As has already been pointed out in Section 3.2.2, good architectural paint should exhibit not only good water resistance, but also good water vapour permeability. This positive combination of properties can be achieved in practice only for paints based on silicone-acrylic binders because silicone polymers are characterized by good permeability of gases due to high mobility of poly(dimethylsiloxane) chains. It was proved in our study that coatings produced from hybrid SIL-ACR and ACR-SIL dispersions showed higher water vapour permeability than those produced from starting ACR dispersions—see Figure 14.

**Figure 14.** Comparison of water vapour permeability determined for starting acrylic/styrene copolymer dispersion (ACR A) and hybrid dispersions ACR-SIL A-1 and SIL-ACR 1-A having the same composition of SIL and ACR parts.

It can be noted from the results presented in Figure 14 that water vapour permeability was better for coatings obtained from SIL-ACR dispersions than from ACR-SIL dispersions, probably because of differences in coating structure that resulted from differences in dispersion particle structure.

#### 3.2.5. Mechanical Properties

If the results of testing the mechanical properties of coatings and films produced from hybrid ACR-SIL and SIL-ACR dispersions presented in Table 2 are compared with mechanical properties of coatings produced from starting ACR dispersions, it is clear that modification with silicone led generally to less brittle coatings, especially in the case of starting dispersion ACR A. The most spectacular difference was in the (direct) impact resistance of coatings—see Figure 15.

For coatings and films produced from starting dispersion ACR B and hybrid coatings and films where ACR B composition of monomers was applied in synthesis of the relevant dispersions, the results of mechanical tests were much less convincing, presumably because *T*g of ACR B was quite high (over 30 ◦C). Cupping test results were good for all coatings and in direct elasticity measurements, only coatings produced from starting dispersion ACR B failed. Elongation at break increased for some films made from hybrid dispersions as compared to films made from starting ACR dispersions and decreased for some others (specifically for these produced from hybrid dispersions with particles having ACR B composition of ACR part) and tensile strength decreased for all films where this could be expected taking into account plasticizing effect of silicone resin. Much higher elongation at break and much lower tensile strength observed for films made from ACR-SIL dispersions than from SIL-ACR dispersions can be explained by a different supramolecular structure of films that results from different morphology of hybrid dispersion particles (see Figure 7) that coalesce to produce these films in the process of air-drying of dispersions.

**Figure 15.** Comparison of impact resistance (direct) determined for coatings obtained from starting acrylic/styrene copolymer dispersion (ACR A) and hybrid dispersions ACR-SIL A-1 and SIL-ACR 1-A having the same composition of SIL and ACR parts.

#### **4. Conclusions**

Simultaneous synthesis of aqueous silicone-acrylic and acrylic-silicone hybrid dispersions (SIL-ACR and ACR-SIL) by (1) emulsion polymerization of acrylic/styrene monomers (BA, ST, KA and MA) mixtures of different composition (ACR A and ACR B) in aqueous dispersions of silicone resins of different composition (SIL 1 and SIL 2) and (2) emulsion polymerization of silicone monomers (D4, VTES and MTES) mixtures of different composition (SIL 1 and SIL 2) in aqueous dispersions of acrylic/styrene copolymers (ACR A and ACR B) was successfully conducted. Hybrid dispersions had good mechanical stability, low minimum film-forming temperature and particle size in the range of 100–150 nm, narrow particle size distribution, and contained very little of coagulate. TEM investigation of hybrid dispersions particle structure revealed that particles of SIL-ACR dispersions exhibited "fruit cake" structure while particles of ACR-SIL dispersions showed "embedded sphere" structure. For most of the dispersions two separate *T*gs of dispersion solids (one for SIL part and the other for ACR part) that were detected by DSC were lower than *T*gs of corresponding starting SIL and ACR dispersions while single *T*g was detected for two of them. These differences were explained by differences in dispersion particle structure.

Most of the hybrid dispersions formed mechanically strong continuous coatings and films. As compared to coatings obtained from starting ACR dispersions, those obtained from hybrid dispersions showed much higher contact angles, much better water resistance and water vapour permeability and exhibited much better impact resistance. Different coating properties were observed when coatings were produced from SIL-ACR and ACR-SIL dispersions having the same composition of ACR and SIL parts, which most probably resulted from different structure of dispersions particles. Films produced from hybrid dispersions were less brittle than those produced from starting ACR dispersions. Determinations of % swell in toluene measured for films produced from hybrid dispersions revealed the difference between crosslinking density of films (i.e., also for coatings) made from ACR-SIL and SIL-ACR dispersions having the same composition of ACR and SIL parts, which was explained by higher possibility of grafting of acrylic/styrene monomers on silicone resin than of grafting VTES on acrylic/styrene copolymer. The authors believe that the selected hybrid dispersions described in this paper can be applied as binders in the formulation of architectural paints that will be characterized by high water resistance and high surface hydrophobicity combined with high water vapour permeability.

**Author Contributions:** Conceptualization, J.K. and J.T.; Methodology, J.T., W.D., I.O.-K., A.K., K.S. and J.P.; Investigation, J.T., W.D. and A.K.; Writing-Original Draft Preparation, J.K.; Writing-Review & Editing, J.T. and W.D.; Visualization, J.P., W.D. and J.T.; Supervision, J.K., D.W. and M.W.; Project Administration and Funding Aquisition, J.K.

**Funding:** This research was funded by Polish State R&D Centre (NCBiR, No. PBS/B1/8/2015).

**Acknowledgments:** The authors wish to thank Piotr Bazarnik from Warsaw University for conducting TEM studies, and Janusz Sobczak from Polish Academy of Sciences for conducting XPS studies. The assistance of Colleagues from the Industrial Chemistry Research Institute in testing mechanical properties of films, water vapour permeability and dispersion particle size distribution, is also acknowledged.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Long-Term Hydrolytic Degradation of the Sizing-Rich Composite Interphase**

#### **Andrey E. Krauklis \*, Abedin I. Gagani and Andreas T. Echtermeyer**

Department of Mechanical and Industrial Engineering, Norwegian University of Science and Technology, 7491 Trondheim, Norway<EMAIL_ADDRESS>(A.I.G<EMAIL_ADDRESS>(A.T.E.)

**\*** Correspondence<EMAIL_ADDRESS>or<EMAIL_ADDRESS>Tel.: +371-268-10-288

Received: 1 April 2019; Accepted: 17 April 2019; Published: 19 April 2019

**Abstract:** Glass fiber-reinforced composites are exposed to hydrolytic degradation in subsea and offshore applications. Fiber-matrix interphase degradation was observed after the matrix was fully saturated with water and typical water absorption tests according to ASTM D5229 were stopped. Due to water-induced dissolution, fiber-matrix interphase flaws were formed, which then lead to increased water uptake. Cutting sample plates from a larger laminate, where the fibers were running parallel to the 1.5 mm long short edge, allowed the hydrolytic degradation process to be studied. The analysis is based on a full mechanistic mass balance approach considering all the composite's constituents: water uptake and leaching of the matrix, dissolution of the glass fibers, and dissolution of the composite interphase. These processes were modeled using a combination of Fickian diffusion and zero-order kinetics. For the composite laminate studied here with a saturated epoxy matrix, the fiber matrix interphase is predicted to be fully degraded after 22 to 30 years.

**Keywords:** composites; sizing; interphase; glass fibers; environmental degradation; aging; model; kinetics; durability; hydrolysis

#### **1. Introduction**

Fiber-reinforced polymer (FRP) composites have experienced a rapid rise in use in the past 50 years due to their high strength, stiffness, relatively light weight and good corrosion resistance, especially when compared with more traditional structural materials such as steel and aluminum [1]. The reason for such superior performance is the synergistic interaction between the constituent materials inside the composite [1]. One such material is the sizing, which is a multi-component coating on the surface of the fibers. During the manufacture of FRPs, this results in the formation of a sizing-rich composite interphase between the reinforcing fibers and the matrix polymer [2]. This composite interphase is of vital importance since the mechanical properties of composite materials are often determined by whether the mechanical stresses can be efficiently transferred from the matrix to the reinforcing fibers [3–5]. The quality of the interfacial interaction is strongly dependent on the adhesional contact and the presence of flaws in the interphase [6]. It is generally agreed that the composite interphase is often the mechanical weak link and a potential source for the initiation of defects in fiber-reinforced composite structures [5].

Composite laminates are often exposed to aqueous and humid environments. Environmental aging is especially interesting for marine, offshore and deep-water applications of composites, such as oil risers and tethers [7–12]. It has been reported that water and humid environments negatively impact the mechanical properties of FRPs partially because of a loss of the interfacial bonding [5,12–15]. Flaws in the interphase can be introduced due to the interaction of the interphase with water taken up from the environment [6]. The removal of the sizing material can also lead to a microcrack initiation at the surface of glass fibers. Furthermore, various sizing components can be extracted by water, resulting

in the loss of the material [16–20]. Quantifying the water-induced aging is especially important for glass fiber-reinforced composites since the glass fibers are highly hygroscopic [5]. The environmental durability is one of the limiting factors in the structural applications [21], since the superior strength and stiffness of such materials are often compromised by the uncertainty of the material's interaction with the environment [22]. Durability is a primary issue because environmental factors such as moisture, temperature and the state of stress to which the material is exposed can degrade interfacial adhesion as well as the properties of the constituent phases. Environmental aging is mainly important at high temperatures, since the dissolution reactions are accelerated at higher temperatures. Therefore, it is of great importance to understand the environmental aging and dissolution kinetics of a sizing-rich composite interphase.

#### *1.1. Sizing and its Composition*

The sizing which forms the interphase, has typically a proprietary composition. Available information about commercial glass fibers tends to contain only one or two sizing-related details. The first is an indication of the chemical compatibility of the sizing with the matrix polymer, e.g., epoxy, as in this case. The second is a value for the loss on ignition (LOI), which indicates the amount of sizing [23]. The key functions of the sizing are: (1) to protect the glass fibers during handling and production; (2) to ensure a high level of stress transfer capability across the fiber-matrix interphase; and (3) to protect the composite matrix interphase against environmental degradation [12].

A typical sizing consists of about 20 chemicals. The most important chemical is an organofunctional silane commonly referred to as a coupling agent [24–26], which is the main component that promotes adhesion and stress-transfer between the polymer matrix and the fiber [12]. It also provides improvements in the interphase strength and hygrothermal resistance of the composite interphase [26–28]. The silane coupling agents have the general structure [*X*–Si(–O–*R*)3] where *R* is a methyl or ethyl group and *X* is a reactive group towards the polymer, in this case, an amine group. When applied to fibers, a silane coupling agent is first hydrolyzed to a silanol in presence of water. It is unstable and further condenses onto the fibers by producing a siloxane/poly(siloxane) network, which then partially becomes covalently bonded to the glass fiber surface. During the composite manufacture, the *X* reactive groups of the silane may react with the thermosetting matrix polymer, leading to a strong network bridging between the fiber and the matrix [12].

Although there are many different silane molecules available, the aminosilanes form the largest proportion of silanes employed in the composites industry [12]. The most common coupling agent is an aminosilane compound called γ-aminopropyltriethoxysilane (γ-APS), also known as APTES, which is the coupling agent in the studied sizing [16]. Usually sizings contain about 10 wt % of the coupling agent [29].

The composition of the sizing also consists of a number of multi-purpose components, such as a film former which holds the filaments together in a strand and protects the filaments from damage through fiber–fiber contact. Film formers are as closely compatible to the polymer matrix as possible. Epoxies, such as in this case, are very common film formers [12]. Usually sizings contain about 70–80 wt % of the film former [12,29].

Much less is known about the other chemicals in the sizing [12]. The sizing may also contain other compounds such as cationic or non-ionic lubricants, antistatic agents, emulsifiers, chopping aids, wetting agents or surfactants, and antioxidants [2,12,30]. Poly(propylene oxide) (PPO) or its co-polymer with poly(ethylene oxide) (PEO) is often used as a surfactant in sizings [2]. Polydimethylsiloxane (PDMS) is a common adhesion promoter, wetting agent, or surface tension reducer [2].

The exact composition of the sizing used in this study was not known to the authors, but based on technical details on the given R-glass fibers elsewhere [16], it is assumed that the sizing is based on the general characteristics described above. The results obtained are compatible with this assumption.

#### *1.2. The Structure of the Sizing-Rich Composite Interphase*

The structure of the sizing-rich composite interphase is very complex [12], as the sizing itself is heterogeneous and not uniform [12,31,32]. Furthermore, it has been observed by various researchers, that sizing is coated on fibers in "islands", "islets" or in patches, meaning that the fiber surface is only partially covered by the sizing, also giving some roughness to the surface [12,33–38]. Thomason and Dwight have concluded that epoxy-compatible sizings cover at least 90% of the glass fiber surface [39]. Mai et al. investigated APTES sizings using atomic force microscopy (AFM) and concluded that sized fibers are rougher than unsized fibers [38]. Also, similar conclusions were drawn by a few other researchers, including Turrión et al., who have shown that thickness of the sizing on the glass fibers varies from some nanometers up to a few hundred nanometers due to roughness [6,31,37].

With regards to the molecular structure of the interphase, APTES forms chemical covalent and physico-chemical hydrogen bonds and van der Waals interactions with the glass fibers and the amine epoxy [12,40]. The majority of APTES molecules which react with the glass surface can only form single Si–O–Si bonds with the glass due to steric limitations, while the vast majority of Si–O–Si bond formation in the silane interphase is due to polymerization—formation of the poly(siloxane) network [12]. A multilayer is formed on the glass fiber surfaces where the amino groups form intramolecular ring structures [32,41].

