**Advanced Coatings for Corrosion Protection**

Editor

**Wolfram F ¨urbeth**

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## **Contents**



## **About the Editor**

**Wolfram F ¨urbeth** has a diploma in Chemistry from the University of Dortmund and received a Ph.D. in Materials Sciences from the University of Erlangen-Nuremberg in 1997. He joined DECHEMA as a research scientist in 1997 and has been head of the corrosion research group since 2005. In 2012, he obtained a habilitation in Materials Technology at RWTH Aachen Technical University. He served as Adjunct Professor at RWTH in 2018, and since March 2017, he has been a member of the Institute Directorate of the DECHEMA Research Institute. Since 2017, he has served as Head of the Scientific Advisory Board of the German Corrosion Society, as well as Chairman of the European Federation of Corrosion Science and Technology Advisory Committee since 2017 and Chairman of Working Party Coatings since 2012.

## *Editorial* **Special Issue: Advanced Coatings for Corrosion Protection**

#### **Wolfram Fürbeth**

DECHEMA Research Institute, Theodor-Heuss-Allee 25, 60486 Frankfurt am Main, Germany; wolfram.fuerbeth@dechema.de; Tel.: +49-69-7564-398

Received: 29 July 2020; Accepted: 30 July 2020; Published: 1 August 2020

**Abstract:** Corrosion is an important issue in many industrial fields. Among others, coatings are by far the most important technology for corrosion protection of metallic surfaces. The special issue "Advanced Coatings for Corrosion Protection" has been launched as a means to present recent developments on any type of advanced coatings for corrosion protection. Fifteen contributions have been collected on metallic, inorganic, polymeric and nanoparticle enhanced coatings providing corrosion protection as well as partly other functionalities.

**Keywords:** metallic coatings; anodizing layers; passivation; polymeric coatings; laser cladding; PVD; superhydrophobic coatings; composite coatings

Corrosion is an important issue in many industrial fields. It leads to high economic losses of 3–4% of the GDP of an industrialized country year by year. Adequate corrosion protection is therefore essential in many applications. Among others, coatings are by far the most important technology for corrosion protection of metallic surfaces.

In the very traditional field of coatings for corrosion protection in the last years a deeper understanding of mechanisms of the protective action and corrosion mechanisms of and below protective coatings has been gained. This was necessary due to upcoming environmental and health issues for some well-established compounds used in former coating systems, e.g., lead or chromates, which have been banned from industrial application. This lead as well to a large amount of research in the field of advanced coating systems for corrosion protection.

This situation is the case for all the different types of protective coatings that are typically used. Novel metallic coatings, e.g., novel zinc alloys are under development, as well as novel pretreatment systems or passivating chemicals avoiding the use of chromates. The upcoming chemical nanotechnology fosters the development of hybrid or inorganic sol-gel coatings, as well as of nanoparticles and nanocapsules to be used as fillers in coating systems. This has also led, in recent years, to the development of novel self-healing and smart coatings. Furthermore, nowadays, bio-based substances are becoming increasingly used for organic coatings. Last but not least, new anodizing processes have also been developed in the frame of an increased use of light metals for light weight construction.

The special issue "Advanced Coatings for Corrosion Protection" has been proposed as a means to present recent developments on any of these types of advanced coatings for corrosion protection. Thus, 15 contributions have been collected on metallic, inorganic, polymeric and nanoparticle enhanced coatings providing corrosion protection as well as partly other functionalities.

Among all of them, inorganic coatings stand out for the number of contributions being submitted to this special issue; however, these are of many different types and for different applications. The thinnest but often quite effective type of an inorganic coating may be a passivating oxide layer. The most commonly used passivating agent to develop such an oxide layer is nitric acid. Lara-Banda et al. investigated an environmentally friendly alternative for the passivation of 15-5 and 17-5PH stainless

steels based on citric acid [1]. It could be shown that, for both types of steel, the passive layer formed in citric acid as passivating solution had very similar characteristics to that formed with nitric acid.

Much thicker oxide layers can be obtained by anodizing techniques especially on light metals. Besides the conventional anodizing treatment at rather low voltages plasma-electrolytic oxidation (PEO) or micro arc oxidation (MAO), leading to ceramic oxide layers, has become an increasingly important alternative, being the topic of three contributions in this special issue [2–4]. Yao et al. prepared different MAO coatings on the magnesium alloy AZ91D [2]. It was found that especially a brown coating doped with Cu is able to significantly reduce the corrosion of magnesium parts in marine environments. Magnesium alloys to be used as biodegradable implant materials are the background of the paper by Anawati et al. [3]. Therefore, they investigated not only the corrosion resistance of PEO coatings, but also their ability to form bone mineral apatite. It was concluded that the alloying element Ca should be limited to 1 wt% as the excess tended to degrade the corrosion resistance and apatite-forming ability of the PEO coating. Kyziol et al. investigated the influence of MAO coatings on the stress corrosion cracking susceptibility of an AlMg6 alloy [4]. The pores in the MAO coating were insignificant and of limited depth. Therefore, the coating could increase the corrosion resistance.

Furthermore, inorganic coatings may be obtained from metal carbides [5], nitrides [6] or borides [7]. Jiang et al. used carburizing by a spark plasma sintering technique to enhance the erosion-corrosion resistance of tungsten in flowing coolant water [5]. W-Cr-C clad tungsten showed a different corrosion behavior than bare tungsten. Ti2AlN coatings were obtained by physical vapor deposition on ferritic steels and submitted to oxidation at a temperature of 700 ◦C by Gröner et al. [6]. The oxide scale of α-alumina was able to reduce the permeability for hydrogen significantly. Hu et al. obtained boride cermet coatings on carbon steel by a laser cladding process to improve the corrosion and wear resistance [7].

Polymeric coatings are widely used in corrosion protection and several contributions deal with this type of coatings as well [8–12]. Aging of the coating in terms of chain scission and phase separation may change the protective properties with time, as shown by Che et al. for polyurea coatings in marine atmosphere [8]. A wide variety of polymeric and hybrid systems can be chosen for protective coatings. Miela et al. showed how materials analysis and molecular dynamics simulation may help to identify the best performing coating system for erosion and corrosion protection of hydraulic water valves [9]. Polymeric coatings may also be used to add further functionalities to a barrier-type protective coating. As such, superhydrophobic properties were generated by Ou et al. on a zinc coating [10] and by Zhang et al. on the aluminum alloy AA5083 [11], providing water-repelling and long-term corrosion resistant surfaces. A very classical application of organic coatings is the insulation of buried pipelines. The paper by Hong et al. addresses this application, focusing on the additional cathodic protection design for a pre-insulated pipeline in a district heating system using computational simulation [12].

Finally, metallic coatings are widely used as noble barrier layers or as sacrificial layers providing cathodic protection to the substrate. One of the latter has been described by Tong et al. [13]. They produced a ZnAl diffusion layer on carbon steel by a mechanical energy aided diffusion method and characterized its corrosion behavior. On the other hand, barrier-type coatings may be reinforced by incorporation of nanoparticles into the coating matrix. Xu et al. demonstrated that the mechanical properties and the corrosion resistance of an aluminum foam can be improved by the electrodeposition of a NiMo coating that has been reinforced by SiC/TiN nanoparticles [14]. Furthermore, Fu et al. studied the effect of doping a NiFeCoP coating with cerium dioxide nanoparticles [15]. With an increased concentration of nano-CeO2 in the composite coating, its corrosion resistance increased as well.

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

#### **References**


© 2020 by the author. 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* **Alternative to Nitric Acid Passivation of 15-5 and 17-4PH Stainless Steel Using Electrochemical Techniques**

**María Lara-Banda 1, Citlalli Gaona-Tiburcio 1, Patricia Zambrano-Robledo 1, Marisol Delgado-E 1, José A. Cabral-Miramontes 1, Demetrio Nieves-Mendoza 2, Erick Maldonado-Bandala 2, Francisco Estupiñan-López 1, José G. Chacón-Nava <sup>3</sup> and Facundo Almeraya-Calderón 1,\***


Received: 5 May 2020; Accepted: 8 June 2020; Published: 24 June 2020

**Abstract:** Increasingly stringent environmental regulations in different sectors of industry, especially the aeronautical sector, suggest the need for more investigations regarding the effect of environmentally friendly corrosion protective processes. Passivation is a finishing process that makes stainless steels more rust resistant, removing free iron from the steel surface resulting from machining operations. This results in the formation of a protective oxide layer that is less likely to react with the environment and cause corrosion. The most commonly used passivating agent is nitric acid. However, it is know that high levels of toxicity can be generated by using this agent. In this work, a study has been carried out into the electrochemical behavior of 15-5PH (precipitation hardening) and 17-4PH stainless steels passivated with (a) citric and (b) nitric acid solutions for 60 and 90 min at 49 ◦C, and subsequently exposed to an environment with chlorides. Two electrochemical techniques were used: electrochemical noise (EN) and potentiodynamic polarization curves (PPC) according to ASTM G199-09 and ASTM G5-13, respectively. The results obtained indicated that, for both types of steel, the passive layer formed in citric acid as passivating solution had very similar characteristics to that formed with nitric acid. Furthermore, after exposure to the chloride-containing solution and according with the localization index (LI) values obtained, the stainless steels passivated in citric acid showed a mixed type of corrosion, whereas the steels passivated in nitric acid showed localized corrosion. Overall, the results of the *Rn* values derived show very low and similar corrosion rates for the stainless steels passivated with both citric and nitric acid solutions.

**Keywords:** stainless steel; passivated; electrochemical noise; precipitation hardening

#### **1. Introduction**

Corrosion in the aeronautical industry remains a major problem that directly affects safety, economic, and logistical issues. Stainless steel alloys have found increasing application in aircraft components that require great strength but can handle the increased weight. The high corrosion and temperature resistances found in stainless steel in harsh environments make it suitable for a range of

aircraft parts such as fasteners, actuators and landing gear components [1–3]. Passivation is a chemical process to remove surface contamination, i.e., small particles of iron-containing shop dirt and iron particles from cutting tools that can act as initiation sites for corrosion. This process also can remove sulfides exposed on the surface of free-machining stainless alloys. In other words, by chemically removing free contaminants from the surface of stainless steel, the passivation process adds a thin oxide layer. More chromium available from a clean surface means a thicker chromium oxide layer at the top of the stainless steel surface. Moreover, this chemically non-reactive surface means more protection against corrosion [4–10].

Precipitation hardening (PH) stainless steels (SS) are a family of corrosion resistant alloys some of which can be heat treated to provide tensile strengths of 850 to 1700 MPa and yield strengths of 520 MPa to over 1500 MPa. These alloys contain 11–18% chromium, 3–4% nickel, and smaller counts of additional metals, including aluminum, niobium, molybdenum, titanium and tungsten. Nevertheless, chromium is the alloying element responsible for the formation of the passive film [11–14]. The family of precipitation hardening stainless steels can be divided into three main types—low carbon martensitic, semi-austenitic and austenitic. These stainless steels are widely used in aerospace structural applications due to its good corrosion resistance and high strength and toughness obtained by the formation of precipitates from age-hardening treatments. Previous investigations on aeronautical-aerospace sector has shown that 15-5PH and 17-4PH steels have good corrosion resistance regarding other stainless steels [15–20].

Back in 1997, the specification QQ-P-35 for passivation of stainless steel parts was withdrawn, and replaced by specification SAE-QQ-P-35, also withdrawn in 2005. The latter was replaced by specification ASTM A967-17. This indicates that both, citric and nitric acid can be used as passivating agents for stainless steels. To be effective, the nitric acid must be highly concentrated. However, many questions has been done regarding the production of harmful to health toxic vapors generated by the use of nitric acid in passivation baths [21,22]. On the other hand, citric acid is a biodegradable alternative that does no generate hazardous waste. Although the citric acid benefits as a passivating agent are well-established, technical information about the passivation process is scarce [23,24]. In 2003, Boeing Company evaluate the use of citric acid as an alternative for steel passivation in the aeronautic industry [6]. In 2008, the National Aeronautics and Space Administration (NASA) began a research program focused on the evaluation of the use of nitric acid in the passivation process of welded parts, using the salt chamber technique [10]. Later, NASA evaluated the use of citric acid on specimens exposed under atmospheric corrosion conditions using adherence tests [21].

It is well know that aggressive ions, especially chloride ions Cl-, affect the protecting nature of the passive film on stainless steels causing its breakdown. This leads to localized attack, mainly pitting corrosion [25,26]. In the study of corrosion mechanisms, a number of electrochemical techniques such as potentiodynamic, potentiostatic, and galvanostatic polarization tests, electrochemical impedance (EIS) and electrochemical noise (EN) are widely used. For instance, the evaluation of important parameters such as passive range, pitting potentials, corrosion rates and transpassive regions are studied using potentiodynamic polarization curves (PPC). Bragaglia et al. [27] studied the potentiodynamic polarization behavior of passivated citric and nitric acid baths) and unpassivated AISI 304 stainless steel samples after 1 h in 3.5 wt. % NaCl solution. The passivation treatment largely increased the pitting potential, particularly in the case of nitric acid. After 24 h exposure, electrochemical behavior for the nitric acid and the citric acid passivated samples were almost identical.

Electrochemical noise is a technique that does not alter the natural state of the system, since no external disturbance is applied [28]. This technique reflect random or spontaneous events of current and/or potential fluctuations. Under open-circuit conditions, these fluctuations appear to be related to variations in the rates of anodic and cathodic reactions causing small transients as a result of stochastic processes such as breakdown and repassivation of passive films and formation and propagation of pits. The fluctuations of current between two nominally identical electrodes as well as their potential versus a reference electrode (three electrode system) are recorded as time series, and by using several methods to analyze noise data, an understanding of the corrosion process occurring can be determined. The EN data can be analyzed by several methods. Perhaps the most commonly used are those related to frequency domain (power density spectral or spectral analysis), time domain (statistical methods as skewness, kurtosis, localization index (LI), and the variation of signal amplitude with time) and time-frequency domains [29,30]. Suresh and Mudali [31] studied the corrosion of UNS S30403 stainless steel in 0.05 M ferric chloride (FeCl3) by spectral, statistical, and wavelet methods to deduce the corrosion mechanism. They found a good correlation of roll-off slopes derived from power spectral analysis and statistical parameters such as standard deviation, localization index (LI), and kurtosis with pitting as the corrosion mechanism. These authors reported a localization index (LI) in the range from 0.7 and 1. LI values of 0.1 to 1 has been attributed to pitting corrosion and hence the mechanism of corrosion was attributed to pitting attack [X]. Ortiz Alonso et al. [32], studied the stress corrosion cracking (SCC) behavior of a supermartensitic stainless steel by EN. They found that the LI value increased during the straining of specimens (in the range from 0.1 to 1), indicating the presence of localized events such as pits or cracks regardless of the susceptibility of the steel to stress corrosion cracking. In spite of some of its drawbacks, other studies also have found a good relationship between the LI parameter and pitting corrosion [33,34]

The aim of the present work is the study of the electrochemical behavior of 15-5PH and 17-4PH stainless steels passivated in nitric and citric acid and exposed to a 5 wt. % NaCl aqueous solution by PPC and EN.

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

#### *2.1. Materials and Samples Preparation*

The materials used in this work were 15-5PH and 17-4PH stainless steels used in the as received condition. The chemical composition of these steels was obtained by atomic absorption spectrometry, see Table 1.


**Table 1.** Chemical composition of the used stainless steels (wt. %).

Stainless steel samples were machined as cylindrical coupons, according to ASTM A380-17 [35]. The specimens were polished with SiC grit paper till 4000 grade, followed by ultrasonic cleaning in ethanol and deionized water for about 10 min each.

#### *2.2. Passivation Process*

The passivation process was carried out under the specification ASTM A967-17 [36]. Gaydos et al. [21] reported that extended passivation treatments give a better protection against corrosion for a series of stainless steels. In the present work, two passivation baths (a) nitric acid (20%v) and (b) citric acid (15%v) solutions were used. A constant temperature of 49 ◦C was maintained along the passivation process. Specimens were immersed in the solutions for 60 and 90 min. Table 2 show the passivation exposure conditions for each type of steel.


**Table 2.** Passivation at a temperature of 49 ◦C.

*2.3. Electrochemical Techniques*

In order to assess the corrosion behavior of passivated specimens (exposed area 4.46 cm2), two electrochemical techniques were used: EN and PPC. The electrolyte was a 5 wt. % NaCl aqueous solution and all tests were carried out at room temperature.

#### 2.3.1. Electrochemical Noise (EN)

This technique was carried out under ASTM G199-09 standard [37]. The experimental setup for EN measurements is schematically depicted in Figure 1. Here, two nominally identical electrodes (passivated stainless steels) as working electrodes (WE1 and WE2) were connected to measure the electrochemical current noise (ECN), whereas the electrochemical potential noise (EPN) was measured by connecting one working electrode to a saturated calomel reference electrode.

**Figure 1.** Experimental set up for electrochemical noise (EN) measurements.

The current and potential electrochemical noise was monitored as a function of time under open circuit condition for each particular electrode–electrolyte combination, using a Gill-AC (Alternating Current) potentiostat/galvanostat/ZRA (Zero Resistance Ammeter) from ACM Instruments. Electrochemical noise measurements started one after the open circuit potential stabilized (about 1 h after immersion in the electrolyte). Since the EN technique involves mostly non-stationary signals, trend removal was carried out. In each experiment, 1024 data were measured with a scanning speed of 1 data/s. The time series in current and potential were visually analyzed to interpret the signal transients and define the behavior of the frequency and amplitude of the fluctuations as a function of time. Resistance noise (*Rn*) data were obtained and used to calculate the corrosion rate according to Equation (1),

$$R\_{\pi} = \frac{\sigma\_{E}}{\sigma\_{I}} \tag{1}$$

where σ*<sup>E</sup>* is the standar deviation of potential noise, and σ*<sup>I</sup>* is the standar deviation of current noise after trend removal. The LI, defined by Equation (2), is a parameter used to estimate, as a first approximation, the type of corrosion occurring in a given system [38–40]. LI values approaching zero, indicates uniform (general) corrosion; values in the range from 0.01 to 0.1 indicates mixed corrosion, whereas values from 0.1 to 1 correspond to pitting corrosion.

$$\text{IL} = \frac{\sigma l}{I\_{\text{RMS}}} \tag{2}$$

where *IRMS* is the root mean square value of the corrosion current noise.

2.3.2. Potentiodynamic Polarization Curves (PPC)

This technique was carried out according to ASTM G5-13 [41] and ASTM G102-89 standards [42]. Here, a conventional three-electrode cell configuration was used, see Figure 2.

**Figure 2.** Conventional three-electrode cell configuration used in the potentiodynamic polarization curves (PPC) tests.

Potentiodynamic polarization curves were recorded in 5 wt. % NaCl aqueous solution at room temperature in a Gill-AC potentiostat/galvanostat from ACM Instruments. The potential scan was carried out from −1000 mV to +1200 mV, at a scan rate of 60 mV/min. A saturated calomel electrode (SCE) and a platinum wire were used as reference electrode and counter electrode, respectively. The working electrode (passivated sample) was hold for about 1 h at open circuit potential before tests.

#### **3. Results**

#### *3.1. Electrochemical Noise*

Figures 3 and 4 show the current and potential time series recorded for 15-5PH and 17-4PH stainless steel passivated in citric and nitric acid solutions at 60 and 90 min, respectively. Figure 2 shows that under passivation conditions at 60 and 90 min in citric acid, the passivated 15-5PH and 17-4PH stainless steel specimens did not present current fluctuations in time, this indicating that the specimens are in passive conditions; also, the potential noise signals remained constant without fluctuations in time (Figure 3a). The 17-4PH sample passivated for 60 min has higher current demand with low amplitude and high frequency transients, while the potential for this alloy has more active potentials (Figure 3d). For both types of stainless steel, the current-potential time series after 1000 s it has a tendency towards passivation.

Windowing analysis of electrochemical current noise between 0 and 200 s (Figure 3b) show no current increase for the 15-5PH samples passivated at 60 and 90 min. The 17-4PH steel passivated for 60 min, shows some transients of low amplitude and frequency, while for the 90 min passivation treatment only one anodic transient of high amplitude and low frequency was recorded 20 s after the start of the test. Another windowing analysis of current noise signal was performed between 900 and 1024 s (Figure 3c). For both types of stainless steel, irrespective of passivation conditions, no current fluctuations were observed. In some way, this behavior indicates stability of the passive layer.

For both types of stainless steel under passivation conditions, windowing analysis from 0 to 200 s and from 900 to 1024 s did not show frequency or amplitude transients, confirming the stability of potentials (Figure 3e,f). It is worth noting that the potentials of the 17-4PH samples are more negative than those recorded for the 15-5PH samples.

**Figure 3.** Electrochemical current and potential noise-time series for 15-5PH and 17-4PH samples passivated in citric acid at 49 ◦C, exposed in a 5 wt.% NaCl solution (**a**,**d**). Windowing of electrochemical current noise (ECN) from 0–200 and 900–1024 s (**b**,**c**); windowing of electrochemical potential noise (EPN) from 0–200 and 900–1024 s (**e**,**f**).

For the 15-5PH and 17-4PH stainless steels passivated in nitric acid, Figure 4 shows the current and potential noise time series recorded. The 17-4PH sample passivated for 90 min, show a decreases in current noise as a function of time; while the potential noise shifts to noble values, indicating stability of the passive layer. A similar behavior was observed for the 15-5PH samples passivated for 60 min. The 15-5PH and 17-4PH samples passivated for 90 and 60 min show a small current demand during the first 300 and 700 s. Afterwards, no significant current or potential fluctuations (transients) were recorded, indicating stabilization of the passive layer (Figure 4a,d).

Windowing analysis of electrochemical current noise between 0 and 200 s (Figure 4b), shows a small increase in in current demand for the 15-5PH and 17-4PH samples passivated for 90 and 60 min, respectively (Figure 4b). From 900 to 1024 s, a windowing analysis of current noise did not show current transients (Figure 4c). Windowing analysis from 0 to 200 s and from 900 to 1024 s did not show frequency or amplitude transients, confirming the stability of potentials (Figure 4e,f). It is interesting to note that, irrespective of the time of passivation treatment, more noble potentials were attained by the 15-5PH stainless steel, in comparison with the 17-4PH steel.

**Figure 4.** Electrochemical current and potential noise-time series for 15-5PH and 17-4PH samples passivated in nitric acid at 49 ◦C, exposed in a 5 wt. % NaCl solution (**a**,**d**). Windowing of ECN from 0–200 and 900–1024 s (**b**,**c**); windowing of EPN from 0–200 and 900–1024 s (**e**,**f**).

The EN parameters derived from the statistical analysis of current and potential time series measurements are shown in Table 3. The *icorr* value obtained from noise resistance (*Rn*) for the samples passivated in citric acid is in the order of 10−<sup>4</sup> (mA/cm2), whereas for the samples passivated in nitric, *icorr* values about 10−<sup>5</sup> (mA/cm2) were recorded. The very low values for *icorr* obtained for both passivating agents indicate that citric acid could be a potential replacement for nitric acid as passivating agent. Information regarding the type of corrosion that could be occurring is given by the LI parameter. As can be seen from Table 3, the stainless steels passivated in citric acid solution mainly show a mixed corrosion type, whereas the stainless steels passivated in nitric acid solution the LI values indicates localized corrosion.


**Table 3.** Electrochemical noise parameters at various conditions in 5 wt. % NaCl at 49 ◦C.

#### *3.2. Potenciodynamic Polarization*

The corrosion kinetic behaviour using potentiodynamic polarization can be observed through cathodic and anodic reactions in polarization curves. Corrosion rate in terms of penetration (mm/sec) is one of the main parameters obtained by potentiodynamic polarization curves, according to Faraday's law (Equation (3)) [40,42–44].

$$\text{Corrosion rate} = K\_1 \frac{i\_{corr}}{\delta} \to \mathcal{W} \tag{3}$$

The potentiodynamic polarization curves obtained for the 15-5PH and 17-4PH stainless steels passivated for 60 min and 90 min in (a) citric acid and (b) nitric acid, and immersed in 5 wt. % NaCl solution are shown in Figure 5.The results for citric acid passivation (Figure 5a) show that the lower Ecorr value was recorded for the 17-4PH sample passivated for 90 min, while the 15-5PH sample passivated for 90 min has the highest Ecorr. Pitting potentials (Epitt) were in the range from 42 mV up to 147 mV. This last value was recorded for the 15-5PH steel passivated for 90 min, this being the best treatment for nitric acid passivation, also corroborated by the lower corrosion rate obtained. For nitric acid passivation conditions, Figure 4b show that the Epitt was largely improved, particularly for the 15-5PH steel passivated for 90 min, and also has the lower corrosion rate in this condition. The lower Epitt value recorded was given by the 17-4PH passivated during 90 min, also giving the highest corrosion rate. On the whole, irrespective of the type of PH stainless steel used, the nitric acid passivation treatment largely increases the pitting potentials compared with the citric acid treatment.

The parameters (Ecorr, Epitt, *icorr*, and corrosion rate (C.R.)) obtained from the polarization potentiodynamic curves are summarised in Table 4. Very low values of corrosion rate (within the same order of magnitude) were recorded for both 15-5PH and 17-4PH steels, irrespective of the passivation treatment conditions.

**Figure 5.** Potentiodynamic polarization curves of 15-5PH and 17-4PH stainless steel passivated in (**a**) citric acid and (**b**) nitric acid, exposed in a 5 wt. % NaCl solution at 49 ◦C.

