In this paper, we introduce a series of recent studies which attempted to obtain the effective physical properties of composites by taking into account the interfacial damage and anisotropy of the matrix, providing a universal formulation for different physical phenomena based on their mathematical analogies. First, with the specific example of elasticity, we introduced the concept of the Eshelby tensor and discussed a how single-inclusion problem can be related to a single inhomogeneity problem, which ultimately is applied to meanfield homogenization to predict the effective properties of composites. Second, based on a mathematical analogy, we show that mean-field homogenization equations for different physical phenomena can be represented by linear equations involving symmetric matrices with different dimensions by adapting Mandel's notation. Third, we extend our discussion by taking into account two common issues in realistic applications: imperfections in a matrix-inhomogeneity interface and anisotropy of the matrix.

Although we have discussed the homogenization method while covering a wide range of issues, there remain a few issues requiring further developments. We categorize the advanced issues of homogenization theory into the six problems of (i) various physical phenomena, (ii) anisotropy of the matrix, (iii) matrix-filler interfacial imperfections, (iv) nonlinear responses, (v) time-dependent responses, and (vi) high volume fractions of reinforcements. Although issues (i)-(iii) have been investigated extensively, we nonetheless cannot take into account important non-spherical reinforcement materials (including ellipsoidal types) such as carbon nanotubes and graphene in the presence of interfacial imperfections due to the non-uniform interior field. Although there have been many scaling up attempts to couple micro/nanoscale simulations and homogenization theory to predict the properties of nanocomposites (Yang et al., 2012, 2013a, 2014; Pan et al., 2013; Shin et al., 2013; Shokrieh et al., 2016), they either rely on an incorrect solution stemming from effective stiffness expressions derived from a few problematic bases, such as violations of Fubini's theorem or inappropriate superpositions, or do not recognize the nonuniform interior field (Barai and Weng, 2011; Yang et al., 2012, 2013b; Pan et al., 2013; Shokrieh et al., 2016). There are a few studies of composites showing highly nonlinear (such as hyperelastic) and time-dependent (such as viscoelastic) responses (Friebel et al., 2006; Miled et al., 2011, 2013; Brassart et al., 2012; Doghri et al., 2016), but most lack consideration of interfacial imperfections, which can play an important role on the macroscale composites during repeated cyclic loading

and/or in nanocomposites including nanoscale reinforcements with large interfacial fractions. Lastly, to account for the high reinforcement volume fraction regime where the interaction among reinforcements becomes important, several heuristic methods have been developed, such as differential schemes (Hashin, 1988) and/or matrix-filler inversion schemes (Friebel et al., 2006) as well as studies of the lower and upper bounds of the effective properties (Hashin and Shtrikman, 1961), leaving room, however, for further improvements. Studies to overcome aforementioned challenges are underway in our research group with the formation of an international collaborative research network.

# AUTHOR CONTRIBUTIONS

SR designed the structure of the review paper. SR and SL wrote the manuscript. SR and SL. contributed the sections on the micromechanics review and the treatment of anisotropy

# REFERENCES

and interfacial imperfections. SR, SL, and JJ contributed to the summary of theoretical derivations, prepared figures, and discussed the mathematical analogy among different physical problems. JL contributed the theoretical derivations. YK contributed the effective inclusion method section.

# FUNDING

This work is supported by the National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT (2016M3D1A1900038, 2016R1C1B2011979, and 2018K2A9A2A12000223).

# SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fmats. 2019.00021/full#supplementary-material

through nature-inspired hierarchical complexity. Acta Mater. 57, 2919–2932. doi: 10.1016/j.actamat.2009.03.003

carbon nanotube polymer matrix composites. Mech. Mater. 58, 1–11. doi: 10.1016/j.mechmat.2012.10.015

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

Copyright © 2019 Ryu, Lee, Jung, Lee and Kim. 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.

## Polymeric and Composite Materials

Federico Carosio

Federico Carosio obtained his PhD in 2012 in Materials Science and Technology at Politecnico di Torino with a thesis titled: "*Layer by Layer assembly of nanostructured coatings for flame retardancy and barrier properties*". At present he is a senior researcher at the Department of Applied Science and Technology of Politecnico di Torino, Italy. His research focuses on polymers and nanostructures, smart fibers and textiles and nanocellulose-based materials with particular attention on their flame retardancy properties achieved through the Layer by Layer deposition technique. He has coauthored more than 70 scientific papers (total citations: 3161, H-index 30, source google scholar) published in international peer reviewed journals and has attended several national and international congresses and workshops.

#### Andrea Dorigato

Andrea Dorigato is an Assistant Professor (RTDB) at the Department of Industrial Engineering at the University of Trento, Italy since 2016. His research activity focuses on the development of multifunctional polymer micro- and nanocomposites, with attention to the investigation of innovative materials for the sustainable development. His most significant research topics are related to the investigation of the viscoelastic and fracture behaviour of polymer nanocomposites and of electro-active multiscale polymer composites. His research also focuses on the development of innovative polymeric materials obtained from renewable resources and recycled materials with low environmental impact, of nanostructured foams for the thermal insulation and energy saving, and of phase transition materials for the thermal energy storage/release. The study of the Life Cycle Assessment (LCA) of engineering materials is also included in his research fields.

#### Gianmarco Griffini

Gianmarco Griffini is currently an Assistant Professor (tenure-track) of Science and Technology of Materials at the Department of Chemistry, Materials and Chemical Engineering "Giulio Natta" of Politecnico di Milano, Italy. He received his PhD degree in Materials Engineering in 2012, from Politecnico di Milano. Since 2014 he serves as a lecturer in the School of Industrial and Information Engineering (MSc and BSc degrees in Materials Engineering and Nanotechnology) and in the School of Design (MSc degree in Design & Engineering and MSc degree in Product Design for Innovation) at Politecnico di Milano. His current research interests are mainly related to the field of science and technology of materials and are focuses particularly on the development and characterization of polymeric materials and on the study of their structure-property relationships.

# Three Organic/Inorganic Nanolayers on Flexible Foam Allow Retaining Superior Flame Retardancy Performance Upon Mechanical Compression Cycles

### Federico Carosio\* and Alberto Fina

Dipartimento di Scienza Applicata e Tecnologia, Politecnico di Torino, Alessandria Campus, Alessandria, Italy

The water-based deposition of flame retardant coatings on flexible polyurethane foams has attracted great interest among the scientific community due to the great performances associated with this technology. Unfortunately, this approach results inefficient as it requires a high number of steps in order to achieve the desired properties. In this paper, we report the production of flame-retardant foams by means of the simple deposition of only three nanoparticles containing layers. The composition and order of the deposited layer has been designed in order to provide specific flame retardant actions, targeting the delayed release of polymer decomposition products to the gas phase, the dilution of these flammable products with water, and the intumescent barrier formation. The morphology of the coated foams after the adsorption of each layer has been investigated by scanning electron microscopy, demonstrating the ability of each adsorbed layer to completely wrap the complex 3D structure of the foam. This three layers-based coating produces a protective exoskeleton that is capable of self-extinguishing the flame in standard flammability tests, leaving the foam almost unaffected (final residue 98%). In forced combustion tests by cone calorimetry, treated foams showed considerably reduced combustion rates, with reduced peak of heat release rate (−50%) as well as consistent reduction in the smoke optical density (−51%) and the total smoke release (−34%). In addition, treated foams have been demonstrated to maintain the ability to self-extinguish the flame as well as reduced combustion rates and smoke production even after being subjected to 100 compression cycles.

Keywords: layer-by-layer, flexible polyurethane foams, boehmite, montmorillonite, ammonium polyphosphate, flame retardancy

# INTRODUCTION

Recent trends in materials science have clearly pointed toward the production of green and environmentally friendly material concepts capable of improving sustainability while also delivering optimized and competitive performances. To this aim, research has been focused on the exploitation of natural resources (e.g., nanocellulose, natural fibers etc.) and water-based processes in different application fields (Mohanty et al., 2002; Klemm et al., 2011). The fire safety of polymeric materials represents an area of concern; indeed, due to their organic nature, polymers can easily ignite after accidental exposure to small flames or irradiation at high heat

### Edited by:

Alessandro Pegoretti, University of Trento, Italy

#### Reviewed by:

José-Marie Lopez-Cuesta, Institut Mines -Télécom Mines Alès, France Mustapha Kaci, University of Béjaïa, Algeria

> \*Correspondence: Federico Carosio<EMAIL_ADDRESS>
#### Specialty section:

This article was submitted to Polymeric and Composite Materials, a section of the journal Frontiers in Materials

> Received: 30 November 2018 Accepted: 04 February 2019 Published: 27 February 2019

#### Citation:

Carosio F and Fina A (2019) Three Organic/Inorganic Nanolayers on Flexible Foam Allow Retaining Superior Flame Retardancy Performance Upon Mechanical Compression Cycles. Front. Mater. 6:20. doi: 10.3389/fmats.2019.00020 fluxes, leading to catastrophic fire events.  This clearly increased societal awareness on the problem of halogen-based FR and the need for safer and better performing alternatives. A possible solution is represented by the exploitation of a nanotechnological approach in which the surface of flammable items (i.e., fabrics and open cell polyurethane foams, PU), is modified with nanostructured coatings deposited by means of green water-based deposition processes (Malucelli et al., 2014; Alongi et al., 2016). Among the possible different approaches, the Layer-by-Layer (LbL) deposition technique represents one of the most widely employed surface modification tool to achieve flame retardancy properties (Holder et al., 2017). Nanostructured coatings are LbL-assembled on the surface following the alternate adsorption of oppositely charged species from aqueous based solutions/suspensions (Decher and Schlenoff, 2012). One of the key advantages of LbL is certainly its versatility in both the components selection and deposition parameters. Indeed, for virtually any combination of water-soluble polyelectrolytes and water-stable nanoparticles, it is possible to find suitable deposition conditions (e.g., time, temperature, concentration, pH) that allow for a steady coating buildup (Hammond, 2004). Thanks to its versatility, the LbL deposition has been applied to various substrates, including textiles, foams and thin films, and the effects of pH, ionic strength and molecular weight on the achieved FR properties have also been investigated (Laufer et al., 2012; Carosio et al., 2015c, 2018a). Although the concept of LbL-assembled FR coatings was first validated on fabrics, open cell foams have recently attracted great interest as a substrate to protect due to their inherently high flammability, which makes them one of the first items to be ignited during a fire (Ahrens, 2008). Different coating compositions and FR strategies have been investigated, mainly targeting the deposition of nanoparticlesrich coatings characterized by brick and mortar structures or the assembly of organic polyelectrolytes with intumescencelike behavior (Carosio et al., 2014b, 2018c; Smith et al., 2017). These two FR concepts have also been combined together using nanoparticles as reinforcing agents either within the intumescent coating or below it in stacked configurations (Cain et al., 2013; Holder et al., 2015). Treated foams usually achieve considerably reduced combustion rates and, in some cases, the ability to stop flame spread during flammability tests. Self-extinguishing typically requires a high number of deposition steps (usually >10), which is currently preventing the scale-up and large scale industrial exploitation of the LbL approach (Richardson et al., 2015; Carosio and Alongi, 2016). In addition, the durability of the LbL coating during repeated foam deformation has never been properly addressed. This represents another strong limitation of the state of the art, as one of the main applications of soft open cell PU foams is in upholstered furniture for buildings and transportations. In the present paper, we aim at the design of a coating composition and its assembly onto a soft PU foam in order to obtain highly performing and durable FR coatings, with only three depositions steps. To this aim, open cell PU foams have been treated, following the procedure schematized in **Figure 1**.

Each deposition step deposits a nanoparticle/polyelectrolyte layer with specific FR characteristics. The hypothesis is that the first layer made of polyacrylic acid and montmorillonite may produce a "brick and mortar" structure to reduce volatile feed to the gas phase, while providing stability to the remaining portion of the coating (Carosio et al., 2015a). In the second layer, boehmite (BOH) nanoparticles deposited within a poly(diallyldimethylammonium chloride) (PDAC) may release water during combustion (with volatile dilution and cooling effects) while also leaving an inorganic protective layer made of aluminum oxide (Rothon and Hornsby, 2014). The outermost layer made of ammonium polyphosphate (APP) and montmorillonite (MMT) acts as inorganic intumescent layer producing an expanded silicoalluminophosphate barrier (Alongi and Carosio, 2016). The resulting coating morphology has been evaluated by scanning electron microscopy while the resulting FR properties have been assessed by horizontal flammability and cone calorimetry tests. Furthermore, in an effort to evaluate the coating durability, treated foams have been subjected to deformation cycles evaluating the effects on coating morphology and FR properties.