The concept of a composite interphase can be represented by a matrix polymer/poly(siloxane)/glass fiber model (shown in Figure 1) [5].

The siloxanes and poly(siloxanes) form covalent bonds with the glass fiber surface, resulting in a two-dimensional interface, the thickness of which is governed by the length of the chemical bonds, and is of an ångstrøm-scale (one tenth of a nanometer) [5].

The composite interphase is a gradient-type blend of the sizing compounds and the bulk matrix polymer, usually being about a micrometer in thickness [5,12,29,42,43]. It was observed, that an interfacial failure occurs at 0.5–4 nm from the glass surface in glass/γ-APS/epoxy interphase, indicating that the interphase region, rather than the two-dimensional interface is the weak link [5].

**Figure 1.** The concept of a polymer–siloxane–glass interphase, after [5]. The dotted line indicates that the sizing is rough [6,31,38].

#### *1.3. The Aim of This Work*

Composites take up water from their surroundings and may release some molecules into the surrounding water. Water uptake curves for composites are not straightforward to interpret since each constituent (matrix, fibers, sizing-rich interphase) interacts differently with the absorbed water. The mass uptake curve presents the combined effect of all these individual interactions.

Testing of water absorption is usually stopped when the composite material's water uptake has reached a maximum. The typical test procedures follow ASTM D5229 [44], where testing is stopped when two subsequent measurements do not differ by more than 0.5% [44]. However, when exposure to water is extended for longer periods, experiments performed in this work showed that degradation of the composite continues and additional mass gain and loss processes are involved. A similar observation was made by Perreux et al. [45], who studied immersion in water of 2.7 mm thick glass fiber epoxy composite plates for up to 10 years. They found that the weight gain of plates aged at 60 ◦C increased strongly after saturation. After about five years in water at 60 ◦C, a composite plate started to continuously lose mass with time [45]. This study will show that these effects can be related to the hydrolytic degradation of the fiber/matrix interphase.

The aim of this manuscript is to describe the degradation of the fiber/matrix interphase with special emphasis on the reaction kinetics.

#### **2. Materials and Methods**

#### *2.1. Materials*

Composite laminates were made with an amine-cured epoxy. The epoxy was prepared by mixing reagents Epikote Resin RIMR135TM (Hexion, Columbus, OH, USA) and amine based Epikure Curing Agent RIMH137TM (Hexion), stoichiometrically, in a ratio of 100:30 by weight. The mixture was degassed in a vacuum chamber for 30 min in order to remove bubbles. The density of the polymer (ρm) was 1.1 g/cm3. Resin and hardener system consisted of the following compounds by composition: 0.63 wt % bisphenol A diglycidyl ether (DGEBA), 0.14 wt % 1,6-hexanediol diglycidyl ether (HDDGE), 0.14 wt % poly(oxypropylene)diamine (POPA) and 0.09 wt % isophorondiamine (IPDA) [46].

A typical glass fiber used for marine and oil and gas applications was selected: boron-free and fluorine-free high-strength, high-modulus 3B HiPer-TexTM W2020 R-glass (3B-the fiberglass company, Hoeilaart, Belgium). Stitch-bonded mats were used. The average fiber diameter was 17 ± 2 μm [47,48]. The density of glass (ρf) was 2.54 g/cm3 [47,48].

Composite laminates 50 mm thick were prepared via vacuum-assisted resin transfer molding (VARTM). Laminates were manufactured using the aforementioned fabrics and epoxy resin. The curing was performed at room temperature for 24 h, continued by post-curing in an air oven (Lehmkuhls Verksteder, Oslo, Norway) at 80 ◦C for 16 h. Full cure was achieved [46,49]. The composite laminates were cut into specimens with dimensions of 50 mm × 50 mm × 1.5 mm. The geometry of the samples and cross section of the fibers is shown in Figure 2. Two configurations C1 and C3 were cut, as shown in Figure 2. Configurations. C1 is representative of a typical composite where fibers are parallel to one of the long sides. The surface area of cut fibers with exposed cross sections is 50 mm × 1.5 mm. Configuration C3 was cut in a way that a maximum number of cut fibers were obtained having exposed cross sections (50 mm × 50 mm). The length of the fibers was just 1.5 mm. This unusual specimen was made to obtain maximum fiber exposure towards the water. The same specimens were also used to measure anisotropic diffusivity in a separate study [49]. The specified dimensions were achieved within 5% tolerance. The thickness was adjusted using a grinding and polishing machine Jean Wirtz PHOENIX 2000 (Jean Wirtz, Dusseldorf, Germany) and SiC discs (Struers, Cleveland, OH, USA; FEPA P500, grain size 30 μm).

Figure 2 also shows a micrograph of a surface with visible cross sections of cut fibers from a specimen with C3 configuration. The micrograph was taken with a confocal microscope InfiniteFocus G4 (Alicona, Graz, Austria).

**Figure 2.** Glass fiber-reinforced epoxy composite plates: (**a**) sample configuration indicating alignment of the fibers in the plate: C3 at the top; C1 at the bottom; (**b**) micrograph of the largest face of the composite plate showing the cross section of the fibers at the surface.

Distilled water (resistivity 0.5–1.0 MΩ·cm) was used for conditioning of the composite samples. It was produced using the water purification system Aquatron A4000 (Cole-Parmer, Vernon Hills, IL, USA). The pH of the distilled water was 5.650 ± 0.010, being lower than neutral due to dissolved CO2 from atmosphere in equilibrium.

#### *2.2. Experimental Methods*

#### 2.2.1. Loss on Ignition

The loss on ignition (LOI) value of the fiber bundles was determined according to the standard practice ASTM D4963 [23]. This technique allows measurement of the weight loss of a sized glass sample. Since the weight loss is due to the burning off of the sizing, the method can be used to determine the amount of sizing on the fiber [12]. According to the LOI measurements, the sizing was 0.64 wt % of the sized fibers. The temperature during the LOI measurement was about 565 ◦C applied for about 5–5.5 h.

The obtained LOI is consistent with literature. LOI of most glass fiber reinforcement products is below 1.2 wt % [12]. For instance, Zinck and Gerard [50] also studied an APTES-based sizing which had a similar LOI value of 0.77 wt %.

#### 2.2.2. Constituent Volume and Mass Fractions of the Composite

The fiber volume fraction of the composite was 59.5% and was determined using the burn-off test, after the ASTM Standard D3171 [51]. The void volume fraction of the composite was 0.44% and was measured by image analysis of optical microscope images, as was described elsewhere by Gagani et al. [49]. Fiber, matrix, interphase and voids volume fractions were 59.5%, 39.2%, 0.9% and 0.44%, respectively. The interphase volume fraction was obtained using the LOI value (0.64 wt %), the mass of sized glass fibers (about 5.6 g), the density of the interphase (1.1 g/cm3) and the mass of the composite (about 7.2 g). Fiber, matrix and interphase mass fractions were 77.2%, 22.3% and 0.5%, respectively. The fiber surface area of one plate was about 0.5 m2 on average. The composite interphase mass fraction (*m*f*i* ) was calculated as:

$$m\_{\rm f\_i} = \frac{\rm LOI \cdot m\_{\rm fibers}}{m\_{\rm comp}} \tag{1}$$

where *m*fibers is the mass of the sized fibers; *m*comp is the mass of the composite.

2.2.3. Conditioning of Composite Plates

Water uptake and hygrothermal aging of the composite laminates was conducted using a batch system. A heated bath with distilled water (60 ± 1 ◦C) was used for conditioning the samples. Samples were weighed using analytical scales AG204 (±0.1 mg; Mettler Toledo, Columbus, OH, USA). Samples were conditioned for a period of about a year. Three parallels were performed.

2.2.4. Specific Surface Area of the Fibers Obtained by N2 Sorption/Desorption and Brunauer–Emmett–Teller (BET) Theory

The specific surface area of the sized and unsized glass fibers was obtained via N2 sorption and desorption. The method uses physical adsorption and desorption of gas molecules based on the Brunauer–Emmett–Teller (BET) theory [52]. The specific surface area was measured using QUADRASORB SI (Quantachrome Instruments, Boynton Beach, FL, USA) equipment. BET tests for specific surface area determination were performed according to the international standard ISO 9277:2010(E) [53]. The method is based on the determination of the amount of adsorptive gas molecules covering the external surface of the solid [53].

Since the sizing's surface is rough [12], the BET tests can provide the specific surface area. The BET theory explains the physical adsorption of gas molecules on a solid surface of a material, and it is the basis for the specific surface area determination.

Due to the roughness of the sizing on the fiber surface, the specific surface area of sized glass fibers measured with BET was 0.180 m2/g (see Figure 3), being higher than the specific surface area of unsized glass fibers of 0.09 m2/g (geometrical considerations as described in Section 2.1) or 0.084 m2/g using the BET method. For the unsized and sized glass fibers, the data with the BET model fit was with a determination coefficient *R*2 of 0.968 and 0.994, respectively.

**Figure 3.** Brunauer–Emmett–Teller (BET) analysis of the specific surface area of the sized fibers.

#### **3. Analytical Model**

#### *3.1. Mass Balance*

When polymers take up water from the environment, their mass is affected by the water uptake itself, leaching and aging mechanisms such as hydrolysis, chain scission or oxidation [46,54]. For the studied epoxy, there is no significant mass loss due to chemical bond scission, since hydrolysis and chain scission are not occurring [46,55].

The combination of the phenomenological perspective and mass balance approach provide a useful tool for analyzing mass uptake/loss processes in composites during hygrothermal aging by breaking down a complex process into constituent-related processes. The processes that affect weight gain or loss of composites are summarized in Table 1.

**Table 1.** Summary of the processes during hygrothermal aging of composites that affect the mass balance.

Gravimetric measurements determine the sample's mass over time during conditioning in water. The mass consists of the following terms:

$$\begin{aligned} m\_{\text{gravitimetric}}(t) &= -m\_{\text{dry}} + m\_{\text{water update}}(t) + m\_{\text{oxidation}}(t) - m\_{\text{leaching}}(t) \\ &- m\_{\text{glass dissociation}}(t) - m\_{\text{interphase dissociation}}(t) \end{aligned} \tag{2}$$

The dissolution of the interphase is then simply given by:

$$\begin{aligned} m\_{\text{interphase dissociation}}(t) &= m\_{\text{dry}} + m\_{\text{water uptake}}(t) + m\_{\text{oxidation}}(t) - m\_{\text{leadching}}(t) \\ &- m\_{\text{glass dissociation}}(t) - m\_{\text{gravimetric}}(t) \end{aligned} \tag{3}$$

The proposed model equation should be a phenomenologically full representation of the interaction between the composite material and the water environment. More details will now be given for each of the terms.

#### *3.2. Water Uptake*

The water uptake for composites includes three sub-processes: the uptake by the polymer matrix, by the interphase, and by the voids [49]. The glass fibers themselves do not absorb any water.

The water taken up by the polymer matrix at any point of time is limited by the diffusivity and the water saturation level [49,54]. The Fickian diffusion model can be used to model the water uptake by the polymer and the interphase [49]. In addition, the effect of voids being filled with water has to be considered [12,49]. The water content at saturation of the studied epoxy is 3.44 wt % if no voids are present [56]. Saturation has been defined as the moment when the difference in two consecutive water absorption measurements is lower than 0.5%, as defined by ASTM D5229 [44]. The composite's saturation water content *M*∞ was determined to be 0.96 wt % [49]. It can be calculated by Equation (4) [49]:

$$M\_{\rm co} = \frac{M\_{\rm co}^{m}(\nu\_{\rm m} + \nu\_{\rm i})\rho\_{\rm m} + M\_{\rm co}^{v}\nu\_{\rm V}\rho\_{\rm water}}{\nu\_{\rm f}\rho\_{\rm f} + (\nu\_{\rm m} + \nu\_{\rm i})\rho\_{\rm m}}\tag{4}$$

where ρm is the matrix density, ρf is the fiber density, ρwater is the water density, νf is the fiber volume fraction, νm is the matrix volume fraction, νi is the interphase volume fraction, νv is the void volume fraction (νf + νm + νi + νv = 1), *Mm* ∞ is the matrix saturation water content (3.44 wt %) and *Mv* ∞ is the void saturation water content (100 wt %). Fiber, matrix, interphase and voids volume fractions are 59.5%, 39.2%, 0.9% and 0.44%, respectively.

The sizing-rich interphase is assumed to have the same saturation water content as the epoxy matrix, since it contains about 70–80 wt % epoxy film-former [12,29,56]. Since the volume of the sizing is very small compared to the composite's volume any deviation from this assumption would have a minimal effect on the water uptake.

It is assumed here that the small voids will be completely filled with water *Mv* ∞ = 1, as was measured experimentally for the composite described here [49].

The water diffusivity of the studied epoxy polymer and the composites C1 and C3 in the thickness direction (with the fibers running transverse and parallel to the thickness direction for C1 and C3, respectively; see Figure 2) are systematized in Table 2 [49]. The higher diffusivity of the composite C3 is due to the fact that the diffusivity of the interphase in the direction parallel to the fibers is almost an order of magnitude higher than that of the polymer, after [49].