**Table 4.** Potentiodynamic polarization parameters in stainless steels passivated at 49 ◦C, in 5 wt. % NaCl.


#### **4. Discussion**

Several EN procedures correlating timed dependent fluctuation of current and potential during the corrosion process have been used to indicate the type of corrosion occurring. For instance, it is well recognized that the main source of electrochemical noise is the passive film breakdown process and repassivation process [45–49].

For the passivated 15-5PH and 17-4PH stainless steels in this work, the electrochemical potential time series recorded under nitric acid passivation show a passive region from 0 to 120 mV, whereas for citric acid passivation conditions, the passive region goes from −300 to 25 mV. Thus, passivation in nitric acid occurs at more positive (noble) potentials that in citric acid solutions. To some extent, this might indicate that the passive layer is more stable at more noble potentials [25]. Hence, higher corrosion resistance (*Rn*) values could be expected for passivation in nitric acid solutions [50]. To some extent, the results of *Rn* in Table 3 confirm this.

As a first approach, to assess the more likely type of corrosion occurring for the 15-5PH and 17-4PH stainless steels passivated in both citric acid and nitric acid solutions, the LI parameter was evaluated from the electrochemical noise data, and was found to be in the range from 0.03 to 0.249, see Table 3. From the LI values obtained for each passivating bath, the corrosion type occurring in citric acid passivation conditions can be attributed to mixed corrosion, whereas for nitric acid passivation conditions, the corrosion type could be attributed to pitting corrosion. LI has been used by several research groups for determining corrosion types under several conditions [51–56].

The use of LI to determine corrosion types has been the subject of many discussions among investigators on the data treatment and interpretations using LI [29,30,38,57–59]. Since the mean of the noise data (detrended) would be negligible, the standard deviation and root mean square current noise would converge to the same value and, hence, the LI evaluated from the data would be unity, irrespective of the corrosion type. Cottis [60] indicated that LI for identification of localization of corrosion is unduly influenced by the mean current and hence less reliable. In the present study, and as a first attempt, the LI parameter was estimated. Of course, it is recognized that in the study of stainless steels such as those in the present work, investigation of procedures based on the frequency domain and time-frequency domain deserves further attention.

The potentiodynamic polarization curves show passivation behavior for the 15-5PH and 17-4PH steels passivated in citric and nitric acid solutions. The passive zones on stainless steels are commonly made up of primary and secondary zones, which are formed before and after transpassivation, respectively. Potentials above Epitt causes a rapid dissolution [61–63]. The passive zone involves the formation of iron and chromium oxide films [61,64,65]. Hence, selective dissolution on the surface of the alloy generates a surface enrichment of Cr3<sup>+</sup> giving rise to Cr(OH)3, as shown in Equation (4). Further dissolution of the hydroxide leads to the formation of a continuum layer of Cr2O3, according Equation (5) [66,67].

$$\rm Cr^{3+} + 3OH^- \rightarrow \rm Cr(OH)\_3 \; + \rm 3e^- \tag{4}$$

$$\rm Cr(OH)\_3 + \rm Cr + 3OH^- \rightarrow \rm Cr\_2O\_3 + 3H\_2O + 3e^- \tag{5}$$

It has been argued that the anodic reactions during the film growth period are mainly from the oxidation of iron and chromium The following equations indicate the oxidation reactions of iron [68]:

$$2Fe + 8OH^- \rightarrow Fe\_3O\_4 + 4H\_2O + 8e^- \tag{6}$$

$$2Fe\_3O\_4 + 2OH^- + 2H\_2O \to 6FeOOH + 2e^- \tag{7}$$

$$2Fe\_3O\_4 + 2OH^- \rightarrow 3Fe\_2O\_3 + H\_2O + 2e^- \tag{8}$$

For the nitric acid passivation treatment, the transpassive region is above 200 mV vs. ECS, whereas for citric acid passivation conditions, the transpassive region is above 50 mV vs. ECS. The passive film formed under nitric acid passivation conditions has higher Epitt values, in comparison with the Epitt values obtained under citric acid passivation conditions. This fact can be seen as a potential disadvantage for the citric acid treatment. Nevertheless, the corrosion rates obtained for both passivation treatments (Table 4) are very low and similar. Thus, for the PH stainless steel used in this work, citric acid passivation treatments can be as effective as nitric acid passivation treatments.

#### **5. Conclusions**

In this work, samples of 15-5PH and 17-4PH stainless steel were passivated in (a) citric acid and (b) nitric acid baths and exposed in a 5 wt. % NaCl solution. Their electrochemical behavior was studied by electrochemical noise and potentiodynamic polarization.

EN results show that, for citric solution passivation baths, the stabilization of the passive layer occurs at more active potentials compared to the stabilization potentials for nitric acid passivation baths. From noise resistance (*Rn*) data, very low corrosion rate values were derived for the PH stainless steels passivated in both (citric and nitric) passivating treatments.

Statistical evaluation of the time record was carried out and the localization index (LI) parameter was evaluated. According to the LI results, the PH stainless steels passivated in citric acid solution mainly show a mixed corrosion type, whereas LI values for the PH stainless steels passivated in nitric acid solution indicates localized corrosion.

In general, potentiodynamic polarization results indicated that, irrespective of the type of PH stainless steel used, the nitric acid passivation treatment largely increases the pitting potentials in comparison with the citric acid treatment. Also, for both passivation treatments, very low corrosion rate values (in the order of 10−<sup>7</sup> mm/year) were recorded for both 15-5PH and 17-4PH steels.

On the whole, citric passivation treatments on PH stainless steels could be a green alternative route to the currently employed nitric passivation treatments.

**Author Contributions:** Conceptualization, M.L.-B., F.A.-C.; Methodology, P.Z.-R., M.D.-E, F.E.-L., E.M.-B.; Data Curation, C.G.-T., J.A.C.-M., and D.N.-M.; Formal analysis, J.G.C.-N., and F.A.-C.; Writing—Review and Editing, M.L.-B., J.G.C.-N., F.A.-C. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Conacyt, Proyect "Estudio de las propiedades electroquímicas y mecanismos de crecimiento de la película de pasivación de aceros inoxidables endurecibles por precipitación en ambientes ácidos", con clave No. A1-S-8882 and UANL (Dirección de Investigación).

**Acknowledgments:** The authors acknowledge The Academic Body UANL—CA-316 "Deterioration and integrity of composite materials".

**Conflicts of Interest:** All authors of this article declare no conflicts of interest.

#### **References**


© 2020 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* **Preparation and Corrosion Behavior in Marine Environment of MAO Coatings on Magnesium Alloy**

**Yuhong Yao 1,\*, Wei Yang 1,\*, Dongjie Liu 2, Wei Gao <sup>1</sup> and Jian Chen <sup>1</sup>**


Received: 1 December 2019; Accepted: 10 January 2020; Published: 12 January 2020

**Abstract:** To improve the corrosion performance of magnesium alloys in the marine environment, the MAO, MAO–Cu2CO3(OH)2·H2O and MAO–Cu2P2O7 ceramic coatings were deposited on AZ91D magnesium alloys in basic electrolyte and the discoloration mechanism of the Cu-doped MAO coatings and the corrosion behavior of the three MAO coatings in the artificial seawater solution were investigated by SEM, EDS and XPS. The results indicated that the formation and discoloration mechanism of the brown MAO ceramic coatings were attributable to the formation of Cu2O in the coatings. Though the three MAO coatings had a certain protective effect against the corrosion of AZ91D substrate in the artificial seawater, the distinctive stratification phenomenon was found on the MAO–Cu2P2O7 coated sample and the corrosion model of the MAO–Cu2P2O7 coatings in the immersion experiment was established. Therefore, the brown Cu-doped MAO coatings were speculated to significantly reduce the risk of the magnesium parts in marine environments.

**Keywords:** magnesium alloy; MAO coating; corrosion behavior; stratification phenomena; marine environments

#### **1. Introduction**

The magnesium alloy is the lightest structure metal material, and is considered as the green engineering material in the 21st century [1]. Now it is widely applied to the electronic industry, aerospace industry, and auto industry [2,3]. However, for its poor corrosion resistance, magnesium alloy, particularly as the magnesium alloy parts for outdoor applications, is confronted with great challenges [4–7]. Many literatures have verified that the surface modification technique, such as chemical conversion coatings, anodic oxidation, micro arc oxidation (MAO), organic coatings, vapor phase processes, etc, is an effective way to change the surface composition and improve its corrosion resistance of magnesium alloy [8–12]. MAO is a newly developed technology for the preparation of ceramic coatings on aluminum, magnesium or titanium alloys to improve the corrosion resistance [10,13–16]. At present, the synthesis of white MAO coatings on magnesium alloy is a mature and universal technology [16], but it has the same disadvantage as the chemical conversion coating technology, which often causes the formation of the light spots on the surfaces of magnesium diecast components and hardly meets the market demand of 3C (computer, communication and consumer) electronic products. It has been reported that the colored MAO coatings can be formed by adding metal salts in the electrolytes [17–24]. The black MAO coating containing V2O3 can be obtained on an aluminum alloy surface by adding ammonium metavanadate into the commonly used (NaPO3)6 (sodium hexametaphosphate) and Na2SiO3 solution [21] and the black MAO coating with excellent properties can also be prepared in the electrolyte with dichromate addition [22]. Moreover, it has been reported that by adding potassium fluotitanate or sodium stannate into the base electrolyte, a yellow or grey MAO coating can be formed on the surface of Mg alloys, respectively [23,24].

*Materials* **2020**, *13*, 345; doi:10.3390/ma13020345 www.mdpi.com 19 /journal/materials

Brown is a very important and common decorative color. Lee et al. [25] has reported that with the addition of 3% and 5% Cu in the base electrolyte, the color of the MAO coating changes from brown to dark brown and the corrosion resistance of the AZ91 alloy is significantly improved after being treated with the micro-arc oxidation process, but the corrosion process and coloring mechanism of this brown MAO coating are still not clear. Furthermore, the magnesium alloy parts with MAO coating for lightweight are sometimes exposed to marine environment and the marine corrosion behavior of the coatings has not been clarified. As a result, the application of magnesium alloy in a marine environment is seriously restricted. In this study, a convenient process to fabricate the MAO coatings with brown color on Mg alloys is introduced by adding alkaline copper carbonate and copper pyrophosphate in the electrolyte and the white and two brown MAO coatings are prepared. Then, the microstructure, formation mechanism and seawater corrosion behaviors of the coatings are investigated in detail. Finally, the seawater corrosion mechanism of the MAO coatings is revealed, which is helpful for the surface protection of magnesium alloy applied in marine environments.

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

AZ91D magnesium alloys were used as the substrate discs in the size of φ30 mm × 5 mm and its nominal chemical composition (in wt. %) was Al 8.5–9.5 %, Zn 0.5–0.9 %, Mn 0.17–0.27 %, Cu ≤0.01, Ni ≤0.01, Si ≤0.01, Fe ≤0.004 and Mg balance. Before the micro arc coatings, the specimens were prepared by means of standard metallographic procedure, such as coarse grinding, accurate grinding, polishing with alumina waterproof abrasive paper up to 1200 grit and then ultrasonically degreased in acetone for 10 min followed by rinsing with distilled water.

The MAO coatings were prepared on the specimen surface by using of micro arc oxidation equipment (JHMAO-60, China) with the constant voltage of 420 V, the frequency of 400 Hz, the duty cycle of 10% and the treatment time of 8 min. The base electrolyte solution was composed of 8.0 g/L sodium silicate (Na2SiO3·9H2O), 5.0 g/L potassium hydroxide (KOH), 5.0 g/L potassium fluoride (KF), 1.0 g/L EDTA (C10H16N2O8) and 3.5 g/L potassium sodium tartrate (C4H4O6KNa·4H2O). The two Cu-doped brown coatings were prepared by respectively adding 2.5 g/L basic copper carbonate (Cu2CO3(OH)2·H2O) and 2.5 g/L copper pyrophosphate (Cu2P2O7) into the base solution and the temperature of the electrotype with pH of about 13 was kept below 35 ◦C during the MAO process.

The thickness of the MAO coatings was measured by using TT240 eddy current thickness meter with an accuracy of 0.1 μm. Six measurements were carried out evenly on the whole sample surface. The surface morphologies and element distribution of the MAO coatings were analyzed by scanning electron microscope (SEM) with Oxford Inca X-Max energy dispersive spectrometry (EDS). X-ray photoelectron spectroscopy (XPS) with Al (mono) Kα irradiation at pass energy of 160 eV (AXIS UTLTRADLD) was used to characterize the chemical bonds of the coatings. The binding energies were referenced to the C 1 s line at 284.6 eV from adventitious carbon. The corrosion behavior of the coated AZ91D magnesium alloy was evaluated by the immersion tests in the artificial seawater, whose composition was shown in Table 1. Before the immersion test, the three MAO-coated specimens were treated with epoxy resin to avoid the effect of defects at the edge of the samples, then immersed in the artificial seawater solution for 14 days and the corrosion morphologies of the samples were observed by SEM.

**Table 1.** Composition of artificial seawater in immersion test.


#### **3. Results and Discussions**

The MAO–Cu2CO3(OH)2·H2O and MAO–Cu2P2O7 coatings on AZ91D magnesium alloys in Figure 1 are prepared with the thicknesses of 8.4 μm, 9.1 μm and 9.6μm, respectively. It reveals that for the more intense micro arc discharge many defects are observed at the edge of the sample in Figure 1b. The microstructures of the three MAO coatings are given in Figure 2. It can be seen from Figure 2 that the surfaces of the three MAO coatings are characterized by lots of micropores with the size range from submicron to several micro scale, and different from the MAO and MAO–Cu2P2O7 coatings, there are some bumps on the MAO–Cu2CO3(OH)2·H2O coating surface, which are formed by the spark discharge and gas bubbles throughout the discharge channels during MAO process [26,27]. So it reveals that the addition of Cu2CO3(OH)2·H2O into the base electrolyte results in a strongly intense micro arc discharge, which promotes the formation of large molten deposits (Figure 2c,d). The elemental compositions of the three MAO coatings are examined by EDS in Table 2. Carbon (C) is considered to be an impurity from the atmosphere or the electrolyte, F and Na are also presumed to originate from the electrolyte or the AZ91D alloy substrate, P and Si are from the electrolyte and Zn comes from the substrate which indicates that the thickness of the three MAO coatings is very thin. Thus the three MAO coatings are mainly composed of Mg, O, Si and a little amount of C, F, Na species. The presence of Si and O reveals that the components of the electrolyte have intensively incorporated into the micro arc oxidation reactions to from the ceramic coatings [16]. Moreover, by addition of Cu2CO3(OH)2·H2O or Cu2P2O7 into the electrolyte, a very small amount of Cu or Cu, P has been respectively doped into the MAO–Cu2CO3(OH)2·H2O and MAO–Cu2P2O7 coatings to make the color of the coatings change from white to brown. The discoloration mechanism of the MAO coatings will be further discussed by using SEM + EDS and XPS.

**Figure 1.** Macrograph of three MAO coatings on AZ91D magnesium alloy, (**a**) MAO, (**b**) MAO– Cu2CO3(OH)2·H2O and (**c**) MAO–Cu2P2O7.

**Figure 2.** *Cont.*

**Figure 2.** Surface morphologies of three MAO coatings on AZ91D magnesium alloy, (**a**) and (**b**) MAO, (**c**) and (**d**) MAO–Cu2CO3(OH)2·H2O and (**e**) and (**f**) MAO–Cu2P2O7.

**Table 2.** EDS results of the MAO, MAO–Cu2CO3(OH)2·H2O and MAO–Cu2P2O7 coatings on AZ91D magnesium alloy (at. %).


The surface morphologies of the MAO–Cu2CO3(OH)2·H2O coatings formed at different oxidation times are shown in Figure 3 and the corresponding EDS analysis results are listed in Table 3. As shown in Figure 3a and Table 3, when the oxidation time is about 70 s, the surface morphology of the AZ91D substrate is inhomogeneous at the moment of starting arc with two distinct regions: region I with a damaged area caused by electrical break-down involving amount of O, F, Si and Cu elements from the electrolyte and Mg and Zn alloying species from the substrate, and region II with a smooth surface morphology and the scratches in the substrate including lower content of O, Si, Cu and much higher content of Mg than those in region I, which indicates that the electrical breakdown phenomenon does not occur in region II. With the prolongation of oxidation time, the surface of the samples presents a typical porous feature and the pore size of the MAO coatings increases with the oxidation time in Figure 3b–d; a similar micro arc process and the mechanism of the pore initiation and the pore development are reported by some literatures [10,21]. However, it is worth noting that the concentrations of O, F, Si, Cu, Mg and Zn elements are similar in the MAO–Cu2CO3(OH)2·H2O coatings, the color of the coatings gradually becomes deeper with oxidation time, so it is very meaningful to analyze the discoloration mechanism of the coatings.

**Figure 3.** Surface morphologies of MAO–Cu2CO3(OH)2·H2O coatings formed with different oxidation time on AZ91D magnesium alloy, (**a**) 70 s, (**b**) 100 s, (**c**) 120 s and (**d**) 180 s.

**Table 3.** EDS results of MAO–Cu2CO3(OH)2·H2O coatings formed with different oxidation time on AZ91D magnesium alloy (at. %).


The chemical states of Cu, Mg and Si are investigated by using XPS in Figure 4. From the wide spectra demonstrated in Figure 4a, Mg, O, Si and F elements are all found in the three MAO coatings and Cu is detected in the MAO coating with MAO–Cu2CO3(OH)2·H2O or Cu2P2O7 addition by XPS. The Cu2p3/2 spectrum consists of two peaks at the binding energies of 932 eV associated with Cu2O and 939.7 eV corresponding to CuF2 in the brown MAO coating with Cu2CO3(OH)2·H2O addition in Figure 4b. Figure 4c,d illustrate that the Mg2p peak at the binding energy of 47.8 eV and the Si2p peak of 100.25 eV are individually assigned to MgO and SiO2. So, it can be concluded that the solute ions (such as Cu and Si) from the electrolyte are involved in the growth process of the MAO coatings, and the same results are found in the growth process of the MAO coatings on Ti substrate [28,29]. Therefore, it is the formation of the red Cu2O in the MAO coatings with MAO–Cu2CO3(OH)2·H2O addition that results in the discoloration of the coatings.

**Figure 4.** (**a**) XPS survey spectra of three MAO coatings, (**b**) Cu2p, (**c**) Mg2p and (**d**) Si2p typical high-resolution XPS spectrum of MAO–Cu2CO3(OH)2·H2O coatings.

Figures 5 and 6 are the macro and micro surface appearances of the three MAO coatings immersed in the artificial seawater solution for 14 days, respectively. Compared with Figure 1, it can be observed from Figure 5 that the color of the three MAO coatings obviously changes and some corrosion products are formed on the immersed coatings. The corrosion is more serious at the edge of the three MAO samples due to the defects caused by micro arc discharge (Figure 1b) or the damage during the embedding. It has been reported that once the corrosion reaction is initiated on the sample through a pit or minute pore, the corrosive medium can come into contact with the substrate to form some corrosion pits [30]. The elemental compositions of the three MAO coatings detected by EDS after the immersion test are listed in Table 4. The three corroded MAO coatings are mainly composed of Mg from the substrate and the artificial seawater and O from the electrolyte during the micro arc discharge, and Si, Cu and F from the electrolytes and Cl, K and Ca elements from the artificial seawater are also found in the corroded MAO coatings, which indicate that the MAO coatings are not completely destroyed. From Figure 6, it is quite clear that the microstructures of the MAO coatings have a significant change before and after the immersion test. The white MAO coating exhibits a relatively uniform surface appearance with a high degree of porosity, some cracks and corrosion products as shown in Figure 6a, which indicates that the AZ91D substrate is still protected by the coatings. The MAO–Cu2CO3(OH)2·H2O coatings are relatively rougher and exhibit a stacking structure with limited number of pores in Figure 6b. For the MAO–Cu2P2O7 coatings, besides many micropores and some micro-cracks, there is a sedimentary layer with a lot of cracks in the coatings in Figure 6c and SEM morphology and EDS analysis results of this MAO coating after the immersion test are shown in detail in Figure 7 and Table 5. It can be seen from Figure 7 that the corroded coatings are divided into three regions: region I (the inner layer near the Mg substrate), region II (the middle layer attached to the

surface coating) and region III (the top layer). From Table 5, Mg, O, F and Si elements are found in the region I and the atom percentage of these elements is similar to those in the MAO coating before the immersion test. Moreover, Cl and Ca elements from the artificial seawater are not observed, indicating that the MAO ceramic coatings in region I have not been destroyed. Region II is a dense layer attached to the MAO coating, where the contents of Mg, Si and F elements dramatically decrease whereas the contents of O element significantly increase compared with that in region I; both Cl and Ca elements have been detected. It indicates that there is the interaction between the MAO coating and the corrosive medium. In the case of region III, this layer is relatively thick and composed of some loose and porous structure sediments, mainly containing O and Ca elements. It is well known that the deposition of corrosion products can hinder the transfer of the charge and increase the inhibition of corrosion, so it is believed that this thick sediment layer is very helpful to improve the corrosion behavior of the MAO–Cu2P2O7 coatings in artificial seawater.

**Figure 5.** Macrograph of three MAO coatings, (**a**) MAO, (**b**) MAO–Cu2CO3(OH)2·H2O and (**c**) MAO–Cu2P2O7, on AZ91D magnesium alloy after 14 days immersion test in the artificial seawater.

**Figure 6.** Micro morphologies of three MAO coatings by SEM after immersion test in the artificial seawater for 14 days, (**a**) MAO, (**b**) MAO–Cu2CO3(OH)2·H2O and (**c**) MAO–Cu2P2O7.


**Table 4.** EDS results of the three MAO coatings immersed in artificial seawater for 14 days (at. %).

**Figure 7.** SEM surface morphologies and EDS results of MAO–Cu2P2O7 coating after immersion test.

**Table 5.** EDS results of the three layers of the MAO–Cu2P2O7 coatings immersed in artificial seawater for 14 days (at. %).


From Figure 7 and Table 5, it has been learnt that the inner layer of the coatings is not destroyed near the substrate during the corrosion process, a middle dense layer is formed in the corrosion solution and the top thick deposition layer is of porous characteristics. Thus, a schematic diagram of the MAO–Cu2P2O7 coatings during the immersion corrosion in the artificial seawater solution is shown in Figure 8. Due to the eruption and condensation of molten materials caused by micro arc discharge, the MAO coating normally has porous characteristics [31]. As a result, some interconnected micro closed-pores exist inside the MAO ceramic layer in Figure 8a. As shown in Figure 8b, the corrosion medium, especially Cl-, can seep into the interface between the electrolyte and the coatings through the micropores during the corrosion process; these micropores are exposed to the corrosive medium due to the oxygen concentration polarization between the interior holes and the interface. Then the corrosion products accumulate in the micropores and form a dense corrosion product layer on the MAO coating with the prolonging of the immersion corrosion time to prevent further corrosion of the coatings. Finally, the further interaction between the substrate and the artificial seawater has effectively been prevented by the corrosion product layer and the top calcium oxide-like thick porous layer forming on the middle dense layer in the following corrosion process in Figure 8c, which can further prevent the corrosion medium into the substrate to enhance its corrosion performance in the artificial seawater. Thus, it can be concluded that the stratification phenomena have been found on the MAO–Cu2P2O7-coated sample and similar results have also been reported in the literatures [32,33]. However, although there is no Mg detected in the top sedimentary layer, for its loose and porous structure there is maybe a risk of corrosion for AZ91D substrate with further extension of the corrosion process.

**Figure 8.** Schematic diagrams of the immersion corrosion of MAO–Cu2P2O7 coating in the artificial seawater. (**a**) Morphology of MAO coating before corrosion, (**b**) Morphology of MAO coating at the initial stage of corrosion, (**c**) Morphology of MAO coating at the latter stage of corrosion.

#### **4. Conclusions**

(1) Brown MAO coatings on magnesium alloy can be prepared in the Na2SiO3 alkaline electrolyte with copper pyrophosphate or copper carbonate as the colorant, and for strongly intense micro arc discharge, some bumps appear on the Cu-doped MAO coatings.

(2) With the increase of reaction time, Cu in the colorant is fully involved in the formation of the MAO–Cu2CO3(OH)2·H2O and MAO–Cu2P2O7 coatings on AZ91D magnesium alloys, and the discoloration mechanism of these coatings are attributed to the formation of Cu2O in the coatings.

(3) After immersion in seawater for 14 days, the stratification corrosion microstructure with the three layers named the inner layer, the middle dense layer and the top calcium oxide-like layer are formed on the MAO–Cu2P2O7 coated sample, which is helpful to prolong the service life of AZ91D magnesium alloys in artificial seawater.

**Author Contributions:** Designing the experiments, Y.Y., and W.Y.; performing the experiments, Y.Y., W.Y., D.L., and W.G.; contributing the reagents, materials and analysis tools, J.C.; analyzing the data, Y.Y., W.Y., D.L., and W.G.; writing—original draft preparation, Y.Y., and W.Y.; writing—review and editing, W.Y, and Y.Y. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Key Research and Development Plan of Shaanxi Province—Industrial project (Grant No. 2018GY-127).