# MATERIALS AND METHODS

# Materials

Open cell polyurethane foams (PU) with a density of 18 g/dm3 and thickness of 18 mm were purchased from a local warehouse. The as received foams were washed by dipping and vigorous squeezing in deionized water in order to remove dust and processing residues and then dried in a ventilated oven at the temperature of 70◦C. Poly(diallyldimethylammonium choride) (PDAC, average Mw 400,000–500,000, 20% wt in H2O), polyacrylic acid (PAA, solution average Mw ∼100,000, 35% wt. in H2O), Sigma-Aldrich. The sodium montmorillonite (MMT) was purchased from Southern Clays. Ammonium polyphosphate (PHOS-CHEKr P30) was purchased from ICL Performance Products, Inc. Bohemite nanoparticles (Dispal 10F4) with average dispersed size 240 nm (as declared by the producer) have been supplied by Sasol. Ultrapure water (18.2 M) supplied by a Q20 Millipore system was employed for the preparation of the solutions and suspensions employed in this work. Both PAA/MMT and PDAC/BOH have been employed to prepare a 3% wt suspension with 1:2 polyelectrolyte/nanoparticle ratio. The suspension was prepared by firstly diluting the polyelectrolyte solution and then adding the required amount of nanoparticles. APP and MMT have been employed with a 2:1 polyelectrolyte/nanoparticle ratio for the preparation of a 3% wt water suspension.

# Layer-by-Layer Deposition

The prewashed PU foam has been treated with a three-step deposition process as schematized in **Figure 1**. First the foam was dipped in the PAA/MMT bath (negative suspension), then

in the PDAC/BOH bath (positive suspension) and eventually in the APP/MMT bath (negative suspension). During each deposition step, the foam is dipped in the bath and vigorously squeezed in order to fill its internal structure with the suspension. After an immersion time of 10 min, the wet foam was removed from the bath, squeezed to remove the excess of suspension and dried in a ventilated oven at 70◦C. The process as shown in **Figure 1** deposits a three layer coating (PAA/MMT)/(PDAC/BOH)/(APP/MMT) on the PU structure. The mass gain, evaluated by weighting the samples before and after the LbL deposition, was 37 ± 3%. In the following, a sample coated with three layers is denoted as PU-3 L.

# Characterization

The morphology of untreated and treated PU foams was investigated using a LEO−1 450VP Scanning Electron Microscope (SEM) operating at 5 kV. Small pieces (∼10 × 10 × 5 mm3 ) of PU foams were positioned on conductive adhesive tape and gold coated prior to observation. Fourier Transformed-Infrared spectra in attenuated total reflectance (FT-IR ATR) were collected at room temperature in the range of 4,000–700 cm−1 (16 scans and 4 cm−1 resolution) using a Perkin Elmer Frontier spectrometer equipped with a diamond crystal. In order to preliminarily evaluate the coating durability during foam deformation, simple compression cycles have been performed on coated PU samples as schematized in **Supplementary Figure 1**. During the test the foam is placed between two 1 mm thick aluminum plates and then manually compressed by the action of a weight placed on top of the plate resulting in 16 kPa pressure1 . The compression is maintained for 5 s after which the weight is removed in order to let the foam recover its original shape. Then, after 10 s the weight is applied again in order to perform a second compression cycle. This cycle has been repeated 50 and 100 times. In the followings, samples coded as PU-3 L @50 c and PU-3 L @100 c refer to a coated foam subjected to 50 and 100 compression cycles, respectively.

Flammability tests were performed in horizontal configuration. 50 × 150 × 15 mm3 specimens positioned on a metallic grid were ignited by applying a 20 mm blue methane flame on the short side of the sample. The flame was applied for 6 s. During the test, the formation of incandescent droplets of molten polymer (melt dripping) was evaluated by placing dry cotton underneath the sample. The time, in seconds, a flame persisted on the sample after the igniting source has been removed was recorded as flaming time. The percentage of mass left at the end of the test is reported as final residue and was evaluated by weighting the sample before and after the test. Three different samples were tested for each formulation. Forced combustion tests were performed by cone calorimetry (Fire Testing Technology). 50 × 50 ×18 mm3 specimens were exposed to a 35 k W/m2 radiative heat flux. Average values concerning Time To

1 (i.e., ∼163 g/cm2 , roughly simulating the weight of a person, despite a full mechanical deformation study would require a more complex setup, well beyond the scope of this study).

Ignition (TTI), peak of Heat Release Rate (pkHRR), Total Heat Release (THR), average Specific Extintion Area (SEA), Total smoke release (TSR) and final residue were evaluated and are presented with their experimental deviations. Measurements were performed four times for each formulation. Prior to flammability and forced combustion tests, all specimens were conditioned in a climatic chamber (23 ± 1 ◦C 50% R.H) for 48 h.

# RESULTS AND DISCUSSION

# Coating Morphology

The changes in surface morphology and chemical composition of small samples of PU foam after each deposition step have been investigated by SEM and ATR spectroscopy. **Figure 2** reports collected micrographs and IR spectra. In **Figure 2** a sample coded as PU 1st layer denotes the deposition of only the (PAA/MMT) layer, PU 2nd layer denotes the deposition of the (PAA/MMT)/(PDAC/BOH) layers and PU 3rd layer denotes the deposition of the (PAA/MMT)/(PDAC/BOH)/(APP/MMT) layers (i.e. the complete coating structure).

The unmodified PU shows a typical open cell structure where cell walls generated during the foaming process are characterized by a smooth and even morphology (**Figure 2A**). The deposition of the PAA/MMT layer (8 ± 1% add-on) completely changes the foam walls morphology. Indeed, the formation of a thin and homogeneous coating containing MMT nanoplatelets embedded within a PAA matrix can be easily observed in **Figure 2B**. This is in agreement with previously reported LbL assemblies comprising polyelectrolytes and MMT (Carosio et al., 2014a). The deposition of the second PDAC/BOH layer (20 ± 3% total add-on) provides an additional change, and BOH nanoparticles are evenly distributed on every surface available producing a granular morphology as reported in high magnification micrographs in **Figure 2C**. The final deposition step adds the APP/MMT layer (37 ± 3% total add-on) and this corresponds to another change in surface morphology that appears to have higher rugosity (**Figure 2D**). It is worth observing that, starting from the first deposition step, the PAA/MMT coating completely covers all the surfaces available without displaying an island growth that is typical of LbL assemblies at early deposition steps. This can be ascribed to the good affinity between the PAA and the PU surface, where extensive hydrogen bonding between the urethane bonds and undissociated carboxylic acids of PAA may occur (Laufer et al., 2013). The stability and homogeneity of this first adsorbed layer is fundamental for the proper depositions of the subsequent layers. The resulting coating creates a complete protective exoskeleton that completely wraps the 3D structure of the foam without altering its open cell nature as reported in low magnification micrographs collected in **Supplementary Figure 2**. ATR spectroscopy has been also employed to evaluate the change in surface chemistry after each adsorption step (**Figure 2E**). Neat PU shows characteristic signals associated to C=O vibrations related to the urethane (broad peak centered at 1,730 cm−1 ), urea group (1,640 cm−1 ), N–H and C–N vibrations of the urethane and urea groups (band centered at 1,544 cm−1 ) and C–O–C asymmetric and

symmetric vibrations (1,226 and 1,094 cm−1 , respectively) (Gómez-Fernández et al., 2016). The neat coating constituents have been also analyzed, and collected spectra are reported in **Supplementary Figure 3**. Pure PAA shows characteristic C=O and C-O-C signals at 1,700 and 1,164 cm−1 , respectively (Dong et al., 1997). MMT displays the strongest peak centered at 1,000 cm−1 and ascribed to Si-O vibrations (Madejová, 2003). PDAC shows the main characteristic bands associated to C-H bonds in CH2 (1,470 cm−1 ) and N-C bonds (1,122 cm−1 ) (Sun et al., 2008). BOH asymmetric and symmetric Al-OH vibrations are found at 1,153 and 1,075 cm−1 , respectively (Haghnazari et al., 2014). This latter peak also accounts for Al-O vibrations while OH groups are evidenced by the peak at 738 cm−1 (Guido et al., 2013). APP shows characteristic peaks associated to P=O (1,245 cm−1 ), P-O (1,062 and 1,022 cm−1 ), P-O− (880 cm−1 ), and N-H of ammonium ions (1,423 cm−1 ) (Socrates, 2006). When the first PAA/MMT layer is deposited on PU foams, the characteristic peaks of the substrate change consistently, evidencing the formation of new peaks at 1,715 and 1,048 cm−1 ascribed to C=O stretching in PAA and Si-O-Si vibrations in MMT, respectively. The deposition of the second layer results in the appearance of signals associated to CH2 vibration (1,476 cm−1 ) in PDAC as well as Al-O/Al-OH (1,072 cm−1 ) and O-H vibration (744 cm−1 ) in BOH. The presence of the last layer is evidenced by signals mainly associated to APP such as P=O (1,243 cm−1 ), P-O (1,062 cm−1 ), P-O− (914 and 875 cm−1 ) and N-H (1442 cm−1 ), while MMT contributes with the signal at 1,030 cm−1 . It is worth mentioning that in the assembled layers many signals associated to polar functional groups of adsorbed species are shifted to different wavenumbers with respect to the neat products. The most apparent is the P-O- signal of APP that appears with two distinct peaks at 914 and 875 cm−1 , likely ascribed to P-O- interactions with PDAC and ammonium ions, respectively. The Si-O signal of MMT appears at higher wavenumbers as a consequence of hydrogen bond interactions with the co-adsorbed polyelectrolytes (i.e., PAA and APP; Chen and Zhang, 2006). In order to evaluate the durability of the coating, the foams treated with three layers have been subjected to repeated compression tests as described in the Materials and Methods section.

The morphology of the coating after 50 and 100 compression cycles has been evaluated by SEM, see **Figure 3** for collected micrographs. The deformation cycles impart strong deformations to the deposited coating that must comply to the change in shape of the substrate. Due to the ionic bonds in between the adsorbed layers and the nanoparticle-based structure, the deposited coating is inevitably rigid and the cyclic deformation results in the formation of cracks where portion of the assembly are partially detached from the substrate (**Figures 3A,B**). Interestingly, no visible differences in terms of damage extent can be appreciated when comparing the coating morphologies after 50 and 100 deformation cycles. This suggests that the cracks that are clearly visible in **Figures 3A,B** are generated during the first compression cycles in order to allow the foam walls to completely collapse upon compression. SEM investigations performed on PU-3 L samples after the first compression cycle support this as reported in **Supplementary Figure 4** where low magnification micrographs clearly show extensive cracks formation with minimal detachment of the coating. This is also confirmed by practical observations of the coated foams that are firmer after the deposition of the coating, but after few deformation cycles no significant qualitative differences can be observed when compared to the untreated foam. Such behavior is in agreement with previously reported literature (Li et al., 2013) where LbLtreated foams partially recovered the original stress/strain curve after few compression cycles. This is particularly important for the application in upholstered furniture and seats, allowing to preserve the comfort feeling for which the pristine foam was designed.

## Flammability

Since PU foams represent one of the first items to be ignited during the start of a fire, it is mandatory to evaluate the flammability properties by means of horizontal flame spread tests. This test simulates the accidental exposure of the foam to a small flame thus allowing for the evaluation of the potential fire starting capability associated with the sample. **Figure 4** summarizes collected flammability parameters as well as snapshots of neat and coated foams during the test. Numerical

parameters with their standard deviation are reported in **Supplementary Table 1**.

The uncoated PU foam can be easily ignited by the methane flame. When the ignition source is removed, the flames grow in size and rapidly spread to the entire length of the sample, completely consuming it (**Figure 4A**). During combustion, a vigorous melt dripping ignites the dry cotton placed underneath the sample. This phenomenon is typical of flexible PU foam and is related to the formation of flaming droplets of molten polymer that pose additional fire risks as not only the foam can be quickly ignited, but also the melt dripping can easily spread the fire to other ignitable materials. The presence of the threelayer coating completely changes this burning behavior. Indeed, upon removal of the ignition source, only a small flame remains confined in the first 2 cm of the foam, constantly reducing in size until it extinguishes within few seconds. No melt dripping is observed, and the residue collected at the end of the test accounts for 98% of the original mass. A nearly identical behavior is observed in coated foams subjected to 50 and 100 deformation cycles. Only a slight increase in the average flaming time was observed for compressed foams, meaning that self-extinguishing takes slightly longer to occur, but the deviation of burning times obtained for different specimens is considerably increased, suggesting a dependency on local density of defects. Nevertheless, all coated foams are capable of stopping flame spread and suppress the melt dripping phenomenon, yielding final residues within 97–98%. These results clearly point out the improved fire safety achieved by the deposited coatings. Indeed, thanks to the engineered structure of the three layers, the protective exoskeleton is able to prevent foam melting and collapsing while at the same time considerably reducing the release of volatile gasses that feed and sustain the flame. The resulting FR action is very efficient as self-extinguishing behavior is achieved with only three deposition steps. In addition, it seems that the formation of cracks during deformation cycles had rather limited effects on the FR performances of the coated foam.

# Cone Calorimetry

Cone calorimetry tests allow to evaluate the reaction of untreated and treated foams when exposed to heat fluxes typically obtained in developing fires, thus assessing the contribution of the samples to fire growth and spread. During the test, samples are exposed to 35 kW/m2 (i.e., heat flux related to early stages of a developing fire; Schartel and Hull, 2007), which triggers thermal degradation processes and the subsequent release of flammable volatiles that, when the proper concentration is reached, can be ignited by a spark igniter leading to flaming combustion. The instrument measures parameters associated to heat and smoke release. **Figure 5** collects HRR plots and histograms of the main parameters evaluated during the test

while numerical values with their standard deviation are reported in **Supplementary Table 2**.

