation of China (Grant Number 61975094).

## ACKNOWLEDGMENTS

We would like to thank Dr Mingsheng Xu, School of Microelectronics, Shandong University for the use of his equipment.

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**Conflict of Interest:** 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 © 2020 Xu, Lu, Liu and Ma. 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.
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# Effect of Co-solution of Carbon Precursor and Activating Agent on the Textural Properties of Highly Porous Activated Carbon Obtained by Chemical Activation of Lignin With $H_3PO_4$

Imane Moulefera<fnref n="1" />,<fnref n="2" />, Francisco J. García-Mateos<fnref n="1" />, Abdelghani Benyoucef<fnref n="2" />, Juana M. Rosas<fnref n="1" />*, José Rodríguez-Mirasol<fnref n="1" /> and Tomás Cordero<fnref n="1" />

## OPEN ACCESS

**Edited by:**

Maria Chiara Bignozzi,
University of Bologna, Italy

**Reviewed by:**

Francisco Carrasco-Marin,
University of Granada, Spain
Alexander M. Puziy,
National Academy of Sciences
of Ukraine (NAN Ukraine), Ukraine

**\*Correspondence:**

Juana M. Rosas
jmrosas@uma.es

**Specialty section:**

This article was submitted to
Carbon-Based Materials,
a section of the journal
Frontiers in Materials

**Received:** 05 March 2020

**Accepted:** 29 April 2020

**Published:** 21 May 2020

**Citation:**

Moulefera I, García-Mateos FJ, Benyoucef A, Rosas JM, Rodríguez-Mirasol J and Cordero T (2020) Effect of Co-solution of Carbon Precursor and Activating Agent on the Textural Properties of Highly Porous Activated Carbon Obtained by Chemical Activation of Lignin With $H_3PO_4$. Front. Mater. 7:153. doi: 10.3389/fmats.2020.00153

This work analyzes the effect of co-solution of carbon precursor and activating agent on the textural and surface chemistry properties of highly nanoporous activated carbons obtained by chemical activation of Alcell lignin with phosphoric acid. The success of this methodology highlights the possibility of directly using the liquors produced in organosolv process (Alcell) to prepare activated carbons by chemical activation with phosphoric acid. Co-solutions of lignin and phosphoric acid were submitted to a two steps thermal treatment, which consisted of a first oxidative treatment in air at 200°C, followed by a thermal treatment in $N_2$ at 400°C, where activation of the oxidized lignin with phosphoric acid took place. A lignin-derived activated carbon with very high apparent surface area (2550 $m^2/g$) and pore volume (1.30 $cm^3/g$) was obtained with an initial phosphoric acid to lignin mass ratio of 2. Up to now, this is one of the highest values of apparent surface area reported not only for activated carbons prepared from lignin, but even for porous carbons prepared by chemical activation of other lignocellulosic materials with phosphoric acid. The use of lignin and phosphoric acid co-solution provided larger and more homogeneous effective interactions between the carbon precursor and the activating agent, by the formation of phosphate esters in the lignin matrix, which seems to be a key factor in the subsequent treatments: promoting crosslinking reactions in the carbonaceous matrix during the oxidative treatment in air at 200°C and enhancing the development of a wide porosity during the following activation thermal treatment.

**Keywords:** Alcell lignin, chemical activation, phosphoric acid, co-solution, activated carbon, phosphate esters, high surface areas

<footnote n="1">¹ Departamento de Ingeniería Química, Facultad de Ciencias, Universidad de Málaga, Andalucía TECH, Málaga, Spain,</footnote>

<footnote n="2">² Laboratory of Technique and Sciences of Water, Mustapha Stambouli University of Mascara, Mascara, Algeria</footnote>
# INTRODUCTION

Within a biorefinery scheme, the conversion of the major components of lignocellulosic biomass (cellulose, hemicellulose, and lignin) into energy and chemicals is essential. However, lignin has been used, up to date, mainly, as a fuel to recover energy and chemical reactants in the pulp and paper industry and only a 2% of the produced lignin is currently commercialized (Higson and Smith, 2011; Smolarski, 2012; Laurichesse and Avérous, 2014; Rosas et al., 2014). There are different pretreatments to fractionate lignocellulosic biomass into their main components, but Organosolv process becomes an ideal candidate in the biorefinery context, given the need of using environmentally friendly treatments. During Organosolv process, lignin is fractionated into smaller parts, dissolved from the lignocellulosic biomass with organic compounds and separated in the form of a liquor rich in phenolic compounds (Xu et al., 2006; Ruiz et al., 2011; Espinoza-Acosta et al., 2014). The next step consists of lignin precipitation from the liquors by partially evaporating the organic solvent, lowering the temperature or a combination of those, followed by a separation step through centrifugation, filtration, etc. The lignin derived from this method presents lower molecular weight and higher purity than others obtained by other processes, such as Kraft one. All these properties make this type of lignin an ideal precursor for the preparation of activated carbon, an added-value material, whose production process would imply both economic and environmental advantages.

Chemical activation with phosphoric acid is a well-known method to prepare activated carbons from different biomass precursors (Jagtoyen and Derbyshire, 1998; Kennedy et al., 2004; Montané et al., 2005; Sych et al., 2012). Many papers have been reported dealing with the optimization of the preparation conditions in order to obtain activated carbons with a high development of the porous texture. The influence of the impregnation ratio and the activation temperature on the textural properties of activated carbons prepared from different biomassic waste has been deeply analyzed (Jagtoyen et al., 1993; Hayashi et al., 2000; Molina-Sabio and Rodríguez-Reinoso, 2004; Rosas et al., 2009; Puziy et al., 2020). However, the effect of the gas atmosphere during the activation step on these properties is not so clear and some contradictions have appeared in this respect in the literature. Generally, activation takes place under inert atmosphere, although some authors have proposed the preparation of activated carbons by chemical activation with phosphoric acid under air atmosphere in order to enhance the textural properties (increase apparent surface area, $A_{BET}$) (Laine et al., 1989). However, (Benaddi et al., 1998) reported no significant influence of the atmosphere on the textural characteristics of activated carbons prepared by chemical activation of wood with phosphoric acid. They obtained values of apparent surface area of approximately 2100 m²/g, independently of the use of air or nitrogen as activation atmosphere. In contrast, Moreno Castilla et al. reported the preparation of activated carbons from olive-mill waste with similar $A_{BET}$ (~ 400 m²/g), using either nitrogen, CO₂ or H₂O atmospheres (Moreno-Castilla et al., 2001), but the $A_{BET}$ value decreased (~ 180 m²/g) when they used air as gas atmosphere.

With regard to lignin, activated carbons with very high $A_{BET}$ values (around 3000 m²/g) by chemical activation of Kraft lignin with KOH has been reported (Fierro et al., 2007a), but these carbon materials presented a very narrow microporosity. Chemical activation of Kraft lignin with zinc chloride and phosphoric acid gave rise to activated carbons with a much wider porosity, depending on the activation temperature and impregnation ratio, but with $A_{BET}$ values lower than 1800 m²/g (Gonzalez-Serrano et al., 1997, 2004; Guillén et al., 2009). However, no much information can be found in the literature regarding the preparation of activated carbons from Alcell lignin (Bedia et al., 2009; Rosas et al., 2012), a co-product that appears dissolved in the liquors produced in one of the most important Organosolv process.

Thus, the objective of this work was to analyze the effect of co-solution of carbon precursor and activating agent on the textural and surface chemistry properties of activated carbons obtained by chemical activation of Alcell lignin with phosphoric acid, in order to assess the possibility of directly using the liquors produced in the Organosolv process (Alcell), which would avoid the lignin precipitation and separation steps in this process. The influence of the use of a low temperature oxidative thermal treatment, previous to the activation process, on the surface chemistry and textural properties of the final lignin-derived carbons obtained was also studied.

# EXPERIMENTAL SECTION

## Activated Carbons Preparation

Alcell lignin (supplied by Repap Technologies Inc., Co.) was used as raw material for the activated carbons preparation. Different procedures were used for the preparation of activated carbons from lignin. Alcell lignin was dissolved in ethanol to simulate the composition of dissolved lignin in the Organosolv liquors. Solid Alcell lignin was directly impregnated with aqueous $H_3PO_4$ 85% (w/w) at room temperature. Co-solution of lignin and the activating agent, $H_3PO_4$ 85% (w/w), was also prepared at room temperature and the ethanol/lignin/$H_3PO_4$ solution was stirred until evaporation of the ethanol (which could be recovered). Both the impregnated lignin and the solid derived from the co-solution of lignin and $H_3PO_4$ were dried for 24 h at 60°C in a vacuum dryer. The amounts of lignin and $H_3PO_4$ used in both cases were those to obtain a final $H_3PO_4$/lignin mass ratio of 2. The as-received lignin and that dissolved in ethanol were denoted by L and LE, respectively. $H_3PO_4$ impregnated lignin and the solid derived from the lignin/ethanol/$H_3PO_4$ solution once dried were denoted as L2P and LE2P, respectively. These two samples were subsequently submitted to a low temperature oxidative treatment at 200°C in air atmosphere for 1.5 h, followed by an activation treatment under inert atmosphere at 400°C, both inside a laboratory horizontal tubular furnace. The heating rate to reach the oxidative and the activation temperature was 3 and 10°C/min, respectively. Once the activation temperature was reached the samples were cooled inside the furnace under inert atmosphere and then washed with distilled water at 60°C until neutral pH. The resulting activated carbons were dried at 60°C.
L2P and LE2P submitted to the oxidative treatment in air at 200°C for 1.5 h were denoted by adding the letter S to the corresponding names (L2PS and LE2PS, respectively). The final activated carbons obtained at 400°C in inert atmosphere were denoted by adding 400 to the respective notations (L2PS400 and LE2PS400, respectively). In case of oxidative treatments at different temperatures, the activated carbons notation will include at the end of the name the temperature of the treatment. For the sake of comparison, some samples were directly activated at 400°C under air atmosphere without being subjected to the previous oxidative low temperature treatment. In these cases, the letter S was not included in their corresponding notations (L2P400 and LE2P400) and the letter A was added at the end of its corresponding name. In order to also evaluate the influence of the gas atmosphere during the low temperature treatment, a sample was treated under inert atmosphere at 200°C for 1.5 h and was denoted by adding the letter N to its corresponding notation.

## Activated Carbons Characterization

The porosity of the samples was characterized by N₂ adsorption–desorption and CO₂ adsorption at –196 and 0°C, respectively, using a micromeritics instrument (ASAP 2020 model). The samples were previously outgassed for at least 8 h at 150°C. From the N₂ isotherm, the apparent surface area (A_BET) was determined by applying the BET equation (Brunauer et al., 1938); micropore volume (V_t) and external surface area (A_t) were calculated using the *t*-method (Lippens and de Boer, 1965); mesopore volume (V_mes) was determined as the difference between the maximum adsorbed volume of N₂ (V_p) at a relative pressure of 0.99 and the micropore volume (V_t) (Rodríguez-Reinoso et al., 1995). From the CO₂ adsorption data, the narrow micropore volume (VDR) and surface area (A_Dr) were calculated using the Dubinin–Radushkevich equation (Dubinin et al., 1947). Pore size distributions were obtaining by applying the 2D-NLDFT heterogeneous surface model from N₂ adsorption isotherm (Jagiello and Olivier, 2013).

Thermal treatments of different samples were carried out in a gravimetric thermobalance system, CI electronics. The thermobalance automatically measures the weight of the sample and the temperature as a function of time. Experiments were carried out in inert atmosphere (N₂) and in air atmosphere, for a total flow rate of 150 cm³ (STP)/min, employing sample mass of approximately 10 mg. The sample temperature was increased at a heating rate of 10°C/min. Differential scanning calorimetry (DSC) experiments were also obtained by a thermal analyzer (Mettler Toledo equipment) coupled to a Mass Spectrometer (Pfeiffer Vacuum model ThermoStar TM GSD 320).

The FTIR spectra of lignin and lignin-derived samples were recorded in a Bruker Optics Vertex 70 FTIR spectrometer, in the 500–4000 cm⁻¹ range, in KBr disks [ca. 1% (w/w)]. A KBr beam splitter and a Golden gate single reflection diamond ATR system detector were used. The spectra were collected for 2 min, with 4 cm⁻¹ resolution.

Solid-state ¹³C and ³¹P-NMR spectra of lignin and lignin-derived samples were recorded, as well, with a NMR spectrometer (400 WB Plus model from Bruker) using the Cross Polarization

Magic Angle Spinning (CPMAS) and high power decoupling (HPDEC) techniques and a 3.2 mm MAS triple-channel probe at a spinning frequency of 15 kHz, with SPINAL-64 ¹H decoupling conditions.

X-ray photoelectron spectroscopy (XPS) analyses of the samples were carried out in a PHI 5700C model Physical Electronics apparatus, with MgKα radiation (1253.6 eV). For the analysis of the XPS peaks, the C1s peak position was set at 284.5 eV and used as reference to position the other peaks. For the deconvolution of the peaks, Gaussian–Lorentzian curves were used. The deconvolution of the phosphorus spectrum was carried out by using two doublet peaks with an area ratio of 0.5 and a separation between peaks of 0.84 eV for each phosphorus groups. TPD experiments were carried out in a customized quartz fixed-bed reactor placed inside an electrical furnace. CO and CO₂, as output gases, were measured by a non-dispersive infrared (NDIR) gas analyzer, Siemens ULTRAMAT 22 (Siemens AG, Munich, Germany). Eighty mg of dried carbon sample was heated from room temperature to 930°C at a heating rate of 10°C/min under N₂ flow (200 cm³/min).

## RESULTS AND DISCUSSION

### Surface Chemistry Characterization of As-Received and Ethanol-Dissolved Lignin

Alcell lignin is a type of lignin derived from Organosolv process, a pulping technique that uses an organic solvent to solubilize lignin and hemicellulose. This type of lignin contains very small amounts of inorganic compounds (Kubo and Kadla, 2004). Delignification is mostly the result of high-temperature uncatalyzed hydrolysis of lignin ether linkages by water. The solvent in the system mainly serves to dissolve the lignin fragments (Girard and van Heiningen, 2000). One of the most used solvents in the Organosolv process is ethanol, which has been the solvent selected in this work to dissolve lignin.