**Conflicts of Interest:** The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

#### **References**


© 2020 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* **Corrosion Resistance and Apatite-Forming Ability of Composite Coatings formed on Mg–Al–Zn–Ca Alloys**

**Anawati Anawati 1,\*, Hidetaka Asoh <sup>2</sup> and Sachiko Ono <sup>2</sup>**


Received: 15 June 2019; Accepted: 10 July 2019; Published: 14 July 2019

**Abstract:** The properties of composite coatings formed by plasma electrolytic oxidation (PEO) were affected by the alloy composition. The corrosion resistance and apatite-forming ability of PEO coatings formed on Mg–6Al–1Zn–xCa alloys with a variation of Ca content were investigated. Potentiodynamic polarization and electrochemical impedance spectroscopy (EIS) measurements showed an order magnitude improvement of corrosion resistance in the AZ61 alloy as a result of the coating. A higher enhancement in polarization resistance was obtained in the Mg–6Al–1Zn–1Ca and Mg–6Al–1Zn–2Ca alloys due to thicker coatings were formed as a result of the incorporation of calcium oxide/hydroxide. However, the underlying substrates were more prone to localized corrosion with increasing Ca content. The microstructure investigation revealed an enlargement in precipitates (Al2Ca, Mg2Ca) sizes with increasing Ca content in the alloys. The growth of larger size precipitates increased the danger to micro galvanic corrosion. Apatite layers were formed on all of the coatings indicating high apatite-forming ability, but the layers formed on the Mg–6Al–1Zn–1Ca and Mg–6Al–1Zn–2Ca alloys contained higher Mg, possibly due to the accumulation of corrosion product, than that on the Mg–6Al–1Zn alloy. The alloying element Ca should be limited to 1 wt.% as the excess tended to degrade the corrosion resistance and apatite-forming ability of the PEO coating.

**Keywords:** magnesium; microstructure; coating; corrosion; polarization; apatite

#### **1. Introduction**

The use of magnesium (Mg) and its alloys for biodegradable materials rely on the surface treatment such as plating, coating, and anodizing, due to the high corrosion rate of Mg in aqueous environments [1–3]. Among the available techniques, plasma electrolytic oxidation (PEO) becomes famous for corrosion protection of Mg alloys considering its flexibility to coat complex geometry and environmentally friendly process [4,5]. The PEO coating that forms as a result of high-voltage (hundreds volt) anodization in an alkaline electrolyte, with the incorporation of the electrolyte species, provides superior corrosion and wear resistance [6–8]. A ceramic-like composite layer consisting of crystalline and amorphous oxides is developed under the exposure of a high-temperature plasma discharge. The corrosion resistance of the PEO coating depends on the processing parameters [9], alloy composition [10], and the type of electrolyte used [11]. Among various alkaline electrolytes, the PEO coating formed in phosphate-based solution exhibited the highest corrosion resistance due to the significant formation of crystalline magnesium phosphate [12]. Magnesium phosphate is a biocompatible compound and has attracted much attention during the past few years as a material for bone replacement [13]. However, both the in vitro [7,14] and in vivo [15,16] tests have indicated that the deposition of bone mineral apatite (Ca10(PO4)6) on the PEO coating was limited. The apatite-forming ability defines the bioactivity of biomaterials.

Various methods have been proposed to accelerate the bioactivity of the PEO coating in physiological solution such as by coating with apatite [17,18], incorporating Ca [19,20] or apatite [21,22] in the coating, and introducing Ca as an alloying element [23]. The unstable Ca compounds easily dissolves in the electrolyte during plasma discharge and therefore adding Ca in the electrolyte bath is not beneficial. Incorporation of apatite in the PEO coating improved the corrosion resistance, but long-term immersion in simulated body fluid (SBF) indicated no further growth of the apatite [19,20]. Apart from the coating modification during the PEO process, the effect of Ca as an alloying element in Mg in modifying the bioactivity of the coating has not been thoroughly investigated. Our previous result [23] showed an increase in the bioactivity of the PEO coating on Mg–6Al–0.26 Mn–xCa alloys with increasing Ca concentration in the alloys. An embodiment of Ca in the coating and its presence in the alloy accelerated the growth of apatite in simulated body fluid (SBF). However, short-term (8 min) grown coatings on Mg–6Al–1Zn–xCa alloys revealed no bioactivity in SBF and significantly high corrosion rate [24]. The reason for such behavior was still unclear. The alloying element Ca might have different effect on the microstructure, and hence the resulting PEO film of the Mg–Al–Zn alloys as compared to that of the Mg–Al–Mn alloys. The development of an Mg based-biodegradable implant included Zn as alloying element to improve its mechanical strength [25–27]. This work aims to throw light on explaining the corrosion behavior and apatite-forming ability of the coatings formed on the Mg–6Al–1Zn–xCa alloys as a result of alteration in the microstructure with a variation of Ca content.

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

#### *2.1. Specimen Preparation*

Rolled plates of Mg–6Al–1Zn–xCa alloys with the composition listed in Table 1 were used as substrates. For the sake of simplicity, the alloy was further designated as AZ61, AZX611, and AZX612 for the alloys containing 0, 1, and 2 wt.% Ca, respectively. The plate was cut into 5 cm<sup>2</sup> working area. For surface observation, the specimen was embedded in a resin and then ground to 1200-grit silicon carbide (SiC) paper. The specimen was then degreased in acetone in an ultrasonic bath for 3 min. The etching treatment was applied on the specimen using 4 vol.% HNO3 in ethanol for 20 s to reveal the microstructure.


**Table 1.** Chemical composition of the alloys.

#### *2.2. Plasma Electrolytic Oxidation*

To dissolve the natural oxide layer on the surface, the specimen was dipped in a mixed acid solution of 8 vol.% HNO3-1 vol.% H3PO4 for 20 s, and then in 5 wt.% NaOH solution at 80 ◦C for 1 min. The PEO coating was developed under a galvanostatic mode at 200 Am−<sup>2</sup> in 0.5 mol·dm−<sup>3</sup> Na3PO4 solution at 25 ◦C for 20 min. A regulated DC power supply was used as the current source. The specimen was placed as an anode while a pair of carbon rod was used as cathode. The voltage output was measured by digital multimeter bench from Keithley series 2700. A coating thickness gauge of dual type (SME-1) from Sanko was used for determining the coating thickness.

#### *2.3. Electrochemical Corrosion*

The electrochemical corrosion measurement was performed by conducting potentiodynamic polarization and electrochemical impedance spectroscopy (EIS) tests in a physiological solution 0.9 wt.% NaCl solution at 37 ◦C based on ASTM G5 [28]. A defined surface with an area of 2.5 cm<sup>2</sup> was exposed to the solution. The electrochemical tests were performed by using a potentiostat instrument from IviumStat. Pt wire was used as a counter electrode, and silver chloride (Ag/AgCl) was used as a reference electrode. The electrochemical cell condition and arrangement were similar to as reported earlier [24]. Potentiodynamic polarization test was conducted from 100 mV below open circuit potential (OCP) and terminated when the current output reached 30 mA at a sweep rate of 1 mV·s−1. The corrosion potential and current density were estimated by using Tafel extrapolation. The EIS was performed on the uncoated substrates at an open circuit potential (OCP) over a frequency ranging from 10−<sup>2</sup> Hz to 10<sup>4</sup> Hz. Prior to the tests, the specimen was left at an OCP for 40 min to stabilize the potential. The coated specimen did not give spectra at low frequency and therefore was analyzed at frequency range 102 to 10<sup>7</sup> Hz.

#### *2.4. Apatite-Forming Ability Test*

The apatite-forming ability test was performed by immersing the specimens individually in SBF at 37 ◦C with a surface-to-volume ratio of 20 mL/cm<sup>2</sup> for 14 days. An SBF10, with the ionic concentration shown in Table 2, was used. The solution preparation followed the Reference [29]. The fresh solution pH was adjusted to 7.4 at 37 ◦C. The solution was replaced after 1, 3, 5, 7, 10, and 12 days, and was collected each time after each replacement for further analysis. The Mg concentration dissolved in the solution was analyzed by using the titration method. A digital titration kit from Hach, which included a digital titrator, buffer solutions, and indicators, was used. The solution was titrated with 0.08 M of sodium ethylenediaminetetraacetic acid (EDTA) using a digital titrator. The method was able to detect Mg concentrations within a 1 ppm margin of error.

**Table 2.** Ionic composition of simulated body fluid (SBF).


#### *2.5. Surface Analyses*

The surface microstructure and the elemental composition was studied by an energy dispersive X-ray spectroscopy (EDS, JEOL EX-54175JMU, Tokyo, Japan) attached to the SEM (JEOL JSM-6380LA, Tokyo, Japan). The crystalline phases in the specimens were analyzed using thin coating X-ray diffraction analysis (TF-XRD, Rint 2000 Rigaku type, Tokyo, Japan) at an incident angle of 1◦. The phases existed in the specimens were analyzed by indexing the peaks in the XRD pattern by referring to the JCPDS cards. Depth profile analysis on the coatings was performed on a circular area with a diameter of 4 mm using glow-discharge optical emission spectroscopy (GDOES, Jobin-Yvon JY5000RF, Horiba, Ltd., Kyoto, Japan). Sputtering was done using Ar<sup>+</sup> ion at 40 W.

#### **3. Results**

#### *3.1. Alloy Microstructure*

Figure 1 shows the X-ray diffraction (XRD) pattern of AZ61, AZX611, and AZX612 alloys and the optical microscope images showing the microstructure of the alloys. All of the alloys consisted of primary α-Mg phase and secondary β-phase (Al12Mg17). Precipitation of Ca-containing intermetallic, Al2Ca and Mg2Ca, was detected in the AZX611 and AZX612 alloys. The peaks for Al2Ca and Mg2Ca phases were quite small and often overlapped with the peaks for other phases. As opposed to the microstructure of AM60 alloy which contained a low number of precipitates [23], the Mg matrix of the

AZ61 alloy exhibited an ultrafine equiaxed grain with the average grain size ~3 μm × 5 μm decorated by numerous spherical intermetallic Al12Mg17 phase with diameters approximately 0.5 μm as shown in Figure 1b. The presence of 1 wt.% Ca expanded the grain size about an order of magnitude, but no further enlargement with increasing Ca content to 2 wt.%. The XRD pattern of AZX611 and AZX612 alloys displayed a higher intensity in the Mg peaks confirming larger metallic grain sizes than that of the base alloy. The equiaxed grains size in both AZX611 and AZX612 alloy was approximately 50 μm × 250 μm. The grain boundaries became thicker with increasing Ca content in the alloys due to segregation of the intermetallic precipitates (Figure 1c,d). The XRD pattern of Ca-containing alloys revealed some additional β-phase peaks emerged at 23.7◦, 40.0◦, 42.0◦, 43.7◦, 61.6◦, 63.3◦, 64.9◦, 66.4◦, and 67.8◦ indicating a higher amount of β phase existed in the alloys.

**Figure 1.** (**a**) X-ray diffraction pattern and surface microstructure of (**b**) AZ61, (**c**) AZX611, and (**d**) AZX612 alloys. The scale bar in image (b) applied to images (c) and (d).

Figure 2 shows the SEM images and the corresponding EDS maps for the elements Mg, Al, and Ca, showing the element composition of the three alloys. In the maps, blue represents Mg, red represents Al, and the green represents Ca. The AZ61 alloy shows a high density of precipitates on its surface. The maps displayed a nearly uniform blue and red maps corresponded to the β precipitate smeared in the matrix. The size of precipitates became larger with increasing Ca content in the alloys, and the matrix became clearer from precipitates. The maps for AZX611 and AZX612 showed that the intermetallic consisted of Mg, Al, and Ca. This analysis proved the unambiguous presence of Al2Ca and Mg2Ca precipitates besides the β phase in both AZX611 and AZX612 alloys. The amount of Al2Ca phase dominated over Mg2Ca phase.

**Figure 2.** Plane-view SEM images and the corresponding EDS maps for Mg (blue), Al (red), and Ca (green), of AZ61, AZX611, and AZX612 alloys. The scale bar applied to all images.

#### *3.2. PEO-Coating Structure and Composition*

The voltage–time curves recorded during the PEO coating on the three alloys are depicted in Figure 3. Above the breakdown voltage (90 V), plasma discharge was developed at a few spots on the specimen surface, indicating the beginning of the electrolytic process. After the discrete plasma formed uniformly on the surface, the voltage stabilized at ~150 V. Strong plasma discharge began to occur at a critical voltage after about 9 min where large oscillation in the voltage (150–250 V) was attained. The presence of 1 wt.% Ca in the alloys shifted the critical voltage towards a longer time of approximately 2 min; however, the curve did not shift any further when the amount of Ca was increased to 2 wt.%.

**Figure 3.** Voltage–time curves during 20 min anodization of AZ61, AZX611, and AZX612 alloys in 0.5 M Na3PO4 solution at 25 ◦C.

Figure 4 displays the XRD pattern of the resulting coatings formed on the three alloys and the cross-sectional FE-SEM images of the coatings. Most of the peaks for substrates no longer appeared in the XRD pattern of the coated specimens indicating an excellent blocking effect of the coatings from the X-ray. The composition of the coatings formed on the three alloys were similar, which were composed of both crystalline and amorphous phases as shown by the appearance of the broad peak between 20◦ and 40◦, which attributed to the amorphous phase, and the small peaks inside the broad peak corresponded to the crystalline oxide phases of Mg3(PO4)2 and Mg(PO3)2. The coatings that formed on all of the specimens exhibited a similar uneven structure, as shown in Figure 4b–d. Pores and cracks, which are the footprint of PEO coating, were observed on both the surface and inner part of the coating. Fine cracks were developed during the PEO process, while heavy cracks occurred unintentionally during specimen preparation. The coating formed on Ca-containing alloys was slightly thicker than that of the base alloy. In agreement with the cross-section images, the average coating thickness measured by coating thickness gauge were 25, 32, and 30 μm for the coating formed on AZ61, AZX611, and AZX612, respectively. The XRD analysis did not detect the presence of Ca in the coating, similar to the earlier result on AM60 [23]. The Ca compounds may have been present in an amorphous state.

**Figure 4.** (**a**) X-ray diffraction pattern of plasma electrolytic oxidation (PEO) coatings and the corresponding cross-section FE-SEM images of the coating on (**b**) AZ61, (**c**) AZX611, and (**d**) AZX612 resulting from 20 min anodization in 0.5 M Na3PO4 solution at 25 ◦C. The scale bar applied to all images.

The coatings composition was further analyzed by depth profile GDOES, and the results are displayed in Figure 5. The dashed line marks the coating–metal interface at which the O and Mg profiles crossed each other. All of the coatings were composed of Mg, P, O, and Al. Confirming XRD results, the primary oxide phases in the coatings consisted of Mg–O compounds. The GDOES profile suggested the presence of other oxide phases, including MgO and Al2O3 in the amorphous state. The incorporation of Ca in the coating of AZX611 and AZX612 alloys was confirmed by the Ca profiles in Figure 5b,c. The intensities of Ca signal in the AZX611 and AZX612 coatings was about half of their bulk values, which verified the presence of calcium oxide/hydroxide phase in the coatings. The coating-metal interface shifted twice towards longer sputtering times from 400 s for AZ61 to 950 s for AZX611 and 900 s for AZX612 specimens. The shift indicated an increase in coating thickness as the thicker coating needs longer sputtering time to reach the bulk substrate.

**Figure 5.** Glow-discharge optical emission spectroscopy (GDOES) elemental depth profiles of PEO coatings on (**a**) AZ61, (**b**) AZX611, and (**c**) AZX612 specimens after anodization in 0.5 mol·dm−<sup>3</sup> Na3PO4 solution at 25 ◦C for 20 min. The dashed lines indicate the oxide–metal interface.

#### *3.3. Electrochemical Corrosion*

The corrosion behavior of the PEO coating was investigated by potentiodynamic polarization test in 0.9 wt.% NaCl solution at 37 ◦C. Figure 6 shows the corrosion potential and current densities data obtained from the polarization curves of the specimens before and after coating. The corrosion potentials of the substrates became slightly nobler from −1.49 to −1.45 VAg/AgCl with increasing Ca content in the alloys. The behavior was preserved after coating that the corrosion potential increased with Ca content in the alloys from −1.66 to −1.64 VAg/AgCl (Figure 6a). The corrosion current densities of the coated specimens were 10 times lower than that of the substrates. The substrates exhibited corrosion current densities in the range of 2.0 <sup>×</sup> <sup>10</sup>−<sup>5</sup> to 3.5 <sup>×</sup> <sup>10</sup>−<sup>5</sup> <sup>A</sup>·cm−<sup>2</sup> while the coated specimens were in the range of 3.5 <sup>×</sup> 10−<sup>6</sup> to 3.6 <sup>×</sup> 10−<sup>6</sup> A·cm−2. Depression of corrosion potential, which was accompanied by a reduction in corrosion current densities as a result of PEO coating, indicated an improvement in corrosion resistance of the alloys. The inhibition of cathodic reaction on the surface suppressed the corrosion potential of the coated specimens to the negative direction.

**Figure 6.** (**a**) Corrosion potentials and (**b**) corrosion current densities of the substrates and coated specimens of AZ61, AZX611, and AZX612 alloys.

The coating resistance was evaluated by EIS measurement in 0.9 wt.% NaCl solution at 37 ◦C. The results are displayed in Figure 7 for uncoated, and Figure 8 for the coated, specimens. The fitted parameters obtained from the EIS data are listed in Table 3. The n1 and n2 refers to constant phase for CPE1 and CPE2, respectively. The maximum n value is 1. The closer to 1 the n1 and n2 value is, indicates more capacitive behavior. The polarization resistance (R1) of the coated specimens was an order of magnitude higher than that of their substrates. The base alloy exhibited the highest polarization resistance in both uncoated and coated conditions relative to the Ca-containing alloys. The

impedance spectrum of AZ61 substrate exhibited one apparent capacitive loop, while the Ca-containing alloys showed an additional inductive loop at low frequency. The corresponding inductive loop is an indication of localized corrosion [5,8]. The presence of inductive loop lowered the n2 values for AZX611 and AZX612 specimens compared to the AZ611 specimen. The presence of a constant phase element (CPE) in parallel with a resistance indicates the presence of a faradaic reaction (charge transfer) and a non-Faradaic reaction (charge accumulation at the interface) occurring at the interface. All of the three PEO coatings exhibited a similar trend of impedance spectrum. The L and RL loops did not exist in the model for the PEO coating (Figure 8), indicating high resistance to metal dissolution. Moreover, a much higher impedance at low and medium frequencies of the Bode plot and an enlarged diameter of the semicircle in the Nyquist plot of the coated specimens are evidence of remarkable higher corrosion resistance than that of the substrates. Moreover, the capacitance values of the coated specimens were a hundred times lower than that of the substrates. The effect of the alloying Ca in the coating impedance was displayed by the Bode plot in Figure 8. The impedance at high frequencies, which represents the outer PEO coating of all the coated specimens, was similar. However, the impedance at low frequencies displayed slightly lower impedance than that of the AZ61 specimen. The results implied that the coating–metal interface of the Ca-containing specimens was more prone to localized corrosion than that of the base specimen.

**Figure 7.** Electrochemical impedance spectra of uncoated specimens of AZ61, AZX611, and AZX612 specimens: (**a**) Nyquist plots, (**b**) bode plots, and (**c**) phase plots, and (**d**) the equivalent circuit.

**Figure 8.** Electrochemical impedance spectra of coated specimens of AZ61, AZX611, and AZX612 specimens: (**a**) Nyquist plots, (**b**) bode plots, and (**c**) phase plots, and (**d**) the equivalent circuit.


#### *3.4. Apatite-Forming Ability*

The apatite-forming ability of the coated specimens was evaluated by immersion test in XRD. The concentration of Mg dissolved from the surface, and the pH solution during the immersion test in SBF was monitored. Figure 9 displays the concentration of Mg detected in the solution and the pH solution during 14 days of immersion. The AZ61 alloy released the lowest Mg concentration in the range of 100–150 mg/L indicating the lowest corrosion rate relative to the two Ca-containing alloys, which released 100–500 mg/L Mg into the solution (Figure 9a). The solution pH of the three alloys increased from 7.4 to 8 after one-day immersion indicating high corrosion activities (Figure 9c). The increase in pH occurred due to the release of hydrogen gas during hydrolysis of Mg ions following the corrosion attack. The pH change was not significant as the solution was buffered. The solution pH of AZ61 alloy decreased, approaching the fresh SBF while that of the AZX611 and AZX612 alloys fluctuated with increasing immersion time. On average, the lowest pH was attained in the AZ61 alloy. The PEO-coated specimens released much lower Mg into the solution relative to their substrates (Figure 9b). In agreement, the solution pH of the coated specimen was relatively lower compared to the substrates during the 14-day immersion. The dissolved Mg in the solution was attributed to both corrosion of the underlying substrate and the coating dissolution. The coated AZ61 specimen released 110 to 200 mg/L Mg at the initial immersion time and then stabilized at day 7 to 12. Dissolution no longer occurred at day 14. The coating on AZX611 specimen exhibited a relatively constant dissolution rate during

1–12 days followed by passivation at the end period of the test. The AZX612 specimen exhibited a dynamic dissolution resulting in a highly fluctuating concentration of the dissolved Mg in the range 300–450 mg/L. All of the coatings exhibited the tendency for passivation at day 14 with solution pH at 7.7, which was presumably due to surface coverage by the apatite layers on their surfaces.

**Figure 9.** Concentration of Mg dissolved during immersion testing in SBF of the (**a**) substrates, (**b**) coated specimens, and the corresponding solution pH for (**c**) the substrates and (**d**) coated specimens.

The presence of apatite on the PEO coating was investigated by observing the surface morphology and analyzing the surface composition after the immersion test. Figure 10 shows the SEM images at different magnifications revealing the presence of uniform apatite layer on all of the three coatings. Most of the pores that existed in the coating of AZ61 specimen were no longer viewed, indicating surface coverage by thick apatite layer (Figure 10a). The layer formed uniformly following the coating structure, as shown clearly in Figure 10b. The spongy structure with scallop type of grain, which was a typical morphology of apatite formed in SBF was seen in a higher magnification image in the inset Figure 10b. The coating remained intact and protected by the apatite layer. Thinner apatite layers were formed on AZX611 and AZX612 specimens as inferred from the existence of open pores and cracks in the coating (Figure 10c,e). The scallop grain of apatite layer in the AZX611 and AZX612 specimens was smaller than that of formed in AZ61 specimen as shown in the inset image of Figure 9d,f. Localized coating damage was recognized at a few spots, pointed by arrows in the images. The AZX611 specimen showed relatively small coating deterioration while larger coating damage on an area approximately 150 μm × 300 μm was observed on AZX612 specimen.

Figure 11 shows the results of EDS analysis from the areas shown in Figure 10 as compared to that of fresh PEO coating. The observation areas for both cases were kept identical by using similar magnification scale. Prior to the immersion test, calcium was not detected in any of the coatings (Figure 11a). The signal attributed to the calcium oxide phase in the coating was weak relative to the other elements. After 14 days immersion in SBF, the Ca signal was detected as high as 8.0, 7.2, and 3.1 at.% in AZ61, AZX611, and AZX612 specimens, respectively, attributed to the apatite layer shown in Figure 10. The Mg concentration increased from 6.8 at.% in the AZ61 to 12.5 at.% in the AZX611 specimens. The concentration of P was relatively not altered at 12–13 at.%. The ratios of Ca/P of the

apatite layers were 0.65, 0.60, and 0.24 for the layers formed on AZ61, AZX611, and AZX612 specimens, respectively. The ratio was half of the stoichiometry apatite, 1.67 [30]. The apatite layer that formed in the coating of AZX611 and AZX612 alloys contained more Mg than that of the AZ61 alloy. Moreover, the strong signal of P derived from the underlying PEO coating lowered the Ca/P ratio.

**Figure 10.** SEM images of PEO coatings on (**a**) and (**b**) AZ61, (**c**) and (**d**) AZX611, and (**e**) and (**f**) AZX612 specimens after 14 days of immersion in SBF showing the corrosion morphology. (b), (d), (f) are higher magnification images of the area inside the rectangle in images (a), (c), (e), respectively. The inset image shows the apatite layer structure.

**Figure 11.** Atomic fraction of elements detected by EDS analysis on PEO coatings of AZ61, AZX611, and AZX612 specimens (**a**) before and (**b**) after 14 days immersion in SBF.

#### **4. Discussion**

Surface characterization showed that the coatings formed on all the three alloys exhibited similar uneven structure decorated with pores and cracks, but the composition and thickness were altered by the presence of Ca in the alloys. The coatings composed of a mixed crystalline and amorphous phase. The Mg3(PO4)2 and Mg(PO3)2 existed in both crystalline and amorphous states as confirmed by XRD analysis, while MgO and Al2O3 were present in an amorphous state as detected by GD-OES and EDS analysis. The presence of Ca in the alloys induced the formation of amorphous calcium oxide/hydroxide phase in the coatings. The depth profile analysis by GD-OES (Figure 5) confirmed the presence of calcium oxide/hydroxide from the surface throughout the coating-metal interface. At identical anodization time, the resulting coating on the AZX611 alloy was 7 μm thicker than that of the base alloy. The coating did not thicken any further in the AZX612 alloy, as no further substantial change in the Ca content of the solid solution matrix.