Upon exposure to the cone calorimetry heat flux, the neat PU foams show a quick ignition (TTI=2 s) followed by the foam structural collapse and the formation of a molten pool of a low viscosity liquid. This latter step produces a steeply increase in heat release rate that reaches the maximum value of about 300 kW/m2 (**Figure 5A**, HRR vs. time plots). The behavior of flexible PU foam during cone testing is well documented and the collapsing and melting stages are normally associated to increased fire spreading risks (Kramer et al., 2010). The presence of the three-layers coating completely changes the burning behavior of coated PU foams. Indeed, upon ignition no structural collapse is observed and the foams maintain their original shape at the end of the test (see **Supplementary Figure 5**). Heat release rates are considerably reduced as demonstrated by a −50% in pkHRR value (**Figure 5B**, pkHRR histograms); on the other hand, the effect on THR values is rather limited, suggesting combustion of the PU is complete, yet delayed in time. Interestingly, a positive effect on smoke related parameters is also observed. Indeed, smoke optical density (SEA) and total smoke produced (TSR) are considerably reduced: −50 and −34%, respectively (**Figures 5D,E**). The final residue is increased to 25% of the starting mass. These results point out that the presence of the coating not only can efficiently reduce the rate of combustion and potential fire spread risk, but can also influence the production of smoke. Such observation is of crucial importance since it is reported that people are normally confronted with smoke and its effects during evacuation (Frantzich, 1994). The presence of dense smoke considerably increases escaping times while also causing loss of ability to react, unconsciousness, slower walking speed and long-term damage such as lung cancer (Gann, 2004). For these reasons, smoke-related effects represent an important cause of fire fatalities. Furthermore, coated foams subjected to 50 and 100 deformation cycles maintain performances similar to the

non-compressed PU-3 L. Despite a slight detrimental effect, proportional to the number of deformation cycles observed, pkHRR values still maintain a reduction ≥40% while average smoke parameters remain within the experimental error of the PU-3 L. It seems that the cracks produced during deformation can affect the efficiency of the coating in reducing volatile release while the smoke diluting action exerted by BOH nanoparticles is mostly unaffected.

The residues collected after cone calorimetry tests have been observed by SEM (**Figure 6**). Low magnification micrographs suggest that the original 3D structure of the PU foam is left almost undamaged (compare **Figure 6** with **Supplementary Figure 2**). By increasing the magnification it is possible to reveal the real structure of the residue that mostly comprises a hollow protective exoskeleton resulting from the deposited coating. The morphology has clearly changed during combustion due to the different processes occurring within the coating (i.e., PAA degradation, BOH water release and APP decomposition) resulting in the formation of a compact barrier. The presence of cracks on the surface can be related to the release of volatiles produced by the PU that is almost completely consumed during combustion as demonstrated by the negligible change in THR values. The achieved results have been compared with other FR systems deposited on open-cell flexible PU foams by the LbL deposition of inorganic nanoparticles, natural polysaccharides or carbon nanotubes, as presented in **Figure 7**, where the best reduction in heat release rate peak achievable with each coating composition is plotted as function of the number of deposited layers (Laufer et al., 2012; Carosio et al., 2014b, 2015b, 2018b; Wang et al., 2014; Haile et al., 2016; Holder et al., 2016; Mu et al., 2017; Smith et al., 2017; Lazar et al., 2018). Red or green symbols refer to flammability and account for the ability to selfextinguish the flame and prevent its spread to the entire surface of the sample.

It can be easily appreciated that most LbL coatings can achieve substantial reductions in pkHRR (in the 40–70% range) regardless of the deposited layer number. However, in order to grant self-extinguishing behavior during flammability tests, a minimum number of 10 deposited layers is needed. This considerably reduces the efficiency of the deposited

LbL coatings and their potential applicability. On the other hand, in this paper we demonstrate that the proper design of the layer composition and the coating structure allows to achieve substantial reductions in pkHRR values (50– 40%) while also obtaining self-extinguishing behavior. In addition, these characteristics are maintained after 100 deformation cycles thus confirming the potentialities of the developed solutions.

# CONCLUSIONS

In this manuscript, we reported the preparation of efficient and durable layer-by-layer coatings based on the deposition of only three layers of nanoparticles containing polyelectrolytes. The deposition procedure produces a protective exoskeleton that completely covers the complex 3D geometry of the foam. The layer constituents and position have been selected on the basis of the desired flame retardant mechanism. The first layer that comprises PAA and MMT nanoplatelets produces a first barrier to volatile release. BOH nanoparticles contained in the second layer can release water with gas phase diluting and cooling effects while leaving a thermally insulating layer made of Al oxide. This latter effect is further enhanced by the third layer where APP and MMT can react together in order to produce an inorganic intumescent layer. The joint action of these three layers can stop flame spread during horizontal flammability tests leaving most of the foam undamaged (final residue up to 98% of the initial mass). Combustion rates are considerably lowered as demonstrated by a 50% reduction in pkHRR during cone calorimetry tests. Smoke production is also efficiently addressed with coated foams showing reduced smoke optical density and total smoke released. This latter aspect can prove fundamental during the evacuation of buildings in case of fire. The durability of the achieved FR performances has been

# REFERENCES

evaluated by means of compression cycles. To this aim, treated foams were compressed cyclically to maximum deformation for 50 and 100 times and then their FR properties have been evaluated again. Only minor differences have been observed for the compressed foams that still exhibited excellent FR properties thus confirming the great efficiency and potential durability of the proposed coating solution. In summary, thanks to an efficient coating composition and by relying on a simple water-based approach it is possible to prepare foams capable of improving the fire safety in several applications where flexible polyurethane foams are used, including upholstered furniture and seats for households, transports, public buildings, theaters/cinemas etc.).

# AUTHOR CONTRIBUTIONS

FC conceived the experiments, carried out the LbL deposition and the characterization. AF contributed to the interpretation of results and the revision of the manuscript. The manuscript was mainly written by FC.

# SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fmats. 2019.00020/full#supplementary-material

treated in the presence of boehmite nanoparticles. Polym. Degrad. Stab. 98, 1609–1616. doi: 10.1016/j.polymdegradstab.2013.06.021

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

Copyright © 2019 Carosio and Fina. 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.

# Thermo-Mechanical Behavior of Novel Wood Laminae-Thermoplastic Starch Biodegradable Composites With Thermal Energy Storage/Release Capability

Andrea Dorigato\*, Giulia Fredi and Alessandro Pegoretti

*Department of Industrial Engineering and INSTM Research Unit, University of Trento, Trento, Italy*

For the first time, multifunctional fully biodegradable composites combining structural and thermal energy storage/release capabilities were prepared and thermo-mechanically characterized. Within this scope, thin beech laminae impregnated with a phase change material (PCM) represented by poly(ethylene glycol) (PEG) were interleaved with thin foils of thermoplastic starch (TPS) and consolidated by hot pressing. From scanning electron microscopy, it was observed that a certain amount of PEG (about 11 wt% of the total laminate) remained entrapped within the wood pores, negatively affecting the interfacial adhesion between wood laminae and TPS foils. The presence of PEG stabilized in the wood laminae was confirmed by differential scanning calorimetry tests, in which a specific melting enthalpy of 27.4 J/g was detected with a melting peak at 55◦C. Wood permeation with PEG was responsible for an increase of the dynamic moduli E′ and E′′, as well as of the tensile and Charpy impact strength of the laminates. Therefore, this paper highlighted the possibility of developing multifunctional fully biodegradable composites capable of combining structural and thermal energy storage properties, in which the selected PCM positively contributed to the mechanical behavior of the laminates.

Keywords: starch, ultrathin wood laminae, composites, poly(ethylene glycol), energy storage

# INTRODUCTION

It is well-known that interest in fiber reinforced polymers (FRPs) has noticeably grown over the last decades (Gay et al., 2003), mostly because of their elevated specific mechanical performance, easy processability and corrosion resistance. On the other hand, various environmental concerns have arisen regarding the use of traditional reinforcing fibers (such as glass, carbon and aramid) in FRPs (Pegoretti, 2007) as life cycle assessments on these materials have evidenced their limits in terms of recyclability and final disposal operations (Dorigato and Pegoretti, 2012). A partial solution to these concerns could be represented by the use of natural fibers (i.e., kenaf, hemp, flax, cellulose, etc.) as reinforcement of both thermoplastic and thermosetting matrices (Wu, 2009). However, even though with these fibers it is possible to reach most of the physical properties of conventional laminates (García et al., 2008; Iwatake et al., 2008; Kunanopparat et al., 2008), it is very difficult from a technological point of view to perform efficient thermo-mechanical recycling of these materials because of the limited thermal degradation resistance of natural reinforcements

Edited by:

*Luca Valentini, University of Perugia, Italy*

#### Reviewed by:

*Debora Puglia, University of Perugia, Italy Loris Giorgini, University of Bologna, Italy*

\*Correspondence: *Andrea Dorigato<EMAIL_ADDRESS>
#### Specialty section:

*This article was submitted to Polymeric and Composite Materials, a section of the journal Frontiers in Materials*

> Received: *30 November 2018* Accepted: *05 April 2019* Published: *26 April 2019*

#### Citation:

*Dorigato A, Fredi G and Pegoretti A (2019) Thermo-Mechanical Behavior of Novel Wood Laminae-Thermoplastic Starch Biodegradable Composites With Thermal Energy Storage/Release Capability. Front. Mater. 6:76. doi: 10.3389/fmats.2019.00076*

**486**

(Houshyar and Shanks, 2003). In this sense, a possible solution to these environmental issues could be represented by the development of fully biodegradable composites, in which both the constituents (i.e., the matrix and the reinforcing phase) can be degraded under certain environmental conditions.

Among the natural reinforcements that can be used for the production of fully biodegradable composites, wood fibers represent one of the most interesting options (Bledzki et al., 2002). Wood is mainly constituted by carbon, hydrogen and oxygen in the form of cellulose (40–50 wt%), hemicellulose (20 wt%), lignin (about 35 wt%) and pectic substances in limited concentrations (Core et al., 1979). From a mechanical point of view, wood behaves like a heterogeneous anisotropic material. Moreover, in the low deformation regime it can be considered as an orthotropic material (Green et al., 1999), and its mechanical performance is strongly influenced by the presence of nodes and microstructural defects. Moreover, considering that wood can absorb up to 25–30 wt% of water, the environmental conditions strongly affect the mechanical response of woodbased materials.

From a microstructural point of view, the cross section of wood-based materials is formed by a regular distribution of cellular lumen with isolated and clustered pores and thick-walled fibers. In the case of beech, large parenchymal rays are generally distended along growth ring boundaries, while the concentration of gum inclusions, salt crystals or lattices is rather limited (Giordano, 1981). In the last years, particular scientific and technological attention was given to wood plastic composites (WPCs), constituted by wood fibers or wood flour compounded with thermoplastic matrices, such as polypropylene (PP), polyethylene (PE), and polyvinylchloride (PVC) (Oksman and Sain, 2008). These materials can be processed through the traditional technologies of thermoplastics (i.e., extrusion, injection molding). However, the mechanical properties of WPCs are generally considerably poorer than those of structural wood components, even at elevated filler amounts, especially under severe environmental conditions (Srimalanon et al., 2016). Quite surprisingly, the possibility to utilize thin wood laminae to develop novel thermoplastic wood composites laminates did not receive much attention in the open literature.

Among biodegradable matrices that could be potentially applied in fully biodegradable composite laminates, thermoplastic starch (TPS) has recently drawn considerable attention. From the chemical point of view, starch is a polysaccharide with a semicrystalline structure, composed by linear amylose and highly branched amylopectin (Bogracheva et al., 2002), and it can be processed through extrusion and injection molding. In order to overcome the problems related to its elevated hydrophilicity, some attempts were made to modify the TPS chemical structure through acetylation (Chi et al., 2008; Ayucitra, 2012) or upon blending with other biopolymers, such as polycaprolactone (PCL) and polyhydroxybutyrate-co-valerate (PHBV) (Averous et al., 2000), or with synthetic plastics, like polyethylene. The most widespread TPS-based products are commercially known as Mater-Bi R (Novamont, Italy) (Bastioli, 1997) and Bioplast R (Biotec, Germany) (Lorcks, 1998).

Two examples of biodegradable laminates have been recently developed by our group. In the first case, polyvinyl alcohol (PVOH)-based composites were prepared by a hand layup procedure, impregnating ultrathin beech laminae with an aqueous solution of PVOH (Dorigato et al., 2018b), while in a second paper thermoplastic composites reinforced with wood laminae were produced through a hot-pressing process by using thermoplastic starch foils (Dorigato et al., 2018a). In both of these papers, it was found that the resulting materials were characterized by a peculiar microstructure, in which the porosity of the beech laminae was occluded by the polymer matrix. The consequence of the matrix permeation within the wood pores was that the resulting laminates showed higher elastic modulus and tensile strength with respect to their constituents.