**Table 2.** Diffusivities in the through-the-thickness direction, after [49].

The following equation links the mass uptake to diffusivity from solving the 1-D Fickian diffusion equation, as described by Crank [58]:

$$M(t) = M\_{\text{co}} \left[ 1 - \left( \frac{8}{\pi^2} \right) \sum\_{i=0}^{\infty} \frac{e^{-\left(2i+1\right)\left(\frac{\pi}{n}\right)^2 Dt}}{\left(2i+1\right)^2} \right] \tag{5}$$

By fitting the exact solution of the diffusion equation to an exponential function, the ASTM standard simplified equation is the following [44]:

$$M(t) = M\_{\infty} \left[ 1 - e^{-7.3 \left( \frac{Dt}{\hbar^2} \right)^{0.75}} \right] \tag{6}$$

where *M*(*t*) is the water content, *M*∞ is the water saturation content, *t* is time, *h* is the thickness and *D* is the diffusivity in the thickness direction of the plate.

More details and 3-D Fickian model calculations can be found elsewhere [49]. 1-D and 3-D Fickian models gave the same result. Thus, for the sake of simplicity, the 1-D diffusion model for water uptake is used in this work.

Experimental gravimetric measurements and modeled water uptake curves using Equation (6) are shown in Figure 4 for a composite C3 with and without voids. It can be clearly seen that the absorption of water in the voids needs to be modeled to get a good fit with the experimental data.

**Figure 4.** Experimental gravimetric measurements of composite C3 plates conditioned in water and modeled water uptake curves using the Fickian 1-D model.

#### *3.3. Oxidation of the Epoxy Matrix*

Photo-oxidation is not present as the material is not exposed to high-energy irradiation [46]. The effect of thermo-oxidation on mass gain due to water uptake is negligible. Thermo-oxidation for the studied epoxy polymer occurs via the carbonyl formation mechanism in the carbon-carbon backbone via nucleophilic radical attack, as is described elsewhere [46].

#### *3.4. Leaching of Molecules out of the Epoxy Matrix*

Water molecules can migrate into the epoxy polymer while at the same time small molecules may leach out of the matrix [59,60]. The leaching phenomenon may occur due to initially present additives, impurities, unreacted hardener or degradation products diffusing out of the epoxy network into the water environment, which is in contact with the polymer. Often leaching follows Fickian-type diffusion [61]. The driving force of this process is due to the difference in concentration of these chemicals inside the polymer, and in the surrounding aqueous environment.

Leaching was determined experimentally using HR-ICP-MS up to about 1100 h in another work for the same epoxy material as used for making the composites in this study [46]. Krauklis and Echtermeyer [46] found that for the studied epoxy polymer there was no leaching of hardener, whilst the leaching occurred of epoxy compounds and impurities, such as epichlorohydrin and inorganic compounds. Based on Fourier transform–near infrared (FT-NIR) spectra (reported in [46]) the leached amount after about 1100 h of conditioning was estimated to be at 54.74 wt % of the initial leachable compounds present in the material. This indicates that more than a half of the small molecules were leached out after the relatively short time of 1100 h. The initial leachable compound content *M*0 leaching was found to be 0.092 wt % (about 1.5 mg) defined as the mass loss due to leaching divided by the initial mass of the polymer (about 1.6 g).

The diffusivity of leached compounds through the epoxy polymer was determined according to 1-D Fickian diffusion [44,61]:

$$M\_{\text{leaching}}(t) = M\_{\text{leaching}}^0 \left[ 1 - e^{-7.3 \left( \frac{D\_{\text{leaching}} t}{h^2} \right)^{0.95}} \right] \tag{7}$$

The diffusivity was obtained by regression analysis of the data performing non-linear Generalized Reduced Gradient (GRG) algorithm, while minimizing the residual sum of squares. The leaching diffusivity *D*leaching obtained in this study was 6.0 × 10−5 mm2/h.

The leached-out compounds were experimentally measured with High-resolution inductively coupled plasma mass spectrometry (HR-ICP-MS) (data from [46]). The modeled leaching behavior from the matrix polymer is shown in Figure 5.

**Figure 5.** Polymer leaching determined experimentally with HR-ICP-MS, after [46], and modeled using Fickian 1-D model, after [61].

#### *3.5. Glass Dissolution*

Glass fibers slowly degrade in water environments via dissolution reactions resulting in a mass loss [47,62,63]. The degradation of glass fibers follows two distinct kinetic regions: short-term non-steady-state (Phase I) and long-term steady-state degradation (Phase II), as described in the dissolving cylinder zero-order kinetic (DCZOK) model for prediction of long-term dissolution of glass from both fiber bundles [47]. During Phase I, the degradation is complex and involves such processes as ion exchange, gel formation and dissolution. When Phase II is reached, the dissolution becomes dominant and the degradation follows zero-order reaction kinetics. For the studied R-glass, the transition from Phase I into Phase II occurs in about a week (166 h) at 60 ◦C and pH 5.65 [28,47]. Elements that are released during degradation of R-glass are Na, K, Ca, Mg, Fe, Al, Si and Cl [47]. The glass mass loss is the cumulative mass loss of all these ions [47]. Si contribution to the total mass loss of the studied R-glass is the largest (56.1 wt %) and seems to govern the dissolution process [47].

The rate of the dissolution depends on the apparent glass dissolution rate constant (*K*∗ 0) and the glass surface area exposed to water (*S*) [28,47]. The glass surface area is proportional to the fiber radius. As the dissolution continues, the radius decreases linearly with time resulting in the mass loss deceleration; the DCZOK model accounts for this effect [47]. Rate constants at various environmental conditions (pH, temperature and stress), as well as more details about the model can be found in other works [28,47,63].

For a thin composite with fibers parallel to the short side through-thickness direction, such as in this work, the dissolution of glass, compared to the free fiber bundles with sizing (not embedded in the composite), is slowed down by 36.84% [28]. The differential mass loss equation for thin composites can be written as [28]:

$$\frac{\partial m}{\partial t} = K\_0^\* \mathcal{S}(t) \tag{8}$$

The *K*∗ 0 includes the effects of diffusion and accumulation of the degradation products inside the composite, the protective effect of the sizing and the availability of water [28,47,63]. The time-dependent parameter is the fiber surface area *S*(*t*).

Considering the two distinct phases of the degradation, the full DCZOK model in the integral form is the following, after [47]:

$$\begin{cases} & t \le t\_{\text{st}}: \; m\_{\text{disolved}} = n\pi \mathfrak{l} \Big( 2r\_0 K\_0^\* \,^I t - \frac{K\_0^{\*\frac{1}{2}}}{\rho\_{\text{f}}} t^2 \Big) \\ & t > t\_{\text{st}}: \; m\_{\text{disolved}} = m\_{\text{disolved}\_{\text{st}}} + n\pi \mathfrak{l} \Big( 2r\_{t\_{\text{st}}} K\_0^\* \,^{II} (t - t\_{\text{st}}) - \frac{K\_0^{\*\.II}}{\rho\_{\text{f}}} \Big( t - t\_{\text{st}} \Big)^2 \Big) \end{cases} \tag{9}$$

where *n* is the number of fibres (6450824); *l* is the length of fibres (1.5 mm); *r*0 is the initial fiber radius (8.5 μm), and ρglass is the density of glass (2.54 g/cm3); *K*∗ *I* 0 and *K*∗ *II* 0 are the apparent dissolution rate constants (g/m2·s) for the short-term non-steady-state (Phase I) and long-term steady-state (Phase II) regions, respectively; *rt*st (m) and *m*dissolved*t*st (g) are the fiber radius and lost mass after time *t*st (s), when steady-state is reached (166 h [47,63]).

Using the composition of dissolving ions reported for the studied R-glass (Si contribution 56.1 wt %) [47], and the composite data after [28], *K*∗ *I* 0 and *K*∗ *II* 0 for the studied composite are 6.91 × 10−6 and 1.54 × 10−6 g/(m2·h), respectively. The dissolution rate constants are systematized in Table 3. The glass mass loss was modeled using the DCZOK Equation (9) as shown in Figure 6. The glass mass loss is normalized by the composite plate's glass fiber surface area (about 0.5 m2).

**Table 3.** Apparent glass dissolution rate constants.

These ions determined with HR-ICP-MS come from both glass material and the sizing-rich interphase. HR-ICP-MS can capture ions from interface and interphase (ionic products of the polysiloxane/siloxane hydrolysis), but ICP does not allow carbon detection due to CO2 in the plasma, thus the organics from sizing-rich interphase are not captured. In other words, the predicted mass loss due to dissolution includes ions coming from the interface and interphase, but does not include organic compounds from the interphase. This is what makes the difference between the HR-ICP-MS determined mass loss and the gravimetric mass loss of the composite.

**Figure 6.** Glass-fiber dissolution modeled using dissolving cylinder zero-order kinetics (DCZOK) for the studied composite, after [28,47,63].

#### *3.6. Interphase Dissolution*

The aging of the sizing-rich composite interphase is the least understood constituent. The small amount of the interphase sizing compared to the composite bulk material makes analysis difficult. The proprietary nature of the sizing's composition allows only general evaluations. For typical sizing formulations, water interacting with the interphase may hydrate the Si–O–Si and Si–O–C bonds [5]. It was found that water molecules adsorbed in the epoxy matrix could migrate towards the sizing/glass fiber interface through the sizing, resulting in the dissolution/decomposition of the polysiloxane [30]. The reaction with water breaks strained Si–O–Si bonds and generates Si–OH sites [12]. Principle silane chemical bonding is reversible in the presence of water, thus the Si–O–Si bonds can be broken due to hydrolysis, as shown in Chemical Reaction (10) [12]:

$$\mathbf{H} \equiv \mathbf{Si}\text{-O-Si} \equiv \text{ +H}\_2\text{O} \leftrightarrow \text{ 2} \equiv \text{Si-OH} \tag{10}$$

In this work, the sizing-rich interphase loss is modeled assuming a simple zero-order kinetic model.

#### **4. Results and Discussion**

The increase of the composite's mass with time within the first few hundred hours could be fairly well described by a standard diffusion approach, as shown in Figure 4. It was important to include the water uptake of the voids in the calculations. However, the diffusion approach would predict a constant mass over time once saturation has been reached (0.96 wt %). The data of C3 show a slight gradual drop in mass after saturation was reached, whereas the mass of C1 is clearly increasing, as shown in Figure 7.

**Figure 7.** Long-term water uptake by composite laminates. Dashed line corresponds to a time when a test following standard practice ASTM D5229 would be stopped [44].

If the water uptake experiments are stopped as suggested by ASTM, then the long-term behavior is not captured. This observation is also consistent with the results of another study on long-term water uptake by composite plates [45]. The diverging behavior of water uptake by C1 and C3 composites can be observed starting only after about 20 h0.5 (about 2 weeks), only after the saturation *M*∞ (0.96 wt %) has been already achieved. The discussion on how the diverging behavior of C1 and C3 can be captured will follow.

#### *4.1. Samples with Short Fibers C3*

Firstly, the gravimetric behavior of C3 is addressed. As described above, a mass loss can be caused by leaching material out of the epoxy and by the glass fibers losing ions. If these effects are added to the mass vs. time curve a fairly good agreement with the experimental data is achieved, as shown in Figure 8. It could be argued that the agreement is sufficient within the experimental scatter. However, a closer look at the data can give some insight in the behavior of the sizing (interphase), although the evaluation is at the limit of what can be analyzed considering the scatter of the results.

Looking at Figure 8, a slightly better fit of the data can be obtained with a curve that has a higher mass loss with increasing time, i.e., is a bit steeper. This extra loss of material could be related to the disintegration of the interphase. The simplest approach is to model the mass loss of the interphase using the zero-order kinetics [64]:

$$\frac{\partial m\_i}{\partial t} = \mathcal{K}\_i^0 S\_i(t) \tag{11}$$

where *mi* is the mass of the interphase, *K*0 *i* is the kinetic coefficient of the interphase dissolution and *Si* is the surface area of the interphase. The solution of this equation for cylindrical fibers is given in Equation (9). For small mass changes and short times, the equation can be approximated by its first linear term with the sizing having a constant surface area *S*0 *i* to be:

$$m\_i(t) = m\_i^0 - K\_i^0 S\_i^0 t \tag{12}$$

The initial mass of the sizing *m*0 *i* (35.7 mg) was determined by the burn-off test to be 0.64 wt % of the sized fibers. Fitting the data in Figure 8 allows finding *K*0 *i S*0 *i* , which basically describes the slightly steeper slope compared to the previous analysis based only on matrix and glass fiber dissolution. Using linear regression, as shown in Figure 9, The best fit for *K*0 *i S*0 *i* = 1.80 × 10−7 g/h.

**Figure 8.** Experimental composite C3 plate mass change during the conditioning in water, shown over a square root of time. Water uptake and mass balance are modeled.

**Figure 9.** Linear regression of the difference between the experimental data and the all modelled terms except the interphase. The regressed line provides insight about the rate of the composite interphase dissolution in water.

Since dissolution is a surface reaction, a surface area of the sizing-rich interphase has to be obtained in order to determine the kinetics of dissolution. Unfortunately, we do not know the exact surface area of the sizing. Using the BET method, it was found that unsized fibers have a surface area of 0.084 m2/g and sized fibers have a surface area of 0.180 m2/g, roughly twice the value of the unsized fibers. As discussed in the introduction, the sizing is rough which creates a larger surface [6,31,37], but it also covers only parts of the fiber [12,32–37]. A typical sizing coverage of 90% of the glass fiber surface is assumed, after [39]. Furthermore, the sizing is bonded to the glass fiber on one side and the epoxy matrix on the other side, which does not create free surfaces at all. Based on the currently available information, the only possibility is to calculate *K*0 *i* for a number of plausible scenarios for the surface area *S*0 *i* .