The microstructure observation revealed that the alloying element Ca decreased the density of intermetallic precipitates in the AZ61 alloys by inducing larger precipitates size and enlarging the metallic grain. The grain size expanded about 30–50 times by the presence of 1 wt.% Ca in the alloy. The substitute Ca atom with atomic diameter larger than Mg atom [31] caused expansion in the matrix grain size. The high number of submicron-sized precipitates in the AZ61 alloy transformed into a low number of 10 micron-sized precipitates as a result of alloying with 1 wt.% Ca (Figure 2). The precipitates grew slightly larger, but the metallic grain size remained the same, with increasing Ca concentration to 2 wt.%. The enlargement of grain size was controlled by the amount of dissolved Ca in the solid solution matrix. The precipitates continued to grow as more excess of Ca existed in the AZX612 alloy. The limited solubility of Ca in Mg, which is only 0.8 wt.% [32], induced the formation of Al2Ca precipitates. With increasing Ca content, the bigger precipitates grew larger at the expense of the smaller ones following Ostwald ripening [33]. Another intermetallic phase Mg2Ca also existed in the AZX611 and AZX612 alloys, but only in a small number. The large misfit in atomic size between Ca and Mg in which the diameter of the Ca atom is 20% larger than Mg, is less favorable for the formation of Mg2Ca phase [31]. The presence of heat-resistant Al2Ca and Mg2Ca phases delayed the formation of strong plasma discharge during electrolytic process resulting in a less crystalline oxide phases in the coatings. Moreover, the presence of Mg2Ca phase is not beneficial from the corrosion viewpoints as it is more anodic than the Mg matrix [2].

During the PEO formation, Mg and Ca in the solid solution matrix was selectively oxidized. From thermodynamic viewpoints, Ca is easier to oxidize as it has lower Gibbs free energy than Mg [33]. However, since the concentration of Ca in the solid solution matrix was limited by its solubility, the resulting calcium oxide/hydroxide in the coating was relatively small compared to the magnesium phosphate/oxide phases. The GD-OES depth profile was proved to be a sensitive technique in detecting the calcium oxide/hydroxide. The calcium oxide/hydroxide phases presumably occurred during the 12 min exposure in fine plasma discharge. The strong discharge, which occurred in the latter stage, tended to dissolve the calcium oxide/hydroxide to the solution. The incorporation of calcium oxide/hydroxide contributed to the thickening of the coating formed on AZX611 and AZX612 alloys relative to the AZ61 alloy.

The higher resistivity offered by the thicker coating formed on the AZX611 and AZX612 alloys was counteracted by the vulnerability of their metal substrates to localized corrosion relative to the AZ61 alloy. The polarization coating resistance of the coated AZ61 specimen was higher than that of the AZX611 and AZX612 specimens. The enlargement of precipitates in the AZX611 and AZX612 alloys increased the contact area of galvanic coupling between the precipitates and the matrix, and therefore increased the danger to micro galvanic corrosion when exposed to the corrosive solution. Ennoblement of corrosion potentials in the polarization curves of the AZX611 and AZX612 alloys due to the increase number of cathodic intermetallic was balanced by the increased susceptibility to localized corrosion. The presence of inductive loop in the EIS spectra of AZX611 and AZX612 alloys (Figure 7a), which was not present in the spectra of AZ61 alloy, was attributed to the localized corrosion. A small amount of big precipitates accelerates corrosion, whereas finely distributed precipitates retarded corrosion [34]. The coating–metal interface in the vicinity of the precipitates was less protective. Once the corrosive solution penetrated the coatings, pockets of the solution were developed at the interface. Corrosion of

the metal substrate occurred immediately as it was exposed to the solution inside the pockets. The increase in intermetallic precipitates size with Ca content in the alloys contributed to the enlargement of the attacked areas (Figure 10).

The higher corrosion rate of the underlying substrate limited the apatite-forming ability of the coatings on the AZX611 and AZX612 specimens. Even though a uniform apatite layer grew on all of the three coatings as a result of immersion in SBF, the layers that formed on the AZX611 and AZX612 coatings was thinner and contained higher Mg than that formed on the AZ61 coating. The incorporation of calcium oxide/hydroxide in the coatings did not necessarily enhance the bioactivity of the coatings, similar to that reported in Reference [19]. The corrosion rate of the underlying substrate is a more important factor in supporting the growth of apatite. An alternate corrosion rate of the underlying substrate, which released fluctuated hydrogen gas, tends to disturb the deposition of apatite on the coatings. In the case of AM60 alloys [23], the corrosion resistance of the PEO-coated specimens were relatively unaltered, and therefore the formation of apatite was not affected by the variation of Ca content in the alloys. During initial immersion time, thinning of the PEO coating occurred selectively at the amorphous part, leaving a rough crystalline oxide phase. The resulting sub-micron surface roughness became a preferential site for nucleation of apatite [14]. The apatite layer was likely to form on the coatings during 7–14 days immersion time where a constant dissolution rate of Mg was attained (Figure 9b). The released substances from the specimens contributed to increasing the degree of supersaturation of the SBF and the local pH, which triggered the deposition of apatite [30].

The following discussion led to the conclusion that the corrosion resistance of the PEO coatings formed on the AZ61 alloys was influenced by the concentration of alloying element Ca, in particular the concentration of the dissolved Ca in the solid solution matrix. Therefore, it is not necessary to add Ca above its solubility in Mg matrix. The Ca-containing precipitates that tended to shorten the lifetime of strong discharge during the PEO process resulted in the formation of lower crystalline oxide phases in the coating. A thinner PEO coating with a high-volume ratio of crystalline to amorphous phases gives much better corrosion resistance than a thicker coating with a lower ratio [12]. Addition of Ca above the solubility limit did not have a beneficial effect on the coating properties. On the contrary, the excess of Ca in the alloys resulted in the enlargement of Al2Ca and Mg2Ca precipitates, which tended to accelerate localized corrosion.

#### **5. Conclusions**

The PEO coating significantly enhanced the corrosion resistance of the Mg-6Al-1Zn-xCa (AZX) alloys relative to their substrates. The presence of 1 and 2 wt.% Ca in the alloys resulted in the formation of thicker coating, which contributed in higher enhancement of the coating resistance. However, the EIS spectra and the surface observation after the immersion test indicated an increased danger to localized corrosion with increasing Ca content in the alloys. The coating formed on AZX alloys exhibited lower polarization resistance than that which formed on the AZ61 alloy. After the 14-day immersion tests in SBF, the coating on AZ61 alloy remained intact while the coatings on AZX611 and AZX612 showed noticeable local damage. The microstructure investigation revealed that the alloying element Ca caused enlargement of metallic grain size and intermetallic precipitates relative to the AZ61 alloy. The increase in corrosion activities underneath the coatings lowered the apatite-forming ability of the AZX coatings, resulting in a thinner apatite layer formed during immersion in SBF.

**Author Contributions:** Conceptualization, A.A.; Methodology, A.A. Software, A.A. and S.O.; Validation, S.O.; Formal Analysis A.A., H.A. and S.O.; Investigation, A.A.; Resources, A.A. and S.O.; Data Curation: A.A.; Writing-Original Draft Preparation, A.A.; Writing-Review & Editing, A.A., H.A. and S.O.; Visualization, A.A.; Supervision, S.O.; Project Administration, A.A. and S.O.; Funding Acquisition, A.A. and S.O.

**Funding:** This research was supported by Direktorat Riset dan Pengabdian Masyarakat (DRPM) Universitas Indonesia through Q1Q2 Research grant 2019 (NKB-0266/UN2.R3.1/HKP.05.00/2019). Grant-in-Aid for Scientific Research from the Japan Society for the Promotion of Science and the Light Metal Education Foundation of Japan is acknowledged. We also thank a Strategic Research Foundation Grant-aided Project for Private Universities matching fund subsidy from the Ministry of Education, Culture, Sports, Science, and Technology of Japan.

**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* **Influence of Micro-Arc Oxidation Coatings on Stress Corrosion of AlMg6 Alloy**

**Lesław Kyzioł 1,\* and Aleksandr Komarov <sup>2</sup>**


Received: 16 December 2019; Accepted: 10 January 2020; Published: 12 January 2020

**Abstract:** This paper shows results of a study on the corrosion behavior of micro-arc oxidation (MAO) coatings sampled from the AlMg6 alloy. The alloy was simultaneously subjected to a corrosive environment and static tensile stress. For comparative purposes, the tests were run for both coated samples and samples without coatings. The research was conducted at a properly prepared stand; the samples were placed in a glass container filled with 3.5% NaCl aqueous solution and stretched. Two levels of tensile stress were accepted for the samples: σ*<sup>1</sup>* = 0.8*R*0.2 σ*<sup>2</sup>* = *R*0.2, and the tests were run for two time intervals: *t1* = 480 h and *t2* = 1000 h. Prolonged stress corrosion tests (lasting up to 1000 h) showed that the samples covered with ceramic coatings demonstrated significantly higher corrosion resistance than the samples without the coatings. Protective properties of the coating could be explained by its structure. Surface pores were insignificant, and their depth was very limited. The porosity level of the main coating layer was 1%. Such a structure of coating and its phase composition provided high protective properties.

**Keywords:** aluminum alloy AlMg6; Al2O3 coating; phase composition; corrosion; stress corrosion; micro-arc oxidation

#### **1. Introduction**

Al-Mg alloys combine good formability, rather high strength, corrosion resistance, and weldability. Therefore, such alloys are used in many structures exposed to weathering, and especially in shipbuilding and offshore structures. It should be noted that studies carried out on samples made of the AlMg6 alloy showed a good resistance to stress corrosion for this alloy, but one much lower than for alloy 5083 [1,2]. However, the strength of alloy 5083 containing 5% Mg is noticeably inferior to the strength of AlMg6 alloy.

Stress corrosion of materials manifests itself through the formation of cracks in the metal when exposed to corrosive environment and static tensile stress. The cracks appearing on the metal surface are perpendicular to the direction of tensile stress and can be of intercrystalline or mixed nature. Studies have shown that before the appearance of pronounced cracks, there is often an incubation period. The intensity of microcracks can be determined on the basis of changes in mechanical properties after time intervals of the stress corrosion test [1–6].

Studies have also shown that Al-Mg alloys with a content of Mg ≤ 3, 5% exhibit a low susceptibility to stress corrosion. This is due to the discontinuity of β-phase molecules at the grain boundaries, which in turn results from the low supersaturation of the solid solution [7–10].

Al-Mg alloys with a content of Mg ≥ 3, 5%, especially above 5% (AlMg5, AlMg6), with a certain state of the structure and under specific external conditions may be resistant to intercrystalline, layer, and stress corrosion [3–5]. The study has shown that Al−Mg alloys with a content of more than 6 wt. % of Mg can be obtained, and they show good weldability and high mechanical properties [1,9–11]. An

increase in the Mg content results in an increase in strength, but leads to higher susceptibility to local corrosion and reduced resistance to corrosion under stress [12]. It is assumed that the grain boundaries are a favorable place for β-phase formation due to their low diffusion barrier (presence of defects such as dislocations and vacancies) [13]. In many cases, overexposure of Al−Mg sheets or other products to a temperature range of 70 ◦C to 200 ◦C may impede production, causing precipitation of the β-phase rich in Mg at the grain boundaries. These alloys are susceptible to intercrystalline, stress, or pitting corrosion, because the β-phase is electrochemically more active than the aluminum matrix [14–18].

Structures with a high Mg content resistant to corrosion are obtained through the use of complex manufacturing methods [4]. The goal of the research was to determine the impact of the protective ceramic coating on the corrosion resistance of the tested alloy. Micro-arc oxidation (MAO) is a very promising process in making tight coatings and curing of metal elements. This is a coating technique capable of forming ceramic coatings on metals such as Al, Mg, Ti, and their alloys [17–22]. This environmental friendly technique allows ceramic layers to be grown, giving the pieces high level mechanical and tribological properties and good corrosion protection in a single step of processing [23–26]. A number of studies have shown that the coatings obtained by micro-arc oxidation are an effective means of protecting aluminum alloys from corrosion [27–33]. The corrosion properties of the MAO coating of the BS 6082 Al alloy were tested against different immersion periods in a 0.5M NaCl solution for up to 48 h [28]. The research has shown the importance of sealing the pores in MAO coatings with the use of the sol–gel technique, enhancing the short-term corrosion resistance in 0.6M NaCl solution. The above investigation also revealed that the MAO coatings improved the corrosion resistance of the Al alloy because of the lack of defects in coatings [29].

Despite numerous studies on the protective properties of coatings, their effectiveness in protecting aluminum−magnesium alloys from stress corrosion has not been scrutinized. The major objective of the present study is to evaluate the overall effectiveness of MAO coatings in terms of resistance to aqueous stress corrosion. Regarding the above objective, MAO coatings were deposited on a AlMg6 alloy and their corrosion behavior was evaluated in 3.5% NaCl solution. In the available literature, there are only a few sources referring to the importance of protective coatings for stress corrosion [34,35]. These tests are very important because often the material is subject simultaneously to corrosive environments and stresses. In this paper the stress corrosion resistance has been evaluated using a change of mechanical properties of samples with coatings and without coatings.

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

Samples of AlMg6 alloy without coating and coated with a ceramic coating by MAO were subjected to stress corrosion. The tests were carried out on identical samples without coatings and with coatings to determine the effectiveness of the ceramic coating that protects the material against stress corrosion.

Samples for the static tensile test and the stress corrosion were made from a sheet with a thickness of g = 3 mm. The samples were cut down in a direction transverse to the rolling direction. The chemical composition and parameters of heat treatment of the AlMg6 alloy plates are presented in Table 1. The shape and dimensions of the samples for the determination of mechanical and stress corrosion properties are shown in Figure 1.


**Table 1.** Chemical composition and heat treatment parameters of the AlMg6 alloy sheet.

**Figure 1.** The shape and dimensions of samples for static and stress corrosion tests of the AlMg6 alloy.

The samples were polished with abrasive paper 600# and degreased with acetone followed by rinsing with distilled water before coating formation. The electrolyte was an aqueous solution of 2 g/L KOH and 5 g/L Na2SiO3. The electrolyte was agitated with compressed air. A forty litter bath from stainless steel was used. The specimen was served as the anode and the bath wall was served as the contrary electrode. The system was cooled by cold water pumped through double walls of the bath. The electrolyte temperature was controlled at 25−30 ◦C throughout the process. The MAO treatment was carried out using a pulsed AC power source. The current density, voltage, frequency, duty cycle, and duration time were 25A/dm2, 280 V, 50 Hz, 50% and 60 min, respectively. After the treatment, the samples were rinsed in distilled water and dried in air.

Then, using abrasive paper 600#, the upper loose porous layer of about 20 μm thickness was removed from the surface of the samples, after which the samples were washed in distilled water with an ultrasonic bath SONOREX for 15 min and dried in air. The thickness of the coating on the samples prepared for testing was 150 ± 5 μm. The surface morphologies and cross-section morphologies of the coatings were examined with an optical microscope Axiovert 25 as well as with a scanning electron microscope (ZEISS, Jena, Germany). X−ray diffraction studies were carried out on a DRON−3M X−ray diffractometer (S. Petersburg, Russia) in scanning mode in 0.1 increments using Cu−Kα radiation in the Bragg–Brentano mode. The time of a set of pulses at a point was 15 s.

Stress corrosion tests of the AlMg6 alloy were carried out on samples for σ = const. The samples were placed in a glass container filled with 3.5% NaCl aqueous solution and stretched. Samples with and without coatings were subjected to stress corrosion tests for two levels of stress:

1—σ<sup>1</sup> = 0.8*R*0.2, 2—σ<sup>2</sup> = 1.0*R*0.2, where *R*0.2 is the average value of the experimentally determined yield strength of samples without coating and with coating equal to 212 MPa and 217 MPa, respectively.

The corrosion tests were conducted for two time intervals: 1—480 h, 2—1000 h. Figure 2 presents a stand for testing stress corrosion resistance.

**Figure 2.** Stand for testing the stress corrosion resistance of AlMg6 alloy, (**a**) test bench scheme for σ = const., 1—container filled with 3.5% of NaCl, 2—sample, 3—lever arm, 4—load, (**b**) container filled with 3.5% of NaCl in which the sample is located.

After stress corrosion tests the samples were subjected to a static tensile test to determine the mechanical changes of the AlMg6 test material. A static tensile test on samples without coatings and with coatings was carried out on a ZwickRoell testing machine (Ulm, Germany). The research was carried out at the Faculty of Mechanical Engineering of the Gda ´nsk University of Technology (Gda ´nsk, Poland).

#### **3. Research Results and Their Analysis**

Figure 3 presents images of samples for mechanical and stress corrosion tests without coating and coated. The formed coating has a light gray color and a uniform surface. The study of the surface morphology of the samples showed that the surface layer of the coating is characterized by the presence of pores up to 10 microns in size (Figure 4). Such a structure is typical for coatings obtained with the use of the micro-arc oxidation method, in which there is an upper loose porous layer and a dense, low porosity basic coating layer [5,6].

**Figure 3.** Images of samples for stress corrosion tests of AlMg6 alloy with sheet thickness g = 6 mm, (**a**) without coating, (**b**) with coating.

**Figure 4.** SEM image of the coating surface prior to testing.

Figure 5 presents images of samples without coating and coated after mechanical and stress corrosion tests. On samples without coatings, there are visible deposits and small corrosion centers (Figure 5a). There is no evidence of corrosion on samples with coating (Figure 5b). Studies performed on the cross-section of the coating after corrosion tests also showed no signs of corrosion (Figure 6).

**Figure 5.** Images of samples after stress corrosion tests of AlMg6 alloy with sheet thickness g = 6 mm, (**a**) without coating, (**b**) with coating, (**c**) without cover after breaking, (**d**) with a coating after breaking.

High protective properties of the coating can be explained by its structure. As follows from the Figure 6, the pores on the surface of the coating have an insignificant size and depth and are limited to a loose surface layer. The main coating layer has a very low porosity (less than 1%), while the permeable porosity is absent. Such coating structures, as well as its phase composition, represented by chemically resistant aluminum oxide in the γ−Al2O3 and α−Al2O3 modifications (Figure 7), provide high protective properties.

On the contrary, the surface of uncoated specimens of AlMg6 alloy was subjected to intense corrosion, as evidenced by the sources of material etching and corrosion products (Figure 5a,c).

**Figure 6.** SEM image of the surface (**a**) and cross-section (**b**) of the sample with coating after stress corrosion tests.

**Figure 7.** A fragment of the X−ray diffraction pattern of coating on a sample for the stress corrosion test.

Table 2 shows results of the test for static tensile stress of the AlMg6 alloy without coatings and with coatings. The obtained results indicate identical mechanical properties of the tested materials. A very thin layer of ceramics (about 150 μm), in relation to the thickness of the sample (gs = 3 mm) has no influence on the mechanical properties of the tested materials.


**Table 2.** Mechanical properties of AlMg6 alloy samples without coatings and with coatings.

The stress corrosion test is very rigorous because the material is simultaneously subject to a corrosive environment and stress. None of the 24 samples tested broke. The results of the tensile test of samples after exposure to stress corrosion for two time intervals and two stress levels are presented in Tables 3 and 4.

Table 3 shows results of the test for static tensile stress in samples with and without coatings at *t* = 480 h exposed to stress corrosion at the level of σ = *0.8R0.2* and σ = *R0.2*.

The results of the tests showed that for the level of stress σ = *0.8R0.2*:


For the level of stress σ = *R0.2*:


For such a high level of stress the samples without coating exhibited a significant decline in mechanical properties caused by simultaneous exposure to stress and a corrosive environment.

In the case of the coated samples, there was an insignificant decline in mechanical properties caused by the high level of stress. The coating very tightly protected the sample surface, and thus the material of the sample was not subjected to corrosive environment.


**Table 3.** Mechanical properties of samples with and without coatings of AlMg6 alloy after exposure to stress corrosion at time *t* = 480 h.

Table 4 shows the results of the samples exposed to stress corrosion at similar levels of stress, i.e., σ = *0.8R0.2* and σ = *R0.2* after the time of t = 1000 h.

For the level of stress σ = *0.8R0.2*:


With the amount of stress corrosion time, a further decline in the mechanical properties of uncoated AlMg6 alloy was observed.

The coated samples showed an insignificant decline in mechanical properties. The decline was not caused by the corrosive environment, but rather by crawling of the material, which was subjected to a high level of stress.

For the level of stress σ = *R0.2*:


For such a high level of stress and a long exposure to a corrosive environment, a further decline in mechanical properties was observed in the samples with no protective coatings. The decline occurred due to a simultaneous concurrence of high stress and corrosion caused by the 3.5% NaCl solution.

Having been exposed to a corrosive environment for *t* = 1000 h and highly stressed for up to σ = *R0.2*, the coated samples sustained an insignificant decrease in mechanical properties. The limited decline is caused by the high level of stress.


**Table 4.** Mechanical properties of samples with and without coatings of AlMg6 alloy after exposure to stress corrosion at time *t* = 1000 h.

Figures 8–10 shows the influence of corrosive environment and stress level on the reduction of mechanical properties of AlMg6 alloy samples with and without a ceramic coating. The research shows that there was almost a 30% reduction in yield strength and tensile strength, and nearly a 60% reduction in the plasticity of AlMg6 alloy samples without a ceramic coating for stress level σ = *R0.2* and corrosion exposure time *t* = 1000 h. For identical corrosion conditions and stress levels, there was a 4% decrease in yield strength and tensile strength as well as a 8% decrease in plasticity of AlMg6 alloy samples coated with ceramic coating.

The samples of AlMg6 alloy without coatings showed a concurrence of static tensile stress and corrosive environment. This kind of corrosion−stress synergy is the very reason for the accelerated degradation of the material. However, the samples with ceramic coatings showed a high resistance to simultaneous corrosion and stress factors as the coatings protected the material against corrosion. An insignificant decrease in mechanical properties occurred due to the high level of stress. Figures 8–10 shows a very narrow area of the decline in mechanical samples of coated AlMg6 alloys.

**Figure 8.** Influence of the corrosive environment and stress level on the yield strength reduction of AlMg6 alloy samples without coatings and with coatings.

**Figure 9.** Influence of the corrosive environment and stress level on the tensile strength reduction of AlMg6 alloy samples without coatings and with coatings.

**Figure 10.** Influence of the corrosive environment and stress level on the relative elongation reduction of AlMg6 alloy samples without coatings and with coatings.

#### **4. Summary**

Studies have been carried out to evaluate the protective properties of the coating obtained by micro-arc oxidation on aluminum alloy AlMg6 from stress corrosion. The protective properties of the coating were evaluated by changing the mechanical characteristics (yield strength *R0.2*, tensile strength *Rm* and relative elongation *A5*) after holding the samples without coating and coated in a 3.5% NaCl solution under a stress of σ*<sup>1</sup>* = *0.8R0.2* and σ*<sup>2</sup>* = *R0.2* for 480 h and 1000 h. The results have showed that the MAO coating provides a fairly effective protection of the aluminum alloy from stress corrosion if the

selected test conditions are met. No corrosive centers were observed on samples with a ceramic coating, but small corrosion centers may be observed on samples without a coating. The effect of corrosion led to a decrease in the mechanical characteristics of the uncoated samples. So, after the 480-hour experiment at both stress levels, there was no noticeable decrease in the mechanical characteristics of coated samples, while the mechanical characteristics of uncoated samples decreased 1.1 (*Rm*) to 1.5 (*A5*) times. After holding the samples for 1000 h, a more noticeable difference in mechanical characteristics was observed. In this case, the decrease in the characteristics of coated samples did not exceed 1.07 times (*Rm*) to 1.12 times (*A5*), while the decrease in the properties of uncoated samples reached 1.3 (*Rm*) to 2.4 times (*A5*). The high protective properties of the coating can be explained by the chemical inertness of its composition (γ−Al2O3 and α−Al2O3) and low porosity (less than 1%) in the absence of through porosity.

**Author Contributions:** Conceptualization, L.K. and A.K.; methodology, L.K. and A.K.; formal analysis, L.K. and A.K.; investigation, L.K. and A.K.; resources, L.K and A.K.; data curation, L.K. and A.K.; writing—original draft preparation, L.K. and A.K.; writing—review and editing, L.K. and A.K.; visualization, L.K and A.K.; supervision, L.K. and A.K.; funding acquisition, L.K. All authors have read and agreed to the published version of the manuscript.

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

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

#### **References**


© 2020 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* **Enhanced Erosion–Corrosion Resistance of Tungsten by Carburizing Using Spark Plasma Sintering Technique**

#### **Yan Jiang 1, Junfeng Yang 2,\*, Zhuoming Xie <sup>2</sup> and Qianfeng Fang 2,\***


Received: 11 May 2020; Accepted: 12 June 2020; Published: 15 June 2020

**Abstract:** The biggest obstacle for the application of tungsten as the target materials in the spallation neutron source is its serious corrosion in the coolant of flowing water. For this reason, W–Cr–C clad tungsten was developed by tungsten carburizing in a spark plasma sintering device, with superior corrosion resistance in the static immersion and electrochemical corrosion test. This work focused on its erosion and corrosion performance in a flowing water system, based upon test parameters simulated under the service conditions. W–Cr–C clad tungsten showed superior corrosion resistance to that of bare tungsten due to the corrosion form changing from the intergranular corrosion of bare tungsten to pitting corrosion of W–Cr–C coating. The corrosion rate of tungsten was as high as tenfold that of the coated sample at 20 ◦C, and at most fourfold at 60 ◦C after testing for 360 h. Effects of water velocity and temperature on pitting and intergranular corrosion were investigated in detail and their corresponding corrosion mechanisms were analyzed and discussed.

**Keywords:** tungsten; W–Cr–C coating; carburization; intergranular corrosion; pitting corrosion

#### **1. Introduction**

Tungsten is being chosen as the most favorable spallation target material for the spallation neutron source (SNS) due to its obvious advantages including high melting point, high neutron yield, and superior thermal conductivity [1,2]. However, on the issue of compatibility with flowing water coolant, tungsten has long been challenged because of its serious corrosion in water, especially in the case of flowing water where both erosion and corrosion occur simultaneously [3–6].