**Figure 1A** shows the normalized CPMAS ¹³C NMR spectra of the as-received Alcell lignin and that dissolved in ethanol. Some differences can be observed, which suggest modifications of the lignin structure by its solution in ethanol. The main difference of both spectra is observed in the range of 15–27 ppm, which can be associated to terminal CH₃ groups. The increase of this signal suggests the accumulation of saturated alkyl structures after dissolving the lignin with ethanol. The predominant signal around 62–68 ppm, which remains practically the same after dissolving the lignin in ethanol, can be assigned to C₅ in xylose internal unit and C_γ, in β-O-4-linked unit; C₄ in xylose non-reducing end and C_γ, in β-O-4-linked unit with C = O. However, the signal around 144–148 ppm slightly increases after solution with ethanol, which can be associated to an increase of C₃ and C₄ in etherified guaiacyl; C₄ and C₄' in etherified 5-5' units and C_α, in cinnamic acid and ester (Almendros et al., 1992), and maybe related to the interaction of -OH groups of ethanol with those of lignin, which would also explain the increase of the signal between 15 and 27 ppm.
<figure><img src="image_59.png" /><figcaption>**FIGURE 1** | Normalized CPMAS ¹³C NMR spectra (A) and FTIR spectra (B) of as-received Alcell lignin and that dissolved in ethanol.</figcaption></figure>

**Figure 1B** represents the corresponding FTIR spectra of both samples. The spectrum of lignin is complicated and most of the bands observed are produced by contribution of various types of vibrations of different functional groups. Therefore, assignment of the bands is possible only in the approximation of the predominant contribution of certain atomic groups (Fițigău et al., 2013). Alcell lignin shows a broad band at 3410–3460 cm⁻¹, attributed to the hydroxyl groups in phenolic and aliphatic structures, which remains the same after dissolving with ethanol. The bands centered around 2935 and 2840 cm⁻¹ can be predominantly associated to CH stretching in aromatic methoxyl groups and in methyl and methylene groups of side chains. In the LE sample also appears a band around 2965 cm⁻¹ related to stretching vibration of aliphatic C-H. This appearance is in concordance to the accumulation of saturated alkyl structures, suggested by CPMAS ¹³C NMR spectra, after dissolving lignin with ethanol. In the carbonyl/carboxyl region, a band around 1700 cm⁻¹ is found, originating from unconjugated carbonyl/carboxyl stretching, which is more intense in the case of the sample LE. Typical bands of lignin are those at 1600, 1515, and 1425 cm⁻¹ related to aromatic skeleton vibrations and at 1460 cm⁻¹ associated to the C-H deformation combined with aromatic ring vibration (Vallejos et al., 2011; Fițigău et al., 2013), which are practically the same in both samples.

The samples show the characteristic vibrations for the guaiacyl unit (G ring and C = O stretch at around 1260 cm⁻¹; CH in-plane deformation at 1150 cm⁻¹; and C-H out-of plane vibrations in position 2, 5, and 6 of guaiacyl units at 830 and 915 cm⁻¹). Syringyl (S) and guaiacyl (G) units are detected by aromatic skeleton vibrations at 1600 and 1325 cm⁻¹ (S), 1515 and 1260 cm⁻¹ (G), and aromatic in plane C-H vibrations at 1110 (S) and 1030 cm⁻¹ (G). This last band is more intense in case of the LE sample (Vallejos et al., 2011; Fițigău et al., 2013). Another band at 880 cm⁻¹, related to C-C asymmetrical stretching mode, is more accused in sample LE, which also evidences the modification of the chemical structure of lignin after dissolving it with ethanol.

Finally, both samples also contain a weak band at 1370–1375 cm⁻¹ originating from phenolic OH and aliphatic C-H in methyl groups and a strong vibration at 1215–1220 cm⁻¹ that can be associated with C–C, C–O, and/or C = O stretching. The aromatic C–H deformation at 1030 cm⁻¹ appears as a complex vibration associated with the C–O, C–C stretching and C–OH bending in polysaccharides (Vallejos et al., 2011; Fițigău et al., 2013).

## Effect of H₃PO₄ and Low Temperature Oxidative Treatment on Lignin Surface Chemistry

**Figure 2** shows the corresponding FTIR spectra of the H₃PO₄ impregnated lignin and the solid derived from the lignin/ethanol/H₃PO₄ solution, L2P and LE2P, respectively. Most of the absorption peaks corresponding to functional groups for these two samples were diminished compared to those corresponding to the samples L and LE. Spectra of samples L2P and LE2P are very similar, showing an important decrease of the intensity in the bands associated to the presence of O-H stretching of hydroxyl functional groups (3500 to 3300 cm⁻¹). This reduction is due to the presence of phosphoric acid, which acts as dehydrating agent immediately after impregnation, even at room temperature. The band around 2965 cm⁻¹ related to stretching vibration of aliphatic C-H is also substantially reduced due to the increase of aromaticity (conversion of aliphatic to aromatic compounds). This loss of aliphatic character takes place at low temperatures and is caused by the cleavage of aliphatic side chains. Jagtoyen and Derbyshire also reported that the primary effects of acid attack are to cleave aryl ether bonds in lignin, accompanied by dehydration, degradation and condensation reactions (Jagtoyen and Derbyshire, 1998), which are the responsible of the decrease observed at 1600, 1515, and 1425 cm⁻¹, related to aromatic skeleton vibrations of phenyl skeletons and, at 1460 cm⁻¹, associated to the C-H deformation, combined with aromatic ring vibration.
<figure><img src="image_60.png" /><figcaption>**FIGURE 2** | FTIR spectra of the different $H_3PO_4$-lignin derived samples.</figcaption></figure>

<figure><img src="image_61.png" /><figcaption>**FIGURE 3** | XPS P2p spectra of the different $H_3PO_4$-lignin derived samples and deconvolution of LE2PS sample (C-O-P groups: continuous line, $P_2O_5$: dotted line and C-PO$_3$: dashed line).</figcaption></figure>

Furthermore, the acid-catalyzed hydrolysis of ether linkages in lignin can lead to the formation of ketones, as can be observed in the characteristic band that appears at $1685 \text{ cm}^{-1}$.

The typical bands for guaiacyl unit are also significantly reduced, those at $1215-1220 \text{ cm}^{-1}$ and at $1030 \text{ cm}^{-1}$, associated to C–C, C–O, and/or C = O stretching and C–O, C–C stretching and C–OH bending in polysaccharides, respectively. However, the appearance of other bands probably related to phosphorus compounds is also observed. In this sense, the broad band at $965 \text{ cm}^{-1}$ and the intense peak at $635 \text{ cm}^{-1}$ could be initially associated to P-O-P and P-C bending of aromatic compounds, respectively, although the peak at $635 \text{ cm}^{-1}$ can be also related to C = CH bonds (Labruquère et al., 1998). In addition, disappearance of this band, at $635 \text{ cm}^{-1}$, is observed for samples submitted to the oxidative treatment in air at $200^{\circ}\text{C}$ (L2PS and LE2PS), which can be related to the formation of more oxidized species of phosphorus or to decomposition of non-stable species. Furthermore, a decrease of the intensity of the bands at 1600, 1515, 1460, and $1425 \text{ cm}^{-1}$ was also detected after this low temperature treatment in air.

In order to analyze if phosphoric acid can act as an oxygen donor to the lignin matrix at this low temperature, sample L2P was treated at $200^{\circ}\text{C}$ for 1.5 h, but under inert ($N_2$) atmosphere (L2PS-N). The FTIR results for L2PS-N are also shown in **Figure 2**. As it can be seen, all the bands already mentioned were also reduced for L2PS-N, which suggests that the presence of oxygen coming from the $H_3PO_4$ is more than enough to produce crosslinking reactions (García-Mateos et al., 2018). In this sense, phosphoric species could be combined with organic species to form phosphate linkages that can serve to connect and crosslink lignin fragments (Jagtoyen and Derbyshire, 1998). These reactions are probably responsible for the reduction of these typical bands of lignin.

The results obtained from XPS analysis confirmed the presence of phosphorus groups on the surface of L2P, LE2P, L2PS,

and LE2PS samples (see **Figure 3**). The XPS P2p spectrum for L2P showed a broad band with a main peak at a binding energy around 134.3 eV, characteristic of pentavalent tetracoordinated phosphorus, as in phosphates and/or polyphosphates (Wu and Radovic, 2006). The obtained P2p broad band indicated the presence of different surface phosphorus species. Lower binding energies (around 133 eV) are characteristic of C-P bonding, as in C-PO$_3$ surface groups, which seem to be practically negligible in these samples (see **Table 1**); whereas higher values (134.2 eV) are ascribed to more oxidized phosphorus species, with C-O-P type bonds, as those in C-O-PO$_3$ surface groups (Wu and Radovic, 2006; Puziy et al., 2008). The contribution of these species was higher than 85%, suggesting that all phosphorus was mainly present in form of C-O-PO$_3$ surface groups. However, due to these samples were not yet activated, the presence of even more oxidized phosphorus species as $P_2O_5$ (135.7 eV) cannot be disregarded. It seems clear that co-solution of lignin with phosphoric acid generated a higher contribution of oxidized P surface groups to the total P surface groups than the traditional impregnation process (LE2P vs. L2P). On the other hand, the lower P content value observed for LE2P, when compared with that for L2P (see **Table 1**), suggests that this sample presents more P in the inner part of the lignin matrix, due probably to a better dispersion and to a more homogeneous distribution of P through the entire lignin matrix obtained by co-solution of the activating agent, phosphoric acid, and the carbon precursor, lignin. After the low temperature oxidative treatment both samples, L2PS and LE2PS, presented a similar surface P content value. However, the amount of surface O was significantly higher for the sample prepared by co-solution of lignin and $H_3PO_4$, LE2PS, which suggests that this preparation procedure enhances the lignin oxidation process at low temperature.

Thermal decomposition of the different lignin-derived samples was studied by TG analyses. **Figure 4A** represents the weight loss of the different lignin-derived samples as a
<table><caption>**TABLE 1** | Mass surface concentration of the different H₃PO₄-lignin derived samples and the corresponding contribution of different phosphorus surface group obtained by XPS analyses.</caption><thead><tr><th rowspan="2">Sample</th><th colspan="3">Surface concentration by XPS</th><th colspan="3">Phosphorus surface groups</th></tr><tr><th>C (%)</th><th>O (%)</th><th>P (%)</th><th>P<sub>2</sub>O<sub>5</sub> (%)</th><th>C-O-P (%)</th><th>C-PO<sub>3</sub> (%)</th></tr></thead><tbody><tr><td>L2P</td><td>30.7</td><td>45.3</td><td>24</td><td>3.9</td><td>89.05</td><td>7.05</td></tr><tr><td>LE2P</td><td>40.1</td><td>39</td><td>20.9</td><td>9.9</td><td>89.6</td><td>0.5</td></tr><tr><td>L2PS</td><td>48.8</td><td>31.3</td><td>19.9</td><td>5.85</td><td>93.7</td><td>0.45</td></tr><tr><td>LE2PS</td><td>42.9</td><td>36.5</td><td>20.6</td><td>9.24</td><td>89.59</td><td>1.17</td></tr></tbody></table>

function of temperature, under inert atmosphere and at a heating rate of 10°C/min. In general, three different ranges can be observed in all the samples. First, there is a small weight loss at temperatures lower than 200°C that can be partially attributed to release of water. Water can come from the phosphoric acid solution and produced by the reaction between phosphoric acid and lignin. As abovementioned, phosphoric acid cleaves aryl ether bonds in lignin, producing many transformations that include dehydration, degradation, and condensation (Jagtoyen and Derbyshire, 1998). At temperatures lower than 150°C, the acid catalyzes the hydrolysis of ether linkages, leading to the formation of ketones. These and other bond cleavage reactions that proceed through ionic mechanism promotes the release of H₂O, CO, CO₂, and CH₄ at these low temperatures, and can contribute to the corresponding weight loss observed in **Figure 4A**.

Between 200 and 500°C a further decrease of mass is observed associated to the release of light compounds from lignin degradation. In this temperature range, crosslinking reactions dominate over bond cleavage and de-polimerization reactions. At 250°C the structures are considered to be small polyaromatic units connected mainly by phosphate and polyphosphate bridges, including polyethylene linkages. As the temperature increases, cyclization and condensation reactions lead to increases in

aromaticity and size of the poliaromatic units, enabled by the scission of P-O-C bonds (Jagtoyen and Derbyshire, 1998).

A change in the slope of the weight-loss curve can be also observed at temperatures higher than 500°C, probably associated to volatilization of carbon-oxygen complexes generated by the activation process, and to a lesser extent, to decomposition of phosphorous-compounds (as P₂O₅), produced by the reaction of phosphoric acid with the carbon matrix (Montané et al., 2005).

The profiles of L2P and LE2P are very similar, suggesting the low impact of using dissolved lignin in ethanol in the thermal decomposition reactions. However, the weight loss curves of the samples treated with air (L2PS and LE2PS) are above those of L2P and LE2P samples. If the difference between the weight loss for a specific temperature of the non-treated and treated samples is represented as a function of temperature (Figure 4B), an important weight loss deviation around 700°C can be noticed for both curves, suggesting that the treatment with air generates significant differences to the samples treated with phosphoric acid. Montané et al. (2005) proposed that the peak observed at 650°C in the DTG curves was attributed to the volatilization of P₂O₅. On the other hand, Martínez De Yuso et al. (2014) found a band around 800°C, which they associated to the more important crosslinking reactions to produce aromatic units. Therefore, the treatment with air seems to stabilize or increase the number of phosphate and polyphosphate bridges, and/or to increase other crosslinking reactions. With this regard, although the weight loss percentage from 800°C is the same for the samples L2PS and LE2PS (around 10%), the remaining weight is not equal, due to a higher weight loss between 750 and 800°C. The weight loss observed in this range temperature coincides with the higher CO evolution observed in the TPD analysis for the LE2PS400 sample, and it is associated to the higher presence of C-O-PO₃ surface groups (Bedia et al., 2009).

Differential scanning calorimetry (DSC) scans of all the lignin-derived samples are shown in **Figure 5**. **Figure 5A** shows a broad endothermic peak near 80°C from absorbed moisture for all the impregnated samples. The impregnated lignin L2P also presents

<figure><img src="image_62.png" /></figure>
<figure><img src="image_63.png" /></figure>

<figure><img src="image_64.png" /><figcaption>FIGURE 5 | (A) DSC curves during pyrolysis of different lignin-derived samples and (B) DSC curves during treatment of L2P and LE2P with air at 200°C.</figcaption></figure>

a broad endotherm peak around 235°C, which can be associated to the decomposition and depolymerization reactions of lignin. Similarly, Karacan and Soy (2013) related an endothermic peak around 214°C to decomposition and depolymerization reactions of cellulose. Another exothermic peak was observed around 300°C. This peak coincides with a maximum in the CO and CO₂ evolution and is due mainly to the crosslinking reactions taking place during the decomposition process (Jagtoyen and Derbyshire, 1998), which could be softened by the endothermic polymerization and condensation reactions that take also place at this temperature range (Zhu et al., 2016). The scan of LE2P is very similar to that of L2P, but a shift of the endothermic peak is observed to 225°C, which suggests an earlier beginning of the decomposition and depolymerization reactions of lignin. In the case of the L2PS and LE2PS samples, the treatment with air results in the reduction of the endothermic peak, probably because the endothermic reaction has already taken place during the treatment with air (at 200°C for 1.5 h). From this temperature, a broad but much less intense exothermic peak can be observed centered on 300°C.