The thickness of the interphase is obtained from the volume of the interphase *Vi* taking geometry and known coverage (i.e., 0.9 or 1) into consideration. The volume of the interphase is known from LOI (0.64 wt %; 35.8 mg) and interphase density (1.1 g/cm3), *Vi* = 0.0325 cm3. The thickness of the interphase is then obtained as follows:

$$\delta\_{\bar{i}} = \frac{V\_{\bar{i}}}{\text{Coverrange} \cdot \mathcal{S}\_{\text{glass}}} \tag{13}$$

where *S*glass is the total glass fiber surface area in a composite plate (about 0.5 m2). For 90% and 100% coverage, a mean interphase thickness is 72 and 65 nm, respectively.

**Scenario 1**. The minimum surface area *S*0 *i* would be just the cross-sectional area of the sizing exposed on the surface of the composite specimen. The fiber fraction was 59.5% and the area of one exposed surface of a C3 specimen was 50 mm × 50 mm. The surface area of fibers on both exposed surfaces is then 2975 mm2. The radius of an individual fiber was 8.5 μm. Based on the burn-off method (LOI 0.64 wt %) and assuming extreme 100% coverage, the average sizing thickness was 65 nm. The ratio of exposed sizing cross sectional area to fiber cross sectional area is then 0.0149 and the exposed sizing area is 44.3 mm2. In this scenario *K*0 *i* = 4.06 × 10−3 g/(m2·h). The sizing would be dissolved along the axis of the fibers while the exposed cross section would remain constant until the sizing is completely dissolved. Equation (12) would accurately describe dissolution in this scenario. For these 1.5 mm-thick samples, the time to dissolve the sizing would be 22.7 years.

**Scenario 2**. The other extreme would be to argue that the epoxy is quickly saturated with water (after about 100 and 81 h for C1 and C3, respectively), The water can then attack and dissolve the sizing. In that case, the exposed area of the sizing would be much bigger. The BET method measured a specific surface area of sized fibers to be 0.180 m2/g. Then, the total surface area of sized fibers (5.6 g fibers) in one plate is 1.01 m2. Since the sizing covers only parts of the fiber, not all of this surface is from the sizing. But to obtain an outer bound *K*0 *i* can be calculated for this maximum surface area (assuming coverage of 100%). In this case using Equation (12), *K*0 *i* = 1.78 × 10−7 g/(m2·h). The *K*0 *i* should be accurately determined by this equation for the relatively small area reduction during the measurement. However, the proper cylindrical Equation (9) taking the surface area reduction with time into account should be used to obtain the long-term dissolving of the sizing. The time to dissolve the sizing would be 30.5 years.

**Scenario 3**. Considering the descriptions of the literature about sizing, a typical sizing covers approximately 90% of the fiber [39]. In that case, the surface area of the sizing would be 0.91 m2. Using the same approach of a cylindrical sizing exposed to water in the epoxy as described for Scenario 2 above the *K*0 *i* for this case would be 1.98 × 10−7 g/(m2·h) and the time to dissolve the sizing would be 30.5 years.

The parameters of the three scenarios are systematized in Table 4.

**Scenario** *K***0** *i S***0** *i* **(g**/**h) Sizing Coverage (%)** δ*i* **(nm)** *S***0** *i* **(m2)** *K***0** *i* **(g**/**(m2**·**h)) Time to Total Dissolution (years) Scenario 1** 1.80 × 10−7 100 65 4.43 × 10−5 4.06 × 10−3 22.7 **Scenario 2** 1.80 × 10−7 100 65 1.01 1.78 × 10−7 30.5 **Scenario 3** 1.80 × 10−7 90, after [39] 72 0.91 1.98 × 10−7 30.5

**Table 4.** Systematized scenarios of the interphase dissolution kinetics.

The mass loss due to long-term gravimetric behavior of composite C3 could be successfully modeled, because the C3 samples did not have a significant accumulation of the degradation products. The C3 plates have a short fiber length (1.5 mm). Once the matrix is saturated with water, the water can attack and degrade the interphase. Any reduction products can be quickly transported along the interphase to the surface of the sample and will be absorbed by the surrounding water.

#### *4.2. Samples with Long Fibers C1*

The C1 samples showed a mass increase with time, see Figure 7, an additional 0.66 wt % of water was taken up after 6673 h of conditioning. Since C1 and C3 samples were made from the same laminate, just cut in a different direction, the change in behavior must be related to the sample's geometry. Compared to the C3 samples the C1 samples have much longer fibers and subsequently much longer fiber matrix interphases (1.5 mm vs. 50 mm).

The matrix of both sample types absorbs water in roughly the same period (see Table 2). The water will attack the interphase between fibers and matrix in the same way. But, it is believed that degradation products (of fibers and interphase) cannot easily move along the interphase and escape into the surrounding water at the composite's surface. Instead, the weakening of the interphase causes the formation of flaws. The degradation products and water can accumulate in these flaws. Thus, the mass of the composite does not decrease with time as for samples C3, but the mass of C1 samples increases with time. Figure 10 shows schematically what such a flaw could look like. Figure 11A shows that such flaws are, indeed, observed in the samples.

**Figure 10.**

Since the laminate absorbed another 0.66 wt % of water, it is possible to estimate the size of flaws needed to accommodate this amount of water. The initial mass of the C1 plate was about 7.36 g. Water in the interphase flaws should thus weigh 48.6 mg, taking up volume of 4.86 × 10−8 m3. Assuming for the moment that all fibers have evenly distributed flaws, the following calculations can be made. Dividing volume necessary to accommodate the extra water by the amount of fibers in a composite C1 plate (193525) and the length of a fiber (50 mm), the cross-sectional area of a water-filled interphase flaw around one fiber is found to be 5.02 × 10−12 m2. The radius of the glass fiber is 8.5 μm, thus the cross-sectional area of the fiber is 2.27 × 10−10 m2. By combining cross-sectional areas of the interphase flaw and the fiber, and deducting the radius of the fiber, an average thickness of a water-filled interphase flaw of 93.5 nm is obtained.

In reality, not all interphase flaws are the same size and not all fiber/matrix interphases are damaged equally, as shown in Figure 11. The weakest links will fail first. Once cracks are formed, stresses are released and more complicated processes follow. However, it is interesting that the first

fiber matrix debondings, as shown in Figure 11A, have dimensions similar to the calculated value of 93.5 nm. Fiber/matrix debondings shown in Figure 11A range from about a 100 nm to a few microns, as was observed experimentally using microscopy after 6673 h of conditioning. The thickness also matches debonding dimensions observed elsewhere for the same composite [65,66].

Three damage mechanisms were observed in the micrographs:

Fiber/matrix debonding appears to be the first failure mechanism, caused by hydrolysis of the interphase. This failure mechanism is described by the observations made for the C3 samples in Section 4.1 When these failure mechanisms accumulate, creating a weakened local region, they can easily combine into a longer "matrix crack" due to a release of curing, thermal and swelling stresses, resulting in a crack formation. The reason for the observed splitting along fibers is less clear. It could be related to the matrix cracks, but it could also be caused by the fibers used for stitching the reinforcing mat. All these flaws (cracks) create volume that can be filled with water and increases the mass of the composite.

**Figure 11.** Micrograph of a composite sample exposed to water for 6673 h at 60 ◦C. The micrograph indicates the (**A**) fiber/matrix debondings; (**B**) matrix transverse cracks; (**C**) splitting along the fibers.

Perreux, Choqueuse and Davies [45] investigated long-term water uptake by 2.7 mm-thick composite plates. The plate was made with an anhydride-based curing agent while this study looked at an epoxy laminate made with an amine-based curing agent. They observed that after ASTM saturation was achieved, there was still a significant continuous mass gain up to about 5 years of conditioning in water at 60 ◦C. After this point, an abrupt and continuous mass loss occurred for the following 5 years until the measurements were stopped. The data is schematically shown in Figure 12.

**Figure 12.** Schematic representation of the long-term water uptake at 20, 40 and 60 ◦C observed by Perreux, Choqueuse and Davies for the 2.7 mm thick composite plates [45].

The results seem to be a combination of what was found for samples C1 and C3 investigated here. An explanation for the behavior observed by Perreux et al. [45] may be given by the findings of this work. Initially flaws form in the composite interphase that is filled with water, resulting in a mass gain as found in C1 samples. At one point, so many flaws have accumulated that an open interpenetrating network with access to the surface of the laminate has formed. This network allows the degradation products and previously absorbed water to diffuse out, creating a mass loss similar to samples C3 (since fiber lengths in C3 were so small the interpenetrating network was present from the beginning). Since the observed mass reduction happened gradually, this means that the network of flaws and cracks is gradually being connected to the sample's surface. The mass drop was not observed for tests carried at lower temperatures. In that case all processes are slower and the samples only increased their mass, indicating the formation of flaws and cracks. But a network of the cracks reaching the surface was not created yet. It should be noted that the matrix of anhydride-based epoxies studied by Perreux et al. [45] is also prone to hydrolysis, so the hydrolysis in their samples may have affected the matrix and the fiber–matrix interphase.

#### *4.3. General Aspects*

For the composite laminates studied here, about 3.5 wt % of the interphase was dissolved in a year's time. The expected total dissolution for the geometry of the C3 sample would occur between 22.7 and 30.5 years, according to the three scenarios at 60 ◦C. At lower temperatures, the processes would be significantly slower, because diffusivities and dissolution rates follow Arrhenius-type temperature dependence [63,67]. Activation energies of these individual processes differ. Thus, it is likely not a straightforward Arrhenius-type influence on the process rate as a whole (summary mass uptake or loss).

The degradation time (22.7 to 30.5 years) should be independent of sample geometry and should be applicable once the matrix has reached saturation. For thick composite laminates the fiber–matrix interphase may only degrade in the surface region, because the matrix in the inside may remain dry. Degradation may also be stopped or slowed down by an accumulation of reaction product, if the degradation reaction is reversible, such as reaction (a). The mass uptake data obtained here showed a slight slowdown of the reaction after 9 months, close to the point when experiments were stopped, see

Figure 7. But it is unclear whether the data really flatten out. The test results from Perreux et al. [45] run over 10 years indicate that the degradation continues all the time.

Damage caused by the hydrolytic aging of the sizing-rich composite interphase very likely leads to a decrease in interfacial strength. For instance, Gagani et al. [68] and Rocha et al. [42] have reported the composite interphase-related deterioration of the mechanical properties due to aging in water. It is likely that the formation of the interphase flaws described in this work is the mechanistic origin of the interfacial strength deterioration of composites.

The authors think that studying the effect of seawater on the hydrolysis of the interphase would be useful, since the composite marine structures are most often used in the seawater environments. The dissolution in seawater conditions is expected to occur slower than in distilled water due to the presence of silica (dissolved from sand and other minerals). The reason for an expected aging rate slowdown in seawater is that the degradation products are already present in the surrounding environment, thus decreasing the driving force—a concentration gradient.

The length of glass fibers should not affect the molecular structure or morphology of the interphase per se. However, it should be added that what is affected by the fiber length is the path (or length) that the hydrolytic degradation products have to travel in order to escape the composite material and diffuse out into the surrounding water. It was shown in this work that water interaction with composites with very short interphase leads to mass loss, whereas for a typical composite an interpenetrating flaw network takes a relatively long time to form in order for degradation products to leave the composite. This leads to another aspect that needs to be studied in more detail: a diffusion of degradation products through the interphase. It is important to understand whether there is a diffusion-controlled aspect.

This paper covers hydrolysis of the composite interphase, but the same approach should be applicable for all other environmental agents and solvents (in general, solvolysis).

#### **5. Conclusions**

Glass fiber composites absorb water with time and the mass of the composites increase subsequently. When measuring diffusivity and saturation level of water according to ASTM D5229 [44] testing is stopped when the mass increase with time stops, i.e., it is reaching a plateau, in this case at about 200 h. However, continuing the tests exposing the laminates to water for longer, the mass of the composite increases again, measured up to 9 months. This additional water uptake was found to be due to the hydrolytic degradation of the sizing-rich fiber matrix interphase. Due to water-induced dissolution interphase flaws being formed which developed further into matrix cracks. The internal volume created by the flaws and cracks can be filled with water leading to the observed mass increase. The microscopically measured size of the flaws matches the order of magnitude of the volume required for obtaining the measured additional mass increase.

The hydrolytic degradation of the fiber matrix interphase could be investigated directly by cutting non-typical specimens from a thick composite laminate. The test specimens were 50 mm × 50 mm × 1.5 mm coupons where all the fibers were running parallel to the short edge. This created specimens with a short fiber–matrix interface length and the interphases being connected to the large sample's surface. When these specimens were conditioned in water, their mass increased during the first 200 h as the typical specimens described above. Continuing the test for longer times leads, however, to a mass loss. For these specimens, the flaws created by the fiber matrix interphase hydrolysis were open towards the surface of the test specimen, since the interphase length (and fiber length) was so short, 1.5 mm. The reaction products of the hydrolysis could migrate into the surrounding water bath leading to a mass drop. This mass loss allowed the product of the dissolution rate constant and the surface area of the interphase *K*0 *i S*0 *i* to be determined. The small specimens tested here would degrade the entire interphase within 22 to 30 years at 60 ◦C. The calculation is based on a full mechanistic mass balance approach considering all the composite's constituents: water uptake and leaching of the matrix, dissolution of the glass fibers, and dissolution of the composite interphase. These processes were modeled using a combination of Fickian diffusion and zero-order kinetics.