It is challenging to impart also multifunctional properties to fully biodegradable composites. In fact, the scientific and industrial interest toward materials able to combine different properties (structural, electrical, magnetic, thermal, etc.) has strongly increased in the latest years (Gibson, 2010). In this scenario, polymer–matrix composites are particularly suitable to be designed as multifunctional materials (Salonitis et al., 2010) because they combine the peculiar properties of polymers, such as the low density, good processability and cheapness, with the elevated stiffness and strength of the reinforcing fibers (Mallick, 2007). According to the authors' opinion, one of the most interesting options could be to prepare fully biodegradable laminates with thermal energy storage (TES) capability. In fact, TES technology is based on the storage of heat and its release when needed, thus reducing the gap between thermal energy availability and request (Pielichowska and Pielichowski, 2014). Therefore, TES systems can be utilized to accumulate excess or waste heat, like in solar thermal power plants, to recover waste industrial heat (Khadiran et al., 2015), to control the temperature inside the buildings (Onder et al., 2008), in smart textiles for body temperature regulation (Sobolciak et al., 2016), or in the cooling systems for electronic devices (Luyt and Krupa, 2009). Multifunctional composites combining good mechanical properties and TES capability could thus be applied were weight and volume saving requirements should be coupled with energy storage capability and/or temperature control operations (Rigotti et al., 2018). TES systems are generally classified depending on the underlying process in sensible, thermochemical, or latent heat-storage systems. Among them, phase change materials (PCMs) undergo an endo/exothermic solid-solid or solid-liquid transition, and can be classified as organic, inorganic or eutectic (Sari et al., 2005). Thanks to their elevated specific phase change enthalpy, narrow transition temperature interval, limited supercooling, lightness and cheapness, organic PCMs, i.e., paraffins, poly(ethylene glycol) (PEG) and fatty acids, are probably the most widely diffused (Dorigato et al., 2017; Fredi et al., 2017). In particular, PEG is a versatile and biodegradable PCM with elevated melting enthalpy, tunable transition temperature range, elevated thermal stability, good biocompatibility, non-toxicity and noncorrosiveness (Sundararajan et al., 2016; Kou et al., 2019). Because of these peculiar properties, researchers have recently put much efforts to synthesize novel PCMs based on PEG, trying to physically blend PEG with other supporting materials, to perform chemical modifications to prepare PEG based solid/solid PCMs or to impregnate PEG in porous materials (Wang et al., 2016; Yang et al., 2016; Zhou et al., 2018). Nowadays, most of the studies on PEG-based PCMs are focused on the preparation of novel PCMs with improved thermal energy storage capability. The problems associated with the poor thermal conductivity of organic PCMs and their need to be confined to limit their leakage at the molten state (Kastiukas et al., 2016) could be partially solved through their encapsulation in organic or inorganic shells (Ozaki et al., 2008), or upon the shape stabilization in a polymer matrix (Yoo et al., 2016), in a layered/porous structure or in a nanostructured filler (Alexandre et al., 2009). Moreover, the use of a conductive nanofiller could also improve the thermal conductivity of the resulting systems. Due to its porous microstructure, it is possible to hypothesize that wood could also be an interesting candidate for the shape stabilization of organic PCMs at the molten state.

Quite interestingly, only a few investigations have been performed on the possibility of preparing multifunctional PMCs combining TES and structural features. For example, our group recently developed novel epoxy/carbon fiber composites containing carbon nanotube (CNT)-stabilized paraffinic wax (Fredi et al., 2017, 2018a), epoxy/discontinuos carbon fibers with microencapsulated paraffin (Dorigato et al., 2019), thermoplastic polyamide 12 (PA12)/glass fiber composites containing both microencapsulated and shape stabilized PCMs, reinforced with long glass fibers (Fredi et al., 2018a) or with discontinuous carbon fibers (Dorigato et al., 2018c; Fredi et al., 2019), or thermoplastic acrylic Elium R /carbon fiber laminates containing microencapsulated paraffin (Fredi et al., in preparation). However, no papers on the production of fully biodegradable composites with good mechanical properties and TES capability can be found in the open literature. On the basis of these considerations, the aim of the present paper is to develop for the first time multifunctional fully biodegradable composite laminates, constituted by a thermoplastic starch matrix reinforced with thin wood laminae, in which the thermal energy storage function can be obtained by impregnating wood laminae with PEG in the molten state. The microstructural, thermal and mechanical behavior of the resulting laminates was investigated, considering also the effect of the orientation of the reinforcing phase within the composites. The properties of the resulting laminates were also compared with those of their constituents.

# EXPERIMENTAL

# Materials

Thermoplastic starch MaterBi R EF51V was supplied by Novamont Spa (Novara, Italy) in the form of foils with a thickness of 100µm. This TPS grade is characterized by a density of 1.2 g/cm3 , a MFI of 4 g/10 min (190◦C, 2.16 kg) and a melting temperature of 167◦C.

The employed poly(ethylene glycol) was PEG 2000 purchased by Alfa Aesar (Karlsruhe, Germany). It has a molecular weight of 1,800–2,200 g/mol and a melting temperature of 51–55◦C, according to the producer's datasheet.

For the preparation of the thin wood laminae (**Figure 1A**), a beech plant with a density of about 0.73 g/cm3 if conditioned at 23◦C and a relative humidity of 50% was selected. The selection of this wood species is motivated by its regular morphology and anatomical structure, as described in the Introduction. Thin wood laminae, with a thickness of about 320µm, were kindly provided by the CNR Ivalsa Institute (San Michele all' Adige,

FIGURE 1 | Representative images of (A) neat wood lamina; (B) wood laminae after PEG impregnation; (C) wood/PEG/starch laminate.

Dorigato et al. Starch-Wood-PEG laminates for TES

Italy). The laminae were obtained through a Marunaka Super Meca-s planer (Marunaka Tekkosho Inc., Shizuoka, Japan) from beech boards, conditioned at a temperature of 23◦C and a relative humidity of 50%. Although this machine allows the production of even thinner laminae, the choice of the thickness was motivated by the need to reach a compromise between the flexibility of the thinner laminae and their PEG storage capability.

# Preparation of the Samples

Wood laminae were cut in square sheets of 120 × 120 mm2 , with the wood fibers oriented either parallel or at 45◦ with respect to the square edge. The sheets were immersed in the molten PEG at 70◦C for 5 min, then taken out and immediately gently blotted with absorbent paper to remove excess molten PEG from the surface (**Figure 1B**). Each wood lamina was weighed before and after immersion in liquid PEG in order to measure the weight fraction of absorbed PEG. Preliminary experiments had shown that a soaking time of 5 min was enough to reach the maximum weight increase due to PEG absorption, as a longer soaking time (up to 24 h) did not result in a greater weight increase. The wood/PEG laminae were interleaved with starch foils of 120 × 120 mm2 and compacted in a hot-plate press at 170◦C for 7 min, with an applied pressure of 1 MPa. Each laminate (**Figure 1C**) consisted of 5 wood/PEG laminae interleaved with 6 starch sheets. The samples were produced following two stacking sequences, i.e., [05] (unidirectional) and [+45/−45/+45/−45/+45] (angle ply). Laminates with virgin wood laminae (without PEG impregnation) were also prepared for comparison. Moreover, the mass of each constituent of each laminate was measured before the hot-pressing stage and compared to the final mass of the prepared laminates to study the PEG leakage during the laminate preparation. Specimens with different orientations were cut out of the prepared laminates for the subsequent characterization. The characterization was performed also on the constituents, i.e., the starch foils, the virgin wood laminae and the PEG-impregnated laminae. The list of the prepared samples is reported in **Table 1**. All the composite laminates with their relative constituents were stored at 23◦C with a relative humidity of 50%. It was decided to condition both the constituents and the laminates at this temperature and humidity level because these are the typical hygrothermal conditions of the laboratory. In this way, the influence of the temperature and of the moisture absorption during the lab tests could be minimized. It could be also interesting to point out that this production process, which consists of embedding the PCM directly in the reinforcing phase and employing matrix foils, allows the preparation of thermoplastic laminates without a melt compounding step, which could partially degrade the PCM (Fredi et al., 2018a). Therefore, the pristine thermal energy storage properties of PEG should be retained.

# Experimental Techniques

The microstructural features of the prepared samples were investigated through a scanning electron microscope (SEM) Jeol IT300, operating at an acceleration voltage of 15 kV. The samples F, FP, FA, and LAM were cryofractured in liquid nitrogen, and the fracture surface was analyzed after Pt-Pd sputtering.

TABLE 1 | List of the prepared samples.

Thermogravimetric analysis (TGA) was performed through a TA Q5000 IR thermobalance. The test was performed on the constituents (A, F and PEG), on the PEG-impregnated wood lamina (FP) and on both the prepared laminate types (FA and LAM). Specimens of ∼20 mg were cut out of all the prepared samples and tested at 10◦C/min up to 700◦C, under a nitrogen flow of 150 ml/min. These tests allowed the measurement of the temperatures corresponding to a mass loss of 1 wt% (T1%) and 5 wt% (T5%), the temperature corresponding to the initial stages of degradation (Tonset), determined with the tangent method, the peak temperature of the mass loss derivative (Td), indicating the maximum degradation kinetics, the initial mass loss before Tonset (mi) and the residual mass after the test (mr).

Differential scanning calorimetry (DSC) was performed with a Mettler DSC30 instrument. The test was performed on the constituents (A, F, and PEG), on the PEG-impregnated wood lamina (FP) and on both the prepared laminate types (FA and LAM). Specimens of ∼10 mg were cut out of all the prepared samples and tested at 10◦C/min in the temperature interval between −20 and 100◦C, under a nitrogen flow of 100 ml/min. All the specimens were subjected to a first heating scan, a cooling scan and a second heating scan. In this way, the melting and crystallization temperatures (Tm, Tc) and enthalpy (1Hm, 1Hc) of the PEG and of the starch phases were determined. An effective weight fraction of PEG inside the laminates was then calculated from the measured phase change enthalpy. The test also evidenced the presence of irreversible phenomena, such as the evaporation of the absorbed water.

The viscoelastic properties of the samples were investigated through dynamic-mechanical analysis (DMA), by using a DMAQ800 machine, provided by TA Instruments (New Castle, DE, USA). The test was performed in tensile mode, between 0 and 150◦C at 3◦C/min, with a maximum deformation of 0.05% applied at a frequency of 1 Hz. The test was performed on the

TABLE 2 | Mass and weight fraction of the constituents of the sample FP and the two laminates (FA and LAM).

samples F, A, FP, FA and LAM, in the L and T directions (where applicable), on rectangular specimens with dimension of 25 × 5 mm2 , while the thickness was that of each sample.

Quasi-static tensile tests were performed with a universal dynamometer Instron 5969, equipped with a 50 kN load cell, according to ISO 527 standard. Rectangular specimens with inplane dimensions of 120 × 10 mm2 were cut from the samples F, A, FP, FA, and LAM and tested at 5 mm/min until failure. The samples FA and LAM were tested in L, T and O directions, while the samples F and FP were tested in the L and T directions. Five specimens were tested for each sample and for each considered direction. The stress-strain curves allowed the determination of the tensile strength (σb), and of the strain at break (εb). While testing the wood laminae with and without PEG (samples F and FP), small plywood blocks were interposed between the specimen and the grips to avoid damaging the sample due to the contact with the hard steel grips. For the samples LAM and FA, five additional specimens for each of the considered directions (L, T, O) were tested at 0.25 mm/min for the measurement of the elastic modulus by using a resistance extensometer Instron 2620 with a gauge length of 12.5 mm. The elastic modulus was measured as the secant modulus of the stress/strain curves between the strain values of 0.05% and 0.25%. The elastic modulus was calculated in the same way also on the stress-strain curve of the other specimens.

Finally, Charpy impact tests were performed on the samples FA and LAM in the L and T directions, by using a Ceast 6549 impactor equipped with a hammer with an arm length of about 25 cm and a mass of 0.48 kg, positioned at a starting angle of 150◦ . In this way, the tests were performed at an impact speed of 2.9 m/s, with a maximum impact energy of 2 J. A thin plasticine layer was used to damp the vibrations to avoid the inertial peak in the force-displacement signal. Tests were performed on rectangular specimens having nominal dimensions of 120 × 10 × 2 mm3 , with a span length of 40 mm. At least five specimens were tested for each sample. In this way, the values of maximum force (Fb) and of the specific energy absorbed under impact conditions (Usp) were determined.

# RESULTS AND DISCUSSION

**Table 2** reports the weight composition of a single wood lamina impregnated with PEG (FP) as well as the nominal mass of each component before the hot-pressing step and the nominal weight composition for each laminate (FA and LAM). It can be observed that the quantity of PEG absorbed by a single lamina is ∼86 wt% of its initial mass. On the other hand, if five of these laminae are alternated to 6 starch foils, the experimental mass of the laminate obtained through hot-pressing (LAM) is ∼25 wt% lower than the sum of its components. This mass loss can be almost entirely attributed to the leakage of PEG out of the wood laminae, as the mass loss of the laminate without PEG (FA) is almost negligible. Therefore, the effective PEG weight fraction in the laminate is ∼11 wt%. This leakage suggests that a considerable fraction of the absorbed PEG has just solidified on the surface of the laminae and is not properly confined in the wood porosity, and the gentle blotting is not sufficient to remove excess PEG. However, the major part of PEG was successfully stabilized into the wood laminae.