A further experiment was carried out trying to simulate the oxidative treatment at 200°C. The corresponding DSC curves are shown in **Figure 5B**. As can be seen, a sharp exothermic peak appears once the temperature is reached. This exothermic peak could be related to that shown in **Figure 5A** at 300°C in samples L2P and LE2P, due to different exothermic crosslinking reactions. In this sense, Jagtoyen and Derbyshire (1998) proposed that crosslinking reactions begin to dominate over depolymerization reactions from 200°C. Therefore, if enough time is provided, the reactions could take place at lower temperatures. Another important aspect is the presence of two endothermic peaks only in the sample LE2P. These two peaks could be related to the splitting of the endothermic peak that appears in **Figure 5A** around 220°C, caused by the superposition of several depolymerization and condensation reactions with different kinetic rates. In this sense, (Montané et al., 2005) already observed with isothermal experiments the existence of several

reactions which appears overlapped during DTG analyses of impregnated samples.

## Characterization of Activated Carbons

Table 2 shows the yield values of each treatment step and the total yield values for the different prepared activated carbons. As it is well known, phosphoric acid limits the formation of tars during the carbonization treatment, increasing the yield of the remaining solid product (Jagtoyen and Derbyshire, 1998). In this case, the yield of the treatment with air at 200°C is, approximately, 90% for both samples derived from co-solution of lignin and phosphoric acid and impregnated with phosphoric acid. The yield of the activation step (mass of activated sample/mass of impregnated or treated with air sample) was around 75% for samples non-treated with air, and about 85% for samples treated with air. There was approximately a difference of 10%, which practically corresponded to the loss of light hydrocarbons from the lignin degradation that was previously observed during the treatment with air at 200°C. The washing yield (mass of washed sample/mass of activated sample) was around 25%, very similar in all cases, and a total yield of around 60% was obtained for all the activated carbons, with no significant differences observed when lignin was co-dissolved or not with the activated agent and/or when the oxidative stabilization step was or was not performed. The total yield values (mass of activated carbon/mass of lignin) are considerably higher than others reported in the literature for lignin and other lignocellulosic materials (Montané et al., 2005; Puziy et al., 2005; Rosas et al., 2008).

**Figure 6** shows the N₂ adsorption-desorption isotherms at -196°C of the different activated carbons, prepared at 400°C and with an H₃PO₄ to lignin mass ratio of 2. The isotherm of the activated carbon L2P400 is a type I isotherm, characteristic of a typical microporous solid, which adsorbs almost all N₂ at low relative pressures. No significant changes in the N₂ isotherm were observed when lignin was co-dissolved with the activating agent. However, the treatment with air at 200°C produced solids with a remarkable increased N₂ adsorbed volume up to 0.4 of
<table><caption>**TABLE 2** | Yield values of the lignin activated carbons.</caption><thead><tr><th>Activated carbon</th><th>Oxidative treatment yield (%)</th><th>Activation yield (%)</th><th>Washing yield (%)</th><th>Total yield (%)</th></tr></thead><tbody><tr><td>L2P400</td><td>–</td><td>75.9</td><td>28.0</td><td>63.8</td></tr><tr><td>LE2P400</td><td>–</td><td>75.9</td><td>25.9</td><td>59.2</td></tr><tr><td>L2PS400</td><td>87.3</td><td>87.8</td><td>27.3</td><td>62.9</td></tr><tr><td>LE2PS400</td><td>90.7</td><td>81.6</td><td>25.2</td><td>63.3</td></tr></tbody></table>

<figure><img src="image_65.png" /><figcaption>**FIGURE 6** | N₂ adsorption-desorption isotherms obtained at –196°C of the different activated carbons obtained at 400°C and with H₃PO₄ to lignin mass ratio of 2.</figcaption></figure>

relative pressure, which indicated a larger presence of wider microporosity in these carbonaceous solids. This N₂ adsorption volume increase was significantly more pronounced when co-solution of lignin and phosphoric acid was carried out previously to the oxidative stabilization step (air at 200°C). The small hysteresis cycle at medium relative pressures shown by the N₂ adsorption isotherms indicates the presence (to a lesser extent) of narrow mesopores.

The textural properties obtained from the N₂ and CO₂ isotherms of the different activated carbons are summarized in **Table 3**. The differences between the micropore volumes obtained from N₂ (Vt) and CO₂ (VDR) isotherms, with higher values for those from N₂ ones, confirmed the presence of a wide microporosity for the lignin-derived activated carbons. However, the mesopore volumes are quite low in all the samples.

The apparent surface area (A<sub>BET</sub>) values increased considerably from 912 m²/g for L2P400 to 1937 m²/g for L2PS400, and to 2551 m²/g for LE2PS400. This last value is much higher than others reported in the literature for lignin-derived activated carbons using other activation methods (Rodríguez-Mirasol et al., 1993a,b; Gonzalez-Serrano et al., 1997; Hayashi et al., 2000; Cotoruelo et al., 2007; Babel and Jurewicz, 2008; Carrott et al., 2008; Sun et al., 2010; Al-Lagtah et al., 2016; Kim et al., 2017; Lin et al., 2019; Rowlandson et al., 2020), including the same activating agent (H₃PO₄), at similar temperatures and impregnation ratios (Gonzalez-Serrano et al., 2004; Guo and Rockstraw, 2006; Fierro et al., 2007b; Bedia et al., 2010; Rosas et al., 2012). Only a few works reported the preparation of activated carbons from lignin with higher apparent surface areas (~3000 m²/g) by using KOH activation, but at much higher activation temperatures than those used in this work (Zou and Han, 2001; Fierro et al., 2007a; Li et al., 2014; Fernandez-Ruiz et al., 2018; Han et al., 2019). In the method here proposed, the use of lower activation temperatures would substantially decrease the cost of the obtained activated carbons, without compromising the high development of porosity.

**Figure 7** collects the micropore size distribution of the different activated carbons obtained at 400°C and H₃PO₄ to lignin mass ratio of 2. The activated carbon prepared by direct impregnation of lignin shows a bimodal micropore size distribution, with a main band with a maximum at around 6.7 Å and a second one (shoulder), less intense, with a maximum at around 16 Å, ending at a pore size value of about 40 Å, indicating the presence of narrow mesopores. The micropore size distribution of the activated carbon when lignin was co-dissolved with the activating agent is practically the same as that of the activated carbon prepared by direct impregnation of lignin. However, the treatment with air at 200°C produces a clear change in the bimodal micropore size distribution. This treatment produced an increase of both bands, with the maximum of the first one slightly shifting to higher micropore sizes (from 6.7 to approximately 8.5 Å). However, the increase of the band corresponding to micropores of about 16 Å was proportionally higher than that corresponding to the lower micropore sizes (6.7 to 8.5 Å), even surpassing it in the case of LE2PS400. These results suggest that the treatment with air at 200°C increased the volume of micropores, especially that of the largest pore size (16 Å), and this effect was clearly enhanced by the co-solution of lignin and phosphoric acid. However, this treatment does not appear to greatly affect the micropore (or mesopore) size of the obtained activated carbons.

<table><caption>**TABLE 3** | Textural properties of the different lignin derived activated carbons obtained from the N₂ and CO₂ isotherms.</caption><thead><tr><th rowspan="2"></th><th colspan="5">N₂ Isotherm</th><th colspan="2">CO₂ Isotherm</th></tr><tr><th>A<sub>BET</sub> (m²/g)</th><th>A<sub>t</sub> (m²/g)</th><th>V<sub>t</sub> (cm³/g)</th><th>V<sub>p</sub> (cm³/g)</th><th>V<sub>mes</sub> (cm³/g)</th><th>A<sub>DR</sub> (m²/g)</th><th>V<sub>DR</sub> (cm³/g)</th></tr></thead><tbody><tr><td>L2P400</td><td>912</td><td>55</td><td>0.412</td><td>0.477</td><td>0.063</td><td>438</td><td>0.176</td></tr><tr><td>LE2P400</td><td>845</td><td>36</td><td>0.395</td><td>0.444</td><td>0.054</td><td>407</td><td>0.163</td></tr><tr><td>L2PS400</td><td>1937</td><td>131</td><td>0.892</td><td>1.044</td><td>0.146</td><td>673</td><td>0.270</td></tr><tr><td>LE2PS400</td><td>2551</td><td>143</td><td>1.158</td><td>1.313</td><td>0.152</td><td>859</td><td>0.344</td></tr></tbody></table>
<figure><img src="image_66.png" /><figcaption>**FIGURE 7** | Micropore size distribution of the different activated carbons obtained at 400°C and with $H_3PO_4$ to lignin mass ratio of 2.</figcaption></figure>

The present results indicate that the use of co-solution of lignin with phosphoric acid in combination with an oxidative treatment at 200°C provides activated carbons with extremely high specific surface area and a large and wide microporosity. These significant differences in the porous texture seem to be the result of larger and more homogeneous interactions between lignin and the activating agent when they are co-dissolved that may increase the formation of effective phosphate esters, which are the responsible of the cross-linking reactions in the carbonaceous matrix during the oxidative stabilization process (Puziy et al., 2005). To support this hypothesis, TPD (Figure 8) and XPS (not shown) analyses were carried out for the different activated carbons. As it was observed in previous works, activated carbons prepared by chemical activation with phosphoric acid presented most of CO evolution at temperatures higher than 700°C, due to the presence of oxygen-containing P surface groups, which include metaphosphates, mainly as $C-O-PO_3$ groups and $C-PO_3$ groups, which decompose as gaseous CO,

producing C-P type surface groups (Bedia et al., 2009; Rosas et al., 2009). In addition, XPS P2p spectra of the different activated carbons, which were very similar to those of the corresponding lignin-derived samples (Figure 3), also evidence the presence of, mainly, C-O-P ones ($CO-PO(OH)_2$ and $(C-O)_3PO$) at 134.2 eV and C-P-O groups (such as $C-PO(OH)_2$ or $C_2-PO(OH)$) at 133 eV. The similarity among the spectra indicated that the contribution of this kind of groups is very similar in the different obtained activated carbons, independently of the preparation method used. In this sense, a similar P surface concentration value of 1.6%wt was obtained for L2P400, L2PS400, and LE2P400. However, the combined effect of the co-solution of lignin and phosphoric acid and the oxidative treatment at low temperature increased this value to 2.4%wt.

TPD CO profiles for the activated carbon prepared from dissolved and not-dissolved lignin are very similar, showing most of their CO release between 700 and 900°C (Figure 8). However, the treatment with air at 200°C and the combined use of this treatment with co-solution of lignin and the activating agent provided a higher release of CO, even at temperatures lower than 700°C. This CO evolution at lower temperatures can be associated to decomposition of mainly phenol and ether surface groups and, to a lesser extent, to anhydride surface groups. As decomposition of this last group also should produce $CO_2$ evolution, the low $CO_2$ evolution for these activated carbons at this temperature range provided evidences of the low presence of anhydride surface groups on those samples. On the other hand, the sample LE2PS400 presented the highest CO evolution between 700 and 800°C, which suggests a higher presence of metaphosphates ($C-O-PO_3$ and/or $C-PO_3$ type) surface groups on this sample (Bedia et al., 2009). These P surface groups, with C-O-P bonds, decompose to gas CO and P surface groups of C-P bond type at these temperatures (Wu and Radovic, 2006). The larger evolution of CO observed in the TPD for LE2PS400 can be associated to a higher generation of phosphate esters during the preparation process, as it was evidenced by FTIR analyses (Figure 2) and by the higher surface P content (2.4%wt) of this activated carbon, revealed by XPS results.

<figure><img src="image_67.png" /><figcaption>**FIGURE 8** | (A) CO and (B) $CO_2$ evolution of the different activated carbons.</figcaption></figure>
## Effect of Temperature of the Oxidative Treatment

Although to our best knowledge the co-solution of the carbon precursor and the activating agent has not been previously reported, some authors have already proposed the activation in two different steps. Toles et al. (1998) analyzed different approaches of acid-activation/oxidation to produce activated carbons with high metals and organics compounds adsorption uptakes. They observed no significant differences on the porosity of samples directly activated in air at 450°C and compared with samples prepared with a further intermediate step at 170°C, obtaining a maximum apparent surface area of $A_{BET}$ = 1662 m²/g. They suggested that the intermediate step was not necessary because the time spent during the ramping stage (averaging approximately 0.8 h) to 450°C was sufficient to allow penetration of the activating agent (phosphoric acid) and any preliminary reactions to take place. Dastgheib and Rockstraw (2001) proposed also a two steps activation process, both of them in air, with the first one up to 215°C, temperature from which ortho-phosphoric acid dehydrates, losing half a molecule of water and forming pyrophosphoric acid, but they do not compare the porosity of activated carbons prepared in other gas atmospheres.

In order to deep in the role of this oxidative treatment, further oxidative treatments were carried out at different temperatures. Table 4 shows the yield values of the lignin derived activated carbons obtained at 400°C and $H_3PO_4$/lignin mass ratio of 2, but previously treated at different temperatures under air atmosphere. As can be seen, the yield of the treatment in air decreases as the temperatures increases, due to a higher release of light hydrocarbons and water. The activation and washing yields remain approximately the same regardless of the treatment previously carried out, since the impregnation ratio and the activation temperature is the same in all the samples. Therefore, the total yield values follow the same sequence found for the treatments with air.