Based on long-term test data from the literature tested for close to 10 years, it seems that composites will initially absorb extra water in the flaws and cracks created by interphase hydrolysis. Eventually these cracks will create a network that is connected to the surface of the composite laminate. When this network is formed reaction products can leave the laminate and the mass will be reduced.

The possible strength degradation due to the flaws in the fiber matrix interface forming within 22 to 30 years (for the tested type of epoxy laminate) in saturated laminates should be taken into account in designs for long lifetimes.

**Author Contributions:** Conceptualization, A.E.K. and A.T.E.; Methodology, A.E.K.; Formal Analysis, A.E.K.; Investigation, A.E.K. and A.I.G.; Resources, A.E.K., A.I.G. and A.T.E.; Data Curation, A.E.K., A.I.G. and A.T.E.; Writing—Original Draft Preparation, A.E.K.; Writing—Review and Editing, A.E.K. and A.T.E.; Validation, A.E.K.; Visualization, A.E.K.; Supervision, A.T.E.; Project Administration, A.T.E.; Funding Acquisition, A.T.E.

**Funding:** This research was funded by The Research Council of Norway (Project 245606/E30 in the Petromaks 2 programme).

**Acknowledgments:** This work is part of the DNV GL led Joint Industry Project "Affordable Composites" with 19 industrial partners and the Norwegian University of Science and Technology (NTNU). The authors would like to express their thanks for the financial support from The Research Council of Norway (Project 245606/E30 in the Petromaks 2 programme). The authors are thankful to Erik Sæter, Valentina Stepanova, Susana Villa Gonzalez and Julie Asmussen. Andrey is especially thankful to Oksana V. Golubova.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **Abbreviations**

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **The Effect of Texturing of the Surface of Concrete Substrate on the Pull-Off Strength of Epoxy Resin Coating**

#### **Kamil Krzywi ´nski and Łukasz Sadowski \***

Faculty of Civil Engineering, Wroclaw University of Science and Technology, Wybrzeze Wyspia ´ ˙ nskiego 27, 50-370 Wrocław, Poland<EMAIL_ADDRESS>
**\*** Correspondence<EMAIL_ADDRESS>Tel.: +48-71-320-37-42

Received: 30 November 2018; Accepted: 18 February 2019; Published: 21 February 2019

**Abstract:** This paper describes a study conducted to evaluate the effect of texturing of the surface of concrete substrate on the pull-off strength (*f* b) of epoxy resin coating. The paper investigates a total of seventeen types of textures: after grooving, imprinting, patch grabbing and brushing. The texture of the surface of the concrete substrate was prepared during the first 15 min after pouring fresh concrete into molds. The epoxy resin coating was laid after 28 days on hardened concrete substrates. To investigate the pull-off strength of the epoxy resin coating to the concrete substrate, the pull-off method was used. The results were compared with the results obtained for a sample prepared by grinding, normative minimal pull-off strength values and the values declared by the manufacturer. During this study twelve out of fifteen tested samples achieved a pull-off strength higher than 1.50 MPa. It was found that one of the imprinting texturing methods was especially beneficial.

**Keywords:** concrete substrate; texturing; adhesion; cohesion; epoxy resin; coating; pull-off strength

#### **1. Introduction**

The construction industry is growing intensively, especially in the field of large area floors. This is mainly due to the growth of the production and transport industries and the fact that larger areas are being adapted for storage. To ensure a suitable floor for more significant loads, epoxy resins are mainly used as concrete surface coatings [1,2]. This usually allows a satisfactory pull-off strength (*f* b) of coatings to be obtained [3]. The destruction of epoxy resin coating occurs during freezing and thawing processes [4], chemical compound aggression [5], erosion and corrosion [6,7], or resistance against thermal shock [8]. As pointed out by Garcia and de Brito [9], the major advantages of epoxy resin coatings are: its high chemical and mechanical resistance, being easy to clean, and its watertightness. Epoxy resin coatings are used to enhance the durability properties of a floor [10,11]. They increase the service life of a floor and decrease its failure [12]. They may also be used as preventive repair [13] and surface protection.

Before the application of epoxy resin coating, the manufacturer recommends preparing the surface of the concrete substrate using sandblasting or grinding, cleaning it, and then using a bonding agent to achieve a guaranteed pull-off strength (usually 2.0 MPa). Usually, grinding is the most effective method of mechanically treating the surface of the concrete substrate before the application of the epoxy resin coating. This was quantified by using 3D roughness parameters of the concrete substrate [14]. However, these steps are labor-intensive and expensive. Moreover, there is a higher probability of failure during these steps. In the authors' opinion there is a need to search for a way to avoid these steps during the construction process of epoxy resin coating floors. One recent example of such a procedure is to modify the composition of the epoxy resin coating using nanosilica [15], carbon nanotubes [16], glass powder [17], polymers [18] or diacrylate monomers [19]. These attempts were successful in

increasing the pull-off strength of the epoxy resin coating to the concrete substrate. However, these kinds of modifications usually have a negative effect on the other important properties of epoxy resins, e.g., mechanical strength and viscosity. Thus, there is a need to find another way to improve the pull-off strength of the epoxy resin coatings [20]. It seems sensible to search for a proper method of treating the concrete substrate before the application of the epoxy resin coating.

Texturing of the surface of the concrete substrate is a promising method [21]. It may be especially effective when carried out on the surface of the fresh and liquid material of the substrate. For example, He et al. [22] used brushing to obtain a satisfactory coarseness of the surface of the concrete substrate. Alternatively, Mirmoghtadaei et al. [23] textured the surface of the concrete substrate using grooving and brushing in order to obtain a satisfactory coarseness. According to the authors, there have been no attempts to investigate the effect of texturing of the surface of concrete substrate on the pull-off strength of the epoxy resin coating. The following questions also remain unanswered: Which method of texturing is the most useful? Will imprinting or patch grabbing be effective?

When considering the above, the purpose of this manuscript is to evaluate the effect of texturing of the surface of concrete substrate on the pull-off strength of the epoxy resin coating. The paper also aims to ensure a proper pull-off strength of the epoxy resin coating by treating the surface of the concrete substrate using different texturing methods. For this purpose, four different texturing methods were used: grooving, imprinting, patch grabbing and brushing. Pull-off strength results were compared with the value obtained for the grinded concrete substrate surface, the value of the normative minimal pull-off strength and the value declared by the manufacturer.

#### **2. Materials and Methods**

#### *2.1. Concrete Substrate*

The concrete substrate samples were prepared in wooden forms measuring 150 × 150 × 40 mm3. To decrease the friction between the wood and sample, internal walls were covered with oil. The 40 mm thick substrate was prepared using a ready mix concrete of class C16/20 (Baumit, Wrocław, Poland). This composition consists of Type 1 Portland cement, quartz aggregate, limestone powder, sand with a grain size of 0–4 mm, and other additives. In this study the weight water-binder ratio of the ready-mix was 0.1 and the mixing time was 3 min. This kind of concrete is commonly used in civil engineering as a concrete substrate for epoxy resin coatings.

#### *2.2. Texturing of the Surface of the Concrete Substrate*

The surface of a freshly laid and liquid concrete mixture of the substrate was textured in four basic ways: grooving, imprinting, patch grabbing and brushing. The names and the descriptions of applied texture methods have been summarized in Table 1. All of these four methods are also shown in Figures 1–4.

Finally, for the grooving, widely available nets (No. 1) or grids (No. 2) with different size holes and thicknesses were used (Figure 1).

The imprinting was created using small cross spacers with thicknesses of 2, 4 and 6 mm. Moreover, two pattern types were designed for the plastic spacers: type "+" with crosses in regular spacing (Nos. 3, 5, 7), and type "x" with an additional rotated (45◦) cross in the middle of four regular spacers (Nos. 4, 6, 8). The general view of the applied imprinting method is shown in Figure 2.

The patch grabbed textures were created with the use of a flooring trowel with different square sizes (Nos. 10–12), and also with a 2 mm rotated corrugated nylon tube along the concrete substrate surface (No. 9). The third designed surface method is shown in Figure 3.

The brushing was prepared with the use of brushes of three hardness levels: delicate bristles (No. 13), medium delicate bristles (No. 14.) and wire bristles (No. 15). After brushing, a texture depth between 2 and 3 mm has been obtained. The fourth method with medium delicate bristles is presented in Figure 4.

**Table 1.** The names and the descriptions of applied textured methods.

**Figure 1.** The applied grooving method for texturing the surface of the concrete substrate (dimensions: *a*1 = 8, 10 mm; *a*2 = 1.0, 8 mm).

**Figure 2.** The applied imprinting method for texturing the surface of the concrete substrate (dimensions: *b*1 = 20, 28 mm; *b*2 = 2, 4, 6 mm).

**Figure 3.** The applied patch grabbing method for texturing the surface of the concrete substrate (dimensions: *c*1 = 2, 4, 6, 12 mm; *c*2 = 1, 4, 6 mm).

**Figure 4.** The applied brushing method for texturing the surface of the concrete substrate.

The exemplary optical views of the surfaces of the concrete substrates after 28 days of maturation are presented in Figure 5.

**Figure 5.** Exemplary optical views of the surfaces of the concrete substrates after 28 days of maturation: (**a**) grooved; (**b**) imprinted; (**c**) patch grabbed; (**d**) brushed.

The manufacturer of the epoxy resin recommended that the concrete surface be treated in two steps: grinding and applying a bonding agent. These actions allow the value of the pull-off strength higher than 2.0 MPa to be obtained after the epoxy resin coating is hardened for seven days. In this study, the concrete substrates were only textured without grinding and applying a bonding agent. For comparative purposes, one sample surface was grinded manually using a grinding stone with ceramic abrasive grain (No. X) in order to compare the obtained values with the textured forms and the pull-off strength declared by the manufacturer (*f* b > 2.0 MPa). The exemplary optical view of the surface of the concrete substrate after grinding has been presented in Figure 6. On this surface no bonding agent has been applied. These values were also compared to the minimum value of the pull-off strength required by the standard EN 1542 [24] (*f* b > 1.5 MPa).

**Figure 6.** Exemplary optical view of the surface of the concrete substrate after grinding.

#### *2.3. Epoxy Resin Coating*

Commercially available epoxy resin (StoPox BB OS, Sto-ispo Sp. z o.o., Wrocław, Poland) was prepared from two components. The first, which is a base, is an epoxy resin based on bisphenol (Component A). The second component is a hardener based on aliphatic polyamines (Component B). The weight ratio of A:B is 100:25. A plastic knife was used to mix the two components together for 3 min in order to obtain a uniform consistency (Figure 7).

**Figure 7.** Epoxy resin coating: (**a**) Components–A & B; (**b**) sample after pouring fresh epoxy resin.

30 min after mixing, the material is suitable to be applied at 20 ◦C. The viscosity of the epoxy resin after mixing with the hardener was in a range between 1400 to 2300 MPa. The epoxy resin was aged in a controlled laboratory environment at the temperature of 20 ± 2 ◦C and a relative humidity less than 65%. The epoxy resin obtains enough strength for the pull-off strength tests seven days after being poured.

#### *2.4. Pull-Off Strength Tests*

The automatic adhesion tester (DY-216, Proceq, Schwerzenbach, Switzerland) was used for the pull-off strength test according to ASTM D4541 [25]. This method has recently become very popular in assessing the pull-off strength of polymer modified coatings [26]. For each sample one specimen was tested in three places. During the test, the load on the fixture was increased in a manner that was as smooth and continuous as possible. The rate of the load was 0.05 MPa/s in order for failure to occur or so that the maximum stress was reached in about 100 s or less. After obtaining the test results, the type of failure and concrete substrate detached thickness were also analyzed (Figure 8).

(**a**) **Figure 8.** *Cont*.

**Figure 8.** Pull-off method: (**a**) scheme of the pull-off strength test; (**b**) Proceq dy-216 measuring the pull-off strength; (**c**) view of the discs glued to the coating; (**d**) cohesive and adhesive failure.

#### **3. Results**

Due to the different texturing methods of the surface of the concrete substrate, the analysis was carried out separately for each method. The results are the mean values of the pull-off strength and the concrete substrate surface detached thickness, which were obtained for each surface. From all 51 tests, one adhesive failure was observed (in this case 15% of the coating was detached). For the rest of the samples, the cohesive failure was observed in the concrete substrate surface (Figure 8d).

#### *3.1. Pull-Off Strength*

It is visible from Figure 9 that only sample No. 1 from the grooving methods obtained values of *f* b higher than 1.5 MPa. On the other hand, for all of the imprinting texturing methods, the values of *f* b were higher than 1.5 MPa. It was observed that the values of the pull-off strength for sample No. 5 were higher than those declared by the manufacturer (*f* b = 2.00 MPa). Half of the patch grabbed samples obtained pull-off strength values higher than 1.5 MPa. The observed values of pull-off strength for brushing were in a range from 1.65 to 1.91 MPa. The pull-off strength result for No. 13 is 1.74 MPa, and for the reference (grinded) surface of the concrete substrate it was equal to 1.82 MPa (No. X).

**Figure 9.** The effect of texturing of the surface of the concrete substrate on the pull-off strength of the epoxy resin coating.

#### *3.2. Detached Thickness*

Figure 10 presents the impact of concrete substrate surface detached thickness on the pull-off strength of the epoxy resin coating for grooving (samples from 1 to 2), imprinting (samples from 3 to 8), patch grabbing (samples from 9 to 12) and brushing (samples from 13 to 15).