To improve the corrosion resistance, tungsten cladding with high temperature and corrosion resistance coatings have been developed as a target material such as HIPed tantalum-clad tungsten [7,8]. However, it suffers from issues of interface defects [7] and the high decay heat of tantalum [8,9]. Alternatively, coatings such as tungsten carbide and chromium carbide have been generally adopted to protect the substrate from corrosion under harsh environments in engineering and industrial fields [10–15]. In the conventional preparation of cemented carbides, cobalt as binder phase is commonly introduced to densify the particles, which is known as a poorly resistant element to corrosion [16]. However, the development of binder-free tungsten carbide coatings needed to solve two problems: density and brittleness. To this end, carburizing by spark plasma sintering (SPS) technology was employed to produce tungsten carbide coatings on tungsten [17]. SPS is a high-temperature and fast-process sintering technique that provides a fast heating/cooling rate, short consolidation time, and controllable pressure, having been widely used in densification fabrication of nanomaterials, gradient functional materials, and ceramic materials. By SPS, the carburized layer over 20 μm thick was

quickly achieved at 1600 ◦C, holding for 10 min with a pressure of 45 MPa [17], and a double thick layer was available when 20 min and 30 min were tried. On this basis, W–xCr–C (x = 0.5, 1, 2, 3, and 6 wt.%) composite layers with higher compactness and uniform grain size distribution were fabricated in the same way [18]. Electrochemical corrosion measurements on a WC clad tungsten system and W–Cr–C clad tungsten system indicated that the W–1%Cr–C clad tungsten sample performed well and exhibited the lowest corrosion current density [19].

As a result of these findings, W–1%Cr–C (W–Cr–C for short in the following text) clad tungsten by the SPS method was chosen as the potential material to undergo the following dynamic corrosion test in flowing water; flowing water (pH: 6~8) was adopted as the coolant in service condition to remove the heat in the target chamber. According to the simulation for the 100 kW neutron spallation source, a flow velocity of about 2 ms−<sup>1</sup> is enough for the heat to be removed from the target, under which the highest target temperature is around 60 ◦C [2,20]. Therefore, this work investigated the corrosion performance of W–Cr–C clad tungsten at such conditions with the main purpose to explore and understand the behavior and mechanism in the erosion–corrosion process. Bare tungsten was taken into comparison.

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

A commercial grade tungsten disc (density: 18.48 g/cm3, Zhuzhou Cemented Carbide Group Co. Ltd., Zhuzhou, China; the main elements included are listed in Table 1) with cylindrical form (Ø16 mm × 3 mm) was embedded in the mixed powder of graphite (99.95%, 8000 mesh, Aladdin) and 1 wt.% chromium (99.5%, ≥325 mesh, Aladdin, Shanghai, China), and then placed in a graphite mold and sintered in a SPS furnace (HPD 5, FCT Systeme GmbH, Rauenstein, Germany), as shown in Figure 1a. A target temperature of 1600 ◦C holding for 10 min with a pressure of 45 MPa was exploited on the graphite die (Ø20 mm). The diameter of the die could change from 10 mm to 100 mm according to the sample dimensions. Pressure was used to bring the powder and disc into full contact. The heating and cooling rates of SPS were both 100 ◦C/min. After sintering, the disc was taken out and cleaned ultrasonically in acetone. A scanning electron microscopy (SEM)(Sirion 200, FEI, Hillsboro, OR, USA) and electron backscatter diffraction (EBSD) detector (Aztec Nordlys Max3, Oxford Instruments, Oxford, UK, incorporated in the SEM device) was applied to characterize the surface morphology, phase composition, and distribution. The samples for the EBSD measurements were mechanically polished with W0.5 diamond paste and then electropolished in 2% NaOH solution. Phase identification was derived from the crystal structure database (HKL, ICSD, and NIST) installed on the EBSD system.

**Figure 1.** The schematic diagram of the sample placement in the spark plasma sintering setup (**a**) and the self-designed water flowing device (**b**).

The erosion–corrosion experiments were conducted in a home-made pipe flow circulating system as illustrated in Figure 1b. Diameter of all water passages was 32 mm. The water-pump was placed into an improved thermostatic water-container (0.6 × 0.6 × 1 m) where the water circulated. PH of the water was about 6~7. The oxygen content in the passages can be regarded as a saturation value since the water-container is a semi-open vessel. The tested samples were side by side fixed on the inner upper wall of the specimen chamber, leaving only one exposed surface (Ø16 mm) parallel to the flow direction.

Table 2 lists the test parameters with reference to the simulated service conditions. The flowing velocity and temperature was controlled by the flow regulator and the thermostatic bath, respectively. The error is also marked out at Table 2. The temperature value 20 ◦C was the entrance temperature of the flowing water. At the relative low temperature of 20 ◦C, a test at the largest velocity of 2 ms−<sup>1</sup> was conducted, in light of the small difference in the corrosion behavior between 1 and 2 ms−1. After the experiments, the corrosion performance was evaluated by analyzing the corrosion rate and the corroded morphology.

The weight was recorded every 12 h in the erosion–corrosion test, using a balance with an accuracy of 0.01 mg. The average corrosion rate <sup>ν</sup> (g/cm2·h) was calculated from the following equation:

$$\mathbf{v} = (\mathbf{m}\_0 - \mathbf{m}\_t) / \mathbf{s} \cdot \mathbf{t} \tag{1}$$

where m0 is the original mass of specimen, mt the mass of specimen after corrosion for t hours, s (cm2) the exposure area before corrosion and t (h) the corrosion time. As the change in s after corrosion is very small, it is considered as a constant. It is worth pointing out in the gravimetric measurements that it was not easy to accurately clean up the corrosion deposits of the corroded specimens without damaging the matrix. Therefore, after the tests, the specimens were directly dried in a vacuum oven at 100 ◦C for 12 h without any surface treatment, and then their mass mt was measured. If ν > 0, it means a mass loss and removal of the corrosion products from the surface. According to the variation of the ν value, one can judge what happened to each specimen in the corrosion process. The corroded surface morphology was studied by means of SEM observations.


**Table 1.** The impurities with concentration over 0.001 wt.% in tungsten bulk.


**Table 2.** The parameters set in the flowing water test.

#### **3. Results**

#### *3.1. Morphology and Phases*

Before corrosion measurement, the morphology, compositions, and phase distribution of the coated sample were characterized by the EBSD system. The fractured surface of the coatings showed good bonding with the substrate and high compact structures with a uniform thickness of around 25 μm (Figure 2a). The coating consisted of external W–Cr–C (hexagonal WC (ICSD [15406]) and WC0.98 (ICSD [77738]) plus dispersed Cr7C3 (ICSD [52289]) (Figure 2b,c) and intermediate hexagonal W2C (ICSD [77567]) (Figure 2e). Close to the carbon and chromium-rich zone, tungsten carbide reached stoichiometry WC (nonstoichiometric WC0.98 is also commonly designated as WC) with an average grain size of 800 nm (Figure 2c,d), and chromium carbide is indexed as Cr7C3. Along the diffusion path, the composition becomes the carbon-depleted W2C and nearly no chromium carbides are detected in the intermediate layer due to the preferential diffusion of carbon. The W2C phase accounts for more than 95 vol.% of the coating due to its much lower Gibbs free energy than WC at high temperature [21]. Furthermore, a small fraction of WC0.98 precipitates were found in the W2C grains (Figure 2e), which were derived from the partial solid-state decomposition of W2C below 1250 ◦C [22]. The vast majority of the W2C phase remained stable due to the fast cooling rate (100 ◦C /min). Figure 2f shows the pole figures and inverse pole figures with respect to the growth direction of the W2C layer. Interestingly, the W2C layer presented obvious crystallographic preferred orientations: the <001>//Y axis (the sintering pressure direction). Strong fiber texture <001> suggests the formation of a columnar structure of W2C coating via spark plasma sintering.

**Figure 2.** (**a**) Scanning electron microscopy (SEM) image of the cross-section of coating; (**b**) phase distribution of the coating surface (an indexing rate of about 91%, the black refers to the unindexed sites, the red to the WC phase, the green to WC0.98, the yellow to Cr7C3); (**c**) Euler image and (**d**) grain size distribution of the WC and WC0.98 phases; (**e**) phase distribution of the inner coating (the blue refers to the W2C phase, the red to the W substrate, the pink dots are the WC0.98 phases; GB: grain boundary; PB: phase boundary); (**f**) pore figure and inverse pore figure of the inner coating.

Such a gradient structure from base to bottom by carburizing through SPS technology could release thermal stress and avoid cracks in the coating. Other technology like HVOF (high velocity oxygen fuel) spray can also be applied in creating nano WC-based coatings made of powders [23]. The spray process acts like the high temperature sintering of the raw powders, and micro-cracks are caused by thermal stress in the spray process. More importantly, like other conventional preparations of a WC coating, the bonding phase of cobalt is commonly introduced to densify the WC, which is known as in-resistant element to corrosion.

#### *3.2. Corrosion Rate*

Under a water temperature of 20 ◦C, the corrosion rate as a function of time is shown in Figure 3. In general, the corrosion rate for tungsten clad with W–Cr–C decreased with increasing time and gradually went to a steady state after many fluctuations. The fluctuations imply that the accumulation and breakaway of the corrosion products occurred alternatively, namely coexistence of chemical and mechanically driven corrosion. The remarkable maximum corrosion rate belongs to bare tungsten, which was located uppermost over the entire time range and finally reached as high as tenfold that of the coated sample, which was relatively stable around the zero line.

**Figure 3.** Variation of the corrosion rate with time for samples at the velocity of 2 ms−<sup>1</sup> under the water temperature of 20 ◦C.

The corrosion rates under the water temperature of 60 ◦C are shown in Figure 4. The change in corrosion rate was positively related to the flow rate. When the flow velocity was 1 ms−<sup>1</sup> and 1.5 ms<sup>−</sup>1, the corrosion rate of the W–Cr–C coated W sample was less than 0.01 mg·cm−2·h−1. The corrosion rate of bare tungsten sample almost doubled and its fluctuation with time was larger. When the flow velocity increased to 2 ms<sup>−</sup>1, the corrosion rates of all samples increased to higher values, varying in the range of 0.25–0.05 mg·cm−2·h−<sup>1</sup> during the experiment time of 360 h.

**Figure 4.** Variation of corrosion rate with time for samples at various flow velocities under the water temperature of 60 ◦C.

#### *3.3. Corrosion Mechanism*

The corroded appearances of tungsten clad with W–Cr–C are shown in Figure 5. The initial surface of the W–Cr–C coating (Figure 5a) presented deficiencies of small scattered pores left by SPS fabrication. After the 360 h test, the morphology changed slightly, but increased pores with increasing velocity or temperature, displaying a dominant feature of pitting corrosion (Figure 5b–e). Depth and the number of pores are difficult to quantify, so Figure 5f counts the area fraction of pores using image software (based upon the difference in primary contrast of images) by collecting six SEM graphs with the same magnification to describe the changes of pitting. Under the water temperature of 60 ◦C, the area fraction of the pores increased from the initial 0.4 ± 0.1% up to 4.5 ± 1% when the flow velocity varied from 0 to 2 ms<sup>−</sup>1. However, there was no single linear or parabolic law between the pore area fraction and flow velocity for a remarkable rise at the flow velocity of 2 ms<sup>−</sup>1. When the water temperature dropped to 20 ◦C, the reduction in pore area fraction reached 60%, almost close to the rate of temperature decline. Pitting area dependent upon velocity and temperature agreed well with the corrosion rate curves.

The corrosion behavior of bare tungsten was totally different, as shown in Figure 6. A rough metallographic structure with bimodal grain-size distribution appeared as a consequence of remarkable intergranular corrosion (Figure 6b–e) as well as slight dissolution of the tungsten grains (Figure 6f). The intergranular corrosion at high flow velocity (2 ms<sup>−</sup>1) was serious, accompanied by the formation of deep groove and micro corrosion pits caused by grain exfoliation, indicating that a mechanical effect would promote intergranular corrosion, weakening the GBs cohesion and loosening the GBs, especially at the triple junction.

**Figure 5.** SEM image of the W–Cr–C coated tungsten samples (the red refers to the pores): (**a**) before corrosion; (**b**) corrosion at 1 ms<sup>−</sup>1, 60 ◦C; (**c**) corrosion at 1.5 ms<sup>−</sup>1, 60 ◦C; (**d**) corrosion at 2 ms<sup>−</sup>1, 60 ◦C; (**e**) corrosion at 2 ms<sup>−</sup>1, 20 ◦C; and (**f**) the bar chart of pitting fraction with velocity and temperature.

**Figure 6.** SEM images of bare tungsten: (**a**) before corrosion; (**b**) corrosion at 1 ms<sup>−</sup>1, 60 ◦C; (**c**) corrosion at 1.5 ms<sup>−</sup>1, 60 ◦C; (**d**) corrosion at 2 ms<sup>−</sup>1, 60 ◦C; (**e**) corrosion at 2 ms<sup>−</sup>1, 20 ◦C; and (**f**) corrosion on grains.

#### **4. Discussion**

Based upon the above results, the corrosion rate of W–Cr–C clad tungsten is much lower than that of bare tungsten due to the corrosion form changing from the intergranular corrosion of bare tungsten to the pitting corrosion of W–Cr–C coating. Pitting corrosion is a unique form of anodic reaction in the electrochemical reaction as well as a local corrosion. It causes less weight loss, so the change in corrosion rate with time gradually moves to a relatively steady state, as shown in Figures 3 and 4. The occurrence of pitting corrosion is known to be closely related to the environment and surface morphology and composition of materials. The medium used in this test containing special ions of oxygen and chlorine become the necessary trigger for pitting corrosion. In the service condition, chlorine and oxygen concentrations need to be controlled to reduce the pitting corrosion. Furthermore,

reducing the porosity of W–Cr–C coating by fine polishing will help decrease the nucleation sites of pitting.

Occurrence of the intergranular corrosion on tungsten is essentially induced by the impurities, as listed in Table 1. It was already understood that the impurities were prone to segregate at grain boundaries (GBs), which resulted in a large composition difference between the GBs and tungsten grains. The nonmetal impurities such as phosphorus, silicon, oxygen, nitrogen, and their compounds acted as intergranular embrittlement elements and mild acidic electrolyte. The impurities iron, aluminum, and silicon, due to lower potential than tungsten, became the anode, which constituted a corrosion cell with tungsten around and induced pitting corrosion at the GBs. With the exacerbation and development of pitting corrosion, it evolved into macroscopic corrosion along the GBs, namely electrochemically-driven intergranular corrosion. In the regions away from GBs, chemically driven dissolution of tungsten proceeds through the interaction with oxygen and water. Understandably, the dissolution rate at the grain boundary is much higher than that on the tungsten grains. Temperature has an accelerated effect on both types of dissolution.

Intergranular corrosion at the triple junction was much more remarkable. From the perspective of defects, the volume percentage of the GBs and defect density at a triple or multiple-junction are much larger than that at conventional GBs, so corrosion would occur more easily at such regions. From the aspect of interface energy, the corrosion behavior of fine grains at the trigeminal boundaries is analogous to the pinning behavior of the second particles occupying similar sites of matrix. The pinning interaction of the second particles with the matrix grains depends on the contact area of the particles with boundaries. The smaller the second particle, the less the boundary area is occupied. For particles with a smaller size, there is a stronger unpinning tendency of the boundaries because of the lower decreased interface energy [24]. Once the pits at the triple junction of GBs are produced, it leads to a noticeable change in weight. This accounts for the large fluctuations of the corrosion rate for tungsten.

Regardless of whether it is intergranular corrosion or pitting corrosion, the effect of flow velocity on corrosion is complex. When flow velocity is elevated, on one hand, corrosion might be further exacerbated on account of sufficient oxygen supply and dissolution of the underlying metal caused by accelerated mechanical damage; on the other hand, corrosion may also be retarded due to the replenishment of inhibitors (corrosion products) to some dead spaces (grooves or pores), which hinder further corrosion [25]. Thus, in the low flow velocity range, the competition between the above-mentioned two effects approaches a balance and the corrosion rate increases slowly with increasing flow velocity. However, when flow velocity exceeds a critical value, the protruding mechanical action overshadows the inhibitors replenishing, and the materials are allowed to deteriorate at a higher corrosion rate, for instance, the corrosion at 2 ms<sup>−</sup>1.

Evidently, temperature imposes a greater and nearly liner impact in promoting intergranular and pitting corrosion than velocity within 60 ◦C. Increase in temperature caused the acceleration of ion mobility and a corresponding rise in electrochemical reaction activity. Despite this, there is a critical corrosion temperature for materials. For instance, the corrosion rate of iron begins to decrease above 80 ◦C [26], while the corrosion rate of P110 steel reaches the maximum at 60 ◦C and then decreases [27]. For tungsten and W–Cr–C coatings in this study, the critical temperatures should be higher than 60 ◦C. Therefore, in the service temperature range, temperature promoted the corrosion rate of both samples.

In the test duration, W–Cr–C endures less weight loss and corrosion rate than intergranular corrosion of tungsten, showing an expected protective role. However, it is worth pointing out that pitting corrosion occurred on the W–Cr–C layer. With an increase in time, the pores will grow gradually, and when the pores stretch to the W2C layer, the columnar crystal structure of W2C will lead the corrosion mechanism to change into intergranular corrosion, which is not conducive to protecting the substrate. Furthermore, the W2C phase is less stable than WC in neutral solution [28]. Therefore, changing the microstructure of the W2C layer and increasing the proportion of the W–Cr–C layer by adjusting the carburizing parameters is the topic of ongoing work.

#### **5. Conclusions**

This work studied the corrosion behavior of tungsten, with and without W–Cr–C cladding prepared by the SPS technique, in a flowing water system, and the main results can be concluded as the follows:


**Author Contributions:** Data Curation and Formal Analysis, Z.X.; Methodology, J.Y.; Writing—original draft, J.Y.; Writing—review & editing, Y.J. and Q.F. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the National Key Research and Development Program of China (Grant No. 2017YFA0402800), the National Natural Science Foundation of China (Grant Nos. 51601189, 11674319, 11475216), the Central Government Guides Local Science and Technology Development Projects (2019ZYD026), and the Sichuan Key Laboratory of Comprehensive Utilization of Vanadium and Titanium Resources (2018FTSZ44).

**Acknowledgments:** The authors sincerely thank the National Natural Science Foundation of China for their support of this work.

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

#### **References**


© 2020 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* **Investigations of the Deuterium Permeability of As-Deposited and Oxidized Ti2AlN Coatings**

#### **Lukas Gröner 1,\*, Lukas Mengis 2, Mathias Galetz 2, Lutz Kirste 3, Philipp Daum 1, Marco Wirth 1, Frank Meyer 1, Alexander Fromm 1, Bernhard Blug <sup>1</sup> and Frank Burmeister <sup>1</sup>**


Received: 31 March 2020; Accepted: 23 April 2020; Published: 1 May 2020

**Abstract:** Aluminum containing Mn+1AXn (MAX) phase materials have attracted increasing attention due to their corrosion resistance, a pronounced self-healing effect and promising diffusion barrier properties for hydrogen. We synthesized Ti2AlN coatings on ferritic steel substrates by physical vapor deposition of alternating Ti- and AlN-layers followed by thermal annealing. The microstructure developed a {0001}-texture with platelet-like shaped grains. To investigate the oxidation behavior, the samples were exposed to a temperature of 700 ◦C in a muffle furnace. Raman spectroscopy and X-ray photoelectron spectroscopy (XPS) depth profiles revealed the formation of oxide scales, which consisted mainly of dense and stable α-Al2O3. The oxide layer thickness increased with a time dependency of ~t1/4. Electron probe micro analysis (EPMA) scans revealed a diffusion of Al from the coating into the substrate. Steel membranes with as-deposited Ti2AlN and partially oxidized Ti2AlN coatings were used for permeation tests. The permeation of deuterium from the gas phase was measured in an ultra-high vacuum (UHV) permeation cell by mass spectrometry at temperatures of 30–400 ◦C. We obtained a permeation reduction factor (PRF) of 45 for a pure Ti2AlN coating and a PRF of ~3700 for the oxidized sample. Thus, protective coatings, which prevent hydrogen-induced corrosion, can be achieved by the proper design of Ti2AlN coatings with suitable oxide scale thicknesses.

**Keywords:** MAX phase; Ti2AlN; PVD coating; oxidation; hydrogen permeation

#### **1. Introduction**

The increasing number of applications in which hydrogen is being used as a storage medium in energy conversion technologies demands the consideration of new construction materials, or at least a profound surface conditioning of established materials to prevent, e.g., hydrogen diffusion induced embrittlement or other forms of corrosion, especially the development of so-called "white etching cracks" [1]. One route for corrosion protection is the development and application of temperature-resistant coatings with excellent barrier properties for hydrogen. Recently performed studies indicate that MAX phase materials might fulfill these requirements [2–6]. The general formula, Mn+1AXn, (short MAX) describes a family of materials consisting of an early transition metal (M), mostly a group 13 or 14 element (A) and nitrogen and/or carbon (X) with the stoichiometry of n = 1,2,3 [7]. The MAX phases crystallize in a hexagonal lattice within the space group D46h (P63/mmc) in which the octahedral Mn+1Xn layers are separated by atomic monolayers of pure A-atoms. MAX phase materials are known to have a good oxidation resistance [3,8,9], a high damage tolerance as well as a high electrical and thermal conductivity [10].

The good oxidation resistance of Al containing MAX phases usually stems from the formation of dense and thermodynamically stable thermal grown oxides (TGO) consisting of α-Al2O3 on the coatings surface at relatively low temperatures of 600–700 ◦C. For comparison, the direct physical vapor deposition (PVD) of an α-Al2O3 in an industrial scale deposition process usually requires temperatures above 1000 ◦C [11]. Lower deposition temperatures of 500–600 ◦C have also been observed but at the expense of a brittle fracture behavior [12,13]. A further advantage of α-Al2O3 oxide scales thermally grown on MAX phase coatings is the well-known self-healing effect whereby small defects or cracks in the coating, which might serve as diffusion pathways, are blocked by oxide growth [2]. For this purpose, the oxidation kinetics of the TGO has to allow for a quick healing and oxidation of the surface, but has to prevent fast oxygen diffusion to the coating–substrate interface. The oxidation kinetics of Ti2AlC at 1200 ◦C were modelled by G.M. Song et al. in [3]. This model contains the growth of oxide grains and assumes that the diffusion paths along the grain boundaries increase with time. This results in a time dependency for the increase in the thickness of the oxide scale *dOx(t)* of:

$$d\_{\rm Ox}(t) \;= \; 2\sqrt{k\_n} \times t^{1/4} \tag{1}$$

The growth factor *kn* = Ω*DGB*Δ*C*3<sup>δ</sup> *<sup>d</sup>*<sup>0</sup> contains a constant prefactor <sup>Ω</sup>, the diffusion coefficient for oxygen along the grain boundaries *DGB*, the size of the grain boundaries δ, the initial lateral grain size *d0* and the gradient in the oxygen concentration Δ*C*.

This model of α-Al2O3-formation, as well as the structural properties of MAX phases, i.e., the sequence of dense MX-layers, motivated the present investigation on their barrier properties against hydrogen diffusion.

Although little information about the diffusion of hydrogen in MAX phases exists, similarly composed carbides or nitrides of early transition metals are already used as diffusion barriers for hydrogen [4,14]. It is expected that the anisotropic structure of MAX phases will induce a directional anisotropy of the hydrogen diffusion. In [5], F. Colonna and C. Elsässer presented the findings of an atomistic simulation of diffusion processes in Ti2AlN using density-functional theory (DFT). Therein, interstitial diffusion paths of hydrogen and oxygen were examined. It was found that, for hydrogen, the migration perpendicular to the basal planes has a maximum barrier of ~3 eV, whereas the migration barrier parallel to the basal plane is one order of magnitude lower. The high migration barrier parallel to the c axis was explained by the presence of the Ti2N double layer, where the interstitial octahedral sites of Al3Ti3 are already occupied by nitrogen atoms.

An experimental study on the hydrogen barrier properties of MAX phase coatings was presented by C. Tang et al. in [6]. Therein, ZrY-4 alloy cylinders were coated with Ti2AlC and Cr2AlC by a multilayer deposition followed by a subsequent annealing step. This process led to a {0001}-textured polycrystalline growth which could also be detected in [15] for Ti2AlN. After loading the specimens in an Ar+H2 atmosphere the cylinders were investigated by neutron radiography. It could be shown that a 5 μm thick Ti2AlC and Cr2AlC reduced the penetration of hydrogen under the detection limit.

To evaluate coatings in terms of their capability to reduce the hydrogen permeation a permeation reduction factor (PRF) can be calculated using the mass specific ion current j:

$$\text{PRF} = \frac{j\_{uncated}}{j\_{coutal}}.\tag{2}$$

D. Levchuk et al. investigated Al-Cr-O coatings [16] and Er2O3 coatings [17] as hydrogen permeation barrier. Both coatings tend to form a dense crystalline structure, which is capable of effectively reducing the hydrogen permeation up to a PRF(Al-Cr-O) = 3500 and PRF(Er2O3) = 800.