**Figures 2A–F** shows the SEM micrographs of the cryofracture surface of the prepared samples. Since the cross section was investigated with this technique, in the samples with longitudinal fibers (L) the wood fibers are perpendicular to the fracture plane, while they are parallel in the samples with transversal fibers (T). In **Figures 2A,B** FESEM micrographs of the wood laminae cut in the transversal and in the longitudinal direction are respectively reported. As already shown in our previous papers on ultra-thin wood laminae impregnated with poly(vinyl alcohol) (Dorigato et al., 2018b), in the longitudinal direction the peculiar structure of the wood can be appreciated, with tracheids, wood fibers and vessels, while the parenchymal cells are less clearly visible. Tracheids and wood fibers are characterized by a mean diameter of about 10µm, while wood vessels present an average size of 40µm. The presence of some tracheids in radial direction can be appreciated in the transversal section. From the micrographs of the laminae impregnated with PEG (**Figures 2C,D**), it can be observed that PEG partially fills the wood porosity. This could be explained by assuming that the volume fraction of the closed porosity is not negligible. Moreover, the viscosity of the molten PEG is too high to allow the smallest pores to be filled in the time window of the impregnation process (up to 24 h). Moreover, the surface of the lamina hosts a considerable amount of PEG, as suggested by the smoothness of the surfaces of the laminae containing PEG (indicated by arrows), if compared to the remarkably rough surfaces of the neat wood laminae. This may complicate the adhesion between the wood laminae and the starch matrix, an issue partially visible from the fracture surface of the two laminates

(**Figures 2E,F**). In the laminate containing PEG (**Figure 2F**) the interlaminar adhesion seems worse than that of the neat laminate, as the cryofracturing operations caused delamination (indicated by the arrow), which is not observed in the neat laminate. However, this apparent worsening in the interlaminar adhesion due to the PCM does not negatively affect the mechanical

TABLE 3 | Results of the TGA tests on the prepared samples.

*T*1%*, T*5% = *temperature corresponding to a mass loss of 1 and 5 wt%.*

*Td* = *temperature corresponding to the maximum degradation rate (peak of the mass loss derivative signal).*

*Tonset* = *temperature corresponding to the initial stages of degradation, after the water evaporation.*

*mi* = *initial mass loss before Tonset.*

*mr* = *residual mass after the test.*

and impact properties, as reported hereafter. In future works on these systems it could be important to perform a detailed characterization of the interlaminar adhesion conditions of the prepared composites, as the interlaminar fracture toughness is one of the key aspects in the assessment of the mechanical behavior of composite materials.

TGA analysis was performed to evaluate the resistance to the thermal degradation of the laminates and of each constituent. **Figure 3** shows the TGA thermograms (with the residual mass and the derivative of the mass loss) of the sample LAM and of its three constituents, i.e., PEG, F, and A, while **Table 3** reports also the results on the sample FP. For the sample F, the initial mass loss (mi) of ∼5.8% is imputable to the absorbed water. This behavior is also observed in the samples FP and LAM, but to a lesser extent, as for these samples the wood laminae are only a fraction of the total mass. After water evaporation, the thermal degradation of the wood phase becomes appreciable above 250◦C and occurs in two main steps. Conversely, no water was detected on the samples A and PEG, which indicates that the total water fraction of the composites is entirely absorbed by the wooden phase. The thermal degradation of these two polymer phases starts to be evident at 320– 350◦C and reaches the maximum rate at ∼400◦C. From these results, it can be concluded that the thermal stability of all the constituents is sufficient to withstand the processing temperature of 170◦C. Moreover, as expected, the degradation thermogram of the sample FP lies between those of its constituents, i.e., F and PEG. Also, the thermal degradation of the sample LAM shows all the degradation steps of its constituents, as it can be better appreciated from the mass loss derivative signal. From TGA/DTG thermograms reported in **Figure 3**, it is also evident that the degradation of PEG and A samples is almost complete at 500◦C, while the F and LAM samples present a last degradation step between 500 and 550◦C, probably due to the degradation of some wood constituents in inert atmosphere. A detailed analysis of this degradation stage, which occurs far above the application temperature interval of the prepared laminates, is out of the scope of this work, and it could be better investigated in the future developments of this work. However, as reported in **Table 3**, the residual mass values at 700◦C (mr) are below 0.4% for all the samples, meaning that all the tested

FIGURE 3 | Results of the TGA tests on the samples F, A, PEG, and LAM: residual mass and mass loss derivative as a function of temperature.

materials underwent a complete thermal degradation even in inert atmosphere.

The DSC tests were performed to investigate the thermal transitions of the constituents and to assess the thermal energy storage capability of the PEG within the laminates. In fact, as reported in many papers in the literature, DSC is the most useful technique to assess the TES capability of these materials (Pielichowska and Pielichowski, 2014). In the present paper, this test allowed the determination of the melting (and crystallization) enthalpy values of the PEG phase within the laminates. The DSC thermograms of the prepared samples are reported in **Figure 4**,

TABLE 4 | Results of the DSC tests on the prepared samples (first heating scan and cooling scan).

*T PEG m , TPEG c* = *melting and crystallization temperatures of PEG.*

1*H PEG m ,* 1*H PEG c* = *melting and crystallization enthalpy values of PEG.*

*T A m, TA c* = *melting and crystallization temperatures of thermoplastic starch.*

1*H A m,* 1*H A c* = *melting and crystallization enthalpy values of thermoplastic starch.*

while the most important results are collected in **Table 4**. The DSC thermogram of the neat starch (A) was not reported here for the sake of brevity, but this test was performed as a preliminary experiment to find a suitable hot-pressing temperature of the laminates (however, the results of this test are reported in **Table 4**). The thermogram of the neat beech (F) shows a broad endothermic peak on the heating scan that is not present in the cooling and second heating scans. This peak, associated with water evaporation, has an enthalpy of 161.5 J/g. Considering that the latent evaporation heat of water is 2,272 J/g, the water present in the neat beech can be estimated at ∼7.1 wt%, which is in good agreement with the mass losses observed by TGA in the low temperature region. The other DSC signals of this sample, such as those associated with the glass transitions of the amorphous wood components, are not clearly visible in this plot, but they have been thoroughly described in a recent work of our group on wood/starch composite laminates (Dorigato et al., 2018a). The signal of water evaporation is partially visible also in the first heating scan on the sample FP, which also presents the signals of the melting and crystallization transitions of PEG, through the presence of narrow peaks at 55 and 25◦C, respectively. The phase change enthalpy measured in the first heating scan is 88.5 J/g. Assuming that the phase change performance and the crystallization behavior of the PEG embedded in the wood are not different from those of the virgin PEG, the total PEG weight fraction in the sample FP can be calculated from the phase change enthalpy of the neat PEG (169 J/g), with the result being 52.4 wt%, in good agreement with the nominal PEG content (see **Table 2**). The peak of water evaporation is also present in the thermograms of the two prepared laminates. The sample FA also clearly shows all the thermal transitions of the starch phase, which are visible in the second heating scan, i.e., a glass transition at ∼60◦C, post-crystallization at 110◦C and a melting peak at 168◦C. The thermal transitions of the starch phase are also visible in the thermogram of the sample LAM (except for the post-crystallization in the second heating scan), although the most intense signal is represented by the phase change transitions of the PEG, whose enthalpy is 27.4 J/g (first heating scan). From this signal, the effective PEG weight fraction in the LAM is 16.2 wt%, slightly higher than that calculated from the masses of the laminates, but in the same order of magnitude. This DSC signal clearly demonstrates that the prepared composites could be potentially applied for TES applications.

**Figures 5A,B**, **6A,B** report the thermograms obtained from DMA on the tested materials, performed to investigate their viscoelastic response below and above the PCM melting temperature. **Figures 5A,B** show the thermograms of the samples A, F, and FP. The latter two samples have been tested with the wood fibers oriented parallel (L) and perpendicular (T) to the applied load. The thermogram of the neat starch (A) has been reported together with the samples in the T directions (for a better comprehension of the plots). The starch phase (A) exhibits a storage modulus at 0◦C of ∼250 MPa, and it shows a considerable drop at ∼60◦C in correspondence of the glass transition, in good agreement with the DSC results. The same phenomenon is detectable through peaks in the E′′ and tanδ thermograms. For the neat wood lamina (F), the value of E′ in the L direction is one order of magnitude higher than that in the T direction, and in both cases it shows a slight decrease at ∼50◦C because of the amorphous fraction in the wood (i.e., lignin) (Kelley et al., 1987). From a comparison between the signals of the neat wood (F) and the PEG-impregnated wood (FP) laminae, it is immediately evident that the presence of PEG significantly contributes to increase the storage modulus of the wood laminae in both directions, probably due to a partial filling of the wood porosity. The major signals in the FP thermograms are those related to the PEG phase transition, as E′ shows a step, while peaks are visible in both the E′′ and tanδ signals, which are evident mostly in the L direction (**Figure 5A**). The positive contribution of PEG in increasing E′ is also transferred to the laminates (**Figures 6A,B**), even though this is accomplished to a lesser extent. Also in this case, the thermograms present all the signals related to the thermal transitions of the PEG and starch phases. For example, the PEG melting is clearly detectable through the DMA thermograms. This effect, which evidences the interesting and unusual possibility to study a melting event through DMA, was also observed by our group in other works related to polymer matrix composites containing a PCM (Fredi et al., 2017, 2018a). Further efforts will be performed in the future to better characterize the PCM melting signals through DMA analysis.

Finally, two tests were performed to investigate the effect of PEG on the mechanical properties of the wood laminae and of

the relative laminates, i.e., quasi-static tensile tests and Charpy impact tests. For the tensile tests, representative stress-strain curves of the sample FA in all the three considered directions are reported in **Figure 7** (the curves for the other samples are not reported for the sake of brevity), while **Figure 8** shows the trends of the elastic modulus and of the tensile strength, respectively. The sample F has a linear stress-strain curve until failure in both the L and T directions, with a rather low strain at break (∼1%). PEG-impregnated laminae (FP) show an initial linear behavior, followed by a region in which the slope decreases until failure. This test also evidences the positive contribution

of PEG on the tensile properties of the wood laminae, especially in the T direction, as the average elastic modulus increases from 120 to 310 MPa and the tensile strength from 60 to 170 MPa, while the properties in the L direction of the two samples are not significantly different from each other. The mechanical properties measured on these thin laminae are remarkably lower than those reported for bulk beech, which exhibits a flexural elastic modulus of 14.5 GPa and a flexural strength of 120 MPa (Giordano, 1981). This effect, already observed in previous works of our group (Dorigato et al., 2018b), is attributable to the considerable destroying of the pristine cellular morphology

to which the ultrathin laminae are subjected during cutting. The stress-strain curve of A is typical of a thermoplastic, highly ductile polymer; the initial linear region terminates with a relative maximum, conventionally regarded as the yield point, followed by a strain-hardening region until failure, which happens at high strain values (i.e., 450%). The properties measured on this starch film are consistent with those reported in the literature for similar materials (Bastioli, 1997; Mark, 1999). From **Figure 7** it is immediate to observe that, as expected, the mechanical properties of the angle ply (O) laminates are intermediate in respect to those of the samples L and T. From the results reported in **Figure 8**, it is also evident that the properties measured on the laminates are considerably higher than those expected from the traditional rule of mixture, largely used to predict the performance of composites. This effect, already observed in previous works (Dorigato et al., 2018a; Fredi et al., 2018b), stems from the partial filling of the wood porosity by the interpenetrated matrix, which reduces the total pore volume fraction. As reported before for the single laminae, this porosity filling effect is also the reason of the enhanced tensile stiffness and strength of the sample LAM compared to FA, in all the considered directions. Considering the elastic modulus results reported in **Figure 8**, it should be pointed out that the laminates (i.e., FA and LAM) and the constituents (i.e., FP, F, and A) were not tested at the same speed. However, elastic modulus results of the laminates and of their constituents could be directly compared, at least from a qualitative point of view. It must also be considered that the most important comparison in the mechanical properties should be performed between the FA and LAM samples in different orientations, which were tested at the same speed.

Similar conclusions can be drawn by analyzing the results of the Charpy impact tests, summarized in **Table 5** (representative force-time curves were not reported for the sake of brevity). The maximum force measured during the test is considerably higher for the LAM samples, and so is the absorbed specific energy, for all the considered fiber orientations. It can be therefore concluded that, contrarily to what has been reported

TABLE 5 | Results of Charpy impact tests on the two prepared laminates, performed in different orientations.

*Fb* = *maximum value of force reached during the test.*

*Usp* = *specific energy absorbed under impact conditions.*

up to now by our group for PCM-containing polymer composites (Fredi et al., 2017, 2018a), in this case the PCM contributes positively to the mechanical properties of the host composite, and the mechanical and TES properties can be considered as synergistic and not parasitic. This is remarkably important in the perspective of developing multifunctional materials combining structural and thermal management functions.