It is visible from Figure 10a that for the grooved samples the surface of the concrete substrate detached thickness was higher for the painting grid (No. 2) than for the building net (No. 1). Figure 10b shows that in the case of imprinted surfaces, the concrete substrate surface detached thickness increases with the thickness of the cross and the number of crosses that were used on one texturing plank. The greatest result of depth after the pull-off strength test was obtained using the cross with a thickness *b*2 = 6 mm. The concrete substrate surface detached thickness for the patch grabbed samples Nos. 10–12. increases proportionately to the size of the textured longitudinal stripes (Figure 10c). The concrete substrate surface detached thickness is almost the same as the pull-off strength for samples No. 13 (13.03 mm) and No. X (13.68 mm). Moreover, for the concrete textured by brushing, a smaller concrete substrate surface detached thickness was observed (Figure 10d). The pull-off strength results for the first three methods are the best when the concrete substrate surface detached thickness value is close to 12 mm.

**Figure 10.** Impact of concrete substrate surface detached thickness on the pull-off strength of the epoxy resin coating for: (**a**) grooving (samples from 1 to 2); (**b**) imprinting (samples from 3 to 8); (**c**) patch grabbing (samples from 9 to 12); (**d**) brushing (samples from 13 to 15).

#### **4. Conclusions**

The purpose of this article was to evaluate the effect of texturing of the surface of concrete substrate on the pull-off strength of the epoxy resin coating. The research also aimed to ensure a proper pull-off strength of the epoxy resin coating by preparing the surface of the concrete substrate using different texturing methods. Based on the performed tests, the following conclusions can be drawn:

The study shows an alternative way to treat the concrete substrate. In further studies, the best concrete substrate texturing methods should be used with modified epoxy resin coating in order to increase the pull-off strength. Some actions should focus on analyzing the concrete substrate surface detached thickness and the impact of texturing depth on the pull-off strength of coatings.

This study also evidenced the technical difficulties for measuring the real adhesion between concrete substrate and epoxy resin. It is proper to note that the pull-off strength test results cannot be used to differentiate the adhesion properties of the epoxy coating to the differently prepared concrete substrates. This is due to the fact that all of the pull-off strength test results resulted in a cohesive failure. Thus, for future studies the development of an adequate method to measure real adhesion between coatings and substrates is required.

**Author Contributions:** Conceptualization, K.K. and Ł.S.; methodology, Ł.S.; validation, K.K.; formal analysis, Ł.S.; investigation, K.K.; resources, K.K.; data curation, K.K.; writing—original draft preparation, K.K.; writing—review and editing, K.K. and Ł.S.; visualization, K.K.; supervision, Ł.S.

**Funding:** This research received no external funding.

**Acknowledgments:** The authors would like to acknowledge the contribution of the COST Action CA15202.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Review* **Electrochemical Strategies for Titanium Implant Polymeric Coatings: The Why and How**

### **Stefania Cometa 1, Maria Addolorata Bonifacio 1,2, Monica Mattioli-Belmonte 3, Luigia Sabbatini 2 and Elvira De Giglio 2,\***

Received: 1 April 2019; Accepted: 18 April 2019; Published: 20 April 2019

**Abstract:** Among the several strategies aimed at polymeric coatings deposition on titanium (Ti) and its alloys, metals commonly used in orthopaedic and orthodontic prosthesis, electrochemical approaches have gained growing interest, thanks to their high versatility. In this review, we will present two main electrochemical procedures to obtain stable, low cost and reliable polymeric coatings: electrochemical polymerization and electrophoretic deposition. Distinction should be made between bioinert films—having mainly the purpose of hindering corrosive processes of the underlying metal—and bioactive films—capable of improving biological compatibility, avoiding inflammation or implant-associated infection processes, and so forth. However, very often, these two objectives have been pursued and achieved contemporaneously. Indeed, the ideal coating is a system in which anti-corrosion, anti-infection and osseointegration can be obtained simultaneously. The ultimate goal of all these coatings is the better control of properties and processes occurring at the titanium interface, with a special emphasis on the cell-coating interactions. Finally, advantages and drawbacks of these electrochemical strategies have been highlighted in the concluding remarks.

**Keywords:** electrochemistry; polymer coatings; titanium implants; corrosion protection; biocompatibility

#### **1. Introduction**

Electrochemical deposition of polymers (ECD) is a relatively new technique for metal modification, even though since ancient times metals have been coated with non-polymeric films by electrochemical processes (e.g., metal plating, anodization and many others).

On the other hand, the synthesis of organic compounds such as polymers has been traditionally accomplished via chemical routes. Alternatively, over the last century, the use of electrochemical methods for polymer synthesis has been investigated at both the laboratory and industrial scale.

ECD has gathered a considerable consensus, since it combines the advantages of an easily controlled and automated technique with the inherent possibility of coating different conducting or semiconducting substrates with polymers having disparate properties [1]. In general, the polymer ECD can be categorized into two separate methods, schematized in Figure 1:

**Figure 1.** Schematic difference between electropolymerization and electrophoretic deposition of polymers to the metal electrode.

The electropolymerization configuration is based on a three-electrode cell (i.e., working, counter and reference electrodes) containing a solution of the monomer (s) and the electrolyte (dopant) in an appropriate solvent. The monomers can be anodically oxidized or cathodically reduced, forming radical anions or cations that react with other monomers, thus forming on the surface of the involved electrode an insoluble polymer layer, which can be conductive, semi conductive or insulating. For insulating polymers, the thicknesses obtainable by electropolymerization are limited to no more than few hundred nanometres, due to the current drop occurring during the film growth. Electropolymerization allows not only to control the film thickness but also to perform in situ characterization of the polymer during its growth with the use of electrochemical and/or spectroscopic methods. Thanks to these properties, the ECD has resulted particularly fascinating in a wide range of application fields, such as corrosion protection in automotive [2,3], electroanalysis [4], electrocatalysis [5], solar cells [6], electronic and microelectronic devices [7], battery technologies, light emitting diodes, electrochromic displays [8] and so forth.

Surely, the greatest impulse to the electropolymerization was given by the development of a sub-class of the chemical sensors, that is, electrochemical sensors and biosensors. The main conducting polymers employed in biosensing are the polyheterocycles, such as polypyrrole (PPy), polythiophene (PT), polyaniline (PANI) and poly(3,4-ethylenedioxythiophene) (PEDOT), developed in the 1980s.

Noteworthy different examples of PPy-based biosensors have been reported in the literature, thanks to its suitability to immobilize enzymes [9]. The interference-free capability of this electrosynthesised polymer was guaranteed by the overoxidation process, which produced a permselective, antifouling membrane able to reject the other typical serum components. PPy films have been electrosynthesized also on mesoporous titanium oxide and resulted able to successfully immobilize glucose oxidase [10].

As far as the electrophoretic deposition process is concerned, it consists of two steps: electrophoresis, that is, the migration of the macromolecules suspended in a solution toward the electrode, attracted by an electrical field and then the formation of a deposit on the electrode surface [11]. Based on the charge of the electrode and of the macromolecules, two electrophoretic deposition processes can be distinguished: cathodic electrophoretic deposition, when the molecules are positively charged and thus attracted to the cathode; anodic electrophoretic deposition, for negatively charged molecules deposited on the anodic surface.

Both these two main ECD processes allow the irreversible, durable and uniform change of the surface properties of metal substrates. Therefore, ECD finds its natural application in various fields of medicine or biomedical engineering, improving wear resistance or enhancing the affinity of metals with living cells and tissues.

Beyond the classical applications of electrochemical techniques in industrial research areas, the past two decades have seen significant achievements of electrochemistry in biomedicine. Indeed, the development of active materials (e.g., medical devices, modified surfaces), implants, sensors and advanced drug delivery systems has benefited from the knowledge of electrochemical processes naturally occurring in living systems [12]. Moreover, in the last 10 years, electrochemical techniques have been further exploited to create complex patterns of macromolecules, accurately guiding protein deposition on conductive or non-conductive substrates [13]. For instance, the atomic-force-controlled capillary electrophoretic printing (ACCEP) has been developed to control protein positioning with high resolution in time and space, envisioning new opportunities for biomedical research [14]. Furthermore, electrochemistry is also giving a valuable contribution to pharmaceutical applications, especially in the development of remote-control drug delivery systems. In this regard, a noteworthy example was given by Huang et al., who fabricated a flexible antiepileptic-delivery system on PET via electrophoretic deposition. Drug elution was triggered by an external magnetic field, leading to tuneable release kinetics [15]. A similar concept has been exploited to enhance the biocompatibility of metallic cardiovascular devices, mainly made of stainless-steel. Innovative drug-eluting stents (DES) have been developed to in situ release anti-restenotic agents, halving in-stent restenosis phenomena with respect to bare metal devices [16].

Investigations of electrochemical methods is also guiding to very attractive findings in neural repair (e.g., in case of peripheral nerve or spinal cord injuries, glial scar treatment or cochlear functionality restoration). Gomez at al. electrosynthesized PPy layers on gold, in presence of nerve growth factor, observing improved neural cell growth [17]. Furthermore, with a similar approach, Quigley et al. managed to guide Schwann cell migration and axonal growth direction electrodepositing PLA-PLGA aligned fibres on gold substrates [18]. These findings display the potential of electrochemically prepared coatings to modify implant surfaces, aiming at precise cell guidance, with applications in regenerative medicine, especially for electrically-active tissues (i.e., muscular and nervous tissues).

To the best of our knowledge, the longest tradition of implant surface modification with electrochemical techniques is related to orthopaedic and orthodontic devices, with a special focus on titanium and its alloys prosthetic elements (artificial hip joints, artificial knee joints, bone plates, screws for fracture fixation, crowns, bridges, overdentures etc.). Hydroxyapatite layers were already electrodeposited on titanium substrates in 1986 [19]. Since then, a plethora of surface functionalization strategies have been proposed, to endow titanium implants with corrosion resistance, biocompatibility, osseointegration and antimicrobial features. The main phenomena that need to be addressed/avoided when an orthopaedic or orthodontic prosthesis is modified with an ECD coating, are schematized in Figure 2.

A pivotal role of coatings endowed with bioactive, anticorrosion or antimicrobial properties is related to the enhancement of an implant's integration with the surrounding tissues. In this respect, several different in vitro approaches were developed to assess the biocompatibility of electropolymerized or electrodeposited titanium coatings, mainly based on osteoblastic cell lines (i.e., MG63 and Saos-2) or primary cells.

This review article deals with the electrochemical strategies to coat titanium implants with polymeric films, both bioinert and bioactive films. Whereas the former class is mainly focused on titanium-based implants protection against corrosion, the latter involves coatings intended to produce an enhanced biological response, in terms of infection prevention, cell adhesion, new bone matrix deposition and so forth.

**Figure 2.** Schematic representation of phenomena occurring at titanium implant-bone interfaces.

In Figure 3, the evolution over time of the number of publications on titanium ECD coatings is reported, suggesting a growing interest toward this research field. On the right of Figure 3, a pie chart shows the distribution of different types of bioactive coatings into three main classes.

**Figure 3.** (**a**) Timeline of electrochemical deposition (ECD)-based coatings on titanium distinguished in bioinert and bioactive polymer films. (**b**) Pie chart relevant to the different bioactive coatings. Data source: 2019 Scopus®.

The State-of-art in this research field is summarized with the aim to gain insight into the future perspectives of ECD in biomedical fields.

#### **2. Bio-Inert Systems: Anticorrosion and Barrier Films**

Corrosion of metallic biomaterials is the gradual degradation of the metal surface by an electrochemical attack, which occurs when the metallic implant is placed in the hostile, highly oxygenated, saline, electrolytic environment of the human body. Blood and other constituents of body fluids are rich of different corrosive substances, including water, amino acids or proteins, plasma, mucin in the case of saliva [20], various anions (such as chloride, phosphate and bicarbonate ions), cations (like sodium, potassium, calcium, magnesium, etc.), organic substances of low-molecular-weight species as well as relatively high molecular-weight polymeric components and dissolved oxygen [21,22]. Changes in the pH values also have a deep impact on metal implant corrosion. Human body pH value is normally near 7.0; however, this value could drastically decrease due to diseases, infections and other factors and after surgery. Indeed, the pH value near the implant could lower up to 5.5 or 5.0. Even if metal implants are routinely pre-passivated prior to final packaging using electrochemical

methods [23], a clinical evidence for the release of metal ions from the implants has been established and this leaching has been ascribed to corrosion processes. Obviously, corrosion attack on metal components of surgical implants is one of the main factor responsible of their short lifetime. This issue is even more severe in the case of prostheses with mobile components. For example, life expectancy of a hip implant is no more than 10 years, mainly due to corrosion phenomena accelerated by the mechanical wear [24].

Commercially pure titanium (CpTi) and its alloys are widely used in orthopaedic and orthodontic applications, also due to their acceptable corrosion resistance [25]. Despite titanium-based implants might not cause any biological adverse reaction, long term stability of these prostheses/implants cannot be totally guaranteed. The lack of surface chemical-physical stability due to corrosive processes could produce many side effects. Indeed, Olmedo and co-workers [26] evidenced the relation between peri-implantitis, a site-specific infectious disease and ion-induced corrosion of the titanium surface.