#### **2. Experimental Details**

#### *2.1. Deposition of Ti2AlN*

The Ti2AlN MAX coatings were deposited on AISI 430 ferritic stainless-steel substrates (Fe81/Cr17/Mn/Si/C/S/P), which were polished (1400 grit) and cleaned in acetone and isopropanol using an ultrasonic bath prior to deposition. A custom build industrial sized magnetron sputter chamber SV400/S3 (FHR Anlagenbau GmbH, Ottendorf-Okrilla, Germany) equipped with rectangular titanium (purity 99.8%) and aluminum targets (99.999%) was utilized. To obtain a pronounced {0001}-texture with a parallel orientation of the basal planes and the substrate surface, we alternately deposited 150 single layers of Ti and AlN on the substrate, beginning with Ti. During the radio frequency-sputtering of the aluminum target, nitrogen (purity 99.9999%) was introduced in the chamber. A final subsequent annealing at 700 ◦C for 1 h in vacuum led to the formation of textured Ti2AlN MAX phase coatings. Details of the deposition process are described elsewhere [15]. The coatings thickness was in the 2 μm–3 μm range.

#### *2.2. Oxidation Procedure and Analysis*

To investigate the oxidation kinetics of the Ti2AlN coatings, comparable samples originating from the same batch were oxidized at 700 ◦C for 5 h, 10 h, 20 h and 100 h in a muffle furnace (Nabertherm GmbH, Lilienthal, Germany) in air. The samples were afterwards removed from the furnace and cooled in air. The crystallographic orientation and phase composition of oxidized and non-oxidized coatings were investigated by X-ray diffractometry (XRD) using a PANalytical Empyrean in parallel beam geometry (Empyrean, PANalytical, Almelo, The Netherlands) and Cu Kα<sup>1</sup> radiation with a 2-bounce Ge 220 monochromator. The samples were irradiated with primary X-rays using a line focus. The diffracted X-rays were detected using a PIXel-3D detector with a 1 mm slit for the phase analysis.

A surface sensitive phase analysis was performed using a confocal Raman spectrometer (Model inVia, Renishaw plc., Gloucestershire, United Kingdom) in backscatter geometry. The excitation wavelength λNd:YAG = 532 nm was used to determine possible changes in the Ti2AlN phase upon thermal treatment whereas the wavelength λHeNe = 633 nm proofed to be suitable for exciting fluorescence bands in the thermally grown AlOx phase. In all measurements, a 100-fold objective focused the laser on the surface to a spot diameter of about 2 μm.

XPS depth profiles were recorded with a PHI 5000 VersaProbe II (Ulvac-PHI, Inc., Chigasaki, Japan) equipped with an argon sputter option using Al Kα-rays. For analyzing the coarse elemental distribution close to the interface of coating and substrate a metallographic cross-section was prepared after an electrochemical deposition of Ni for protective purposes. The electron probe micro analysis (EPMA) was performed utilizing a JXA-8100 (Jeol, Akishima, Japan). The measurements were performed with an acceleration voltage of 15 kV and a dwell time of 30 ms.

#### *2.3. Deuterium Permeation Setup*

To investigate the diffusion of deuterium from the gas phase through coated and oxidized membranes, a permeation setup was developed following the works of C. Frank et al. [18], J. Gorman et al. [19] and D. Levchuk et al. [20]. The test rig consisted of two chambers separated by a thin steel membrane (see Figure 1). The high pressure side is filled with the diffusional species or the purging gas, the low pressure side is evacuated by a turbomolecular pump and an ion getter pump down to pressures of ~10-6 Pa. The latter is also equipped with a quadrupole mass spectrometer (Model PrismaPlusTM QMG 220, Pfeiffer Vacuum Technology AG, Aßlar, Germany) to determine the gas composition as well as the mass and time dependent ion current, which is detected by a secondary electron multiplier. With infrared transmissible windows on both sides, the membrane was heated by a focused halogen radiation heater and the temperature as well as the temperature distribution was recorded by a heat sensitive camera. The membranes, illustrated in Figure 1b, were water jet cut (Ø = 30 mm) from a 0.2 mm thick steel foil (AISI 430) and coated as described before. The uncoated

back sides were corundum blasted to increase the absorption of infrared radiation. The membrane was mounted with conical copper gaskets with the coating facing the high pressure side. After a minimum base pressure of 1 <sup>×</sup> <sup>10</sup>−<sup>5</sup> Pa was reached, the measurement was started.

**Figure 1.** (**a**) Schematic illustration of the hydrogen permeation test rig with quadrupole mass spectrometer (QMS), turbomolecular pump (TMP) and ion getter pump (IGP). (**b**) Schematic illustration of the coated ferritic steel membrane.

To investigate the hydrogen barrier properties, the isotope deuterium was employed in order to avoid interpretation ambiguities due to contaminations with residual gases or water molecules. The permeation measurements were performed close to thermodynamic equilibrium. First, deuterium was injected on the atmospheric pressure side. Then the membrane temperature was set to a maximum and was reduced stepwise when a constant ion current was reached. The ion current of the atomic mass m(D2) = 4 was recorded. The permeation reduction factor was calculated by (2) using the steady state values of the ion currents of a non-coated sample as a reference.

#### **3. Results and Discussion**

#### *3.1. Oxidation*

To investigate the influence of the oxygen exposure at high temperatures on the phase composition, XRD diffractograms were recorded for different exposure times and compared to the pristine sample. In Figure 2a, the diffractogram of the as-synthesized coating reveals an almost phase pure Ti2AlN coating having a strong {0001}-texture. The peak at 42.5◦ is attributed to the (200) lattice plane of TiN. The Fe-bcc peaks at 44.6◦ and 65.0◦ are assigned to the steel substrate. Diffractograms of the oxidized samples are depicted in Figure 2 with an offset for better visibility.

These phase compositions appear almost unaltered upon thermal exposure. Only in the 2Θ-region between 42◦–43◦ a slight change in the peak position is visible. This region is depicted in detail in Figure 2b. Due to broad peak widths and weak angle dependent interferences, the signals from TiN and α-Al2O3 cannot be clearly distinguished. Further ambiguities arise due to the small TGO layer thickness and its possibly amorphous structure. Hence, further surface sensitive Raman analysis was performed (see Figure 3). The Raman spectra of the coatings in Figure 3a still feature the characteristic Raman peaks for Ti2AlN despite the oxidized surface, though an increase in the background is detected. The broad background and the peak between 500 cm<sup>−</sup>1– 600 cm−<sup>1</sup> might stem from the formation of surface oxides and/or oxycarbides [21] as well as from the formation of TiN close to the surface due to Al depletion [22]. Titanium oxides like anatase and rutile, which were reported in [23] after the oxidation of Ti2AlN coatings at 750 ◦C, are not detected.

**Figure 2.** XRD diffractograms of as-synthesized and oxidized Ti2AlN coatings on ferritic steel samples: (**a**) overview and (**b**) enlarged region around 2θ◦ ≈ 40◦.

**Figure 3.** (**a**) Raman spectra (λNd:YAG = 532 nm) of the Ti2AlN coatings before and after oxidation. (**b**) Raman fluorescence spectra (λHeNec = 633 nm) of the Ti2AlN coatings before and after oxidation.

The fluorescence spectra in Figure 3b exhibit an increase in the background with an increasing oxidation time. Particularly after 100 h of oxidation, the formation of distinct peaks close to 1379 cm−<sup>1</sup> and 1402 cm−<sup>1</sup> can be detected. According to literature, these peaks are attributed to the fluorescence of Cr3<sup>+</sup> and Fe3<sup>+</sup> impurities in an α-Al2O3 environment [24–26].

After oxidation, XPS depth profiles of all coatings were created by convoluting the distribution of the binding energies. The underlying XPS spectra are not shown herein. In Figure 4 profiles of a sample oxidized for 100 h (a) and of an uncoated substrate (b) are presented for comparison. In the case of the MAX phase coating Figure 4a, oxidic Al2p bonds with a maximum in the energy of 74.3 eV were detected and ascribed to the formation of Al2O3 at the sample surface. To a smaller extend of about 8 at%, oxidic Ti2p bonds with a maximum in the energy of 458.2 eV were detected, which were distributed over two regional maxima located in a depth of ~45 nm in the Al2O3 scale and at the interface of Ti2AlN/TGO. At the interface, the shift of the Ti2p-signal to nitridic binding energies of 454.3 eV and the shift of the Al2p-signal to metallic binding energies of 72.3 eV revealed the transition to the Ti2AlN phase. With only ~16 at% of Al2p bonds and sustaining increasing signals at a depth of 200 nm, a transition regime close to the interface is observed where an Al depletion exists. The thickness dOx of the TGO scale was determined by the decrease of the O1s signal, and set to the sputter depth where the O-signal fell below 50% of the original ratio.

The numerical fit of the values for dOx according to (1) is plotted in Figure 5. The errors of the oxide thicknesses were estimated to 10 nm resulting from a temporal variation in the sputter rate during XPS measurements. The growth factor of the TGO was calculated to *kn* = <sup>402</sup> <sup>±</sup> <sup>17</sup> mol <sup>m</sup>×<sup>s</sup> with a quality factor of *R*<sup>2</sup> = 0.9794. The quality of the fit argues for the suitability of the mathematical description by (1 for the oxidation kinetics. However, no conclusion can be drawn so far as to whether the O or the Al diffuses along the grain boundaries to the oxidizing interface according to the above described model of G.M. Song et al.

The spectra of an uncoated ferritic steel substrate after 100h at 700 ◦C in Figure 4b revealed the formation of a TGO consisting of of Cr-, Mn- and Fe-oxides. The thickness of the TGO was calculated to 540 nm, which compares to approximately 100 nm in the case of a coated substrate.

**Figure 4.** XPS depth profiles after oxidation for 100 h at 700 ◦C of (**a**) Ti2AlN coating, and (**b**) uncoated ferritic steel substrate.

**Figure 5.** Measured oxide thicknesses by XPS depth profiles of the Ti2AlN coatings after oxidation. Underlying fit was performed using (1).

The EPMA images of the sample oxidized for 100 h at 700 ◦C depicted in Figure 6 represent the elemental distributions of Ti (b), Al (c), N (d), O (e), Ni (f) and Fe (g), where the colors indicate the normalized elemental concentration. The measured distribution of Ti, Al and N across the coating thickness features a depletion in Ti and Al at the interface of Ti2AlN/TGO and in the subsurface region. Accordingly, the substrate is locally enriched by Al and N and the formation of precipitates perpendicular to the surface is visible. In such areas, only a minor Fe-concentration is measured, as the microprobe signal is always to normalized 100% for all elements. The inward diffusion of Al and N is accompanying the outward diffusion of Fe into the coating according to the Fe elemental distribution map. Besides the thin oxide scale, which formed on top of the MAX phase coating, oxygen can be detected within the Ni-plating. This is caused by the formation of a longitudinal crack within the Ni-plating during the preparation of the cross-section.

The strong interdiffusion of the weakly bound A-element of MAX phases with the substrate is known to be a crucial aspect, when it comes to the chemical stability in high temperature applications [23]. Therefore, the interdiffusion should be suppressed by applying additional barrier films against Al diffusion between the substrate and coating. Further loss of the A-element also occurs during oxidation

and annealing in vacuum due to Al2O3 formation and evaporation. In [27], Zhang et al. calculated that the Ti2AlN MAX phase lattice structure is capable of accommodating Al vacancies down to a Ti2Al0.75N stoichiometry.

**Figure 6.** Scanning electron microscope image (**a**) and electron probe micro analysis (EPMA) scans of the cross-section after oxidation for 100 h at 700 ◦C for Ti (**b**), Al (**c**), N (**d**), O (**e**), Ni (**f**) and Fe (**g**).

#### *3.2. Hydrogen Permeation*

The D2 ion current was measured by secondary ion mass spectrometry using the setup illustrated in Figure 1. The influence of the coatings on the permeation was investigated in a state close to the thermodynamic equilibrium. Three different membranes were investigated: the uncoated substrate material (substrate), the substrate coated with 2.7 μm of Ti2AlN (substrate+Ti2AlN) and the substrate coated with 2.7 μm of Ti2AlN with a subsequent oxidation for 20 h at 700 ◦C (substrate+Ti2AlN+TGO). In Figure 7, the D2 ion currents are plotted in the temperature range from about 50 ◦C to 400 ◦C. Three measuring cycles were performed for each membrane. The results of the different cycles are denoted by the different symbols in the graph (square, circle and triangle). The small variance in the data points confirms the reproducibility of the permeation induced ion current.

**Figure 7.** Arrhenius Plot of D2 ion currents of ferritic membranes with and without Ti2AlN or Ti2AlN+thermal grown oxide (TGO) coating. The different symbols on each line refer to results of consecutive measuring cycles. The quasi-linear fit is performed to illustrate the Arrhenius type behavior.

The diffusion through all three membranes follows an Arrhenius type behavior, which confirms the assumption of a diffusion-controlled permeation. The deposition of a 2.7 μm thick Ti2AlN coating already reduces the permeation of deuterium. The PRF at 300 ◦C is calculated to a factor of 45. As optical investigations on the coating after the measurements still revealed some minor cracks in the

coating, the PRF might still be lower for a defect-free MAX-phase coating. However, the formation of the oxide scale leads to a further, significant reduction of the permeation. With an oxidation of the coated steel membrane for 20 h at 700 ◦C, the formation of TGO of 80 nm thickness is expected, compare Figure 5. With respect to the uncoated steel membrane, a PRF of about 3700 was achieved. This reduction can be explained by the low solubility of hydrogen in the α-Al2O3 phase as well as by the potential healing of small defects in Ti2AlN, which blocks alternative migration paths with low energy barriers. The reduction of three orders of magnitude strongly supports the initially assumed suitability of Ti2AlN coatings as high temperature hydrogen diffusion barriers.

#### **4. Conclusions and Outlook**

Ti2AlN coatings were synthesized on ferritic steel samples by a repeated deposition of Ti/AlN double layers and a subsequent annealing in vacuum. The oxidation experiments at 700 ◦C in air revealed the formation of a thin TGO at the sample surface consisting mainly of α-Al2O3. By analyzing the thickness of the TGO, the kinetic confirms the findings of G.M. Song et al., which are described by a growth in thickness with a time dependency of ~t1/4. Thereby, a thin protective oxide is quickly formed after exposure to air but further growth is strongly hindered by a slow diffusion of migrating particles through the dense oxide. It was shown that the TGO not only serves as an effective protective layer against further oxidation, but also serves as a diffusion barrier against hydrogen. Whereas a 2.7 μm thin Ti2AlN coating reduces the permeation of deuterium to a factor of 45, the formation of an α-Al2O3 scale further reduces the permeation within three orders of magnitude. The healing of coating defects like pores and cracks at elevated temperatures upon oxidation is seen as an additional advantage of thermally grown diffusion barriers in comparison to directly deposited barrier coatings.

Further investigations need to focus on the interdiffusion process at the interface of coating and substrate in order to reduce the loss of the Al, which is required for the formation of α-Al2O3. Finding an optimum thickness of the TGO, which significantly reduces the hydrogen permeation and at the same time exhibits a sufficient thermal and mechanical stability is the crucial task for the utilization of Ti2AlN as protective coatings in industrial applications.

In summary, Al containing MAX phase coatings, which tend to form a dense α-Al2O3 on the surface upon oxidation, seems to be effective protective coatings in high temperature applications, where oxygen and hydrogen corrode the substrate material.

**Author Contributions:** Conceptualization, L.G.; investigation, L.K., L.M., P.D. and L.G.; technical support, M.W.; writing—original draft preparation, L.G.; writing—review and editing, L.M., M.G., F.M. and B.B.; supervision, F.B.; funding acquisition, A.F. and F.B. All authors have read and agreed to the published version of the manuscript.

**Funding:** Financial support by the Baden-Württemberg-Stiftung gGmbH in the context of "CleanTech" (project CT-6 "LamiMat") is gratefully acknowledged.

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

#### **References**


© 2020 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 E**ff**ect of Laser Power on the Properties of M3B2-Type Boride-Based Cermet Coatings Prepared by Laser Cladding Synthesis**

#### **Zhaowei Hu, Wenge Li \* and Yuantao Zhao \***

Merchant Marine College, Shanghai Maritime University, Shanghai 201306, China; huzw0731@sina.com **\*** Correspondence: wgli@shmtu.edu.cn (W.L.); zhaoyt@shmtu.edu.cn (Y.Z.); Tel.: +86-1391-799-6912 (W.L.)

Received: 12 March 2020; Accepted: 14 April 2020; Published: 16 April 2020

**Abstract:** Boride-based cermet can serve as a good protective coating for low-corrosion and wear-resistant materials, such as carbon steels, due to their mechanical and chemical properties. In this study, M3B2 (M: Mo, Ni, Fe, and Cr) boride-based cermet coatings were fabricated on Q235 steel with mixed powders of Mo, B, Ni60, and Cr by laser cladding synthesis, and the effects of laser power on the properties of the cermet layer were investigated. Three laser powers (2200, 2500, and 2800 W) were used at the same scanning speed. The X-ray diffraction (XRD), scanning electron microscopy (SEM), and energy-dispersive X-ray spectroscopy (EDS) analysis confirmed that all the coatings were composed of M3B2-type borides and {Fe, Ni} alloys. The micro-hardness, corrosion, and frictional experiments showed that the cermet coatings enhanced the corresponding performances of the Q235 steels at the three laser powers. However, the micro-hardness of the coatings decreased as the power increased, and the maximum micro-hardness value was 1166.3 HV (Vickers Hardness). The results of the corrosion and frictional experiments showed that the best performance was obtained at a laser power of 2500 W, followed by 2800 and 2200 W.

**Keywords:** boride-based cermet; laser cladding synthesis; laser power; microstructure; hardness; corrosion resistance; wear resistance

#### **1. Introduction**

In recent years, wear- and corrosion-resistant materials have been used in various industrial areas. However, the slurry flow unit breaks easily due to erosive wear and erosion–corrosion, thus a material that has comprehensive properties needs to be used. Boride-based cermet composite coatings consisting of a transition metal base matrix with dispersed hard phases, such as Mo2FeB2, WCoB, MoCoB, and Mo2NiB2, possess high hardness, high melting points, good wear and corrosion resistance, and good optical and thermal properties [1–6]. These cermet coatings have been applied to injection molding machine parts, cutters for heat sealers, bearings for sea water pumps, offshore engineering parts, and slurry flow units in the coal-mining industry [7,8]. In particular, as emerging and promising materials with superior strength, hardness, wear, and corrosion resistance, M3B2-type boride-based cermet coatings, with Mo2FeB2, Mo2NiB2, and W2NiB2 as hard phases, are qualified candidates for applications that require wear and corrosion resistance [9,10].

Kenichi Takagi prepared the ternary boride using the sintering method, and systematically studied the effects of the Mo/B atomic ratio and additional elements on the mechanical properties and structure of boride cermets, and the mechanism of boride formation [11–17]. Cr and V addition resulted in improvement of the mechanical properties, such as TRS (transverse rupture strength) and hardness associated with a simultaneous phase transformation of complex boride from orthorhombic to tetragonal. The addition of Fe, Co, Ti, Mn, Zr, Nb, and W did not result in structural changes, but did degrade the mechanical properties [11]. Adding 3.5 wt.% Cr and 11.5 wt.% V, the TRS of boride

increased when the Mo/B atomic ratio increased and then decreased, while the hardness and density increased as the ratio increased.

The maximum TRS and hardness were 2.95 GPa and 90.5 RA (Rockwell A hardness) [12]. The sintering mechanism of boride was analyzed by studying the microstructure of adding 10 wt.% Cr and 12.5 wt.% V, respectively, at different sintering temperatures, and the mechanical properties were discussed. The results showed that Mo2NiB2 was formed in the compact by a reaction of 2MoB + Ni = Mo2NiB2, and the addition of V yielded better mechanical properties [13]. The addition of Mn had an effect on the mechanical properties and microstructure of Mo2NiB2-based cermets in the powder composition of Mo, Ni, B, V, and x wt.% Mn. The hardness increased with Mn addition, while the TRS increased then decreased; the maximum TRS was 3.5 GPa at 2.5 wt.% Mn [14]. Adding 12.5 wt.% V and 2.5 wt.% Mn, the TRS of boride increased with the increasing Mo/B atomic ratio and then decreased, while the hardness increased as the ratio increased.

The maximum TRS and hardness were 3.25 GPa and 90.8 RA [15]. Adding V can greatly improve the mechanical properties of cermets. Corrosion tests in the molten fluorocarbon resin revealed that the Mo2NiB2 boride cermets had far better corrosion resistance than high-speed steel and SUS 304 [16]. The effects of Cr on the properties of Mo2NiB2 ternary boride were studied, while the Mo2NiB2 molar formula was assumed as (Mo2−xNi1−xCr2x)B2 for Cr-substituted Mo and Ni. Cr was doped into Mo2NiB2 ternary borides. In particular, 10 and 15 wt.% Cr exhibited high hardness and high elastic moduli, which makes them a suitable alternative material for wear-resistant hard materials, such as WC [17]. Yuan et al. [18] prepared Mo2NiB2 cermets and found that the maximum bending strength and hardness of Mo2NiB2 cermets reached 1.85 ± 0.04 GPa and 85.7 ± 0.1 RA, respectively.

In Lei Zhang's study, Mo2NiB2-Ni cremets with different Ni contents and ball milling times were fabricated by reaction boronizing sintering [19,20]. The cermets when molar ratio of Ni/B is 1.1 had the best mechanical properties and the lowest wear rates. The cermets with a milling time of 11 h exhibited a maximum hardness and bending strength of 87.6 HRA and 1367.3 MPa, respectively. Further, the high-temperature compressive strength and tribological behaviors of the cermets were investigated [21]. Sevinch [22] prepared Mo2NiB2 by self-propagating high-temperature synthesis using a mixture of MoO3, NiO, B2O3, and Al powders.

Most research has focused on the sintering method. Wu et al. [23,24] found that continuous, dense, and adherent Mo2NiB2 cermet coatings were obtained by a laser surface cladding technique, which was chosen as laser cladding possesses some advantages over other methods, such as its convenience, high efficiency, high energy, cost effectiveness, and the fact that it is environmentally friendly [25]. However, the influence of laser parameters has not been deeply discussed in laser cladding synthesis. The laser power, as one of the most important variants, should be primarily investigated, as it would increase the scope of applications of Mo2NiB2 coatings.

Herein, several various laser powers were studied to reveal the influences of power on the microstructure and properties of M3B2 boride-based cermet coatings on Q235 steels. Then, the microstructures of the coatings were analyzed by X-ray diffraction (XRD), scanning electron microscopy (SEM), and energy-dispersive X-ray spectroscopy (EDS). The hardness, corrosion resistance, and wear resistance of the coatings were evaluated. How the laser power affected the microstructure and properties of the coatings was revealed.

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

#### *2.1. Laser Cladding Synthesis*

Commercially available Mo (Tianjiu, Changsha, China, 99.9% pure, particle size < 48 μm), B (Tianjiu, Changsha, China, 99% pure, particle size < 48 μm), Cr (Tianjiu, Changsha, China, 99% pure, particle size < 48 μm), and Ni60 (Kenna Metal, Pittsburgh, PA, USA, particle size < 45 μm) powders were used as raw materials. The composition of mixed powders is listed in Table 1. The Q235 steel

(C 0.14%–0.22%, Mn 0.3%–0.65%, Si ≤ 0.3%, S ≤ 0.05%, and P ≤ 0.045%) was used as substrate with a size of 100 mm × 80 mm × 10 mm.


**Table 1.** Composition of the powder.

The raw powders were ball milled for 12 h in a planetary ball mill machine. To remove rust, dust, and oils, the substrates of Q235 steel were sandblasted for 10 min, and then cleaned with 99.8% alcohol in an ultrasonic cleaner for 15 min. As one of the common methods, the pre-placed powder method has been utilized here to investigate the effect of laser power on the properties of coatings. The mixed powders were placed on the surface of the substrate using polyvinyl butyral (PVB) binder. The thickness of the placed layer was about 1 mm. Then, the substrate layer was dried naturally in a drying cabinet. A Rofin DC030 laser (Rofin, Hamburg, Germany) was used as the laser source. The wavelength was 10.6 μm with a rectangular spot of 6 mm × 2 mm. The substrate with the layer was scanned by a laser after drying, as shown in Figure 1. During the laser cladding synthesis process, argon was used as a protective gas to prevent oxidation. As clearly seen in Figure 1a, the argon is blown out of the nozzle shielding oxygen entering the molten pool. The laser scanning sequences are shown in Figure 1b, and the laser spot was overlapped.

**Figure 1.** Laser cladding synthesis: (**a**) X–Z direction view; (**b**) X–Y direction view.

Here, the laser powers were 2200, 2500, and 2800 W, respectively, which was decided by good surface conditions in the forgone investigations, while the scanning speed was 1 mm/s, and overlap rate was 33%. The samples of 2200, 2500, and 2800 W were marked as No.1, No.2, and No.3, respectively.

#### *2.2. Microstructure Analysis*

The phase compositions of the cermet coatings were analyzed by X-ray diffraction (XRD, Rigaku Ultima IV, Tokyo, Japan) with Cu-Kα radiation (λ = 1.54 A) operated at 40 kV and 30 mA. The detected diffraction angle (2θ) was scanned from 20◦ to 100◦ and the scanning speed was 5◦/min. The results were obtained using the XRD analysis software X'Pert HighScore Plus (v1.1, Malvern Panalytical, Etten Leur, Netherlands). The microstructure and composition distribution of the cermet coatings were characterized by scanning electron microscopy (SEM, Hitachi TM3030, Tokyo, Japan) equipped with an energy-dispersive X-ray spectroscopy (EDS, Oxford Swift 3000, Oxford, UK) machine.