**Figure 9** shows the $N_2$ adsorption-desorption isotherms at -196°C of the different activated carbons prepared at 400°C and $H_3PO_4$/lignin mass ratio of 2, treated at different oxidative temperatures. The isotherms of the series are all of them type I isotherms. The results indicate that the oxidative treatment at 100°C did not generate large porosity development compared to that of the sample without any oxidative treatment. The highest amount of $N_2$ adsorbed volume was observed for the sample treated at 200°C. An increase of this temperature to 300°C

slightly reduced the $N_2$ adsorbed volumes. Table 5 collects the textural properties derived from the $N_2$ and $CO_2$ isotherms of the different lignin derived activated carbons obtained at 400°C, a $H_3PO_4$/lignin mass ratio of 2 and treated at different oxidative temperatures. As can be seen, the textural properties of the sample treated at 100°C are very similar to those reported in Table 3 for the sample L2P400, which was not submitted to any oxidative treatment. It can also be observed that maximum pore volume (1.31 cm³/g) and apparent surface area (2551 m²/g) were obtained with an oxidative treatment at 200°C, while a relative decrease in the apparent surface area to 2128 m²/g was noticed when the oxidative treatment was performed at 300°C.

The treatment in air at low temperature seems to enhance the formation of phosphate and polyphosphate bridges, which are responsible for connecting and crosslinking the lignin fragments. These phosphorus groups are inserted in the carbon matrix, separating the organic species. These groups generate an expansive process that, after removal of the acid, leaves the matrix in an expanded state, with a high pore development structure. In this case, the highest porosity development is observed when the oxidative treatment is performed at intermediate temperatures (200°C). The treatment at 100°C does not produce any enhancement in the porosity, and at 300°C some reduction begins to be noticed. These differences can be explained, according to the findings of Jagtoyen and Derbyshire (1998) by the fact that bond cleavage reactions that proceed through ionic mechanisms are taking place at low temperature and cross-linking reaction dominates over bond cleavage and de-polymerization reactions from 150°C. Therefore, the results seem to indicate that the presence of oxygen promotes preferentially the cross-linking reactions around 200°C and from this temperature some over-oxidation could take place.

<figure><img src="image_68.png" /><figcaption>**FIGURE 9** | $N_2$ adsorption-desorption isotherms obtained at -196°C of the different activated carbons obtained at 400°C and $H_3PO_4$/lignin mass ratio of 2, treated at different temperatures.</figcaption></figure>

<table><caption>**TABLE 4** | Yield values of the lignin-derived activated carbons obtained at 400°C and $H_3PO_4$/lignin mass ratio of 2, treated at different temperatures.</caption><thead><tr><th>Sample</th><th>Oxidative treatment yield (%)</th><th>Activation yield (%)</th><th>Washing yield (%)</th><th>Total yield (%)</th></tr></thead><tbody><tr><td>LE2PS400-100°C</td><td>98.1</td><td>85.8</td><td>25.7</td><td>65.0</td></tr><tr><td>LE2PS400-200°C</td><td>90.7</td><td>81.6</td><td>22.2</td><td>63.3</td></tr><tr><td>LE2PS400-300°C</td><td>87.7</td><td>88.7</td><td>24.0</td><td>56.0</td></tr><tr><td>LE2P400-A</td><td>-</td><td>80.1</td><td>27.2</td><td>64.9</td></tr><tr><td>LE2PS(N)400</td><td>88.8</td><td>87.3</td><td>26.2</td><td>60.9</td></tr></tbody></table>
<table><caption>**TABLE 5** | Textural properties of the different lignin activated carbons obtained at 400°C and H₃PO₄/lignin mass ratio of 2, treated at different temperatures under air atmosphere, obtained from the N₂ and CO₂ isotherms.</caption><thead><tr><th rowspan="2">Sample</th><th colspan="5">N₂ isotherm</th><th colspan="2">CO₂ isotherm</th></tr><tr><th>A<sub>BET</sub> (m²/g)</th><th>A<sub>t</sub> (m²/g)</th><th>V<sub>t</sub>(cm³/g)</th><th>V<sub>p</sub>(cm³/g)</th><th>V<sub>mes</sub> (cm³/g)</th><th>A<sub>DR</sub>(m²/g)</th><th>V<sub>DR</sub>(cm³/g)</th></tr></thead><tbody><tr><td>LE2PS400-100°C</td><td>877</td><td>61</td><td>0.386</td><td>0.459</td><td>0.066</td><td>420</td><td>0.168</td></tr><tr><td>LE2PS400-200°C</td><td>2551</td><td>143</td><td>1.158</td><td>1.313</td><td>0.152</td><td>859</td><td>0.344</td></tr><tr><td>LE2PS400-300°C</td><td>2128</td><td>181</td><td>1.036</td><td>1.237</td><td>0.175</td><td>570</td><td>0.228</td></tr><tr><td>LE2P400-A</td><td>707</td><td>26</td><td>0.306</td><td>0.339</td><td>0.031</td><td>501</td><td>0.201</td></tr><tr><td>LE2PS(N)400</td><td>1263</td><td>111</td><td>0.575</td><td>0.706</td><td>0.127</td><td>483</td><td>0.194</td></tr></tbody></table>

In addition, other experiment was carried out but treating the lignin/H₃PO₄ samples at 200°C under inert (N₂) atmosphere (instead of air), followed by activation also in N₂ at 400°C (LE2PS(N)400). The amount of N₂ adsorbed volume over the entire range of relative pressure for the sample thus obtained was significantly higher than that observed for the LE2P400 sample, evidencing the presence of larger and wider microporosity. These results suggest that some reactions are being promoted at 200°C even under inert atmosphere, when sufficient time is provided. According to DSC results, it must be related to crosslinking reactions taking place during the decomposition process (Karacan and Soy, 2013) and/or to the polymerization and condensation reaction with H₃PO₄ (Zhu et al., 2016). Based on these results, it seems that, in agreement with the hypothesis of Toles et al. (1998) there are some kinetic limitations, as can be deduced from the increase of the microporosity with the treatment at 200°C under inert atmosphere, mainly associated to some crosslinking reactions. However, the use of an oxidative atmosphere becomes necessary to boost the porosity development, as pointed out by Dastgheib and Rockstraw (2001), but at adequate experimental conditions, as those previously observed.

In order to provide evidences of a possible over-oxidation of the activated carbon, which avoids the porosity development, a further activation of the LE2P sample at 400°C but directly in air was carried out. The isotherm of the corresponding activated carbon, LE2P400-A, (shown in **Figure 9**) presented lower N₂ adsorption volumes than the ones for the corresponding carbon materials obtained by activation under inert atmosphere. The same tendency was found for the sample L2P also activated in air at 400°C (not shown). In particular, a reduction in the apparent surface area of almost 200 m²/g (**Table 5**) was observed, suggesting that some over-oxidation has occurred, causing some pores to collapse (Wang et al., 2014). In this line, several authors prepared activated carbons under different gas atmospheres in order to analyze their influence on the porous texture development. However, some discrepancies can be found in these works. Laine et al. (1989) observed an increase of the surface area of activated carbons from coconut shell at 400°C, when forced air was used, varying from 180 (in N₂) to 860 (in air) m²/g. Molina-Sabio et al. (1995) suggested that the presence of oxygen at 450°C, at low impregnation ratios (R < 0.32), inhibits the aromatization process, thus reducing the influence of the phosphoric acid in both promoting the

development of porosity and reducing the extent of contraction produced during the heat treatment of peach stones. At higher impregnation ratios (R > 0.32), they proposed that the access of impregnation will be more restricted and will have smaller effect on the process. Benaddi et al. (1998) reported very similar micropores volumes and surface areas independently of the gas atmosphere (O₂ or N₂) used to prepare activated carbons from wood at 480°C. Puziy et al. (2003) associated a maximum of porosity in activated carbons prepared from a styrene-divinylbenzene copolymer in air, at 900°C, to a further air gasification. Wang et al. (2014) mixed ammonium persulfate with phosphoric acid. They proposed that the oxygen generated from decomposition of ammonium persulfate enhanced both the formation of phosphate esters and the formation of CO and CO₂ by gasification. At low oxygen dosages, the formation of phosphate esters predominated, hence promoting the crosslinking between phosphoric acid and lignin. They also pointed out that lignin presents higher carbon content than wood and can thus produce more phosphate esters. However, the excessive amount of persulfate leads to over-oxidation and carbon deletion. Furthermore, the over-oxidation also can cause some pores to collapse (Wang et al., 2014).

Based on our results and taking into account the statements of all these authors, it seems that in our experimental conditions, the presence of oxygen at intermediate temperatures, where cross-linking reactions dominate, promote the porosity development due to the creation of more phosphate esters, but at higher temperatures, where the aromatization process preferentially takes place, an excess of oxygen was counter-productive, generating an over-oxidation that would inhibit the increase of the porosity.

This new methodology can be an interesting alternative for the valorization of Alcell lignin into highly porous activated carbons. The high apparent surface area obtained and the presence of surface phosphorus functional groups, which provides to the activated carbons with surface acid character (Valero-Romero et al., 2017) and oxidation resistance (Rosas et al., 2012), open the possibilities for different applications (Puziy et al., 2020). Specifically, these activated carbons can be used in adsorption processes (gas and liquid phases), like adsorbent for basic compounds (Guo et al., 2005; László, 2005; Suhas et al., 2007), in heterogeneous catalytic processes, like catalyst or catalytic support (Bedia et al., 2009, 2010) and even in electrochemistry applications (Hulicova-Jurcakova et al., 2009).
## CONCLUSION

This work analyzes the effect of co-solution of carbon precursor and activating agent on the textural and surface chemistry properties of highly nanoporous activated carbons obtained by chemical activation of Alcell lignin with phosphoric acid. The success of this methodology highlights the possibility of directly using the liquors produced in organosolv process (Alcell) to prepare activated carbons by chemical activation with phosphoric acid. Co-solutions of lignin and phosphoric acid were submitted to a two steps thermal treatment, which consisted of a first oxidative stabilization in air at 200°C, followed by a thermal treatment in N₂ at 400°C, where activation of the oxidized lignin with phosphoric acid took place. A lignin-derived activated carbon with very high apparent surface area (2550 m²/g) and pore volume (1.30 cm³/g) was obtained with an initial phosphoric acid to lignin mass ratio of 2. Up to now, this is one of the highest values of apparent surface area reported not only for activated carbons prepared from lignin, but even for porous carbons prepared by chemical activation of other lignocellulosic materials with phosphoric acid. The use of lignin and phosphoric acid solutions provided more homogeneous and effective interactions between the carbon precursor and the activating agent, by the formation of phosphate esters in the lignin matrix, which seems to be a key factor in the subsequent treatments: promoting

crosslinking reactions in the carbonaceous matrix during the oxidative treatment in air at 200°C and enhancing the development of a wide porosity during the followed activation thermal treatment.

## DATA AVAILABILITY STATEMENT

All datasets generated for this study are included in the article/supplementary material.

## AUTHOR CONTRIBUTIONS

IM and FG-M prepared the activated carbons. IM, FG-M, and JR performed the characterization of the samples and prepared the manuscript. JR, AB, JR-M, and TC developed the synthesis concept and planned the experiments. All authors discussed the result and commented on the manuscript.

## FUNDING

We gratefully thank MICINN and FEDER (Project RTI2018-097555-B-I00) and Junta de Andalucía (Project UMA18-FEDERJA-110) for financial support.

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**Conflict of Interest:** 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.

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<figure><img src="image_69.png" /></figure>

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# Mechanically Robust Free-Standing Single-Walled Carbon Nanotube Thin Films With Uniform Mesh-Structure by Blade Coating

Yuichi Kato, Atsuko Sekiguchi*, Kazufumi Kobashi, Rajyashree Sundaram, Takeo Yamada and Kenji Hata

National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba, Japan

Carbon nanotubes (CNTs) have garnered tremendous attention as building blocks for self-supporting membranes owing to remarkable developments in the manufacturing technology of high-quality CNT films. CNT films are expected to be applied in a wide range of applications, such as ultrafiltration membranes and as switching or sensing elements in microelectromechanical systems. However, the main challenge has been in fabricating CNT films by versatile and scalable processes suitable for industrial production while retaining lab-scale high performance. In this work, we succeed in fabricating macroscale (10 cm²) free-standing CNT films with thicknesses as low as 200 nm showing tensile strengths of ~166 MPa by simple, versatile, and scalable blade-coating of CNT suspensions. Our study demonstrates that it is possible to control CNT film bundle size distribution and pore structure by controlling the CNT dispersibility and entanglement in the suspensions. We find that controlling bundle size distribution, pore structure uniformity, and packing can lead to five-fold, four-fold, and three-fold higher tensile strengths, fracture strain, and Young's modulus, respectively, compared to films with poorer uniformity and packing.

**Keywords:** carbon nanotube, free-standing film, tensile strength, fracture strain, bundle thickness, pore size

**OPEN ACCESS**

**Edited by:**
Ming Xu,
Huazhong University of Science
and Technology, China

**Reviewed by:**
Qingwen Li,
Suzhou Institute of Nano-Tech
and Nano-Bionics (CAS), China
Renyuan Zhang,
Tongji University, China

*Correspondence:
Atsuko Sekiguchi
atsuko-sekiguchi@aist.go.jp

**Specialty section:**
This article was submitted to
Carbon-Based Materials,
a section of the journal
Frontiers in Materials

**Received:** 15 May 2020
**Accepted:** 14 August 2020
**Published:** 08 October 2020

**Citation:**
Kato Y, Sekiguchi A, Kobashi K, Sundaram R, Yamada T and Hata K
(2020) Mechanically Robust Free-Standing Single-Walled Carbon Nanotube Thin Films With Uniform Mesh-Structure by Blade Coating.
Front. Mater. 7:562455.
doi: 10.3389/fmats.2020.562455

## INTRODUCTION

Thin and robust carbon nanotube (CNT)-based free-standing films are in high demand as switching or sensing elements in microelectromechanical systems (MEMS), ultrafiltration membranes, or liquid film supports for water purification, metal ion refinement, etc. For sensing/switching elements in MEMS, sufficiently thin substrate-free films are necessary to be sensitive and rapidly respond to external stimuli, such as fluid pressure, static electricity by elastic deformation, etc. Ultrafiltration membranes and liquid film supports (Parhi, 2013) require free-standing films that readily allow liquid passage while being mechanically strong to withstand liquid influx backpressure. Such robust and thin ultrafiltration membranes obviate the need for support films and save space in liquid separation modules, increasing filtration efficiency and throughput. Thinner films increase permeation flow velocity, which could also result in enhanced separation efficiencies of liquid film supports. Among the various materials developed for self-supporting films (such as polymer-based, organic-inorganic hybrids), CNTs offer the advantages of high
thermal and chemical stability and mechanical strength as well as easy formation of mesh-like structures with high surface areas (Narisawa, 1982; Hecht et al., 2011; Jiang et al., 2011; Kobashi et al., 2013; Janas and Koziol, 2014).