# CONCLUSIONS

This work aimed at developing novel multifunctional fully biodegradable composite laminates with thermal energy storage/release capability. With this aim, thin wood beech laminates were impregnated with PEG at the molten state, and the resulting reinforcements were interleaved to foils of thermoplastic starch through hot compaction. The microstructural and thermo-mechanical behavior of the resulting laminates was thoroughly investigated and compared with that of the relative constituents.

Even though not all the PEG introduced in the samples is effectively confined within the wood pores, a PEG fraction of about 11 wt% of the whole laminate remained entrapped within wood laminae. From SEM micrographs it was found that PEG is able to partially fill the wood porosity, but the presence of PEG on the surface of wood laminae could negatively affect the adhesion and the compenetration between the wood laminae and the starch matrix, thereby worsening the delamination resistance of the resulting composites. From DSC tests it was evident that the stabilization of PEG within the wood pores was responsible for the presence of an endothermic signal at 55◦C, with a specific melting enthalpy of 27.4 J/g. DMA analysis revealed that the presence of PEG significantly contributed to increasing the dynamic moduli of the laminates because of the partial filling of the wood porosity. For the same reason, the quasi-static properties (stiffness and strength) measured on the laminates were considerably higher than those expected from the traditional rule of mixture, and Charpy tests also showed that the impact resistance of the composites was systematically higher than that of their constituents and of the neat wood/starch laminate.

# REFERENCES

This work demonstrated the possibility to prepare fully biodegradable multifunctional materials by combining structural and thermal management functions, in which the insertion of a biodegradable PCM had a positive effect on the mechanical performance of the host composite.

# AUTHOR CONTRIBUTIONS

AD performed part of the experimental work and performed the editing of the manuscript. GF performed part of the experimental work and performed the editing of the manuscript. AP coordinated the editing of the manuscript.

# ACKNOWLEDGMENTS

Mr. Federico Vespa is gratefully acknowledged for his support to the experimental work.

storage medium: a review. Solar Energy Mater. Solar Cells. 143, 78–98. doi: 10.1016/j.solmat.2015.06.039

polyethylene, paraffin wax and expanded graphite. Renew. Energy. 88:372. doi: 10.1016/j.renene.2015.11.056

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

Copyright © 2019 Dorigato, Fredi and Pegoretti. 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.

# Host Matrix Materials for Luminescent Solar Concentrators: Recent Achievements and Forthcoming Challenges

#### Gianmarco Griffini\*

Department of Chemistry, Materials and Chemical Engineering "Giulio Natta", Politecnico di Milano, Milan, Italy

Luminescent solar concentrators (LSCs) have attracted increasing attention in the past few years as appealing solar energy technology for the seamless integration of photovoltaic (PV) systems into the built environment. Traditionally, research in this field has focused on two main aspects: the optimization of the device assembly, in the quest for more efficient architectures to maximize collection, transport, and conversion of photons into usable electrical energy; the development of novel, highly emissive luminescent species, to ensure broad light collection and efficient photon emission. Only recently, the attention has also been directed toward the selection and development of suitable host matrix/waveguide materials with appropriate optical properties, sufficient chemical compatibility with the guest luminescent species, good processability for easy device fabrication and prolonged durability in outdoor operation. In addition to consolidated polymeric systems based on polyacrylates or polycarbonates, in recent years different examples of alternative host matrix systems have been proposed, characterized by peculiar chemical, physical and optical characteristics specifically designed to meet the stringent requirements of the LSC technology. This mini-review will focus on recent developments in the design of new host matrix materials for LSC applications. An overview of the most recent examples of novel LSC host matrices will be provided here, mainly focusing on new polymers, polymer-based organic-inorganic hybrids and multifunctional organic systems. Finally, opportunities and challenges in the field will be considered in view of the effective exploitation of the LSC technology in real application scenarios.

Keywords: luminescent solar concentrator (LSC), photovoltaic (PV), solar cells (SC), waveguide (WG), polymer, organic-inorganic hybrid, multifunctional devices

# INTRODUCTION

Luminescent solar concentrators (LSCs) have been regarded as a promising, accessible solar energy technology for reducing architectural barriers to the integration of photovoltaic (PV) systems into the built environment (Meinardi et al., 2017).

Originally proposed in the late 70's (Weber and Lambe, 1976; Goetzberger and Greube, 1977), LSC systems typically consist in bulk slabs or thin films/coatings of a transparent host matrix material embedding luminescent species that act as light-harvesting and spectral-conversion

Edited by:

Alessandro Pegoretti, University of Trento, Italy

#### Reviewed by:

Luca Beverina, Università Degli Studi di Milano Bicocca, Italy Rodrigo Cercená, Universidade do Extremo Sul Catarinense, Brazil

> \*Correspondence: Gianmarco Griffini<EMAIL_ADDRESS>
#### Specialty section:

This article was submitted to Polymeric and Composite Materials, a section of the journal Frontiers in Materials

> Received: 13 January 2019 Accepted: 12 February 2019 Published: 04 March 2019

#### Citation:

Griffini G (2019) Host Matrix Materials for Luminescent Solar Concentrators: Recent Achievements and Forthcoming Challenges. Front. Mater. 6:29. doi: 10.3389/fmats.2019.00029 centers of the incident light. Upon light absorption, photons are re-emitted via radiative decay by these luminophores and a fraction of the downshifted light is confined within the host matrix as a result of the difference in refractive indexes between the LSC material and the surrounding (air) medium. Total internal reflection (TIR) allows such trapped photons to be conveyed in waveguide mode toward the edges of the LSC device, where they become concentrated and can be collected by small-area optically-coupled solar cells for photonto-electron conversion

In the past decade, research has pushed forward the performance of LSCs making them closer to practical deployability on the market.

In this mini-review, the focus will be on recent developments in the design of new host/waveguide materials for LSC applications. This discussion will include new polymers, polymer-based organic-inorganic hybrids and multifunctional systems. Opportunities and challenges in the field will also be considered in view of the effective exploitation of the LSC technology in real-life scenarios.

# LSC OPERATION AND HOST MATERIAL REQUIREMENTS

In general terms, the performance of LSC systems can be quantified based on the so-called optical efficiency of the LSC ηopt, which is a measure of the optical power effectively delivered at the edges of the LSC (P opt out) with respect to the optical power incident on its top surface (P opt in ):

$$
\eta\_{opt} = P\_{out}^{opt} / P\_{in}^{opt} \tag{1}
$$

This parameter provides a direct estimation of the fundamental properties of a LSC (i.e., to collect and concentrate incident light) as it is directly influenced by all intrinsic loss mechanisms occurring to photons prior to being successfully collected at the edges of the waveguide (Debije and Verbunt, 2012; Moraitis et al., 2018). To tackle some of these loss mechanisms, research in the field of LSCs has been traditionally driven by the development of highly emissive luminescent species and by the engineering of more efficient device assemblies, as recently summarized in some excellent review articles (Debije and Verbunt, 2012; McKenna and Evans, 2017; Mazzaro and Vomiero, 2018; Pucci, 2018).

The choice of the host/waveguide material employed for the LSC assembly has also a direct effect on efficiency as it can greatly impact the processes of photon harvesting (reflection and scattering losses are matrix-dependent), absorption/emission (photoluminescence quantum yield (PLQY) of the luminophore can be affected by its solubility in the host matrix material), waveguiding/collection at the LSC perimeter (TIR and parasitic absorption by the host matrix are determined by its optical properties) (Kastelijn et al., 2009). Host materials for LSC applications should be designed so as to provide simultaneously a number of desirable features (Zettl et al., 2017), including: suitable refractive index; high transmittance in the visible and near-infrared range; high thermodynamic solubility for the embedded luminescent species; suitable thermal properties (namely, thermo-oxidative stability, and glass transition temperature Tg) compatible with the conditions used for materials processing and with the target application; long-term durability in outdoor contexts.

Glass was originally considered a reference material for LSCs due to its excellent optical properties as well as unbeatable chemical, photochemical, and environmental stability. However, the high processing temperatures needed for the production of conventional glass together with its high specific weight have shifted the attention toward polymeric systems. In this area, commodity polymers such as poly(methyl methacrylate) (PMMA) or polycarbonate (PC) have been traditionally employed as inexpensive host waveguide materials in bulk LSCs or as easily processable host carriers in thin-film LSCs. To overcome some of the limitations of these conventional systems and to provide additional functionalities to the resulting LSC system, in recent years a surge of new alternative matrix systems has been proposed with chemical, physical, and optical characteristics appropriately designed to meet the specific requirements of the LSC technology, as will be discussed in the next sections.

# POLYMERIC MATRICES

PMMA is by far the most widely employed host polymer for LSCs because of its excellent transparency to visible light and its suitable refractive index (n ≈ 1.5) (Li et al., 2019). Recently, its use as host matrix material in bulk LSCs was benchmarked against other commodity thermoplastics with similarly suitable thermal characteristics [polystyrene (PS), styrene-acrylonitrile (SAN) and different classes of PC] in terms of optical properties, compatibility with the target luminescent species (Lumogen F 305 Red by BASF, LF305 from here on) and resulting optical efficiency (Zettl et al., 2017). It was found that PMMA and PC (compact-disc grade) displayed the highest output irradiance level from the LSC edges, in both cases yielding ηopt values of 8.9% with dye doping concentration of 500 ppm. This result was correlated with the high internal transmittance of these matrices, which enables strong absorption of incoming photons by the embedded fluorescent dye.

However, PMMA usefulness is limited by the poor solubility that efficient luminophores feature in such a matrix (Al-Kaysi et al., 2006; Colby et al., 2010; Yoo et al., 2010), as well as by its limited photostability, especially toward high energy UV light (Kaczmarek et al., 2000; Griffini et al., 2013a). Improving the stability of PMMA as host polymeric matrix for LSCs has been a key focus for many research groups in the last few years. In an early study (Mansour et al., 2005), a copolymer made of methyl methacrylate and styrene [poly(STY-co-MMA)] doped with three different organic dyes was developed and its photodegradative response was benchmarked against PMMA. It was shown that all tested copolymer/dye systems displayed much enhanced stability compared with PMMA homopolymer, irrespectively from the type of dye employed, due to more favorable dye/matrix interactions.

More recently, crosslinked systems based on partially fluorinated polymers were presented as potential alternative host matrix materials for the fabrication of durable polymer-based LSC devices (Griffini et al., 2013b). The new host matrix systems were obtained by thermally crosslinking a functional chlorotrifluoro-ethylenevinyl-ether (CTFE–VE) copolymer with different types of aliphatic polyisocyanates. The resulting LSCs in thin-film configuration (with LF305 as fluorophore) were shown to yield superior long-term operational stability compared to PMMA-based devices without compromising the efficiency. Further work by the same group demonstrated that different crosslinking agents (isocyanate vs. melamine) play a role in the evolution of chemical, physical, and morphological properties of the CTFE–VE-based LSC coatings during weathering (Griffini et al., 2014).

In the field of cross-linked systems, an unsaturated polyester (UP) resin (a mixture of UP and STY with methyl ethyl ketone peroxide as thermal initiator) was proposed as a possible alternative LSC host material to more conventional PMMA (Lim et al., 2012). To enhance the optical properties of the host matrix, a blend between UP and MMA was proposed, embedding coumarin 6 as luminescent species and resulting in satisfactory absorption and emission response.

A different, interesting approach for the development of novel polymeric host matrix systems was recently proposed based on the use of poly(p-xylylene) (or parylene) (Maggioni et al., 2013). By employing a suitably developed vacuum-based co-sublimation process, the parylene-based host material and

FIGURE 1 | (A) Schematic representation of optical fiber LSCs with the used luminescent active layers and corresponding photographs of the obtained demonstrators. Adapted with permission from Correia et al. (2016). Copyright 2016, John and Wiley and Sons. (B) Schematic representation of the chemical mechanism for the formation of planar di-ureasil LSCs doped with LR305. Reproduced with permission from Kaniyoor et al. (2016). Copyright 2015, WILEY–VCH Verlag GmbH & Co. KGaA. (C) Synthetic route for the preparation of ureasil-based LSCs embedding a conjugated polyelectrolyte as luminophore. Reproduced with permission from Meazzini et al. (2017). Copyright 2017, The Royal Society of Chemistry.

the guest fluorophore (LR305) were deposited simultaneously from the solid state on the target substrate (glass or PMMA), thereby allowing the in-situ formation of a parylene polymeric film embedding a fine dispersion of the dye within the matrix, while maintaining its surface roughness in the nanometer range. Tests on LSCs obtained by depositing fluorescent parylene films on glass and polymer waveguides showed a maximum ηopt as high as 22.6%. Further studies by the same group focused on widening the applicability of vacuum-deposited parylenebased matrices in combination with another luminescent species based on the rare-earth complex tris(2-thenoyl trifluoroacetone) mono (1,10-phenantroline) europium (III) (Eu(tta)3phen) (Tonezzer et al., 2015).