#### *2.3. Properties Analysis*

The microhardness of the cross-sections of the cermet coatings was measured with a Vickers hardness tester (HXD-1000TMC/LCD, Shanghai TaiMing, Shanghai, China) under a load of 0.98 N and a dwell time of 15 s at room temperature. The corrosion resistance of the cermet coatings was evaluated by an electrochemical workstation (Autolab PGSTAT302N, Herisau, Switzerland) in 3.5 wt.% NaCl solution at room temperature. A standard three-electrode system was used. The reference electrode was a silver chloride electrode (Ag/AgCl), the counter electrode was platinum, and the working electrode was coated with specimens with a surface area of 1 cm2. Before each corrosion experiment, the cermet coatings were immersed in 3.5 wt.% NaCl solution for 1.5 h. The open circuit potential (OCP) was measured for 1 h.

The electrochemical impedance spectra (EIS) was measured at the OCP with a potential amplitude of 5 mV and a frequency of 0.01 to 100,000 Hz. Potentiodynamic polarization experiments were scanned at 1 mV/s in the range of OCP ± 0.8 V, and the corrosion potential (Ecorr) and corrosion current (icorr) were obtained by the Tafel extrapolation method [26]. The corrosion results were analyzed by the software Nova 2.1.4 (version 2.1.4, Metrohm Autolab, Utrecht, The Netherlands). The wear resistance of the cermet coatings was measured by tribometer (UMT TriboLab, Bruker, Campbell, CA, USA). The cermet coatings reciprocated sliding against a Si3N4 ball with φ8 mm at room temperature under dry sliding conditions. The experiments were continuously applied with a load of 45 N lasting 1 h and a sliding frequency of 5 Hz for a sliding distance of 6 mm. The surface features of the cermet coatings were measured using a 3D optical microscope (ContourGT-1, Bruker, Campbell, CA, USA). After wear testing, the worn morphologies were observed by SEM.

#### **3. Results**

#### *3.1. Microstructure and Composition of Cermet Coatings*

The XRD patterns of cermet coatings were recorded. The laser power was increased to analyze the change in crystal structures. Figure 2. shows the XRD patterns of the coatings. The (M2M)3B2-type boride phase, {Fe, Ni} phase and unknown phase were detected in Sample 1, Sample 2, and Sample 3. The M in the (M2M)3B2 phase represents Mo, Ni, Cr, and Fe. They also represent the sorting of the intensity of the M3B2 and {Fe, Ni} phases, respectively, in a standard card with the black numbers and green numbers. The intensity of M3B2 in Sample 1 was much stronger than in Sample 2, while the intensity of Sample 2 was a little stronger than that of Sample 3. The intensity of {Fe, Ni} in Sample 1 was much stronger than in Sample 2 and 3 at position (1). The intensities of Sample 1 and Sample 2 were approximately the same at position (2) but much stronger than that of Sample 3. Position (1) was the strongest peak and position (2) was the second strongest peak in the standard card. The peak intensities of all samples were similar at the other three positions. The M3B2 phase intensity decreased as the laser power increased. The intensity of the {Fe, Ni} phase also decreased, but not sharply, as the laser power increased.

Figure 3a,b show a macro view of the coating after cladding. It can be seen that the overall coating after lapping is relatively flat and there are no visible macro cracks. Figure 3c,d are SEM sample images after aqua regia corrosion is taken from the coating. The light gray is the substrate and the dark gray is the coating. The coating and the substrate are well combined, and the coating can be seen with detailed overlapping patterns.

The cross-section morphological images of cermet coatings were recorded with the increase of laser power, as shown in Figures 4–6. There is an interface between the substrate and the coating in the morphological images. Distributed in cermet coating Sample 1, Sample 2, and Sample 3, there are white and gray phases, which are the M3B2 hardness phases and {Fe, Ni} binder phases, respectively. The curve radian of the interface at the molten pool area is evidently larger than that of the overlap area shown in Figure 4a. The white phase is long and uniformly distributed in the coating, as shown in Figure 4, while the laser power was 2200 W. Compared with the gray phase in the coating, the white phase is more abundant. The smaller grain size is about 2.8 μm, the larger grain size is about 10 μm, and the much longer grains will break, as shown in Figure 4b.

**Figure 2.** X-ray diffraction (XRD) patterns of the cermet coatings at various laser powers: (**a**) laser power of 2200 W; (**b**) laser power of 2500 W; (**c**) laser power of 2800 W.

**Figure 3.** The macro morphology of the coatings: (**a**) 2500 W; (**b**) 2800 W; (**c**) cross section at 2500 W; (**d**) cross section at 2800 W.

**Figure 4.** Scanning electron microscopy (SEM) morphology of the cermet coatings at 2200 W: (**a**) low magnification; (**b**) high magnification.

**Figure 5.** SEM morphology of the cermet coatings at 2500 W: (**a**) low magnification; (**b**) bonding zone at high magnification; (**c**) surface area at high magnification; (**d**) middle area at high magnification.

**Figure 6.** SEMmorphology of the cermet coatings at 2800W: (**a**) lowmagnification; (**b**) highmagnification.

Figure 5 shows the SEM images that depict the cross-sectional morphology of the cermet coatings at the laser power of 2500 W, and shows enlarged views of the bonding zone (Figure 5b), the surface area (Figure 5c), and the middle area (Figure 5d), respectively. The curve radian of the interface at the laser power of 2500 W is larger than that of 2200 W. The curve radian of the bonding zone at the molten pool area is clearly larger than that of the overlap area. Figure 5a shows that the white phase is evenly distributed in the molten pool area, and less distributed in the overlap area and the area near the interface. Compared with the grain at 2200 W, the grain size at 2500 W was significantly smaller and more regular. However, the distribution of grains in the entire cross-section of the coating was more uniform at 2200 W.

It can be seen from Figures 4a and 5a that the white phase was present in a lower quantity at 2500 W than that at 2200 W. The morphology of the gray phase is different in the bonding zone and the surface area. In the surface area, the gray phase is constituted by crystals that are shaped as triangles, squares, bars, and other shapes, and the edges have burrs similar to feathers. In the interface area, most of the gray phases contain crystals that are square-shaped, and some of them contain thin, long white phases. It can be seen from Figure 5d that the shape of crystals in the white phase in the middle area is basically square, the small grain size is approximately 1.8 μm, and the largest grain size is approximately 4.9 μm.

Figure 6 shows the SEM images depicting the cross-sectional morphologies of the cermet coating at the laser power of 2800 W and shows enlarged views of the middle area (Figure 6b). As in Figures 4a and 5a, the curve radian of the interface at the molten pool area is larger than that of the overlap area in Figure 6a. The white phase was present in a lower quantity, and it was unevenly scattered in the coating at 2800 W. There are more white phases distributed at the bottom of molten pool than other areas. The morphology of the white phase is irregular, as shown in Figure 6b. The smaller grain size was approximately 5.1 μm, the larger grain size was approximately 29 μm, and the long grain size could reach up to 50 μm.

The compositions of the white and gray phases—marked as spectrums 1 to 5 in Figures 4–6—were evaluated, as shown in Figure 7, with the increase of laser power. In order to elucidate the effects of laser power on the elements of the white phases, the EDS spectrums 1, 3, and 5 were employed in three kinds of cermet coatings, respectively. Meanwhile, spectrums 2, 4, and 6 were also carried out to depict the variation of gray phases in the cermet coatings. Evidently, the element contents of the white phase and the gray phase are completely different, as shown in Figure 7. The contents of Mo are much higher in the white phases than in the gray phases, whereas the contents of Fe are greatly increased in the gray phase. Thus, the gray phases can be considered as the binder of {Fe, Ni} with few Mo, and the white phases are M3B2 with B aggregation, in which M represents Mo, Ni, Cr, and Fe. The high concentration of Fe originates from the dilution of substrate steels, which is a common phenomenon in laser cladding [27]. The detailed element contents are given in Table 2.

**Figure 7.** Energy-dispersive X-ray spectroscopy (EDS) analysis of the white and gray phases in the cermet coatings at various laser powers.



#### *3.2. Hardness of M3B2-Based Boride Cermet Coatings*

The positions at which the micro Vickers hardness tests of the cermet coatings were carried out are shown in Figure 8. Seven points were tested on the cross-section, and the first point was taken as close as possible to the coating surface. The distance between the test points was one fifth of the thickness of the coating. The distance between the test points was different in the three samples as the coating thicknesses of the samples were different.

**Figure 8.** Positions of the hardness measurements.

The results of the hardness test are shown in Figure 9. This indicated that the micro-hardness decreased along the surface of the substrate at different laser powers. The hardness of the coating decreased with increasing laser power. The hardness at 2500 W was slightly greater than the hardness at 2800 W, but significantly less than the hardness at 2200 W. The hardness suddenly changed at a certain position in the middle area of the coating when the power was 2800 W; it was not only higher than the hardness of the two positions before and after it, but also higher than the hardness of the other two samples at the same position. The maximum hardness value in all test points was 1166.3 HV, when the power was 2200 W at the surface position.

**Figure 9.** Vickers hardness of the cermet coatings at various laser powers.

#### *3.3. Corrosion Tests of Cermet Coatings*

The substrate and all cermet coatings were immersed in 3.5 wt.% NaCl solution for 1.5 h before the OCP test. The results of the OCP were listed in Table 3. Compared to the coatings and substrate, the OCP shifted positively. The difference of the OCP between 2200 and 2800 W was not great, but the OCP reached −0.3161 V at 2500 W with a significant increase.

**Table 3.** The open circuit potential (OCP) of the substrate and cermet coatings immersed in NaCl solution.


As is known, the EIS technique has been extensively used to characterize surface coatings in a semiquantitative way [28]. Nyquist plots recorded with samples immersed in NaCl solution are shown in Figure 10. All impedance spectra of the samples displayed depressed semicircles. The diameter of the impedance spectra reflected the rate of the electrochemical reaction [29]. As shown in Figure 10, the diameter of the impedance spectra arc of coatings increased compared to the substrate. The laser power had a great effect on the impedance spectra of the coating, and the diameter of the impedance spectra arc increased first and then decreased with increasing power. When the power was 2500 W, the corrosion rate was minimum.

**Figure 10.** Nyquist diagrams for the substrate and cermet coatings at various laser powers.

Moreover, Figure 11 displays the electrochemical data of the samples as Bode plots. They exhibited the same features as the Nyquist plots. The Z value was the smallest for the substrate and the largest in Sample 2 among the three coating samples at the lowest frequency, which indicated that the coating had better corrosion resistance than the substrate and had the best corrosion resistance at 2500 W.

**Figure 11.** Bode diagrams for the substrate and cermet coatings at various laser powers.

The Bode diagrams with a wide phase angle in Figure 11 indicated that there are at least two time constants [30]. The equivalent circuit of the two time constants in Figure 12 was used to fit the EIS data using Nova software. Here, Rs is the solution resistance from the reference electrode to the working electrode, Rf is the surface coating resistance, and Rct is the constant phase elements of the charge transfer resistance. The impedance of the regular phase-angle element Q is related to the angular frequency (ω) by the following relation: ZCPE = Y0 <sup>−</sup>1(jω) <sup>−</sup>n, where Y0 is a proportionality factor and n is the deviation parameter, which reflects the roughness of the surface; Qf is the constant phase angle element (CPE) between the solution and the surface, and Qdl is the CPE of the interface double-layer between the solution and the matrix. The fitting lines are also plotted in Figures 10 and 11, and the fitting results are in good agreement with the experimental results, indicating that the equivalent circuit is feasible.

**Figure 12.** Equivalent circuit model of the substrate and cermet coatings.

The fitting results of the EIS data of all samples immersed in 3.5 wt.% NaCl solution with different processing parameters are listed in Table 4. The solution resistance Rs did not change significantly, indicating that the environmental conductivity was relatively stable. Compared with the substrate, the Rf of coatings increased, and the Rf at 2200 W was much larger than at 2500 W and 2800 W, which indicated that laser power had an effect on the surface coating resistance. The polarization resistance Rp (Rp = Rf + Rct) is used to evaluate the corrosion resistance of materials [31]: the greater the Rp value, the better the corrosion resistance. The Rp of Sample 2 was 4115.6 Ωcm2, more than twice that of Sample 1 and Sample 3, and the Rp of Sample 3 was slightly larger than that of Sample 1. The results are consistent with the previous analysis of the Nyquist and Bode plots.


**Table 4.** Fitted values of the equivalent circuit of the electrochemical impedance spectra (EIS) diagram.

The polarization curves of the substrate and cermet coatings immersed in 3.5 wt.% NaCl solution are presented in Figure 13. It is evident from Figure 13 that both the substrate and the coatings have undergone the transition process from the activated state to the passivated state and ultimately, to overpassivation; however, the duration of each process varied. The self-corrosion potentials of Samples 2 and 3 shifted positively compared to the substrate, but the self-corrosion potentials of Sample 1 shifted negatively, while the OCPs of coatings were all positively shifted. The passivation state of the substrate and Sample 3 was short, quickly changing to the overpassivation state, but Sample 3 experienced a brief secondary passivation state. Samples 1 and 2 maintained a long passivation state, and Sample 2 entered a secondary passivation state after overpassivation. The polarization curves show that the coating at 2500 W had very good corrosion resistance.

**Figure 13.** Polarization curves of the substrate and cermet coatings.

According to the relationship between the current density i and the polarization potential E of the Butter–Volmer equation (Equation (1) and (2)), the polarization resistance Rp can be calculated by Equation (3) [32].

$$\mathbf{i} = \mathbf{i}\_{\text{corr}} \text{[expr(2.303\Delta E)/\mathbf{b}\_a - \exp(-2.303\Delta E)/\mathbf{b}\_c]},\tag{1}$$

$$
\Delta \mathbf{E} = \mathbf{E} - \mathbf{E}\mathbf{corr} \tag{2}
$$

$$\mathbf{R\_{p}} = \mathbf{b\_{a}} \mathbf{b\_{c}} / \text{(2.303(b\_{a} + b\_{c})i\_{\text{corr}})}.\tag{3}$$

The corrosion parameters related to the polarization curves are calculated and listed in Table 5. icorr is the corrosion current density, Ecorr is the self-corrosion potential, ba is the anode Tafel constant, and bc is the cathode Tafel constant. Ecorr,Obs in Table 5 is the observed value, and Ecorr,Calc is the calculated value. The icorr decreased from 16.9 to 2.93 μA, which indicated that the coating on the substrate greatly improved the corrosion performance of the substrate. As the laser power increased, icorr decreased and then increased, and the polarization resistance showed the opposite trend. The results of the analysis are consistent with the previous results.

**Table 5.** The electrochemical parameters obtained from the polarization curves of the substrate and coatings.


#### *3.4. Wear Tests of Cermet Coatings*

The frictional coefficient curves of the substrate and cermet coatings are shown in Figure 14. The frictional coefficients of the substrate and coatings rose rapidly at the start of the test, and then the substrate and Samples 2 and 3 entered the steady wear stage after a breaking-in process of approximately 2.5 min, whereas Sample 1 took approximately 7.5 min. After entering the steady state, the frictional coefficient of the substrate and Sample 1 displayed some fluctuations and showed a slight upward trend, while the fluctuations in the frictional coefficient of Samples 2 and 3 were small and showed a downward trend. The order of the frictional coefficients of the substrate and the coating tended to be consistent after sliding for 30 min. The frictional coefficient of the substrate was larger than that of the coatings, and this decreased as the power increased. The frictional coefficient of coatings was the smallest at 2800 W.

**Figure 14.** Friction coefficient of the substrate and cermet coatings.

The wear volumes of the substrate and cermet coatings are shown in Figure 15. The volume loss of the substrate was approximately 2.6 to 3.1 times that of the coatings. The influence of laser power on the volume loss of coatings was different from the frictional coefficient. The volume loss increased as the frictional coefficient increased. The frictional coefficient of Sample 2 was larger than that of Sample 3, but the volume loss of Sample 2 was smaller. The coating with the best wear resistance was obtained at a laser power of 2500 W.

**Figure 15.** Volume losses of the substrate and cermet coatings at various laser powers.

The three-dimensional morphologies of the wear tracks of the substrate and coatings are shown in Figure 16. The plow-furrows can be clearly identified in the worn area of the substrate, and some pits were found on the surface of Sample 1, while only fine scratches were found on the surfaces of Samples 2 and 3. As seen from Figure 16, the wear tracks of coatings had a lower wear depth than that of the substrate (31.74 μm). The wear depth of Sample 2 (9.41 μm) was similar to that of Sample 3 (9.69 μm), which was much smaller than that of Sample 1 (13.17 μm). This further confirmed that the coating had the best wear resistance at 2500 W.

The worn morphologies of the substrate and cermet coatings were observed by SEM, as shown in Figure 17. The worn surface of the substrate displayed large plastic deformation and furrows, and many adhered materials and microcracks parallel to the ploughing direction (Figure 17a). The worn surface of Sample 1 did not have obvious ploughing furrows, but had distorted plastic deformation, peeling and adhered materials, and micro-cracks appeared in the peeling layer (Figure 17b). The worn surface of Sample 2 was relatively flat, with small plastic deformation and peeling. Micro-cracks were

observed in the peeling layer, and no obvious adhered materials were seen (Figure 17c). The worn surface of Sample 3 had shallow furrows and more peeling and adhered materials, and there were micro-cracks in the dark area (Figure 17d). The analysis is consistent with the previous results.

**Figure 16.** Three-dimensional topography of the wear tracks: (**a**) substrate; (**b**) laser power of 2200 W; (**c**) laser power of 2500 W; (**d**) laser power of 2800 W.

**Figure 17.** The SEM morphology of the worn substrate and cermet coatings: (**a**) substrate; (**b**) laser power of 2200 W; (**c**) laser power of 2500 W; (**d**) laser power of 2800 W.

#### **4. Discussion**

The XRD analysis showed that there were at least two phases in the coating. By comparison with the standard card, the two phases were identified as M3B2 and {Fe, Ni}. The absolute intensity of M3B2 in Sample 1 was much stronger than in Sample 2, while the intensity of Sample 2 was slightly greater than that of Sample 3. The absolute intensity of {Fe, Ni} in Sample 1 was also much stronger than in Sample 2, but the intensity of Sample 2 was a little stronger than that of Sample 3. From the XRD diffraction peaks of the samples, their integrated intensity changes were the largest at 2200 W, the second largest at 2500 W, and the smallest at 2800 W. Therefore, according to the XRD phase calculation method, the change in the content of the M3B2 phase was the largest at 2200 W, followed by 2500 W, and the smallest change was observed at 2800 W. It can be seen from the SEM that the coatings consisted of a white phase and gray phase. The results of EDS analysis of the white and gray phases indicated that the white phase was M3B2 and the gray phase was {Fe, Ni}. Fe and Cr occupied the positions of the Ni atoms in the white phase M3B2 [21,24].

The M3B2 phase was long and uniformly distributed in the coating at 2200 W. Compared with the gray phase in the coating, the white phase was present in a greater quantity. The M3B2 phase was evenly distributed in the molten pool area, and less distributed in the overlap area and the area near the interface at 2500 W. The M3B2 phase was present in a smaller quantity, and it was unevenly scattered in the coating at 2800 W, while there was a greater amount of M3B2 phase distributed in the bottom of the molten pool than other areas. The smaller grain size was approximately 2.8 μm, and the larger grain size was approximately 10 μm at 2200 W. The small grain size was approximately 1.8 μm, and the largest grain size was approximately 4.9 μm at 2500 W. The smaller grain size was approximately 5.1 μm, the larger grain size was approximately 29 μm, and the long grain size could reach 46.5 μm at 2800 W.

When the laser power was changed, the energy obtained per unit time, the depth, and disturbance of the molten pool, and the solidification time changed as a result. When the power was 2500 W, the molten pool became deeper, and Fe in the substrate floated to the top of the molten pool; however, as the disturbance became larger and the cooling time became longer, the grain size became smaller. When the power was 2800 W, a large amount of Fe floated to the top of the molten pool, which changed the atomic ratio of the elements in the molten pool, resulting in a reduction in the quantity of M3B2 phase that was formed. The complex spatial microstructures of the coatings including the complex-shaped M3B2 and {Fe, Ni} phases were caused by the non-equilibrium solidification behavior of the laser cladding coating in the quick-cooling temperature filed, which has been investigated comprehensively in other works [33,34].

The improvement of the hardness of the cermet coatings can be explained by particle dispersion strengthening [35]. The M3B2 boride particles in the composite coatings acted as a strengthening phase that was distributed in the solid solution phase. M3B2 is a hard and brittle phase that is difficult to deform. The crystal structure of M3B2 is different from the {Fe, Ni} phase; therefore, when the dislocation cut through the particles, the slip resistance increased, which resulted in an increase in the hardness of the coating. It was found from Figure 9 that the hardness of the coating decreased as the laser power increased. The main reason for this result was the different volume fraction of the reinforcing particles in the coating. The volume fraction can be clearly seen in Figure 4a, 5a and 6a. The results of the hardness test show that the hardness gradually decreased from the surface to the bonding zone at 2200 W. Figure 4a shows the microstructure of the coating at 2200 W, the grain size in the surface area was smaller than that in the middle area, and fewer particles were present in the bonding area. The grain refinement can improve the performance according to the Hall–Petch strengthening formula (Equation (4)) [36].

$$
\sigma\_{\rm s} = \sigma\_0 + \mathbf{K} \mathbf{d}^{-1/2},
\tag{4}
$$

where σ<sup>s</sup> is the yield strength of the materials, σ<sup>0</sup> is the deformation resistance of the grain, K is the influence coefficient of the grain on deformation, and d is the average size of the grain. The hardness of the coating at 2500 W decreased slightly from the surface area to the middle area, and decreased significantly in the bonding zone. Figure 5a shows the microstructure of the coating at 2500 W, the volume fraction and size of the grains did not change greatly in the surface and middle area but occupied a small proportion in the bonding area. Figure 9 shows the hardness of the coating at 2800 W gradually decreased from the surface area to the bonding area, but suddenly increased in the middle area approaching the bonding area. The sudden increase in hardness caused by a large amount of M3B2 phase gathered at that area, which can be evidently seen in Figure 6a. Figure 6a shows the microstructure of the coating at 2800 W, the grain size and volume fraction of the second phase in

the surface and middle area were small, but they were large at test position 4. The second phase distribution in the coating was different, due to the differences of melting depth, disturbance, and cooling time of the molten pool under different laser powers.

The EIS test results are shown in Figures 10 and 11, which show that the corrosion performances of the coatings were much better than that of the substrate. The corrosion resistance of the coating was different at different laser powers; it was highest at 2500 W, followed by at 2800 W and lowest at 2200 W. The EIS plots were fitted by the Nova software, and the model Rs(Qf(Rf(QdlRct))) was applied to analyze the variations in corrosion. Both Rf and Rct increased, indicating that the surface properties of the substrate were changed with the coating, and the corrosion performance was improved. The Rf coating at 2200 W was much larger than that of coatings 2 and 3, due to the difference of the boride ceramic phase volume fraction. The law of Rf at various powers was consistent with the law of the volume fraction of boride ceramic phase M3B2 in the coating. In the EIS diagram, the high-frequency part is related to the charge transfer, which is associated with the effect of an electric double-layer [37]. The Rct of the coating at 2500 W was larger than the others, which means that the electric double-layer formed by the structure of the coating at 2500 W was more likely to hinder the electron transfer in 3.5 wt.% NaCl solution.

The polarization curves show that both the substrate and the coatings underwent the transition process from the activated state to the passivated state and ultimately, to overpassivation; however, the duration of each process varied. The corrosion resistance was improved by the coatings compared to the bare substrate. The corrosion current density icorr, the Tafel constant for an anode ba, and the Tafel constant for a cathode bc, which reflect the corrosion mechanism and kinetics are considered as electrochemical corrosion parameters [38]. The absolute values of ba are greater than those of bc, indicating that the anode is the control step of the whole electrochemical reaction. The corrosion mechanism of substrate and coatings in 3.5 wt.% NaCl solution can be deduced as follows.

Anodic reaction:

$$\mathbf{M} = M^{n+} + \mathbf{n}e^{-},\tag{5}$$

Cathodic reaction: *<sup>n</sup>*

$$\text{m}\frac{\text{n}}{2}\text{H}\_{2}\text{O} + \frac{\text{n}}{4}\text{O}\_{2} + \text{ne}^{-} = \text{n}\text{OH}^{-}\tag{6}$$

where M is the metal elements in the coating (according to the XRD and SEM analyses, M is Fe and Ni), and n<sup>+</sup> is the valence of elements. The formed passivation film can reduce the average speed of the anode dissolution process, which was characterized in Hamadou's study [39]. The polarization curves show that the passivation film formed on the substrate and the coatings, but the electrode reaction of the substrate was faster, and the film was quickly dissolved. The electrode material and surface state had a significant effect on the reaction speed of the electrode. The microstructure of the coatings was different, which caused the differences in the coatings' abilities to hinder electron transfer, and the coating at 2500 W was the strongest according to the EIS results.

The volume fraction of {Fe, Ni} was large in the surface area of Sample 3, and the short maintenance time of the passivation state was similar to that of the substrate. However, the M3B2 phase in the coating caused the potential to shift positively, and the corrosion resistance was improved. The volume fraction of M3B2 was larger in the surface area of Coatings 1 and 2, the passivation state was maintained in this microstructure. However, the dense M3B2 phase was distributed in the {Fe, Ni} phase, which made it easier to form galvanic cells between them, which accelerated the corrosion process.