Resistance to corrosion of titanium implants can be increased by alloying it with aluminium and vanadium or aluminium and niobium (even if also other metals such as molybdenum, zirconium, rhenium, chromium, nickel or manganese have been used). The most frequent use of the Ti-6Al-4V alloy for either orthopaedic or orthodontic implants is linked to an interesting combination of resistance to corrosion, durability, low elasticity module and high osseointegration [27–29]. However, several issues can be related to the side effects of the alloy's components. Detailed studies have shown that they lead to long term ill effects such as peripheral neuropathy, osteomalacia and Alzheimer disease due to the release of aluminium and vanadium ions from the alloy. In addition, vanadium, which is present both in the elemental state and in oxides (V2O5) is highly toxic [30].

In recent years, various researchers have studied different strategies to mitigate titanium-based implants corrosion. In general, this could be attempted by surface modification techniques [31]. Some examples of surface modification processes are physical and chemical vapour deposition [32], laser treatments [33–35], thermal oxidation [36,37] and thermal spraying [38], plasma spray [39,40], ion implantation [41], micro-arc oxidation [42], sandblasting [43,44] and electrochemical treatment [45].

Among these surface modification techniques, the manufacture of multi-materials obtained by deposition of coatings on metals can be a successful way for the mitigation of the metal corrosion process. Polymers, in particular, represent an optimal building block for the achievement of a barrier film on the metal surface.

Different coating strategies can be adopted, such as plasma assisted deposition [46], solvent evaporation [47], sol-gel dip coating [48], dip coating [49] and so forth. As a matter of fact, the application of polymeric protective coatings via electrochemical processes on titanium represents one of the most powerful strategies to obtain a superior adhesion of the polymer layer to the metal substrate with high resistance to mechanical wear.

Although this approach has been widely used in the development of anti-corrosion coatings on various metals (iron, steel, copper, etc.) in different sectors such as naval, aerospace, automotive and so forth [50–52], in the biomaterials field the development of electrochemical polymeric coatings on metals and, in particular, on titanium based implants, for corrosion protection remains a niche study.

In this respect, one of the first studies was carried out by our research group, which achieved the electrosynthesis of a polypyrrole (PPy) coating in aqueous media on both titanium and TiAlV alloy [53]. It is worth noting that, while PPy films are easily synthesized at inert anodes, the electropolymerization at oxidizable metals can occur only under electrochemical conditions that strongly passivate the electrodes without preventing the electropolymerization growth. Indeed, PPy coating is poorly adherent to the oxidizable metals (Fe, Zn, Al, Ti) because of the competition between two simultaneous anodic processes: PPy formation and metal oxidation. Interestingly, in this work it was demonstrated that a thin layer of the native oxide film did not hinder an efficient PPy polymerization on the titanium surface. This polymer is particularly versatile in that it can be either easily functionalized with biologically active molecules, able to stimulate positive interactions with bone tissue or act as an efficient protective barrier. For example, PPy was polymerized starting from the relevant

monomer solution on anodically polarized NiTi electrode, demonstrating that the polymer improves the corrosion performance at the open circuit potential and at potentials where the bare substrate suffers pitting attack [54]. An improvement in the corrosion resistance property was attempted by the electro-co-deposition of PPy and PEG (composite coating) on Ti-6Al-7Nb alloy. The authors in this paper substantiated the crucial role of PEG in the corrosion tests carried out in Hank's balanced salt solution [55]. The biological response of G292 human OBs was tested onto these PPy/PEG films, showing a good cell viability and proliferation. Furthermore, G292 preserved their morphology within the first 24 h of culture. Overall, the study suggested that the proposed film offers a microenvironment with an increased bioactivity, improved corrosion resistance and biocompatibility. Finally, another paper proposed the modification of Ti-based implants surfaces through incorporation of torularhodin, a natural compound with antimicrobial effect, by means of a polypyrrole film. The results showed that PPy–torularhodin composite film, besides showing antibacterial activity and no harmful effect on cell viability, acted also as an anticorrosion coating [56]. These observations, as well as other studies with different cytotypes [57], strengthened the role of PPy as a material for different biomedical applications.

Another class of polymers showing interesting barrier properties are the polyacrylates. In 2005, our research group focused on the study of performances of poly(methyl methacrylate) (PMMA) coatings, as such or modified by an annealing process, as barrier films against corrosion of titanium-based orthopaedic implants [58]. In this work, the electrosynthesis of MMA on titanium substrates was performed for the first time by an electro-reductive process from aqueous solutions. This study evidenced that the presence of PMMA coatings produced a decrease in ion release from Ti alloys. Moreover, the annealing treatment considerably reduced the ion dissolution rate, leading to very efficient protective coatings. As a further improvement of this research, De Giglio et al. carried out the electrosynthesis of poly(acrylic acid) (PAA) films on pure Ti or Ti-6Al-4V sheets [59]. The idea was to obtain a versatile coating for titanium-based orthopaedic implants acting both as an effective anti-corrosion barrier and as a bioactive surface, thanks to the presence of carboxylic groups in PAA that can be functionalized with bioactive molecules. Also in the case of PAA coatings, the annealing procedure resulted in a more compact film able to strongly inhibit the ion release, as demonstrated in simulating tests. Since PAA is a poly-anion, the pH dependency of the polymer barrier properties was also studied [60], indicating that PAA barrier properties were optimal at a pH value near the polymer pKa (4.9); this could be ascribed to a good compromise between the number of charged groups and the polymer swelling.

More recently, Meng et al. [61] synthesized a copolymer Poly(2-Hydroxyethyl methacrylate*co*-2-(Dimethylamino) ethyl methacrylate-*co*-7-hydroxy-4-methylcoumarin methacrylate) (PHDC))–by free radical polymerization and then self-assembled into colloidal particles and immobilized on the NiTi alloy by a simple one-step electrophoretic deposition. This study demonstrated that the copolymer coatings could significantly decrease the release of nickel ions into the environment. As far as the coating's biological evaluation is concerned, two different cytocompatibility approaches were tested. Indeed, NIH 3T3 fibroblasts were used to assess both the effect of the release of nickel ions through the physical coating barrier and the effect of the direct interaction between the cells and the coating.

Another promising polymer in the biomaterials field is poly(ether ether ketone) (PEEK). The most important properties of PEEK are low density, high mechanical and tribological properties, as well as good chemical and thermal stability and irradiation sterilization resistance. It was also exploited to develop coatings, as such or reinforced with different fibres or particles, by electrodeposition on Ti-13Nb-13Zr [62,63]. Interesting results were obtained in terms of improvement of tribological properties, such as wear resistance, friction coefficient and adhesion of the coatings to the substrate, as well as corrosion resistance. Indeed, scratch tests have been performed on these coatings, revealing no delamination even up to the maximal load of 30 N (Figure 4a). Cracks at the borders and within the scratch track were visible at high magnifications (Figure 4b). Overall, this analysis evidenced an excellent adhesion of PEEK coatings to Ti-13Nb-13Zr.

**Figure 4.** Scanning electron microscope (SEM) images of the poly(ether ether ketone) (PEEK) coating on Ti-13Nb-13Zr after scratch test at a load of 30 N (**a**) and 25 N (**b**). (Reprinted with permission from [63] 2016 Elsevier, copyright n.<PHONE_NUMBER>187).

Composite coatings, based on bioglass/chitosan and sol-gel glass/chitosan, were electrophoretically deposited on a near-β Ti-13Nb-13Zr alloy [64]. It was found that both types of coating improve the electrochemical corrosion resistance of the Ti-13Nb-13Zr alloy in Ringer's solution. Furthermore, a good biocompatibility was demonstrated with MG63 osteoblast-like cells.

A recent work discussed the electropolymerization of poly(3,4-ethylenedioxythiophene) (PEDOT) coating on near-β Ti-20Nb-13Zr (TNZ), which was pre-treated with three different surface treatments to ameliorate the substrate in terms of morphology, topography and hydrophobicity as well as to facilitate the formation of a compact PEDOT film [65]. The results from electrochemical corrosion studies clearly evidenced that the PEDOT coatings on the surface-treated TNZ substrates improved barrier protection performances in simulated body fluid (SBF).

A further example of protective polymer film electrodeposited on titanium substrates was reported by Bosh et al. [66], which deposited on Ti an anti-corrosion coating based on Halar®, an ethylene chlorotrifluoroethylene (ECTFE) thermoplastic polymer, followed by post heat treatment. It was concluded that this polymer could be applied by electrophoresis as protective coating to improve the physical and mechanical properties of the metal substrate and to reduce the stress shielding effect.

Finally, a recent paper reviewed the corrosion resistance of the main types of biocompatible metals, declaring that metals will reasonably continue to be used as biomaterials due to their unmatched mechanical excellence [67]. Therefore, electrochemical polymeric coatings can still provide new solutions for improving the stability of titanium prostheses.

#### **3. Bioactive Coatings: From Drug-Eluting Systems to Antimicrobial Surfaces**

The previous section dealt with the strategies to protect titanium implants from unwelcome phenomena (e.g., corrosion) due to the interaction between the biomaterial and the surrounding tissues. Besides, several researchers are also focusing on a complementary issue, that is, implant integration through the enhancement of cell adhesion and new matrix deposition.

The interface between a biomaterial and the host microenvironment is the main player in the process that determines implant success or failure. Indeed, several strategies are focusing on titanium surface modification in order to improve cell colonization [68].

An effective method consists in providing titanium with topographical features that stimulate cell adhesion. As an example, Wang and co-workers combined sandblasting, anodic oxidation and acid etching to fabricate micro- and nanoscale structures on titanium implants, studying the effect on cell behaviour [69]. Several other papers [70–73] proposed different techniques (ion implantation, plasma spray, lithography, electric field-aided casting) to engineer metallic surfaces, creating hierarchical patterns to guide cell adhesion. Jeon et al. discussed the impact of pattern shapes (e.g., disordered squares, grooves, pillars) on different cell types, providing new perspectives on implant fabrication to finely tune cells' activities [74].

Moreover, beyond the choice of a specific surface pattern, the achievement of a bioactive implant could take advantage from the modification of its surface's physico-chemical features (e.g., wettability, protein adsorption features, degradation rate). Indeed, bare titanium and its alloys, as well as other metallic biomaterials, do not display excellent "biofunctions" intended as the abilities to attract cells and being spontaneously integrated in the surrounding tissues [75]. In this context, the development of bioactive polymeric coatings represents one of the most effective approaches available at reasonable costs. The electrochemical process allows the direct growth of polymeric coatings on titanium, leading to tightly adherent films with adjustable thickness [76]. In this respect, our research group prepared poly(2-hydroxyethyl methacrylate) (PHEMA) coatings on different metallic substrates, using cyclic voltammetry. By simply varying the number of cycles or adding ethylene glycol dimethacrylate (EGDMA) as crosslinker, the authors highlighted the opportunity to tune the polymer thickness, porosity and swelling characteristics, affecting cell adhesion and morphology. The biocompatibility of PHEMA was assessed with fibroblasts, evidencing that hydrogel surfaces had a strong effect on cell adhesion, thus proving useful for both biomedical and diagnostic devices. Indeed, the crosslinked PHEMA-EGDMA coating displayed a smooth and continuous surface, which promoted fibroblasts elongation, while the uncrosslinked PHEMA coating, with its more porous structure, improved cell adhesion and spreading. More recently, Bhattarai et al. electrosynthesized a poly(aniline) coating via cyclic voltammetry, using a titanium-anodized substrate with nanotubular geometry [77]. Their work demonstrated the role of the polymeric coating in enhancing pre-osteoblasts attachment, spreading and osteogenic differentiation. Moreover, Popescu et al. afforded the deposition of a PPy coating through the bioinspired self-polymerization of an adhesive poly(dopamine) layer, which was able to improve the stability and duration of the PPy coating [78].

Beyond the electrosynthesis of polymers starting from their relevant monomers, electrophoretic deposition is an effective alternative, widely used to prepare films based on nature-derived polymers. In this respect, Kamata et al. developed a collagen-electrodeposited coating on titanium, discussing the superior homogeneity of electrochemically assisted deposition as compared to a conventional dipping technique [79]. Moreover, Zhuang et al. recently managed to align chitosan nanofibers on titanium using a magnetically assisted chemical electrodeposition that involved the incorporation of iron oxide nanoparticles (IOPs) and the application of an external magnetic field during ECD. This aligned collagen coating positively guided bone marrow mesenchymal stem cells (BMSCs) growth and elongation, favouring cell differentiation toward an osteogenic phenotype (e.g., expression of *m*RNA for ALP, COLL I and OCN) [80].

In addition to the already discussed benefits, coating of metallic implants with a polymeric, bioactive film involves the advantage to use it as a versatile platform to load and deliver bioactive molecules (e.g., drugs, growth factors or proteins). Through electrochemical techniques, the bioactive molecules could be entrapped during the coating formation and their release kinetics could be controlled varying the coating's features, as well as the electrochemical parameters. Moreover, since biomolecules are often sensitive to harsh temperatures, extreme pH conditions and UV exposure, the electrochemical route could preserve them from denaturation or chemical rearrangement.

Following these principles, in 2008 our research group electrosynthesized three acrylate-based hydrogels in presence of a model drug (caffeine) and a model protein (bovine serum albumin), pointing out that their release could be tuned by controlling the nature and/or the amount of the coating's crosslinker [81].

Several classes of bioactive molecules (e.g., anti-inflammatory and anticancer drugs, antimicrobials, growth factors) have been loaded on titanium and its alloys, exploiting polymeric coatings to modulate their release kinetics. As an example, indomethacin, an anti-inflammatory drug, was loaded on titanium, modified with an array of titania nanotubes (TNT) and coated with a thin, biodegradable layer of chitosan and poly(lactide-co-glycolide) [82]. The presence of the polymeric coating enabled to

extend the drug release up to one month and to adjust it, playing with the film thickness. Furthermore, the model drug ibuprofen was also loaded in an electrochemically deposited chitosan coating on titanium by means of an inorganic carrier (i.e., mesoporous silica nanoparticles) [83] (Figure 5).