The increasing demand for environmentally-friendly materials has recently led to the development of novel biobased polymeric host matrices that may offer some ecological advantages compared with conventional oil-based systems. Within this framework, a poly-L-lactic acid (PLA) matrix was proposed as an environmentally friendly alternative to PMMA in LSC applications (Fattori et al., 2011). The smaller index of refraction of PLA vs. PMMA led to a minor reduction in maximum light trapped. On the other hand, when a luminescent species was added to the PLA matrix, the PLQY resulted to be higher than in PMMA at similar loading. Further studies from the same research group reported on another class of biopolymer-based eco-friendly matrix from silk fibroin (SF) (Melucci et al., 2012), with promising optical properties.

Another potential source of useful bio-based materials is represented by lignocellulosic biomass, in which cellulose is present as major constituent (Tuck et al., 2012). The possibility to exploit cellulose nanocrystals (CNCs) as host matrix in thin-film LSCs was recently assessed in comparison to a more conventional

FIGURE 2 | (A) Surface wettability behavior of fluorinated LSCs compared with bare glass as observed from movie stills depicting water droplets falling at increasing times. Adapted with permission from Pintossi et al. (2017). Copyright 2017, The Royal Society of Chemistry. (B) Scheme illustrating the fabrication process to realize superhydrophobic LSCs, SEM images of the obtained surface features and photographs of 3D-printed templates (Upper) with corresponding LSCs devices (Lower) obtained by replica molding. Adapted with permission from Credi et al. (2017). Copyright 2017, Elsevier. (C) Photographs of a thermo-responsive LSC demonstrating the color change as the surrounding temperature is increased from 40 to 70◦C, with changes in the optical response measured in terms of absorbance spectrum, edge emission spectrum, and ratio between edge emission intensities at 623 and 510 nm. Adapted with permission from Sol et al. (2018a) under the terms of the CC BY-NC license. (D) Photographs of the smart LSC window at increasing bias with a schematic representation of LCs and dyes alignment, and the emission direction. Adapted with permission from Sol et al. (2018b). Copyright 2018, WILEY-VCH Verlag GmbH & Co. KGaA.

acrylic polymer emulsion, in both cases using rhodamine 6G as fluorescent doping species (Chowdhury et al., 2017). It was shown that the CNC matrix may be considered suitable as host matrix material because of the enhanced absorption properties compared to the acrylic polymer emulsion. However, the latter behaved better as waveguide maintaining high efficiency also for long light propagating distances.

Very recently (Geervliet et al., 2018), two different types of renewable polyesters obtained via a catalyzed two-step melt-polycondensation reaction (a homopolymer of diethyl 2,3:4,5-di-O-methylene galactarate (GxMe) and isosorbide and a random copolymer of GxMe with 1,3-propanediol and dimethyl terephthalate) were proposed as host matrix materials in thin film LSCs. The spectroscopic characterization and the ηopt of resulting LSCs (embedding LR305 or an aggregation-induced emission molecule as guest luminophores) evidenced a performance similar or superior to that of reference LSC thin films based on PMMA/LR305, thus further indicating the potential of bio-based polymers as valuable renewable matrices for highperformance LSCs.

# ORGANIC-INORGANIC HYBRID MATRICES

In addition to purely polymeric materials, also hybrid organicinorganic systems have been thoroughly investigated as viable option for LSC host matrices. In these materials the high refractive index and enhanced stability typical of inorganic glasses are combined with the good processability and accessible chemical functionality characteristic of polymers (Pandey and Mishra, 2011; Reisfeld et al., 2011).

An interesting demonstration of LSCs based on a hybrid system employed a polysiloxane-rubber waveguide as flexible host material for LSC applications doped with LR305 (Buffa et al., 2012). It was shown that for low dye concentrations efficiencies comparable to those obtained by using conventional PC waveguides could be achieved, as a result of the good dispersibility of the fluorophore in these conditions. On the contrary, at higher concentrations luminescence quenching was observed, associated to the lower solubility of the luminophore in an apolar environment. LSC devices obtained by coupling this doped polysiloxane rubber with Si or GaAs PV cells demonstrated the viability of this material as flexible host matrix system for LSCs.

In the field of glassy hybrids, bridged silsesquioxanes containing europium(III) salts and 2-thenoyltrifluoroacetone were more recently proposed as integrated host-guest system for LSC device applications (Freitas et al., 2015). It was shown that high ηopt could be reached due to the peculiar optical properties resulting from the combination of the Eu3+ complex and the hybrid matrix. This allowed the easy fabrication of free-standing thick films (440µm) and highly luminescent coatings (200– 400 nm) on glass substrates exhibiting acceptable PLQY (∼60%). Further studies concentrated on the use of different functional bridged silsesquioxanes in combination with various trivalent lanthanide ions (Ln3+), demonstrating the high versatility of these systems for the production of defect- and crack-free films with controlled nanostructure via the sol-gel chemistry (Graffion et al., 2011, 2012).

Another important class of sol-gel-based hybrid materials widely investigated as host matrices for LSCs is that of ureasil systems. These compounds are generally prepared from the reaction between a commercially available polyetheramine and an organo-silica precursor, followed by hydrolysis and condensation of the silica network using sol-gel chemistry (McKenna and Evans, 2017). One interesting application based on such ureasil matrices is their use to fabricate cylindrical LSCs. In a first proof of concept study, a commercial PMMA optical fiber was coated with an organic-inorganic ureasil hybrid layer doped with a Eu3+ complex as luminescent species (Correia et al., 2015). Optimized fiber-based LSC devices could attain ηopt > 20%, calculated in the absorption spectral region of the luminescent material (300–380 nm). Using an optical-fiber drawing facility, the same group extended their previous studies by scaling up large-area LSCs (up to 2.5 m) based on bulk and hollow-core optical fibers (**Figure 1A**), using rhodamine 6G- or Eu3+-doped hybrids based on two different matrices, namely a di-ureasil and a tripodal tri-ureasil system (Correia et al., 2016). Efforts on the use of tri- and di-ureasil compounds as host matrices in LSCs have also focused on their combination with different doping species in planar devices (Nolasco et al., 2013; Rondão et al., 2017; Frias et al., 2018).

The ureasil family was recently further exploited as precursor for the formation of host matrix systems for planar LSC devices in bulk configuration and embedding LR305 as luminescent species (Kaniyoor et al., 2016). The host matrix was obtained by the reaction between a polyetheramine of the Jeffamine family and an isocyanate-functionalized organo-silane with the formation of urea bonds (**Figure 1B**). From the optical analysis it was found that this material exhibits enhanced transmittance in the whole UV-visible range, with optical clarity (and refractive index) comparable to glass or PMMA. Förster resonance energy transfer (FRET) was demonstrated between the matrix and the organic dye, which enabled an extension of the light-harvesting window of the LSC device. For optimized devices, ηopt = 14.5% (emission: 300–800 nm, G = 3.3) was reported. In a later work (Meazzini et al., 2017), the same class of host matrix system was tested in the presence of a poly(fluorene-alt-phenylene) copolymer containing on-chain perylenediimide chromophore units as the luminophore (**Figure 1C**), further confirming the key role played by the interplay between host matrix material and guest luminescent species in determining the optical performance of the resulting LSC.

# MULTIFUNCTIONAL MATRICES

Recent efforts in the field of LSCs have been addressed to the development of multifunctional systems in which the matrix material not only acts as host environment for the luminophore, but also displays added functionalities such as extended photochemical durability, peculiar surface properties and/or stimuli-responsiveness. Within this context (Pintossi et al., 2017), a novel photopolymeric matrix for thinfilm LSCs was recently developed for higher outdoor durability but also displaying a highly hydrophobic and oleophobic surface character (**Figure 2A**). Such system was based on a blend of three different UV-curable fluorinated oligomers to be coreacted upon UV-light exposure with a suitably functionalized perylene-based organic luminophore bearing lateral carbon double bonds, to yield a solid crosslinked LSC thin film. The long-term stability of such photocurable LSC system was demonstrated by full retention of its initial performance, as opposed to the decline observed in reference host/guest luminescent systems. The highly hydrophobic character (water contact angle ∼120◦ ) and moderate oleophobicity (mineral-oil contact angle > 90◦ ) resulting from the highly perfluorinated nature of the polymeric matrix was shown to impart easy cleanability to the LSC coating. In a successive work by the same group, superhydrophobicity of the LSC surface could be achieved by combining a soft-lithographic process together with an additive manufacturing approach based on stereolithography (Credi et al., 2017). Surface-microstructured superhydrophobic LSCs could be obtained using a suitably developed UV-curable formulation incorporating perfluorinated monomers doped with LR305 (**Figure 2B**), leading to water contact angles >160◦ .

Another interesting approach to multifunctionality in LSCs is focused on providing the host matrix with thermo-responsive properties. In this respect, the possibility to have liquid crystals (LC) as host matrix was recently explored, with the incorporation of donor–acceptor fluorescent dye pairs which could yield fluorescence through a FRET mechanism (Sol et al., 2018a). In particular, at low temperature one of the two dyes is in its aggregated form, preventing any type of fluorescence, while the second is emitting in the green region of the electromagnetic spectrum. Upon increasing temperature, the solubility of both dyes in the LC matrix increases and FRET phenomena are allowed to occur, shifting the emission of the device from green toward red. This system (**Figure 2C**) offers the opportunity to tune both absorption and emission colors, which may be a particularly interesting esthetic feature with great potential for LSCs and smart windows in general. Furthermore, positive and

# REFERENCES

negative dichroic dyes can be exploited to change the appearance of a window as a function of applied voltage (**Figure 2D**) (Verbunt et al., 2009; Debije et al., 2014; Khandelwal et al., 2015; Sol et al., 2018b).

# CONCLUSIONS AND PERSPECTIVES

The importance of the selection of appropriate host matrix materials for LSCs has driven research to the development of novel systems working alongside more conventional host matrix platforms such as glass or commodity polymers. Within this framework, the extensive array of properties achievable by the use of polymers has made them the most widely exploited class of materials for use in LSC devices. Together with polymers, research has also focused on the development of hybrid systems to simultaneously exploit the beneficial properties of both the glassy inorganic phase (chemical inertness, optical clarity) and the organic phase (chemical tunability, light-weight). Only recently, the attention has been also shifted toward matrix materials capable of imparting added functionalities to the LSC ensemble. In this respect, examples of (super)hydrophobic, oleophobic, or stimuli (thermally, electrically) responsive host matrices have been proposed and successfully demonstrated.

Considering that research on alternative matrix materials for LSCs has been addressed in detail only in recent years, the examples presented in this mini-review clearly demonstrate that there are enormous opportunities for enhanced LSC performance by developing suitably tailored host materials. In addition, the development of multifunctional matrix systems will be expected to pave the path to novel and truly smart devices with potential application in a variety of fields, thus anticipating a wider permeation of the LSC technology into the market.

# AUTHOR CONTRIBUTIONS

The author confirms being the sole contributor of this work and has approved it for publication.

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

Copyright © 2019 Griffini. 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.

## Smart Materials

Xian-Xu 'Frank' Bai

Dr Bai received a PhD degree in Instrument Science and Technology from Chongqing University in 2013. He joined the Hefei University of Technology in 2013 and founded the Laboratory for Adaptive Structures and Intelligent Systems (LASIS) in 2016. His research interests are focused in two areas. (i) Design, optimization, dynamics, and control of smart structures based on smart materials, including magnetorheological fluids/elastomers and magnetostrictive materials, applied to automotive and aerospace systems, and (ii) New mechatronics-based vehicle dynamics and control with an emphasis on intelligent/unmanned vehicles. Dr Bai has authored over 50 international journal and conference articles. He is an inventor on 16 issued Chinese patents and two PCT US patents (pending). Currently, Dr Bai serves as an Associate Editor of the Journal of Intelligent Material Systems and Structures. He is a Committee Member of the Adaptive Structures and Materials System Branch of the Aerospace Division of ASME. He is a peer reviewer of over 30 internal journals. Dr Bai is a member of ASME (S'09, M'12), SAE-China (M'16) and IEEE (M'17).

### Ali El Wahed

Ali K. El Wahed was awarded a PhD degree in Mechanical Engineering by the University of Liverpool in 1991. He has been active in various areas of research, especially in Smart Materials and in Applied Electrostatics. He is currently a Reader in Mechanical Engineering, University of Dundee where he leads the Smart Materials Group. He has published over 60 papers in technical literature.

#### Yancheng Li

Dr Yancheng Li obtained his PhD degree from Nanjing University of Science and Technology in 2007, majoring in Mechanical Engineering. After graduating, he joined the School of Civil and Environmental Engineering at the University of Technology Sydney in 2008 and has been working at the university as a Postdoctoral Research Associate, Chancellor's Research Fellow, Lecturer, and Senior Lecturer. He has been a Professor in Structural Dynamics at the College of Civil Engineering, Nanjing Tech University since 2018. His research interest include to innovatively explore the applications of various smart materials into building protection of civil structures, including magnetorheological materials (fluids, elastomer, gel and grease) and piezoelectric materials. Dr Li has published around 100 peer-reviewed journal and conference papers on the topic of smart materials and structures.