Two broad strategies have evolved to fabricate thin and robust CNT free-standing films: (i) dry process and (ii) wet process. Dry process involves either spinning from vertically aligned nanotube arrays or depositing directly from a floating catalyst chemical vapor deposition reactor (Jiang et al., 2011; Nasibulin et al., 2011). In general, spinning from vertically aligned nanotube arrays yields an oriented CNT film. On the other hand, depositing directly from a floating catalyst chemical vapor deposition reactor leads to randomly aligned CNT films. With appropriate CNT synthesis methods, nanotube structure and entanglement can be controlled, leading to ≤ 200-nm-thick free-standing films with tensile strengths of ~65–360 MPa (Zhang et al., 2005; Ma et al., 2007; Liu et al., 2011; Nasibulin et al., 2011; Zhou et al., 2012). Dry-processed freestanding thin films have even been mass produced and are available commercially. Wet processing, the second route to fabricate free-standing CNT films, involves dispersing and debundling CNTs in a liquid medium to form suspensions that are subsequently converted into films (Hu et al., 2010; Hecht et al., 2011; Kobashi et al., 2013; Janas and Koziol, 2014; Wang and Bao, 2015). To make the CNT suspension, ultrasonic cavitation, shearing, or turbulence force are applied to CNT powders suspended in a liquid medium to enable nanotube de-bundling. The CNT suspensions are converted into films by methods such as spin-coating, blade-coating, dip-coating, vacuum filtration, etc. The primary advantage of wet process is fabrication and structural control flexibility. CNTs (including commercial samples) with regulated structures (wall number, diameters, lengths, crystallinity, chirality) and purity can be chosen for wet-process film fabrication (Hu et al., 2010; Wang and Bao, 2015). In contrast, in dry process, CNT structural and purity control have to be exercised at the synthesis stage.

Fabricating free-standing CNT films with thicknesses ≤ 200 nm by wet process is a major challenge (Sreekumar et al., 2003; Gu and Swager, 2008; Jo et al., 2010; Liu et al., 2010; Shi et al., 2011; Janas and Koziol, 2014; Ibañez et al., 2016; Janas et al., 2017). Achieving thinness in combination with high strength by wet process with minimal CNT damage during the dispersion step have been widely investigated. Shi et al. show promising results on wet-processed CNT films (Shi et al., 2011). They obtain 20-nm-thick free-standing films with tensile strength and Young's modulus of 860 MPa and 36 GPa, respectively, from suspensions of ultralong CNTs. In their process, the CNT films are fabricated by CNT suspension filtration. Shi et al. (2011) claim their filtration process to be scalable as the process is used in commercial paper production. However, the filter material needs to have appropriate porous structure, and solvents for the suspension need to be designed to have good permeability for the filtering process. Therefore, there are still unresolved issues in devising versatile and industrially compatible CNT film fabrication methods by wet processing to enable utility in a wide range of applications.

Here, we report macroscale (10-cm-square) freestanding CNT films with thicknesses ~200 nm and tensile strengths

of ~166 MPa fabricated by wet process involving simple and scalable blade-coating of nanotube suspensions. In addition, we demonstrate the possibility of controlling the pore and bundle size in the CNT film mesh structure by regulating CNT dispersibility and entanglement in the suspensions. Further, we find this control in bundle and pore structure to be critical for achieving high CNT film tensile strengths, fracture strain, and Young's moduli.

## MATERIALS AND METHODS

Our film fabrication protocol by wet process involves suspending CNT powders in a liquid medium aided by ultrasonication and blade-coating the suspensions into films. For fabricating 200-nm-thin free-standing CNT film, we peeled off the CNT film by dissolving the Al sacrificial layer and scooped up the film by a photo flame.

The CNTs we used in this study were synthesized by water-assisted chemical vapor deposition (“Super Growth”). The CNTs were grown as a vertical array on a Si substrate using iron nanoparticles as catalyst and $C_2H_4$ as carbon source with water vapor as growth enhancer (Hata et al., 2004). The as-grown CNTs were removed from the substrate to obtain powders, which were used as starting materials for film preparation.

We prepared four types of CNT suspensions for film fabrication using four types of liquid media (Table 1 and Supplementary Table S1). In suspension 1, we used 10-dodecyl-7,8-dimethyl-10H-benzo[g]pteridine-2,4-dione (Flavin) (Ju et al., 2009; Kato et al., 2015), as the surfactant and toluene as the solvent to debundle the CNT powders. In suspension 2, we used sodium deoxycholate (SDOC) as the surfactant in aqueous solution. In suspension 3 and 4, N-methylpyrrolidone (NMP) and propylene glycol (PG) were used as the solvents, respectively (without any surfactant). The CNT and surfactant concentrations were optimized to make suspensions suitable for blade-coating. We made the CNT suspensions according to procedures reported elsewhere (Sakurai et al., 2018).

Details of the procedures we used for CNT film characterization—tensile testing, scanning electron microscopy (SEM), sheet resistance measurement, mercury porosimetry, gas adsorption, CNT film thickness measurement, Raman spectroscopy, thermogravimetric analysis (TGA), and CNT suspension characterization (by flow image analysis) are provided in the **Supplementary Material**.

## RESULTS AND DISCUSSION

### Manufacturing of Free-Standing CNT Films and Mechanical Properties

Our fabrication process is designed to be scalable, easy, and versatile, showing compatibility to various types of substrates and solvents. We fabricated our suspension blade-coated CNT films on 4-inch wafer scale (Figures 1A,B) and as 10-cm-square free-standing samples with <200 nm thickness (Figures 1C,D). As shown in **Figure 1D**, our free-standing films are translucent and
<table><caption>**TABLE 1** | Four types of CNT films prepared using four types of suspensions and their mechanical properties by tensile test – tensile strength, fracture strain, and Young's modulus.</caption><thead><tr><th rowspan="2">Sample name</th><th colspan="2">Suspension</th><th colspan="3">Mechanical properties</th></tr><tr><th>Solvent</th><th>Surfactant</th><th>Tensile strength [MPa]</th><th>Fracture strain [%]</th><th>Young's modulus [GPa]</th></tr></thead><tbody><tr><td>Film 1</td><td>Toluene</td><td>Flavin</td><td>166</td><td>2.4</td><td>7.6</td></tr><tr><td>Film 2</td><td>Water</td><td>SDOC</td><td>71</td><td>2.2</td><td>2.9</td></tr><tr><td>Film 3</td><td>NMP</td><td>No surfactant</td><td>55</td><td>1.5</td><td>8.4</td></tr><tr><td>Film 4</td><td>PG</td><td>No surfactant</td><td>33</td><td>0.6</td><td>2.4</td></tr></tbody></table>

<figure><img src="image_70.png" /></figure>

uniform in appearance without agglomerates over the entire span of the membrane. Optical microscopy images of the freestanding films also suggest that the film is uniform at the microscale (Figure 1E). Among the 4 types of samples prepared (Table 1), free-standing films with submicrometer thickness and > 10 cm in size could only be obtained from CNT/flavin/toluene suspensions (film 1). Samples other than film 1 ruptured during the transfer process, suggesting that film 1 is more suitable for applications requiring free-standing films, such as switching/sensing elements in MEMS or ultrafiltration membranes.

Differences in transferability observed for the four types of samples stem from differences in the CNT film mechanical properties. It should be noted that film 1 showed markedly better mechanical properties than the other samples (films 2–4).

Figure 2, which shows the stress-strain curves of the CNT films, indicates that film 1 shows the highest tensile strength (166 MPa), fracture strain (~2.4%), and Young's modulus (7.6 GPa) among the four samples. The tensile strength, fracture strain, and Young's modulus of each sample calculated from the stress-strain curve (summarized in Table 1) show distinct differences in mechanical performances of the four samples. Film 1 shows nearly three times higher tensile strength than the other three samples (films 2–4). With regard to fracture strains, films 1 and 2 show similar values - nearly 1.5 times the values shown by the other two samples (films 3 and 4). In terms of Young's modulus, films 1 and 3 show similar values, which are more than twice the values shown by the other two films (films 2 and 4). We think the different mechanical properties observed are
<figure><img src="image_71.png" /><figcaption>FIGURE 2 | Tensile stress-strain curves of the CNT films.</figcaption></figure>

related to differences in the bundle and pore structures of each sample. Therefore, we systematically characterized and studied the influence of these structural factors. The results of the film structure characterization and their correlation to the mechanical performances are presented below.

We note that the tensile strength of film 1 (166 MPa) is comparable to thin free-standing films obtained by dry process (Zhang et al., 2005; Ma et al., 2007; Liu et al., 2011; Nasibulin et al., 2011; Zhou et al., 2012). The mechanical robustness of film 1 contributes to free-standing film fabrication in two ways.

1. It helps the film support its own weight, making free-standing submicrometer-thick samples possible.

2. It maintains film integrity during lift-off in the wet-process protocol.

We also measured film sheet resistances (film 1: 18.2 ± 1.9Ω, film 2: 18.3 ± 0.4Ω, film 3: 14.6 ± 1.1Ω, and film 4: 18.1 ± 1.7Ω). Since the values are similar and do not show a clear correlation with the film structure, we focused only on the mechanical properties.

## Characterization of the CNT Film Mesh Structure

As reported in previous papers (Sakurai et al., 2018), bundled CNTs entangle in blade-coated CNT films, forming a mesh structure. Two factors could affect the film strength: (A) mesh structure of CNT bundles constituting the film (including attributes such as bundle thickness and inter-and intra-bundle pore sizes) and (B) characteristics of the CNTs themselves (such as length, diameter, wall number, crystallinity, functionalization level, etc.). In this paper, we fixed factor (B) by using identical CNT starting materials and processing methodology (sonication-based dispersion and blade-coating) for fabricating all our film samples. Our CNTs are typically single-walled nanotubes (average wall number ~1.2), 100–300 µm in length, and 1–5 nm in diameter

(average ~3.7 nm) (Kobashi et al., 2019). Therefore, we assume that CNT attributes in our samples are comparable, which is supported by the similar Raman G/D ratios of our films (0.9, 1.9, 2.2, and 2.3, respectively, for films 1–4). Hence, we have attempted to explain the exceptional tensile strength of film 1 compared to the rest of the samples based on differences in the CNT bundle mesh structure. We characterize the mesh structure of the four films through relative standard deviation (RSD) of bundle thicknesses measured by SEM, Brunauer-Emmett-Teller (BET)-specific surface area (SSA) by gas adsorption analysis, bulk density, and G/D ratio by Raman spectroscopy (summarized in Table 2).

Our analysis of the mesh structure indicates that film 1 shows narrower bundle thickness and pore-size distributions than the other samples, which could contribute to the high mechanical strengths observed. **Figure 3** compares the mesh structure uniformity of the strongest and weakest CNT films (films 1 and 4, respectively) in terms of SEM images and bundle thickness distribution. The CNTs are randomly aligned in both samples, which is typical for CNTs in films prepared by wet process. However, film 1 shows smaller bundles with a narrower thickness distribution (bundle thickness 10–90th percentiles: 7–11 nm, RSD: 20%) than film 4 (bundle thickness 10–90th percentiles: 19–109 nm, RSD: 77%). Consequently, film 1 contains CNT bundles of similar thicknesses forming a dense mesh with similar-sized small pores distributed uniformly throughout (Figure 3a). This narrow bundle thickness distribution of film 1 is a key feature contributing to its structural uniformity and, hence, high strength. In contrast, in film 4 (Figure 3b), bundles of various sizes entangle to result in pores of different dimensions. The presence of bundles of various thicknesses usually leads to the formation of macropores, which reduces structural uniformity and thereby lowers film mechanical strength as observed in film 4.

Quantitative pore size distributions obtained by mercury porosimetry support SEM results on the presence of smaller pores with more size uniformity in film 1 compared to film 4 (Figure 4). Figure 4, which plots pore volume (normalized by sample weight) versus pore diameter, shows that film 1 contains fewer pores of sizes >300 nm compared to film 4. The widths of the pore volume versus diameter plots can be used as a measure of the pore-size distribution and uniformity. Film 1 shows a smaller pore volume versus diameter plot width than film 4, indicating higher pore-size uniformity in film 1 than film 4. The overall structural uniformity of film 1 (arising from similarly sized pores and bundles) enables even stress distribution, precluding stress concentration, leading to high tensile strength.

In addition to structural uniformity, bulk density is a well-known factor affecting mechanical strength: the higher the density, the better the CNT packing and stress transfer and the higher the tensile properties (Ma et al., 2007; Jiang et al., 2011; Zhou et al., 2012). In addition to better structural uniformity, film 1 also has a higher bulk density (0.84 ± 0.06 g/cm³) than film 4 (0.47 ± 0.14 g/cm³), which additionally contributes to the higher strength values observed. We emphasize that the
<table><caption>**TABLE 2** | RSD of bundle thickness measured by SEM, SSA by gas adsorption analysis, bulk density, and G/D ratio by Raman spectroscopy of the four types of CNT films prepared using four types of suspensions.</caption><thead><tr><th>Sample name</th><th>G/D ratio</th><th>Bundle thickness (RSD) [%]</th><th>Bulk density [g/cm³]</th><th>SSA [m²/g]</th></tr></thead><tbody><tr><td>Film 1</td><td>0.9</td><td>20</td><td>0.84</td><td>766</td></tr><tr><td>Film 2</td><td>1.9</td><td>35</td><td>0.69</td><td>595</td></tr><tr><td>Film 3</td><td>2.2</td><td>71</td><td>0.82</td><td>500</td></tr><tr><td>Film 4</td><td>2.3</td><td>77</td><td>0.47</td><td>1130</td></tr></tbody></table>

<figure><img src="image_72.png" /><figcaption>**FIGURE 3** | SEM image and CNT bundle thickness distribution of (a) film 1, (b) film 4.</figcaption></figure>

mesh structure uniformity in terms of bundle thickness and pore-size uniformity more profoundly affects film mechanical strength than the bulk density. The sample with high bulk density (film 3, 0.82 ± 0.12 g/cm³) shows a much lower mechanical strength than film 1. We attribute this to the mesh structure of film 3, which, similar to film 4, is non-uniform with diverse bundle thicknesses and pore sizes (bundle thickness 10–90th percentiles: 19–111 nm, RSD: 71%) (see **Supplementary Figure S1**).