This system demonstrated different drug release profiles in response to pH and electrical stimuli, providing new opportunities to trigger the desired drug release with externally controlled signals. In addition, to address the complex issue of poor vascularization around the implant, other authors managed to develop a VEGF-loaded coating on titanium implants, able to promote both mineralization and in situ angiogenesis, two processes strictly interrelated during osseointegration. MG63 human osteoblast-like cell behaviour was assessed up to 21 days, demonstrating their adhesion, viability and maintenance of osteoblastic phenotype [84,85].

Even if the controlled release of a bioactive molecule is the goal to achieve in several applications, a wide range of surface functionalization techniques is also devoted to the development of stable, bioactive surfaces via grafting of chemical groups (e.g., moieties to be used as organic bridges to improve titanium reactivity, adhesive molecules, anti-thrombogenic drugs, etc.). In this respect, De Giglio's research group focused on the electrochemical coating of titanium with a tightly adherent PPy film, modified with L-cysteine as terminal residue [53,86].

**Figure 5.** Optical picture (**a**) and SEMimage (**b**) of electrodeposited chitosanwithoutibuprofen-mesoporous silica nanoparticles. Optical picture (**c**) and SEM image (**d**) of the chitosan coating embedding ibuprofen-mesoporous silica nanoparticles. (Reprinted with permission from [83] 2014 Elsevier, copyright n.<PHONE_NUMBER>536).

Furthermore, this single amino acid was exploited to covalently bind a peptide sequence containing the RGD moiety, to enhance osteoblasts attachment on the implant surface [87]. Indeed, the authors observed an adhesion improvement of newborn rat calvaria cells equal to 230% on the modified coating, correlated to RGD surface density. Other authors grafted the RGD sequence on PEG coatings, exploiting different techniques (plasma polymerization, electrodeposition, silanization). Their aim was to achieve an antifouling system able control osteoblasts adhesion while inhibiting random protein adsorption and biofilm formation. Morphological observation showed that RGD improved cells adhesion, whilst no differences were detected for cells spreading and morphology. Authors concluded that the impaired cell adhesion due to the antifouling feature of the PEG coatings was improved by the immobilization of the RGD cell adhesive motive, while maintaining the effect on the bacterial adhesion [88].

Another successful strategy consists in modifying the relevant monomers for electropolymerization, in order to confer further reactivity to the coating. In this regard, a carboxylic acid-substituted PPy (PPy-3-acetic acid) was electrosynthesized on titanium, allowing the subsequent grafting of adhesive sequences [89]. Cell adhesion, which is essential in cell growth, migration and differentiation, as well as *m*RNA levels of Alkaline Phosphates (ALP), Collagen Type I (COLL I) and osteocalcin (OCN) were studied up to 4 days of cell culture. The viability of osteoblasts (OBs), obtained from mouse bone marrow, was not affected by the developed coating. Moreover, OBs adhered and proliferated on the electrosynthesized PPy-3-acetic film, preserving their osteoblastic phenotype.

Covalent binding techniques are also being investigated to graft anti-thrombogenic molecules on metallic implants, hindering platelets activation. As an example, Li et al. proposed a fibronectin-heparin coating on titanium, able to accelerate the endothelialization process, improving blood compatibility of metallic implants such as heart valves, stents or ventricular pumps [90]. In addition, Wagner and co-workers explored the opportunity to lower the thromboembolic risk associated with titanium implants used as ventricular assist devices. These authors modified titanium surfaces with poly(2-methacryloyloxyethyl phosphorylcholine) (MPC), a phospholipidic polymer. They demonstrated superior in vitro blood compatibility, opening new perspectives for milder intravenous anticoagulant therapies [91].

All the examples cited so far demonstrate that the combination of polymeric coatings with metallic biomaterials helps overcoming the drawbacks related to bare metals. As far as the bioactivity of a metallic implant is concerned, one of the most successful strategies is based on the development of composite biomaterials, able to mimic natural tissues. In this respect, composite biomaterials are particularly useful for orthopaedic applications, given the intrinsic dual nature of bone, with its organic (mainly collagen) and inorganic (hydroxyapatite) phases tightly intertwined. Several attempts to recreate native bone microenvironment on titanium implants have been reported, consisting in immobilizing the main components of the surrounding tissue on the biomaterial surface [92–94]. As an example, collagen was embedded during PPy electropolymerization, adjusting the applied potential to allow the collagen fibres, positively charged, to be attracted by a titanium electrode. Then, hydroxyapatite was sprayed on the PPy coating leading to a bone-mimicking composite surface [95]. More recently, other authors observed that the use of bioinspired materials (e.g., nacre, bioactive glasses, silk fibroin, etc.) could stimulate implant integration better than the precise reproduction of native bone. Therefore, several research groups studied hydroxyapatite substitutes to trigger an improved osteointegration of metallic implants [64]. Similarly, collagen was replaced by other low-cost nature-derived polymers, such as chitosan. Recently, Zhang et al. reported the successful deposition of calcium phosphate/chitosan/gentamicin films on titanium alloys, enhancing bone regrowth at levels significantly higher than the normal bone growth rates [96]. Moreover, Boccaccini's research group developed a chitosan/bioactive glass composite on titanium alloys, previously treated by grit blasting in order to tune the surface topography, roughness and wettability. The authors discussed the opportunity to change the applied voltage during electrophoretic deposition to modify the coatings' morphology [97]. With the same perspective, Metoki et al. chemisorbed on Ti-6Al-4V several self-assembled monolayers (SAM) with different end-group charge, length of the chains and anchoring groups, observing their impact on the electrodeposition of calcium phosphate (CaP). The morphology of the coating was mainly affected by the end-group type of the SAM, affecting cell colonization, spreading and biomineralization around the implant [98].

The reviewed strategies to promote biomaterials' osseointegration are often combined with tools to prevent implant infections. Indeed, a delicate, dynamic equilibrium between the risk of bacterial colonization and host's cells adhesion is immediately established after implantation, a process which is called "the race for the surface" [99]. Several works have been addressed to investigate if polymeric coating endowed with antimicrobial properties can be tethered to biomaterial surfaces without losing functionality as well as maintain compatibility with the surrounding environment. Such strategies allow the administration of the antimicrobial agents with the benefit of providing direct interaction with the site of infection. Different strategies could be pursued, such as antibiotics loading into swellable polymers, release during polymer degradation and/or the use of antibacterial molecules directly as coating components. An example of antimicrobial compounds loaded into a swellable polymer is reported in reference [100].

In this respect, electropolymerized coatings based on PHEMA swelling were designed to deliver antimicrobial drugs. Cell adhesion and viability were evaluated up to 120 h of culture, considering the burst release of the ciprofloxacin from the proposed system. Both cytoskeletal organization and cell morphology were unaffected by the presence or release of ciprofloxacin as evidenced by immunofluorescence and SEM observations (Figure 6). Moreover, Cometa et al. electrosynthesized a poly(ethylene glycol diacrylate) (PEGDA) coating on titanium, entrapping vancomycin or ceftriaxone onto its surface. The coatings inhibited methicillin-resistant *Staphylococcus aureus* in vitro, providing a useful tool for in situ prevention of dental or orthopaedic infections [101].

**Figure 6.** Fluorescence detection of F-actin stress fibres (**a**,**b**) and SEM micrographs (**c**,**d**) of MG63 cells cultured on ciprofloxacin-loaded PHEMA coatings on titanium at 24 h (**a**,**c**) and 120 h (**b**,**d**) of culture. Arrows in (**a**) indicate local contacts, in (**c**,**d**) filopodia and a star-shaped morphology. Main figure scale bar 20 μm; inset scale bar 5 μm. (Reprinted in part with permission from [100] 2011 Elsevier, copyright n.<PHONE_NUMBER>304).

More recently, Raj et al. developed a titanium alloy coating based on TiO2-SiO2 mixtures, chitosan-lysine biopolymers and electrodeposited gentamicin sulphate (Figure 7). The antimicrobial experiments against *S. aureus* and *E. coli* demonstrated that the coating could be useful to eradicate bone infections caused by Gram-negative and Gram-positive bacteria, while repairing bone loss subsequent to osteomyelitis [102].

Beyond conventional antibiotic drugs, antimicrobial peptides (AMP) represent an intriguing alternative to address the concerning issue of drug resistance. In this regard, Hoyos-Nogués et al. prepared a trifunctional coating based on a PEG anti-fouling polymer, embedding both cell adhesive

molecules and an AMP. The modified titanium surface hindered protein adsorption and *S. sanguinis* adhesion, while promoting the attachment and spreading of osteoblasts. Data reported are within the 4 h after seeding and are promising for future application in bone replacing approaches [103].

**Figure 7.** FESEM images of TiO2–SiO2 coating (**a**), TiO2–SiO2/Chitosan-Lysine coating (**b**) and hydroxyapatite formation on Drug loaded on TiO2–SiO2/Chitosan-Lysine coating (**c**). (Reprinted with permission from [102] 2018 Elsevier, copyright n.<PHONE_NUMBER>469).

Even antimicrobial metals, in the form of ions or nanoparticles, have been extensively studied as antibiotics substitutes, because of their ability to inhibit microbial proliferation and biofilm formation [104,105]. Several studies pointed out the effectiveness of silver ions and/or nanoparticles, even if cytocompatibility was not always considered [106]. In 2013, De Giglio et al. electrosynthesized a PEGDA-co-acrylic acid coating, embedding silver nanoparticles obtained by a green procedure. The developed system displayed in vitro antibacterial effectiveness on two *S. aureus* clinical isolates. The presence of AgNPs showed no significant toxic effects on osteoblast-like cells in a week of exposure to the examined coatings [107]. More recently, the same authors prepared composite polymeric coatings based on poly(acrylic acid) and chitosan, modified with silver or gallium ions to merge antibacterial effectiveness with osteointegration enhancement [108,109]. Authors evidenced that the viability of MG63 was slightly affected by the gallium-loaded bilayer, whilst no differences were present for cell adhesion, morphology and maintenance of phenotype. Interestingly, changes in coating topography due to the chemical composition represented an instructive pattern for cell arrangement. As far as the silver-loaded bilayer was concerned, cells were affected by the presence of Ag within the polymeric matrix only during the first hours of contact, whilst good cell adhesion, morphology and viability were detected after 7 days of culture. This suggests an acceptable compatibility of the silver-modified coatings with osteoblast-like cells. Overall, the biological observations strengthened the effectiveness

of these bilayer coatings for a potential application in orthopaedic and/or dental field. An agent-based modelling approach was also adopted to simulate *S. aureus* interactions with silver-loaded titanium, supporting the choice of the least effective amount of the antibacterial agent [110]. Furthermore, Kim and co-workers electropolymerized poly(dopamine) films, achieving the spontaneous reduction of silver, thus leading to an effective antimicrobial coating [111]. The electrophoretic deposition approach was also exploited by Erakovi´c et al. to attract a silver-doped hydroxyapatite, dispersed in lignin, onto a titanium electrode. Applying constant voltage, it was possible to obtain a system with immediate and continuous release of silver ions, enabling the inhibition of *S. aureus* growth. Moreover, the coating did not show toxicity against peripheral blood mononuclear cells [112].

The growing research on metal-based antimicrobial coatings suggests that the latter would be potentially coupled with conventional antibiotics to overtake drug resistance, exploiting the effectiveness of metals while limiting their toxicity.

#### **4. Conclusions and Future Perspectives**

ECD is a fascinating technique for the modification of different metal surfaces and, in particular, it has been successfully applied in biomaterials research, thanks to the possibility of using metals currently employed in both orthopaedic and orthodontic fields as working electrodes.

Not so many techniques are characterized by the main advantages of the ECD, such as low costs, high purity of products, high deposition rates, homogenous covering distribution, capability of combining different monomers to obtain copolymers with a wide range of composition (in the case of electropolymerization), possibility of employing complex geometries of substrate, reduced waste materials, ease of process control and automation.

Moreover, polymer coatings with customized morphology and assemblies can be simply obtained by controlling ECD operating conditions (i.e., initial pH, electrolytic solution composition, current density, deposition time and temperature, typology of electrochemical cell or experimental design etc.).

Finally, ECD can be performed at room temperature from water-based electrolytes, avoiding the use of organic solvents. Hence, ECD can be considered a green chemistry approach to be employed without concerns in biomaterials field, where it is mandatory to avoid the use of potentially toxic substances.

However, there are still some issues to be solved, that represent the current research challenges. Mainly, both electropolymerization and electrodeposition usually produce a fairly dense polymer. Hence, the growth of a passivating polymer film prevents a further polymer deposition (in electrodeposition) or monomer migration (in electropolymerization) on the electrode surface, leading only to very thin coatings. Consequently, other techniques are nowadays required to develop highly porous structures or to obtain thick coatings on titanium.

For all these reasons, when ECD strategies can be exploited, then relatively small areas of conductive substrates could be coated with polymers. Therefore, rather than the scale up of these procedures to coat metal surfaces of large size, the future perspectives of ECD will be related with miniaturisation. Certainly, for the biomedical research, as well as for other technological sectors, it is predictable that the major future trend will be the miniaturisation, rather than the use of macro-devices. In this respect, electrochemical deposition of polymers on metals represents an intriguing strategy, due to the possibility to operate with microelectrodes.