Jong-Seok Oh°

Jong-Seok Oh received a BSc degree, MSc degree, and PhD degree in Mechanical Engineering from Inha University in 2009, 2011, and 2015, respectively. He is currently an Assistant Professor at Kongju National University. His research interests include robust controller design and the control of various systems using smart actuators such as magnetorheological (MR), electrorheological (ER) fluids, and piezo actuators.

#### Ubaid Ubaidillah

Ubaidillah currently works at the Mechanical Engineering Department, Faculty of Engineering, at the Universitas Sebelas Maret. He obtained a PhD degree in 2016 from the Universiti Teknologi Malaysia; a MSc degree in 2010 from the Universiti Teknikal Malaysia Melaka; and a B.Eng. degree in 2007 from the Institut Teknologi Sepuluh Nopember (ITS). His research interest include smart materials development and actuators based on smart materials. He started to do research in magnetorheology in 2007 and has published more than 30 ISI papers. He has also started to produce an MR damper for lateral damper, seismic damper, and a secondary suspension for train etc. Some types of MR brakes were fabricated for ankle foot orthoses braking systems.

# On the Hysteresis Mechanism of Magnetorheological Fluids

#### Xian-Xu Bai 1 \* and Peng Chen1,2

*1 Laboratory for Adaptive Structures and Intelligent Systems, Department of Vehicle Engineering, Hefei University of Technology, Hefei, China, 2 GAC Automotive Research and Development Center, Guangzhou Automobile Group Co., Ltd., Guangzhou, China*

In this paper, hysteresis of magnetorheological (MR) fluids is identified from experimental tests following the mechanism of rate-independence and further studied to explore the hysteresis mechanism. The theoretical analysis based on the dipole model is provided to reveal the hysteresis mechanism of MR fluids. Specifically, the performance tests of a self-developed double-rod MR damper under different excitations show that instead of the typical force-velocity plot, the relationship between the force and displacement meets the requirements of rate-independence of the hysteresis well. A critical concept of "yield displacement" is defined and analyzed in the force-displacement plot to illustrate the hysteresis characteristics. In addition, the relationship of the normalized restoring force vs. strain is derived for a single chain from the dipole model. Then a stress-strain curve is developed for the multi-chain structure with an assumption of the dynamic equilibrium between the rupture and reconstruction of the chains. Sequentially, the hysteresis mechanism is established based on the force-displacement characteristics under reciprocating excitations. The consistency between the yield displacement in experiment and yield strain in theory verifies the effectiveness of the hypothetical hysteresis mechanism. The analysis results provide a guideline for the structural design of MR devices to enhance/reduce the hysteresis effect. The hysteresis mechanism-based further extension for the stress-strain properties of MR elastomers and the response time of MR fluids are provided as well.

 Hysteresis is an inherent property of MR fluids and MR fluid-based actuators. Such hysteretic non-linearity increases the control complexity of MR actuators, which in turn limits the MR applications. Understanding the working mechanism of the hysteresis characteristics of MR fluids is significant for studying the hysteresis characteristics of MR actuators and implementing high-efficiency control for MR semi-active systems.

 Experimental tests of MR dampers in valve or shear mode show that the hysteresis only occurs around the direction-change points under sinusoidal displacement excitations

#### Edited by:

*Seung-Bok Choi, Inha University, South Korea*

### Reviewed by:

*Yang Yu, University of Technology Sydney, Australia Xuan Shouhu, University of Science and Technology of China, China U. Ubaidillah, Sebelas Maret University, Indonesia Fitrian Imaduddin, Faculty of Engineering, Sebelas Maret University, Indonesia*

#### \*Correspondence:

*Xian-Xu Bai<EMAIL_ADDRESS>www.lasiser.com*

#### Specialty section:

*This article was submitted to Smart Materials, a section of the journal Frontiers in Materials*

Received: *11 January 2019* Accepted: *18 February 2019* Published: *12 March 2019*

#### Citation:

*Bai X-X and Chen P (2019) On the Hysteresis Mechanism of Magnetorheological Fluids. Front. Mater. 6:36. doi: 10.3389/fmats.2019.00036*

**509**

(Wang and Liao, 2011; Bai et al., 2015). Coincidentally, the associated velocities in this area are relatively small, and the hysteresis displays more distinctly in the force-velocity plot. However, from a view of rate-independence of hysteresis (Visintin, 2013; Chen et al., 2018), the experimental results (Wang and Liao, 2011; Bai et al., 2015; Sapinski et al., 2017; Yu et al., 2017; Shou et al., 2018) show that the hysteresis behaviors of MR fluids occur in force-displacement plot while not in the typical force-velocity plot. The "rate-independence" concept is defined by the path of the couple [with the input z(t) and the output F(t) of the system] that is invariant with respect to any increasing time, i.e., the output is not dependent on the derivative of the input z (Visintin, 2013). It should be noted that the hysteresis loops are not regarded as an essential feature of hysteresis (Visintin, 2013). The rate-independence, as the critical property of the hysteresis, provides an index to identify the input/output in any hysteresis phenomenon. In addition, the modeling analysis of MR fluids operated in squeeze mode indicates that the hysteresis in squeeze mode cannot be described by the normal force-velocity plot-based hysteresis models, but can be well-described by the force-displacement relation-based model–Bouc-Wen model (Chen et al., 2017a) and resistor-capacitor (RC) operator-based models (Chen et al., 2018; Bai et al., 2019). Similarly, as a derivative of MR fluids, the macroscopic mechanical properties of MR elastomers can also be modeled by the Bouc-Wen model or RC operator-based models. It indicates that MR fluids have the similar hysteresis characteristics to MR elastomers (Behrooz et al., 2014). Based on the rate-independence property of hysteresis, the macroscopic mechanical properties of MR actuators are actually a coupling of Duhem-type hysteresis and viscous properties which is ratedependent (Visintin, 2013; Chen et al., 2018).

Compared with the experimental investigations of the hysteresis characteristics of MR actuators, the research on its working mechanism is rare to be found. Theoretical and experimental studies show that the compressibility of MR fluids and the compressibility of the accumulator (if exists) can be effectively used to establish the mechanical model of the hysteresis characteristics of MR dampers in valve mode (Wang and Gordaninejad, 2007; Guo et al., 2014). However, the compressibility of MR fluids fails to explain the inherent hysteresis of MR fluid- and elastomer-based actuators (Weber, 2013; Behrooz et al., 2014) working in shear mode.

Dipole model can be used to theoretically explain the working mechanism of the controllable mechanical properties of MR fluids and MR elastomers from the microscopic level. Based on the simply dipole model, researchers have deduced the enhanced dipole model considering the magnetic interaction, the local magnetic field interaction and the magnetic chain interaction in the whole chain (Jolly et al., 1996; Furst and Gast, 2000; Xu et al., 2016). A corresponding stress-strain curve is derived under the assumption that all chains maintain consistent deformation.

The main contribution and work of this paper is to theoretically analyze the experimental performance of MR dampers and explain the hysteresis mechanism of MR fluids. In order to avoid the influence of the compressibility of the accumulator and viscous damping on the hysteresis characteristics, a self-developed double-rod MR damper will be tested under triangle displacement excitations. Applying normalization method, the expression of hysteresis of MR fluids in force-displacement plot is studied. Yield displacement concept is proposed to quantitatively analyze the experimental results. Theoretically, using the dipole model, under the assumption of the dynamic equilibrium between the rupture and reconstruction, the expression of hysteresis of MR fluids via the stress-strain curve will be derived.

# EXPERIMENTAL TESTS OF HYSTERESIS CHARACTERISTICS OF MR FLUIDS

The hysteresis characteristics of MR fluids are tested on a selfdeveloped double-rod MR damper (Bai et al., 2013, 2018a,b), as shown in **Figure 1A**, using the servo-hydraulic testing system (Type: LFH-LFV3068, SAGINOMIYA Inc.), as shown in **Figure 1B**. Observing **Figure 1B**, the experimental setup is composed of the servo-hydraulic testing system, the MR damper, and a power supply (Type: IT6861A, from ITECH Company). The servo-hydraulic testing system is integrated with a shaker actuator driven by a servo-hydraulic system, data acquisition and processing system with force sensors and displacement sensors, and a controller (i.e., a host computer, not shown in the figure). In order to avoid the influence from viscosity characteristics under different excitation velocities, the mechanical behaviors of the MR damper are tested under triangle displacement excitations provided by the servo-hydraulic testing system. The force and displacement signals are acquired by the data acquisition and processing system with a sample rate of 1 kHz. The geometrical dimensions and parameters of the self-developed double-rod MR damper are listed in **Table 1**. Such specifically developed double-rod MR damper will effectively eliminate the influence from the compressibility of the accumulator. As a reference, the experimental results of a commercial MR damper with an accumulator are provided in **Appendix A**.

**Figure 2** presents the test results of the MR damper under a triangle displacement excitation with a frequency of 1 Hz and an amplitude of 25 mm. As shown in the figure, the damping force of the MR damper changes rapidly and converges to a stable value in a region where the excitation direction changes and the excitation velocity is constant. Moreover, the tendency of the damping forces under different currents shows a good consistency in this region.

**Figure 3** presents the force-displacement plot of the MR damper under different excitation frequencies to further verify the rate-independence of the hysteresis characteristics. As seen from **Figure 3**, the force-displacement plots also show a very good consistency. Combining **Figures 2**, **3**, the forcedisplacement hysteresis characteristics of the MR damper/fluids are verified. Such phenomenon of the hysteresis being presented by the force-displacement plot is also consistent with the preyield and post-yield definitions of the stress-strain curve of MR fluids (De Vicente et al., 2011). Sequentially, the preyield and post-yield regions are defined in **Figures 2**, **3**. As known, the unavoidable friction for the MR damper shown

in **Figure 1** also behaves as a kind of Duhem hysteresis. The macroscopic mechanical characteristics of MR damper are the coupling of the hysteresis characteristics of the MR fluids and the hysteresis characteristics of friction. As seen from **Figure 2**, in field-off state (i.e., 0 A, where only the friction shows the hysteresis characteristics), the length of pre-yield is significantly different from that in fieldon states. It indicates that the hysteresis characteristics of MR fluids is an important factor for the hysteresis of the MR damper, while not the friction. In addition, in **Figures 2**, **3**, the consistence of the pre-yield lengths under different currents provides a clue for the study of the hysteresis mechanism.

To study the hysteresis characteristics of MR fluids, a concept of normalized damping force is defined as

$$
\hat{F}\_{I,i} = \frac{F\_{I,i} - F\_{0,i}}{F\_{I,0} - F\_{0,0}} \tag{1}
$$

where FI,i is the damping force of the MR damper with current of I at displacement i when under triangle displacement excitation; FI,0 is the damping force when under forward excitation and the displacement excitation at position zero; F0,i and F0,0 are the damping force of the MR damper with no current at displacement i and the damping force of the MR damper with no current at the position zero, respectively; Fˆ I,i is the normalized damping force.

On one hand, the influence of the friction on the hysteresis characteristics of the MR damper (i.e., FI,i − F0,i) can be effectively excluded by using the normalization concept. On the other hand, it is convenient to study the commonality of the hysteresis characteristics of MR fluids with different currents (magnetic fields). **Figure 4** presents the normalized force-displacement plot of the MR damper with different currents under triangle displacement excitation with a frequency of 1 Hz and an amplitude of 25 mm. As seen from **Figure 4**, the normalized force-displacement plots under different currents show a very good consistence. The distance from the displacement turning point to the convergence point of the damping force is defined as the yield displacement in this study. The yield displacements of the MR damper are nearly equal to each other under different currents, as seen from **Figure 4**. The yield displacement increases little with the increase of the current, due to the compressibility of the MR fluids. In other words, the compressibility of the MR fluids is not the essential factor for the hysteresis characteristics of MR actuators/fluids.

# THEORETICAL ANALYSIS OF HYSTERESIS CHARACTERISTICS OF MR FLUIDS

Test results and analysis have shown that the friction is only a very limited part of the hysteresis characteristics of the MR damper, and the compressibility of the MR fluids cannot fully explain the hysteresis characteristics of the MR damper. We will try to theoretically explain the hysteresis characteristics of MR damper/fluids via a theoretical perspective with the magnetic interaction of ferromagnetic particles.

# Structural Evolution of Chain Clusters of MR Fluids

Dipole model can be used to essentially explain the MR effect in microscopic point. Based on the dipole model, the deformation of the chains or columns of MR fluids leads to the restoring force

FIGURE 3 | Mechanical behavior of the MR damper with 1.8 A current under triangle displacement excitations with different frequencies and an amplitude of 25 mm.

in the presence of the magnetic field. Further, the expression of the restoring force between two particles is derived with relation to the particle distance and the polar angle between the line of centers and the direction of the magnetic field (Klingenberg et al., 1990). This model can be applied to analyze the yield stress of MR fluids in shear mode (Bossis and Lemaire, 1991; Jolly et al., 1996). Following this idea, modified dipole models that account for the particle interaction along a single chain have been proposed (Furst and Gast, 2000; Tao, 2001; Si et al., 2008). In such models, the chains are assumed to deform in shear mode, Chen et al. proved that the restoring force would be varied with different kinds of structural evolution and the deformation of the chains would be affected by the hydraulic inaction when under valve or squeeze mode (Chen et al., 2017b).