The structure and performance of film-2 also support that structural uniformity is important for film robustness. Film-2 shows similar mesh-structures as film 1 (see **Supplementary Figure S1**). The bundle thickness distribution for film 2 (bundle thickness 10–90th percentiles: 8–19 nm, RSD: 35%) is comparable to that for film 1 (bundle thickness 10–90th percentiles: 7–11 nm, RSD: 20%). Although film 2 (strength: 71 MPa) is weaker than film 1 (strength: 166 MPa), film 2 is stronger than film 4 (strength: 33 MPa) and film 3 (strength: 55 MPa). The discrepancy in the film strengths between films 1 and 2 indicates other factors are in play, affecting mechanical robustness. The poor tensile performance of film 2 can be explained based on two factors.

First, film 2 shows a smaller bulk density (0.69 ± 0.10 g/cm³) than film 1 (0.84 ± 0.06 g/cm³), which translates to poorer CNT packing and stress transfer. Second, film 2 might contain a higher amount of residual surfactant (SDOC), which essentially acts as defects and, therefore, stress centers, undermining the tensile performance of the sample. We observed 8.6 wt% of TGA (air) residue for film 2 (see **Supplementary Figure S2**). On the other hand, very little residual flavin is left as residue in film 1. We recovered and detected at least 99% of flavin in rinsing solutions, as confirmed by ultraviolet-visible spectrophotometry (see **Supplementary Figure S3**). In addition, the virtual absence of residual surfactant may also be attributed to the higher (more than two times) Young's modulus of film 1 compared to film 2. However, it should be also noted that Young's modulus is different between film 3 and film 4 by more than two times despite similar mesh structures and being unaffected by surfactant residues. This suggests that CNT packing is more important for Young's modulus than uniformity. In fact, gas adsorption/desorption isotherm analyses (see **Supplementary Figure S4** and **Table 2**) indicate that film 4 contains more micropores and shows higher SSA compared to film 3, suggesting
<figure><img src="image_73.png" /><figcaption>**FIGURE 4** | Pore-size distribution of films 1 (red) and 4 (black) measured by mercury porosimetry.</figcaption></figure>

that the existence of inter- or intra-bundle pores is an important factor affecting the Young's modulus of the CNT films.

Finally, the correlation of CNT film mechanical properties with mesh-structure uniformity and CNT packing is summarized in **Figure 5**. **Figure 5** plots the tensile strength, fracture strain, Young's modulus versus bundle size distribution and bulk density for the four types of samples. These figures show a clear dependence of tensile strength and fracture

strain on bundle size distribution, which represents network structure uniformity. This suggests that controlling bundle size and, therefore, bundle and pore structures leads to five-fold and four-fold higher tensile strengths and fracture strains, respectively, compared to films with more bundle size variation. On the other hand, Young's modulus depends more on bulk density, which represents the number of pores and degree of CNT packing. Therefore, films with higher bulk density with better CNT packing show three-fold higher Young's modulus than those with the least packing (and, therefore, bulk density).

We believe that the differences in film mesh structures observed are related to differences in the CNT aggregate sizes and their distributions in the parent suspensions. Suspending CNTs in a liquid medium by ultrasonication is typically known to result in nanotube bundle agglomerations of sizes ranging to tens of microns (Kobashi et al., 2013; Sakurai et al., 2018). These aggregates, in turn, form higher order structures, (Kobashi et al., 2013) such as our films, when suspensions are processed by techniques such as blade-coating, drop-casting, etc. We observe clear differences in CNT aggregate size distributions of the different suspensions we made for film fabrication as seen in the number frequency versus aggregate size plots (Figure 6). The parent suspensions of films 1 and 2 show a noticeably higher number fraction (>95%) of smaller aggregates (<10-micron-diameter) than that observed in the parent suspensions of films 3 and 4 (78.1 and 54.5%, respectively). The higher fraction of smaller aggregates in parent suspension 1 and 2 is indicative of superior nanotube

<figure><img src="image_74.png" /></figure>

<figure><img src="image_75.png" /></figure>

<figure><img src="image_76.png" /></figure>

<figure><img src="image_77.png" /><figcaption>**FIGURE 5** | Plots correlating mechanical strength and bundle network structure. (A,C) Tensile strength (blue), fracture strain (red), and (B,D) Young's modulus (green) versus bundle thickness distribution and bulk density of the CNT films.</figcaption></figure>
<figure><img src="image_78.png" /><figcaption>**FIGURE 6** | Size distribution of CNT aggregates measured by flow image analyzer of (A), suspension 1, (B) suspension 2, (C) suspension 3, (D) suspension 4.</figcaption></figure>

debundling during the dispersion process, which leads to smaller and uniform-sized CNT bundles that ultimately result in the uniform mesh-structured films 1 and 2. The superior debundling of suspensions 1 and 2 can be attributed to the presence of a surfactant, i.e., flavin and SDOC, respectively, that are known to stabilize CNT aggregates and prevent their agglomeration by strong steric/electrostatic repulsions. In contrast, suspensions 3 and 4 are made without a surfactant by sonicating CNT powders in a neat solvent (i.e., PG and NMP, respectively) that stabilize the aggregates only by weak solvation effects. To sum up, well-stabilized CNT suspensions with a large fraction of small aggregates are key to making films with highly uniform mesh structures, in turn, vital to achieving high film mechanical strengths. Therefore, our study draws a holistic connection between the suspension quality, mesh structure of the films obtained, and their mechanical performances. We believe this bottom-up tailoring of robust free-standing ultrathin CNT films by wet-process

blade-coating demonstrated here provides a pathway to the nanocarbon community for custom-building free-standing high-performance films of required properties. For instance, our work could be extended to CNTs with any predefined set of attributes (wall numbers, diameters, chirality, lengths, etc.) to improve electrical performances, optical properties, and so on because the suspension technique and bladed-coating can be adapted to various CNTs easily.

## CONCLUSION

In summary, we fabricated large (10-cm-square) 200-nm-thick free-standing robust CNT films by a scalable, wet process involving blade-coating of CNT/surfactant suspensions. We demonstrate that a uniform mesh structure with narrow bundle and pore size distribution is vital to achieving robust thin films with strengths as high as 166 MPa. Further, we correlate the mesh structure with CNT aggregate sizes (and its distribution) in the parent suspension. We find that well-stabilized suspensions using surfactants with higher fractions of smaller aggregates lead to CNT films with uniform mesh structures that exhibit high mechanical performances. Our work presents a bottom-up approach of tailoring robust ultrathin free-standing CNT films by wet process through a versatile and potentially scalable fabrication methodology.

## DATA AVAILABILITY STATEMENT

The raw data supporting the conclusions of this article will be made available by the authors, without undue reservation.

## AUTHOR CONTRIBUTIONS

YK performed sample preparation, measurement, and analysis. YK and AS conceived the idea and designed the experiments. KK contributed to flow image analysis. YK, AS, and RS prepared the draft of the manuscript. TY and KH conceived the idea and supervised the project. All authors contributed to the article and approved the submitted version.

## ACKNOWLEDGMENTS

We thank Dr. T. Morimoto and Dr. T. Okazaki for useful discussion. M. Nishimura, R. Shiina, F. Kamada, K. Miura, M. Fujii, and S. Nemoto are acknowledged for their technical support.

## SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fmats.2020.562455/full#supplementary-material
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**Conflict of Interest:** 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 © 2020 Kato, Sekiguchi, Kobashi, Sundaram, Yamada and Hata. 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.
Ceramics and Glass

<figure><img src="image_79.png" /><figcaption>Francesca Tittarelli</figcaption></figure>

Prof. Francesca Tittarelli has a Ms. Degree in Chemistry, a Ph.D. in Materials Engineering, and is an Associate Professor in Materials Science and Technology at Università Politecnica delle Marche (UNIVPM) for the "Civil and Environmental Engineering" and "Building Engineering and Architecture" courses. She is a member of the RILEM Technical Committee HDB "Hygrothermal behaviour and Durability of Bio-aggregate based building materials" and of the Directive Board of the National Interuniversity Consortium of Materials Science and Technology (INSTM). She is the author of more than 220 papers about the development, characterization, durability, and sustainability of building materials and the co-inventor of two national patents.

<figure><img src="image_80.png" /><figcaption>Chiara Giosuè</figcaption></figure>

Dr. Chiara Giosuè, Eng, Ph.D., is a Research Fellow at the SIMAU department (UNIVPM). Her research focuses on the interaction between environment and materials, such as waste valorization in sustainable construction and depolluting materials, alternative binders/aggregates for the construction industry, functionalization of materials to improve indoor air quality, material characterization and materials for conservation of cultural heritage, and the management and optimization of critical raw materials in the context of a circular economy. She is involved in teaching activities in Materials Science and Technology as a supervisor/co-supervisor of Bachelor's and Master's theses, the author/co-author of more than 40 scientific publications in international journals, and a reviewer of several of those same journals.
<figure><img src="image_81.png" /></figure>

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# Multifunctional Lightweight Mortars for Indoor Applications to Improve Comfort and Health of Occupants: Thermal Properties and Photocatalytic Efficiency

Chiara Giosuè<fnref n="1" />*, Mattia Pierpaoli², Alessandra Mobili<fnref n="1" />, Maria Letizia Ruello<fnref n="1" /> and Francesca Tittarelli<fnref n="1" />,³*

## OPEN ACCESS

**Edited by:**

Emilia Morallon,
University of Alicante, Spain

**Reviewed by:**

Maria Angeles Lillo-Rodenas,
University of Alicante, Spain
Inês Flores-Colen,
Instituto Superior Técnico, Portugal
Mohammed Quzzine,
Université Sultan Moulay Slimane,
Morocco

*Correspondence:

Chiara Giosuè
c.giosue@univpm.it
Francesca Tittarelli
f.tittarelli@univpm.it

**Specialty section:**

This article was submitted to
Ceramics and Glass,
a section of the journal
Frontiers in Materials

**Received:** 11 May 2020

**Accepted:** 14 July 2020

**Published:** 26 August 2020

**Citation:**

Giosuè C, Pierpaoli M, Mobili A,
Ruello ML and Tittarelli F (2020)
Multifunctional Lightweight Mortars
for Indoor Applications to Improve
Comfort and Health of Occupants:
Thermal Properties and Photocatalytic
Efficiency. Front. Mater. 7:255.
doi: 10.3389/fmats.2020.00255

A new generation of smart building materials, able to passively improve the indoor environment and the comfort of occupants owing to their interaction with the surrounding environment, can be addressed. This paper investigates the use of three highly porous aggregates to manufacture hydraulic lime-based multifunctional mortars to be used as indoor finishes. The same water/binder ratio was used for each mortar mix, and conventional calcareous sand was totally replaced by volume with zeolite, silica gel, and activated carbon. Nanosized titanium dioxide (TiO₂) was added to award a photocatalytic behavior under UV radiation to the mortars. Results show that, as expected, when highly porous aggregates are used, mortars absorb more water by capillary suction. However, even though the mortars manufactured with lightweight aggregates have a lower density, the mechanical behavior of zeolite and activated carbon mortars is comparable or even higher than that of sand mortars, thanks to an optimum interfacial transition zone (ITZ) between the binder paste and the aggregate. The photocatalytic activity, in terms of photocatalytic NOₓ degradation efficiency and selectivity of unwanted produced NO₂, results to be optimal when silica gel-based mortar is tested. Additionally, the thermal-insulation properties are enhanced up to 40% by using all the unconventional aggregates.

**Keywords:** Indoor Air Quality, lightweight mortar, titanium dioxide photocatalytic oxidation, capillary water absorption, thermal conductivity

## INTRODUCTION

A remarkable issue of modern time is the sustainability of the building sector since it contributes 40% on global carbon dioxide emissions (Wang et al., 2012). For this reason, nowadays European laws and directives are becoming stricter in terms of energy efficiency and, consequentially, more sealed buildings are assembled and not enough air changing can be guaranteed (Wolkoff, 2013). In this condition, the concentration of chemical, physical, and biological pollutants, such as volatile

<footnote n="1">¹ Department of Materials, Environmental Sciences and Urban Planning (SIMAU), Università Politecnica delle Marche – INSTM Research Unit, Ancona, Italy, ² Faculty of Electronics, Telecommunications and Informatics, Gdañsk University of Technology, Gdañsk, Poland, ³ Institute of Atmospheric Sciences and Climate, National Research Council (ISAC-CNR), Bologna, Italy</footnote>
organic compounds (VOCs), anhydrides, particulate matters, and molds, can be higher inside than outside, impairing the Indoor Air Quality (IAQ) of living environments. Currently, people are spending about 90% of their time indoor (Frey et al., 2015); thus, it is of utmost importance that the indoor living environment is wholesome to guarantee the health and comfort of occupants (Pierpaoli and Ruello, 2018).

Several strategies are used to improve the IAQ also in terms of thermal comfort; the conventional ones are source control, dilution by ventilation, and active engineered control systems (in heating, ventilation, and air-conditioning, HVAC) or stand-alone air purifiers (Matsumoto et al., 2009). Active control systems are generally more effective than passive systems for the control of IAQ, but they have a substantial energy loss that can be reduced by the addition of passive systems in the environment (Trník et al., 2016). Moreover, the combination of multiple techniques to improve IAQ, and in particular to remove pollutants, may result in a more efficient synergistic effect than the sum of the individual ones (Pierpaoli et al., 2017, 2019a). To remove pollutants, titanium dioxide (TiO₂) is the most studied photocatalytic agent, because of its chemical inertness, non-toxicity, and low cost. Therefore, there is an increasing interest in using such catalyst in cementitious materials (Ballari et al., 2011; Folli et al., 2012).

Smart materials are based on the innovative concept of designing the material composition by introducing functional components into the matrix (Han et al., 2016) able to give different functionalities to the same material such as good mechanical performances, lightness, and ability to improve IAQ in terms of low thermal conductivity, to act as an hygroscopic buffer (Senff et al., 2018), and/or to remove pollutants.

In this case, the addition of highly porous materials with adsorbent properties is the key aspect to providing the multifunctionality of mortars, as successfully demonstrated in previous works where they were added to commercial finishes (Giosuè et al., 2017a). Several studies have tested the NOₓ photocatalytic degradation, where TiO₂, added at 1–4% by binder mass, gives a removal efficiency ranging from 30 to 80% (Senff et al., 2013, 2018), also under visible irradiation (Giosuè et al., 2018).

In the current study, the photocatalytic activity is associated with the adsorption of unconventional highly porous aggregates added to different mixes. The adsorption on a solid surface can be influenced by forces of electrostatic nature due to differences in polarity and differences in molecular weight and molecular geometry which induce selective deposition on the solid surface (Masel, 1996). The adsorption process is preferred on solids with a high specific surface since adsorption is a superficial process (Matsumoto et al., 2009). In mortars, the effectiveness of adsorbent aggregates to remove VOCs can reach 65%, as it has been previously tested (Tittarelli et al., 2015; Giosuè et al., 2019b). Adsorbent fillers/aggregates have been also tested to remove NOₓ in cementitious systems with photocatalytic agents under UV-A radiation (Horgnies et al., 2012; Pierpaoli et al., 2019b).

The use of unconventional highly porous aggregates can also be employed to develop lightweight mortars with low sound transmission (Branco and Godinho, 2013) and minimal thermal

transmittance (Benfratello et al., 2013; Branco and Godinho, 2013; Giosuè et al., 2019b) to the advantage of thermal efficiency and indoor sound comfort. It is well known that thermal conductivity can be reduced by the addition of aggregates with low bulk density and/or low thermal conductivity (Benfratello et al., 2013). Another positive effect is the reduction of the material density with a consequent lowering of dead loads on the structure. Nevertheless, in materials the reduction of density is often related to the reduction in compressive strength (Demirboga and Gül, 2003; Yu et al., 2015).

The current work is aimed to develop innovative multifunctional mortars for indoor finishes/ renders able to passively improve IAQ, besides fulfilling the ordinary requirements (according to the standard EN 998). The main objective is not to find the mortar with the best

<table><caption>TABLE 1 | Properties of different aggregates.</caption><thead><tr><th>Mix</th><th>Image</th><th>ID</th><th>Density kg/m³</th><th>Water absorption %</th></tr></thead><tbody><tr><td>Calcareous sand</td><td></td><td>S</td><td>2650</td><td>5</td></tr><tr><td>Zeolite</td><td></td><td>A1</td><td>1600</td><td>20</td></tr><tr><td>Silica gel</td><td></td><td>A2</td><td>1310</td><td>86</td></tr><tr><td>Activated carbon</td><td></td><td>A3</td><td>530</td><td>30</td></tr></tbody></table>

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<figure><img src="image_84.png" /></figure>

<figure><img src="image_85.png" /></figure>

<figure><img src="image_86.png" /><figcaption>FIGURE 1 | X-ray diffraction (XRD) patterns of TiO₂, P-25 where anatase (A), rutile (R), and the amorphous phase (Am) are detected.</figcaption></figure>
<table><caption>TABLE 2 | Mix proportions (g/L), density (kg/m³), and accessible total porosity (%) of mortars.</caption><thead><tr><th rowspan="2">Mix</th><th colspan="7">Mix proportions</th><th rowspan="2">Density kg/m³</th><th rowspan="2">Accessible porosity %</th></tr><tr><th>Water</th><th>HL</th><th>S</th><th>A1</th><th>A2</th><th>A3</th><th>TiO<sub>2</sub></th></tr></thead><tbody><tr><td>M0-S</td><td>255</td><td>440</td><td>1535</td><td>-</td><td>-</td><td>-</td><td>-</td><td>1800</td><td>32</td></tr><tr><td>M0-S TiO<sub>2</sub></td><td>255</td><td>414</td><td>1535</td><td>-</td><td>-</td><td>-</td><td>26</td><td></td><td></td></tr><tr><td>M1-A1</td><td>255</td><td>440</td><td>-</td><td>927</td><td>-</td><td>-</td><td>-</td><td>1390</td><td>41</td></tr><tr><td>M1-A1 TiO<sub>2</sub></td><td>255</td><td>414</td><td>-</td><td>927</td><td>-</td><td>-</td><td>26</td><td></td><td></td></tr><tr><td>M2-A2</td><td>255</td><td>440</td><td>-</td><td>-</td><td>759</td><td>-</td><td>-</td><td>910</td><td>47</td></tr><tr><td>M2-A2 TiO<sub>2</sub></td><td>255</td><td>414</td><td>-</td><td>-</td><td>759</td><td>-</td><td>26</td><td></td><td></td></tr><tr><td>M3-A3</td><td>255</td><td>440</td><td>-</td><td>-</td><td>-</td><td>683</td><td>-</td><td>1210</td><td>35</td></tr><tr><td>M3-A3 TiO<sub>2</sub></td><td>255</td><td>414</td><td>-</td><td>-</td><td>-</td><td>683</td><td>26</td><td></td><td></td></tr></tbody></table>

mechanical properties, but to compare mortars that can act as multifunctional renders/finishes being able to improve also IAQ.

For this purpose, hydraulic lime is used as binder with and without the addition of TiO₂. Lime-based mortars are extensively used as renders/plasters in indoor applications. Moreover, hydraulic lime mortars are commonly used for rehabilitation of historical buildings where cement-based mortars are not allowed, and lime is a more sustainable binder than ordinary Portland cement (Barcelo et al., 2013; Lucas et al., 2013). It has been also demonstrated that hydraulic lime provides to mortars a higher depolluting activity than cement; this is due not only for the lower presence of mineralogical components that can hide the catalyst (Sugrañez et al., 2013; Kaja et al., 2019) but also for the higher porosity of its hardened microstructure (Karatasios et al., 2010; Giosuè et al., 2018).

The conventional sand is substituted by highly porous and adsorbent materials usually employed in filters for water-air cleaning processes as zeolite, silica gel, and activated carbon. As reference, a conventional sand-based mortar is used.

The mechanical strength and porosity of these mortars have been measured in a previous work (Giosuè et al., 2020). Here, the different compressive strength values are correlated with the corresponding microstructures observed by scanning electronic microscopy (SEM). Moreover, the behavior under capillary water absorption is investigated. In order to find the best mortar for improving the indoor comfort and health of indoor occupants, the corresponding thermal conductivity and photocatalytic activity are also tested. The previous work has demonstrated the optimum hygroscopic behavior of these mortars in terms of moisture buffering capacity and water vapor permeability but also a photocatalytic activity hidden by the adsorption process until saturation of specimens (Giosuè et al., 2020). Therefore, in this paper, to demonstrate the effectiveness of photocatalytic activity in batch, all mortars are tested in saturated conditions.

<figure><img src="image_87.png" /><figcaption>FIGURE 2 | Compressive strength of hydraulic lime-based mortars.</figcaption></figure>

# MATERIALS AND METHODS

## Materials

The binder is a commercial hydraulic lime LIC 3.0, according to UNI EN 15368, with a density of 2650 kg/m³. As aggregates, sand

(S) is fully substituted by volume with selected unconventional aggregates, namely, natural clinoptilolite zeolite (A1), silica gel (A2), and activated carbon (A3). During the cast, all aggregates are added in saturate surface dry (SSD) condition, which means that all voids of aggregates are saturated by water but the surface is dry; hence, when the water necessary for the cast is added, no water exchange between the aggregate and the paste takes place. The density and water absorption of all aggregates to reach the SSD condition are evaluated and summarized in Table 1.

A commercially available titanium dioxide, TiO₂, P-25 Aeroxide by Evonik, is used as photocatalytic agent, added by replacing 5% of the binder weight. TiO₂, in powder form, was added to the dry cement mixture prior to the water addition, carefully mixed manually for 3 min and then with a mortar mixer.

As reported in the data sheet, its chemical composition is a mixture of anatase–rutile–amorphous phases, 78–14–8% in weight, respectively. Particles have nano-size dimension around 20–50 nm. The specific surface, measured by BET, is 35–65 m²/g,
<figure><img src="image_88.png" /><figcaption>**FIGURE 3** | SEM images of mortars manufactured with different aggregates and TiO₂. (a) HL-S TiO₂, (b) HL-A1 TiO₂, (c) HL-A2 TiO₂, (d) HL-A3 TiO₂.</figcaption></figure>

and density is equal to 3100 kg/m³. The pH value of 4% dispersion in water is 3.5–4.5. The X-ray diffraction (XRD) pattern is shown in **Figure 1** where the crystalline phases are detected.

The mixes are prepared by keeping constant the water to binder ratio at 0.6. Mix proportions, density, and accessible total porosity of cured mortars without TiO₂ (since these properties were not significantly affected by TiO₂ nanoparticles) are reported in **Table 2** (Giosuè et al., 2020).

Since the retail price in Euro/kg is about 0.10 for calcareous sand, 2.02 for zeolite, 0.59 for silica gel, 5.00 for activated carbon, 44 for TiO₂, and 0.14 for HL, the price in Euro/m³ of the photocatalytic mortars with zeolite and activated carbon is about double that with conventional sand, whereas the price in Euro/m³ of the silica gel photocatalytic mortar is about half that of conventional sand photocatalytic mortar.

## Methods

Mortars are casted and then cured for 28 days according to the UNI EN 1015-11 procedure: at temperature (T) of 20°C and relative humidity (RH) of 95% for 7 days and then at 20°C T and 65% RH. After 28 days of curing, mortars are tested according to the current standards.

Properties such as mechanical performance, density, and porosity of these mortars have been already investigated in

Giosuè et al. (2020) and are briefly summarized in the current paper. Compressive strength is measured by means of a hydraulic press (Galdabini) after 28 days of curing. Density is evaluated by the ratio between the weight of the specimen and its dimension, measured with a caliper. Each measurement is repeated three times, and the average values are reported. The dry weight is evaluated after placing the specimens in oven at T = 105°C until a constant weight is reached (when the change in the weight of the test specimen over 24 h period is less than 0.1%). Porosity is evaluated by means of MIP (Thermo Fisher, Pascal series 240) on samples after 28 days of curing on three fragments per mortar of 1 cm³ volume, and the average results are reported.

In this study, the mechanical strengths are correlated with the microstructures observed by SEM. Moreover, the capillary water absorption, the photocatalytic behavior by NOₓ degradation test, and thermal conductivity are further investigated.

The morphology is analyzed by means of the SEM ZEISS 1530 (Carl Zeiss, Oberkochen, Germany) equipped with a Schottky emitter, with two different secondary electrons detectors (the inc-lens and the Everhart-Thornley) and operating at 10 keV (EDAX probe). Specimens are fragments of about 1 cm³.

Capillary water absorption is investigated according to UNI EN 15801. Capillary water absorption is affected not only by the pores volume but also by the distribution and size of the pores, since the lower the pore diameter, the higher the level
<figure><img src="image_89.png" /><figcaption>FIGURE 4 | SEM images (a) and EDAX (b) performed on HL-S TiO₂ mortar.</figcaption></figure>

reached by the rising water, according to the Washburn's equation (Washburn, 1921). Test specimens are prisms 160 × 40 × 40 mm broken in two halves. Specimens are dried in an oven at 105°C until the constant mass is reached (difference in the weight values measured after 24 h should be lower than 0.1%). The specimens are placed on a water-saturated filter paper, and the capillary water absorbed by the tested surface (16 cm²) is monitored over time by weighing the specimens periodically. The results are expressed as water absorbed per unit area $Q_i$ (kg/m²) at time $t_i$ (s⁰.⁵).

Concerning photocatalytic efficiency, after curing, the pH of mortars is measured in order to demonstrate the lowering of the basicity of the mixes. It is well known that CO₂ can penetrate within the matrix as easily as the porosity increases, interacting with the matrix through the carbonation process. This process causes the decrease of the pH and the densification of the matrix owing to the precipitation of calcite. In turn, calcite can cause the occlusion of the catalysts (Sugrañez et al., 2013). Also, the redox process that occurs at the TiO₂ surface is influenced by the pH, due to the Nernst relationship (MacPhee and Folli, 2016). 1 g of material is collected on a glass slide; then, a drop of water (about 1 ml) is added and the pH of the solution is measured.

The test adopted for assessing the photocatalytic property of the mortars follows the UNI 11247 procedure, which assesses the photocatalytic activity of inorganic photocatalytic materials by determining the degradation rate of nitrogen oxides (NOx) in air. Therefore, the NOx photocatalytic activity is expressed by the ratio of removed NOx over total NOx flux under plug flow conditions. Briefly, a NOx flux is directed inside the reactor by a NOx tank, 499 ppb NO (SAPIO S.r.l., Monza, Italy), and it is kept constant with a dilution system (Calibrator 8188, Rancon Instruments S.p.A., Milan, Italy) by mixing with atmospheric air at $T = 25 \pm 2°C$ and RH between 40 and 50%. NO and NO₂ concentrations are continuously monitored by a chemiluminescence NOx analyzer (nitrogen oxide analyzer model 8841, Monitor Labs, Englewood, CO, United States) at the reactor outlet. The photoreactor consists of a 3-L borosilicate glass cylinder, in the center of which the specimens (8 cm diameter and 0.8 cm height) are placed. A UV metal halogen lamp, having a measured irradiance of 20 W/m², is placed outside the reactor over the specimen and placed at 25 cm from the

specimen' surface. The irradiation of the sample is guaranteed until stable conditions are reached, usually 30 min.

The thermal conductivity ($\lambda$, in W/mK) of specimens is measured at room temperature ($T = 20°C$) and RH = 50%. Specimens are cylinders of 14 cm diameter and 3 cm height. The test is performed according to the UNI EN 12664 standard and $\lambda$ calculated by the following equation (Eq. 1):

$$ \lambda = \frac{Jd}{(T_1 - T_2)} \qquad (1) $$

where J is the heat flux (W/m²), d is the distance between thermocouples (m), and $T_1$ and $T_2$ are the temperatures at the two different sides of the sample (K). Before being tested, specimens have been conditioned at $T = 20°C$ and RH = 50% until constant weight is reached (when the change in the weight of the test specimen over a 24 h period is less than 0.1%).

# RESULTS AND DISCUSSION

## Compressive Strength and Microstructural Observations

The compressive strength values already measured in Giosuè et al. (2020) and the morphological observations obtained by SEM are reported in Figures 2, 3. For the sake of brevity, in Figure 3, only images of specimens manufactured with different aggregates and TiO₂ are shown.

Since the main aim of this paper is to compare the effects of different types of aggregates on the ability of mortars to improve IAQ in terms of thermal conductivity and photocatalytic activity, the mixes are prepared by keeping constant both the water amount and the same water-to-binder ratio. For this fact, it is worth underlining that while the resulting compressing strength values for A2 mortars are good for applications as finishes/ renders, those for A1 and A3 mortars are too high and not suitable for this type of application.

However, the good mechanical behavior of the mortar manufactured with sand (S) can be explained by the high mechanical strength of natural sand, due to its low porosity and the good adhesion between sand and hydraulic lime paste
<figure><img src="image_90.png" /><figcaption>**FIGURE 5** | Capillary water absorption in time of hydraulic lime-based mortars.</figcaption></figure>

(Figure 3a). EDAX analysis (Figure 4) evidences Ca of the aggregate and of the binder with Si and S present in the hydraulic-lime paste. TiO₂ particles are not visible in SEM images, confirming their optimum dispersion in the mix; anyhow, their presence is evident by EDAX analysis performed on the paste (Figure 4) due to the appearance of the Ti peak.

When the mortar is manufactured with zeolite (A1), even though zeolite is more porous, lighter, and therefore less resistant than natural sand, the mechanical resistance of the mortar remains high thanks to the pozzolanic activity of zeolite that forms additional hydration products, well visible as indicated in Figure 3b (Caputo et al., 2008; Uzal et al., 2010). Moreover, thanks to the pozzolanic reaction, the interfacial transition zone (ITZ) between zeolite and the binder paste can be hardly recognized, highlighting the optimum adhesion between the two components.

On the other hand, in Figure 3c, the ITZ between silica gel and hydraulic lime is clearly visible due to an evident detachment between the two components. As a matter of fact, the adhesion between hydraulic-lime paste and this type of aggregate appears very poor, due to the smooth surface of silica gel particles. This issue explains the worst mechanical behavior of silica gel-based mortars.

The ITZ between hydraulic lime paste and activated carbon is better than that with silica gel. In this case, even though activated carbon is more porous, lighter, and therefore less mechanical resistant than sand, the mechanical strength of mortars is even greater than that of sand-based mortars, probably thanks to its surface roughness which improves the adhesion with the

binder paste, guaranteeing an interlocking mechanism. This behavior was detected in other carbon-based materials with a porous surface that, acting as reservoirs, were able to provide a migration of water from their inside to the surrounding binder paste during the curing period, guaranteeing an optimum ITZ (Mrad and Chehab, 2019).

## Capillary Water Absorption

The open porosity is one of the most influencing parameters for water absorption of mortars: the higher the open porosity, the higher the water quantity that can fill the pores. Therefore, the use of unconventional aggregates increases the water uptake of mortars because they increase the porosity of the mortars compared to conventional sand (Giosuè et al., 2017b). As already demonstrated for these mixes (Giosuè et al., 2020) and reported in Table 2, the mortars with the highest value of open porosity are those manufactured with silica gel, followed by zeolite-based mortars and finally activated carbon-based mortars. The reference mortar has obviously the lowest porosity value.

In the absence of TiO₂, as expected, the reference mortar, being the less porous, absorbs the lowest amount of water (Figure 5). Zeolite (A1)-based mortars adsorb an amount of water at 60%, higher than sand-based mortars. Silica gel (A2) mortars have the highest water absorption (80% higher than sand-based mortar) since they are the most porous.

Activated carbon-based mortars absorb lower (at the beginning of the test) and then higher (40%) amounts of water than sand-based mortars. This behavior is mainly related to the characteristics of the aggregate: activated carbon is hydrophobic
<figure><img src="image_91.png" /><figcaption>FIGURE 6 | Photocatalytic efficiency of mortars under UV-A radiation.</figcaption></figure>

(Yang, 1987), with low affinity to water, at the early stage since it is a non-polar absorbent material, but it becomes hydrophilic for longer periods of contact with water and/or when some water has already been adsorbed (Giosuè et al., 2017b).

When the TiO₂ agent is added, generally the water absorption increases. This phenomenon, already shown in a previous study (Giosuè et al., 2017b), is due to the presence of TiO₂, since the activation of the photocatalyst gives hydrophilicity to mortars (Banerjee et al., 2015; Zhang et al., 2015).

## Photocatalytic Oxidation

The behavior of mortars under photocatalytic radiation is reported in **Figure 6**. After the curing period, the pH of the specimens is evaluated as indicated in section “Methods.” After the cast, the pH value of the cementitious product is around 13.5 (Tittarelli et al., 2018). The results show that the finishes lose their initial strong basicity since the detected pH values are equal to 8 for all mortars, apart from those of the M3 series (M3-A3 and M3-A3 TiO₂) with a pH of 8.5. For all TiO₂ specimens, a NO removal efficiency ranging from 20 to 40% is found. The lowest efficiency, equal to 20%, is registered in

mortars manufactured with activated carbon; their dark color can reduce the reflectance of radiation and consequently influences the photocatalytic properties of the substrate (Chen and Poon, 2009; Daoud, 2013).

In the TiO₂-mediated NO photocatalytic oxidation, NO₂ is produced as an unwanted product of the reaction. For this reason, the photocatalytic mortar exhibiting a lower NO₂ production should be preferred. In this regard, the ratio between NO₂ and NO is considered: the higher the ratio, the higher the production of unwanted NO₂ during the reaction (Pierpaoli et al., 2018). Whereas the ratio between the produced NO₂ and the removed NO is constant and around 16% for reference, zeolite, and activated carbon specimens, it is significantly lower when silica gel is used. This synergistic effect was already found in previous studies (Pierpaoli et al., 2018, 2019b) and may be attributed to an enhanced reactivity with the Ti-OH via NO₂ disproportion (Eq. 9) caused by a more available adsorbed water in the vicinity of the photocatalyst (Dalton et al., 2002) and by the subsequent reaction of the nitrates with alkali or alkaline earth metals ions (Eq. 10).

### TiO₂ photocatalysis

$$ \mathrm{TiO_2} + \mathrm{h\nu} \rightarrow \mathrm{TiO_2}^* (\mathrm{h}_{\mathrm{vb}}^{+} + \mathrm{e}_{\mathrm{cb}}^{-}) \quad (2) $$

$$ \mathrm{TiO_2}^* (\mathrm{h}_{\mathrm{vb}}^{+}) + \mathrm{H_2O}_{\mathrm{ads}} \rightarrow \mathrm{TiO_2} + \mathrm{OH}_{\mathrm{ads}}^{\bullet} + \mathrm{H}^{+} \quad (3) $$

$$ \mathrm{TiO_2}^* (\mathrm{e}_{\mathrm{cb}}^{-}) + \mathrm{O}_{2\mathrm{ads}} \rightarrow \mathrm{TiO_2} + \mathrm{O}_{2\mathrm{ads}}^{-} \quad (4) $$

### Oxidation via OH⁻

$$ \mathrm{NO}_{\mathrm{ads}} \xrightarrow{\mathrm{OH}_{\mathrm{ads}}^{\bullet}} \mathrm{HNO}_{2\mathrm{ads}} \xrightarrow{\mathrm{OH}_{\mathrm{ads}}^{\bullet}} \mathrm{NO}_{2\mathrm{ads}} + \mathrm{H_2O} \quad (5) $$

$$ \mathrm{HNO}_{2\mathrm{ads}} \rightleftharpoons \mathrm{NO}_{2\mathrm{ads}}^{-} + \mathrm{H}^{+} \quad (6) $$

$$ \mathrm{NO}_{\mathrm{ads}} \xrightarrow{\mathrm{OH}_{\mathrm{ads}}^{\bullet}} \mathrm{NO}_{3\mathrm{ads}}^{-} + \mathrm{H}_{\mathrm{ads}}^{+} \quad (7) $$

### Oxidation via O₂⁻

$$ \mathrm{NO}_{\mathrm{ads}} \xrightarrow{\mathrm{O}_{2\mathrm{ads}}^{-}} \mathrm{NO}_{3\mathrm{ads}}^{-} \quad (8) $$

<figure><img src="image_92.png" /><figcaption>FIGURE 7 | Relation between (A) thermal conductivity and density and (B) thermal conductivity and total pore volume.</figcaption></figure>
Disproportionation of NO₂

$$3\text{NO}_{2\text{ads}} + \text{H}_2\text{O} \rightarrow 2\text{HNO}_3 + \text{NO} \quad (9)$$

$$\text{HNO}_3 + \text{H}_2\text{O} \rightarrow \text{H}_3\text{O}_{\text{aq}}^{+} + \text{NO}_3^{-}_{\text{aq}} \quad (10)$$

## Thermal Conductivity

The enhancement of the indoor comfort of occupants is strictly related to the thermal insulation of the environment: if building materials with low thermal conductivity are used, the thermal comfort can be guaranteed with a lower energy consumption by active systems. According to the standard UNI EN 998-1, the threshold limits for λ to classify a mortar as a T1 or T2 thermal mortar are 0.100 W/mK and 0.200 W/mK, respectively.

In the current study, almost all mortars can be classified as T2 mortar since their λ is lower than 0.200 W/mK. In particular, the reference mortar has the highest value (λ = 0.142 W/mK), followed by activated carbon-based mortar and zeolite-based mortar with λ values of 0.110 and 0.110 W/mK, respectively. The lowest value of λ is detected in silica gel-based mortars (0.079 W/mK), which can be classified as a T1 mortar.

The λ of a mortar is related to both its density and porosity. The lower the density, the lower the thermal conductivity, and the higher the porosity, the lower the thermal conductivity (Demirboga and Gül, 2003). As expected, the use of highly porous aggregates permits to obtain lightweight mortars (Giosuè et al., 2018). Thanks to the total replacement of calcareous sand with silica gel (A2) and activated carbon (A3), lightweight finishes, highly appreciated for non-structural materials, can be obtained (Demirboga, 2003; Branco and Godinho, 2013).

Results of porosity and density of specimens are already shown in Giosuè et al. (2020), but in this paper the correlation between these properties and thermal conductivity is proposed and reported in Figure 7. Similar to the findings already published in other studies, a very good linear correlation between thermal conductivity and density (Coppola et al., 2018; Giosuè et al., 2019a) is obtained, as shown in Figure 7A. Usually, mortars with a high percentage of voids have strong insulating properties and can be used as plastering joints, building blocks, and panels able to guarantee the thermal comfort. Again, a linear correlation (but in this case with a lower coefficient) can be found between thermal conductivity and the total pore volume, as reported in Figure 7B.

## CONCLUSION

In the present work, hydraulic lime-based mortars manufactured with unconventional highly porous aggregates are studied to be used as multifunctional indoor finishes. Calcareous sand (S) is replaced entirely by volume with three different lightweight aggregates, namely, zeolite, silica gel, and activated carbon. Mortars are prepared with and without the addition of TiO₂. The compressive strength values of these mortars are related to their microstructure observed by SEM. Then, the capillary water absorption, the photocatalytic oxidation of NOₓ, and the thermal conductivity are also investigated.

From the current research, it is possible to stand that:

* The dispersion of TiO₂ nanoparticles in the different analyzed matrixes is good, as shown by SEM analysis;

* Even though the mortars manufactured with lightweight aggregates have a lower density, the mechanical behavior of zeolite and activated carbon mortars is comparable or even higher than that of sand mortars, thanks to an optimum ITZ between the binder paste and the aggregate. In the case of silica gel-based mortars, the worst mechanical strength is due to the bad adhesion observed between this aggregate and the paste;

* Zeolite, silica gel, and activated carbon-based mortar adsorb a total amount of water 60%, 80%, and 40% higher than that of the reference mortar, respectively, because of their higher total porosity;

* The mortar prepared with silica gel exhibits the highest photocatalytic activity and the lowest selectivity toward NO₂ because of the higher availability of adsorbed water, which reflects into an enhanced production of OH radicals due to an enhanced reactivity with Ti-OH via disproportion caused by a more available adsorbed water in the photocatalyst vicinity.

* An optimal linear correlation is confirmed between thermal conductivity and density, whereas a good linear correlation is found between porosity and density.

Further studies will be carried out in order to analyze the photocatalytic behavior of TiO₂ active under visible light.

## Permission to Reuse and Copyright

Figures, tables, and images will be published under a Creative Commons CC-BY license, and permission must be obtained for use of copyrighted material from other sources (including republished/adapted/modified/partial figures and images from the internet). It is the responsibility of the authors to acquire the licenses, to follow any citation instructions requested by third-party rights holders, and to cover any supplementary charges.

## DATA AVAILABILITY STATEMENT

The raw data supporting the conclusions of this article will be made available by the authors, without undue reservation.

## AUTHOR CONTRIBUTIONS

CG: original research plan, methodology and investigation, material characterization, specimen preparation, original manuscript draft, and preliminary data curation. MP: collaboration in data curation of photocatalytic activities and manuscript editing. AM: collaboration in data curation of thermal conductivity, and manuscript review and editing. FT: coordination and supervision of the project, analysis of data, and manuscript draft and editing. MR: supervision of tests to evaluate the depolluting properties of mortars, interpretation of relative data, and manuscript editing. All authors contributed to the article and approved the submitted version.
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**Conflict of Interest:** 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 © 2020 Giosuè, Pierpaoli, Mobili, Ruello and Tittarelli. 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.
Colloidal Materials and Interfaces

<figure><img src="image_93.png" /><figcaption>Lourdes Basabe-Desmonts</figcaption></figure>

Prof. Lourdes Basabe-Desmonts has been a Research Professor since 2012, the Director of the Microfluidics & BIOMICs Cluster UPV/EHU at the University of the Basque Country since 2019, and a Young Member of Jakiunde, the Academy of Science and Arts of the Basque Country, since 2016. Lourdes studied chemistry at the Universidad Autónoma de Madrid and completed a Ph.D. at the MESA+ Institute for Nanotechnology in the University of Twente, the Netherlands. In 2006, she joined the Biomedical Diagnostics Institute in Dublin, Ireland, and she moved to Spain in 2012. Her research interest is focused on the development lab-on-a-chip analytical microsystems.

<figure><img src="image_94.png" /><figcaption>Fernando Benito-Lopez</figcaption></figure>

Prof. Fernando Benito-Lopez studied chemistry at the Universidad Autonoma de Madrid and completed his Master's studies in the Department of Inorganic Chemistry in 2002. He obtained his Ph.D. at the University of Twente, the Netherlands, in 2007. He carried out his postdoctoral research in the group of Prof. Dermot Diamond at Dublin City University, Dublin, where, in 2010, he became Team Leader in polymer microfluidics. In 2012, he moved to CIC microGUNE—a Research Centre working in Microtechnology in Spain. From 2015, he has been a Ramón y Cajal Fellow and leader of the Microfluidics Cluster UPV/EHU at the University of the Basque Country, Spain. In 2019, he became an Associate Research Professor and leader of the Microfluidics and Biomics Cluster UPV/EHU Group.

<figure><img src="image_95.png" /><figcaption>Monica Francesca Pucci</figcaption></figure>
Dr. Monica Francesca Pucci, Ph.D. in mechanical engineering, is an assistant professor at IMT MINES Alès. She has background in mechanical engineering and has dedicated a significant part of her research to wetting and wetting dynamics. Those works have made her a renowned researcher in the field of impregnation of fibrous media driven by capillary forces for composite manufacturing. Her research activities are focused on the characterization of material microstructure and interface properties by means of Atomic Force Microscopy. In 2020, she was awarded the Daniel Valentin prize by the French association for composite materials (AMAC).