The tribological properties are not inherent properties of the materials and depend on the mechanical properties and some other factors, such as the tribo-phase, roughness, and toughness [40]. The substrate and coatings were ground under the same process conditions to ensure consistent roughness. The decrease of the frictional coefficient and wear rate of the coatings compared to the substrate due to the hardness improved according to Archard's principle [41].

The substrate underwent abrasive wear and adhesion wear demonstrated by the furrows and adhered materials on the worn track (Figure 17a). The M3B2 phase in the coating resisted deformation, so that there were no obvious furrows in the wear track, and the abrasive wear was reduced. The hardness of Sample 1 was greater than that of Samples 2 and 3; however, the frictional coefficient and volume loss were greater, because the volume fraction of the M3B2 phase in Sample 1 was large, and a small amount of the bonding phase was insufficient to protect the hard particles during the friction process. Then, the particles were peeled off (Figure 17b).

The frictional coefficient of the coating at 2500 W was greater than at 2800 W in Figure 14, but the volume loss was smaller in Figure 15. Sample 3 displayed a large area of peeling during the friction process, but only a small amount of the peeling layer remained in the wear track, so the frictional coefficient was not large; however, the volume loss was large (Figure 17d). Fine, square-shaped M3B2 particles were uniformly distributed in Sample 2, which improved the hardness of the coating, and a sufficient bonding phase protected the M3B2 particles, thus, there were no obvious deformations, furrows, or adhesive materials in the wear track, only fewer peeling phenomena (Figure 17c).

#### **5. Conclusions**

In this study, M3B2 boride-based cermet coatings were developed on Q235 steel by laser cladding synthesis using different laser powers. The laser power had significant effects on the microstructure, and the properties of the coatings were studied. The coating consisted of a white, hard M3B2 phase and a gray bonding phase {Fe, Ni}. The laser power increased and the grains in the coating were refined; however, the uniformity of the distribution decreased due to the increase of the molten pool disturbance. When the power was further increased, the dilution rate increased, and a large amount of Fe entered the coating from the steel. The thickness of the coating increased, but the amount of hard M3B2 phase in the coating decreased.

The hardness of the coatings demonstrated a linear relationship with the power, that is, the hardness decreased with the increase of power, because the increase in power caused the precipitation of Fe in the steel and then the reduction of the proportion of the hard phase in the coating. The maximum hardness of the coating was 1166.3 HV, which was obtained at 2200 W.

From the EIS and polarization curve analysis, it was shown that the corrosion performance of the coating was much better than that of the substrate. The corrosion current of the coatings was reduced by an order of magnitude compared to the substrate, and the polarization resistance was increased by an order of magnitude. The corrosion performance of the coatings first increased and then decreased with increasing power. The coating with the best corrosion performance was obtained at a power of 2500 W, at which time the corrosion current was 2.93 μA, and the polarization resistance was 11949 Ω.

The frictional coefficient and wear loss of the coatings were lower than those of the substrate. The wear performance was also better. Under the same conditions, the volume loss of the substrate was 19.29 μm3, and the maximum volume loss of the coating was 7.34 μm3, which greatly improved the wear resistance. The frictional coefficient decreased with increasing power, while the volume loss was reduced and then increased. The volume loss showed that the coating had the best wear resistance when the power was 2500 W.

The prepared coating greatly improved the overall performance of the substrate. The key process parameter—laser power—should not be too low or too high, but the performance at high power was better than that at low power. At 2500 W, the coating demonstrated not only the best corrosion resistance but also the best wear resistance. These cermet coatings can be used in the manufacture and repair of mining machinery parts, seawater pump parts, offshore engineering parts, and slurry flow units.

**Author Contributions:** Conceptualization, W.L. and Z.H.; methodology, Z.H.; validation, Z.H., Y.Z., and W.L.; formal analysis, Z.H. and Y.Z.; investigation, Z.H. and Y.Z..; resources, W.L.; data curation, Z.H.; writing—original draft preparation, Z.H.; writing—review and editing, Z.H. and Y.Z.; visualization, Z.H.; supervision, W.L.; project administration, W.L. and Y.Z.; funding acquisition, W.L. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the National Natural Science Foundation of China, grant number 51572168.

**Acknowledgments:** The authors acknowledge Lei Hong from Logistics Engineering College of Shanghai Maritime University for the laser cladding experiments, Shuxiang Ma from Shanghai Maritime University for the SEM measurements, and Dongsheng Wang from Shanghai Maritime University for the frictional experiments.

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

#### **References**


© 2020 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* **Investigations on Aging Behavior and Mechanism of Polyurea Coating in Marine Atmosphere**

#### **Kaiyuan Che, Ping Lyu \*, Fei Wan \* and Mingliang Ma**

School of Civil Engineering, Qingdao university of technology, Qingdao 266033, China; 18764811850@163.com (K.C.); mamingliang@qut.edu.cn (M.M.)

Received: 5 September 2019; Accepted: 1 November 2019; Published: 5 November 2019

**Abstract:** In this investigation, the aging behaviors of polyurea coating exposed to marine atmosphere for 150 days were studied and the mechanism was analyzed. The influences on surface and mechanical properties, surface morphology, thermal stability behavior, as well as chemical changes evolution of the coating were investigated. By attenuated total reflectance fourier transform infrared spectroscopy (ATR–FTIR) and X-ray photoelectron spectroscopy (XPS), changes in the chemical properties of polyurea coatings before (PCB) and after 150 d (PCA) of aging were analyzed, and emphasis was given to the effect of aging on functional group change, the hydrogen bonding behavior, and phase separated morphology. The results displayed prominent chain scission during aging, such as N–H, C=O, and C–O–C and the hydrogen bonded urea carbonyl content showed a decrease trend. The relative content of soft and hard segments showed a significant change, which increased the degree of phase separation.

**Keywords:** polyurea; aging mechanism; morphology; chemical properties; phase separation; hydrogen bond

#### **1. Introduction**

Polyurea as a thermoset elastomer has excellent performance for its special phase structure and physical crosslinks, which is formed by reacting a diisocyanate with an amine terminated compound by a step growth polymerization process [1–4]. Due to its exceptional mechanical and physical properties, chemical and moisture resistance, polyurea has been extensively applied in a wide number of applications. [1,3] For example, polyurea has been found to change the failure and fragmentation behavior upon impact by hypervelocity projectiles and is applied in military fields such as bulletproof vests [5]. In recent years, polyurea has been increasingly used in marine engineering, such as dock steel piles, sea-crossing bridges, and drilling platforms. Polyurea has also been studied as a binder for marine joint structures exposed to impact loads of varying amplitudes and strain rates. Thus, polyurea's moisture resistance properties and its impact mitigation capacity over a wide range of frequencies make it an ideal candidate for applications that involve exposure to or submergence in sea water [6].

As is known, there are aggressive factors such as ultraviolet radiation, floating and pushing fouling, aggressive chemicals, mechanical stresses, moisture and so on in the ocean atmosphere, which can not only cause failure of coating, but also pose a threat to property and life safety [7,8]. Based on this background, research work on the constitutive properties of polyurea are being given increasingly more interest [9–13]. Youssef [14] investigated the effect of UV radiation on the dynamic mechanical properties of polyurea samples and found that the dynamic creep modulus increased with an increase in the UV exposure duration. Whitten [15] studied the color change from white to deep tan after extended exposure to ultraviolet radiation. Youssef [16] investigated the effect of UV radiation on the ultrasonic

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properties of polyurea and concluded that the acoustic wave speeds and p-wave attenuation were found to exhibit minimal change, while shear attenuation showed convergence as the temperature and ultraviolet exposure duration increased. Gupta [17] presented a new laser-generated stress-wave-based test method to measure polyurea characteristics under high strain rates and limited strains.

In recent years, many scholars have done a lot of research on the chemical composition of coatings. [18–24]. For example, S. Bhargava et al. [25] have investigated the UV aging mechanism of waterborne polyurethane coatings. The result showed that there exists scission in functional groups such as C–O–C, C=O, C–H, and CO–NH. Chain scission of the polyurethane binder resulted in the appearance of N–H groups. Boubakri [26] irradiated thermoplastic polyurethane (TPU) coatings for 140 hours of ultraviolet light and showed that glass transition temperature (Tg) of TPU coating decreased in the beginning and then increased with UV-exposure time, then drew the conclusion that there were competition between TPU chain breakage and cross-linking. Liu P et al. [27] have characterized the alkyd coating and polyurethane coating after accelerated UV aging. The results showed that the apparent destruction mode of the alkyd system and the polyurethane system were the same, but the failure mechanism was different. Zhu [28] researched the aging behavior of aliphatic polyurethane coatings and acrylic polyurethane coatings under ultraviolet light irradiation, and concluded that the acrylic urethane was mainly C–O bond cleavage, while the aliphatic urethane was mainly C–N bond cleavage. Rossi [29] showed that the main degradation mechanism of the polyurethane under UV exposure was produced by oxidation of the carbon atom at the alpha position of the nitrogen atom of the urethane group. In conclusion, past research mainly focused on the effects of surface morphology, functional group changes, and the thermodynamic properties of coating.

Polyurea has a unique microphase separation structure composed of a soft segment and a hard segment [26,30,31], of which the soft segment is composed of long aliphatic polyether chains and form amorphous domains whereas the hard segment is composed of a urea bond (NH–CO–NH), forming carbonyl to amino hydrogen bonds [32]. Research indicated that the microphase separation structure of polyurea results in the formation of good physical properties, such as high tensile strength, higher elongation, and so on. Iqbal [33] et al. investigated the effect of soft segment length on their mechanical properties and concluded that the relative content of soft and hard segments had a high correlation with its mechanical strength and modulus. Hydrogen bond plays a pronounced role in defining the macroscopic properties of polyurea, which can promote physical cross-linking between macromolecules [34]. As is known, natural exposure has a vital impact on the macro-properties and constitutive properties of polyurea coating [35–37], but there are only a few studies on the effect of exposure in marine atmosphere on the chemical properties of especially hydrogen bond behavior and phase separation morphology.

In this investigation, emphasis was given to the effect of aging on the functional group change, hydrogen bonding behavior, and phase separated morphology in marine atmosphere. The present study evaluated the morphological, chemical, thermal, and mechanical behavior of polyurea during exposure in marine atmosphere and the aging mechanism was analyzed. Spectrometer and contact angle tester, as well as mechanical tests were performed to characterize the degradation. On further studying the SEM, AFM revealed changes in micromorphology. By ATR–FTIR and XPS, changes in the chemical properties of PCB and PCA were analyzed, and emphasis was given to the effect of aging on hydrogen bonding behavior and phase separated morphology.

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

#### *2.1. Materials*

#### 2.1.1. Raw Materials

The polyurea coating was manufactured by Qingdao Shamu Advanced Material Co., Ltd (Qingdao, China), which was a two-component elastomeric material formed by the reaction of component A and B. Component A was a semi-prepolymer of terminal NCO groups formed from diphenylmethane-4,4' diisocyanate (MDI) and a polyether polyol and component B consisted of a terminal amino polyether Jeffamine® D2000 and an amine chain extender. The molecular structures are shown in Figure 1. The two components were, respectively, placed in a raw material tank and mixed, and then thoroughly reacted in a ratio of 1:1 before spraying. The spray equipment was PHX-40 proportioner (PMC Global, Inc., Branford, CA, USA) and AP-2 gun (PMC Global, Inc., Branford, CA, USA). The spray temperature and pressure were 65 ◦C and 2500 psi, respectively. The coating film thickness was 2 mm with 7 days to be fully cured.

**Figure 1.** Molecular structures of polyurea precursors.

#### 2.1.2. Properties of Polyurea Coating

The polyurea used in this investigation is pure sheet-like coating without spraying on any substrate, and is mainly applied to marine steel structure protection in engineering. It contains zero Volatile Organic Content, which is extremely friendly to the environment. Due to the fast curing speed and insensitivity to moisture and temperature, it can be continuously sprayed on a sloping or vertical surface without sagging or running.

#### *2.2. Experimental Method*

The experimental site is located at the Qingdao University of Technology, which is situated 4 kilometers away from the actual coastline. Qingdao is situated in the southeastern part of the Shandong Peninsula, borders the Yellow Sea, and is of a temperate monsoon climate, where the annual average relative humidity, temperature, as well as precipitation are 70%, 12.3 ◦C, and 680 mm, respectively. The polyurea coating used in the experiment should be hung on the exposed frame in the artificial sea pool in accordance with the requirements of GB/T9276-1996 standard. The exposure frame is oriented 45◦ to the south, ensuring that it is fully exposed to natural sunlight, as shown in Figure 2. The exposure time is 150 days, from mid-March to mid-August.

**Figure 2.** Qingdao University of Technology exposure site.

#### *2.3. Characterization*

#### 2.3.1. Contact Angle Measurements

Static contact angle measurements were carried out at room temperature on an SDC-200 contact angle goniometer (Dongguan Shengding Precision instrument Co., Ltd., Dongguan, China), using a syringe with a needle to drop a droplet of distilled water on the surface of the sample to be tested, which was then stopped for a few seconds, waiting for the drop to stabilize, and then digital images of the droplet silhouette were captured. The contact angle was measured within 30 s. The polar and dispersive components of the surface energy were calculated using the Fowkes method. Before this measurement, the sample surface before and after aging was ultrasonically cleaned with alcohol to remove impurities and oil.

#### 2.3.2. Mechanical Properties

Tensile properties were conducted using MZ-4000D1 universal testing machine (Jiangsu Mingzhu Test Machinery Co., Ltd., Yangzhou, China) at a constant speed of 500 mm/min and at room temperature (25 ◦C). The dimensions of tensile test samples were taken according to the ASTM D 412 (Standard Test Methods for Vulcanized Rubber and Thermoplastic Elastomers—Tension).

#### 2.3.3. Scanning Electron Microscope

The microstructure of the sample before and after aging was observed by a JSM-7500F scanning electron microscope produced by JEOL (Beijing, China). The resolutions were 1.0 nm (15 kV) and 1.4 nm (1 KV), respectively. The accelerating voltage range was from 0.1KV to 30 KV, and the electron gun used a tungsten filament lamp with magnifications ranging from 25 times to 1,000,000 times. Before the test, the sample was gently wiped with anhydrous ethanol to remove surface dust and was then dried. Then, the sample was sputtered-coated with gold to form a conductive film to improve the image quality and resolution.

#### 2.3.4. Atomic Force Microscopy

Tapping mode atomic force microscopy (AFM) was performed with an Ntegra Prima Scanning Probe system from NT–MDT Prima (NT–MDT Spectrum Instruments Corp., Moscow, Russia). Topographic (height) and phase images were recorded simultaneously under ambient conditions. A silicon cantilever probe with a force constant of 5 N·m−<sup>1</sup> and a resonance frequency of 70 kHz was used to work on the tapping mode. Before the test, a 15 mm × 15 mm square specimen of PCB and PCA were prepared, and the cleaned samples were obtained by absolute ethanol to remove surface dust.

#### 2.3.5. Thermogravimetric Analysis

Thermogravimetric analysis (TGA) was performed on a STA449C (Netzsch Gerateball, Selb, Germany) under the N2 atmosphere at 20 mL/min. In the experiment, a sample weighing approximately 8 mg was heated at 10 ◦C/min from room temperature to 700 ◦C.

#### 2.3.6. Attenuated Total Reflectance Fourier Transform Infrared Spectroscopy

ATR–FTIR (Attenuated Total Reflectance) spectra of PCB and PCA were investigated with VERTEX 70 FTIR spectrometer (Bruker Optics, Inc., Ettlingen, Germany) with 32 scans and a resolution of 4 cm<sup>−</sup>1.The spectral region was from 4000 to 500 cm−<sup>1</sup> and the surface of samples for measurements was cleaned by anhydrous ethanol. The degree of degradation could be detected by detecting changes in the dipole moment associated with the stretching vibration of the functional group.

#### 2.3.7. X-ray Photoelectron Spectroscopic

The X-ray photoelectron spectroscopic (XPS) experiments were performed on Thermo ESCALAB 250XI system (AXIS SUPRA DLD, Shimadzu Kratos Inc., Kyoto, Japan) with Al K<sup>α</sup> (hv = 1486.6 eV) radiation. The binding energy (BE) scale was regulated by setting the C1s transition at 284.8 eV. The accuracy of the BE values was ± 0.2 eV. The surface of the measurement samples was cleaned by anhydrous ethanol and then naturally dried.

#### **3. Results**

#### *3.1. Surface and Mechanical Properties*

Gloss refers to the ability of a surface to reflect light projected thereon and can be used to indicate the degree of aging of the coating. Figure 3 showed the change of gloss of polyurea coating for 150 days. With extended exposure time, the glossiness value showed a decreasing tendency. As the aging time reached 150 days, the glossiness value decreased 91.95%. Literature [38] stated that the greater the roughness of the coating surface, the lower the reflectivity and gloss value obtained, so the decrease of the coating gloss corresponded well to the cracks of the coating surface.

**Figure 3.** Change in gloss with aging time.

Determining the water contact angle can provide information about the change in the surface topography through the wetting properties [39,40]. The contact angle results and calculated surface energies of the polyurea coatings were presented in Figure 4. As observed, a steady decrease in the value of the contact angle was evident. The contact angle value of the coating before aging was about 78◦ ± 1.2◦ and dropped to 37.7◦ ± 2.1◦, as the aging time reached 150 days, which decreased by about 51.28%. The result indicated a significant change in its wettability on the sample. As the value decreased, the levels of hydrophilicity increased, which could increase the adhesion of contaminants [41]. This was one of reasons for the decrease in gloss.

**Figure 4.** Change in contact angle with aging time.

In this section, the aging impact on the tensile properties of the polyurea coating was investigated. Figure 5 showed the result of a polyurea coating during natural exposure to the ocean atmosphere for 150 days. A slight increase in the tensile strength values was observed in the sample during the early stage of aging. This related to the enhancement of molecular cross-linking in the early aging period, and in this period, the effect of molecular cross-linking was greater than molecular bond rupture [26]. However, the tensile strength presented the descent trend on the whole, in the long term, and the tensile strength reduced by 7.44% after exposure for 150 days.

**Figure 5.** Tensile strength of polyurea coating.

#### *3.2. Surface Morphology and Topography*

The morphology changes on the surface of samples before and after ocean atmosphere exposure for 30 d, 60 d, 90 d, 120 d, and 150 d were obtained by SEM in Figure 6. As shown in Figure 6, the surface of the unaged coating was relatively smooth and homogeneous, and no obvious holes and cracks appeared. With an increase in the aging time, the coating surface generated more cracks, which resulted in the increase of the surface roughness. The defect area calculated by the image J software is shown in Figure 7. The increase in area defects correlated to the exposure duration, which was consistent with research conducted by Youssef [14]. The contact angle decreased with an increase in the defect area, and the two were basically negatively correlated. The appearance of cracks on the coating was related to the breakage of molecular bonds. As found in the literature [38], under the combined action of violet radiation and other influencing factors, the molecules in the coating could degrade to form micropores, and then microcracks had developed. The degradation products might also leave the surface, creating a rough surface.

**Figure 6.** SEM of polyurea—(**a**) 0-day aged; (**b**) 30-days aged; (**c**) 60-days aged; (**d**) 90-days aged; (**e**) 120-days aged; and (**f**) 150-days aged.

**Figure 7.** Correlation between the defect area and the contact angle.

The surface morphology of PCB and PCA were probed using tapping mode AFM, and the phase contrast, and topographic images were shown in Figures 8 and 9. Figure 8 provided direct visual evidence of two distinct phases—the darker areas correspond to the soft segment domain and the lighter areas correspond to the hard segment domain [42], the hard segment dispersed randomly within the soft segment, forming a microphase separation structure. Compared to the PCB, the PBA had a higher ordinate, indicating a greater degree of microphase separation. Figure 9 showed AFM 2D (Figure 9a,b) and 3D (Figure 9c,d) topographic image representations of PCB and PCA. We can clearly observe in Figure 9c,d, using a qualitative approach, that an increase in surface roughness was achieved after exposure in the marine atmosphere. According to the NOVA software analysis, the root mean square roughness (Rrms) value for the original sample was 6.03 nm and the value of average roughness (Ra) was 3.24 nm. The surface topography was relatively flat. As the exposure time reach 150 days, an obvious increase in the roughness degree values could be observed, the Rrms value reached 49.50 nm while the Ra was close to 36.3 nm. Surface roughness was the main reason for reducing the contact angle [41] which explains the reason for the above-mentioned contact angle reduction. The surface height of PCB and PCA was analyzed by the NOVA software, and the results are shown in Figure 9e. As observed, there are two different roughness profile zones—some very rough zones with a surface height from 50 to 300 nm and other zones had very low roughness values in the range of less than 50 nm.

**Figure 8.** Atomic force microscopy (AFM) of polyurea coating (**a**) 3D phase contrast image of PCB; and (**b**) 3D phase contrast image of PCA.

**Figure 9.** AFM of polyurea coating (**a**) 2D topographic image of PCB; (**b**) 3D topographic image of PCB; (**c**) 2D topographic image of PCA; (**d**) 3D topographic image of PCA; and (**e**) the height of surface before and after 150 d aging.

#### *3.3. Thermal Stabilities*

Thermal gravity (TG) weight loss curves and the differential thermal gravity (DTG) for PCB and PCA are shown in Figure 10. The onset degradation temperatures are defined by the temperatures of 5% weight loss in TGA curves, whereas temperatures of the maximum degradation rate were evaluated by the peaks in the DTG curves. It can be seen that PCB and PCA both exhibited only one degradation step in nitrogen atmosphere during 250–500 ◦C. Compared to PCB, PCA had a better thermal stability, for it started to degrade at as high as 297 ◦C and the maximum weight loss rate occurred at 413 ◦C, while the coating before aging started to degrade and the maximum weight loss rate occurred at the temperature of 272 ◦C and 403 ◦C, respectively. The results of this experiment had the following implications. First, it is possible for PCA to degrade into small molecules, which has a better heat resistance than macromolecules. Second, the soft segment content is reduced, and the coating with a high hard segment content exhibits better heat resistance [42]. Practical reasons need further experimentation.

**Figure 10.** TG and DTG curves for polyurea before and after aging.

#### *3.4. Chemical Changes*

The FTIR spectra of PCB and PCA are shown in Figure 11. The spectral bands corresponding to the hard segment domains were observed at 3284–3287 cm−<sup>1</sup> for N–H stretching vibrations and at 1600–1700 cm−<sup>1</sup> attributed to C=O stretching vibrations. The absorption bands pertaining to the C–N stretching vibrations were at around 1540 cm<sup>−</sup>1, all above were associated with the existence of urea bond in the interior of the composites. The Figure 11 shows that the major difference between the spectra corresponding to PCB and PCA was their intensity of bands. A decrease in the intensity signal of the aged sample indicated the beginning of chain scission or the disappearance of a certain group. For example, the absorbance peaks at 1020–1100 cm-1 decreased, confirming the cleavage of the C–O–C bond at exposure aging, and at the same time, the peak of C=O absorption decreased because of the scission of C=O band. The secondary amino N–H stretching vibration peak near 3287 cm−<sup>1</sup> was weakened and broadened, and shifted towards a lower frequency, which indicated the existence of N–H bond cleavage and the reduction in the degree of hydrogen bonding. The C–H stretching vibration in the 2860–2970 cm−<sup>1</sup> range was obviously weakened, probably the C–H bond in the D2000 side methyl and methylene had fractured [43]. The results coincided perfectly with the evolution of the surface energy and the tensile properties.

**Figure 11.** FTIR spectra at the region of 4000–500 cm−<sup>1</sup> for unaged and aged polyurea.

Hydrogen bonding characteristic was one of the most notable features of the polyurea materials, especially polyurea materials with a high hard segment content [44]. As major objectives of hydrogen bonding behavior research, amino, and carbonyl largely determine the hard segment structure and affect the structure and properties of the material [30]. In order to investigate the effect of aging on hydrogen bonding degree, emphasis was given to the amino region and the urea carbonyl region of the spectrum.

The peak at about 3440 cm−<sup>1</sup> was assigned to the amino that had not formed a hydrogen bond. In the N–H region, the stretching vibration band of N–H of PCB and PCA showed in the frequency of 3287.28 cm−<sup>1</sup> and 3284.46 cm−1, respectively. The magnitude of the shift of the infrared absorption peak number was the indicator of the strength of the hydrogen bond, thus, the length of hydrogen bond can be calculated through the following Equation (1) [45].

$$R = 3.21 - \frac{\Delta V}{0.548 \times 10^3} \tag{1}$$

where ΔV represented the shift distance from the amino band without hydrogen bonding to the existing amino band, while R denote the length of hydrogen bond. From the above formula, the lengths of the amino hydrogen bonds before and after aging were both calculated to be about 2.93 Å, thus, aging has little effect on the hydrogen bond length in the N–H region.

A detailed analysis of hydrogen bonding characteristics was carried out by curve fitting the broad band in the carbonyl region of PCB and PCA. The band deconvolution analysis was performed on the carbonyl region at 1620–1690cm−<sup>1</sup> using the OMNIC 8.2 software (Thermo nicolet, Madison, WI, USA). Hydrogen bonded carbonyl ureas were characterized by the presence of disordered and ordered urea carbonyls, consistent with previous reports [46,47]. The peaks due to the urea domain were characterized by bands at 1626–1645 cm−<sup>1</sup> and 1651–1669 cm−1, attributed to the ordered and disordered hydrogen-bonded urea carbonyls, respectively [43,47,48]. The absorption bands arising around 1673–1685 cm−<sup>1</sup> corresponded to free hydrogen-bonded urea carbonyl. The wavenumber assignment areas as well as the hydrogen bonding degree of the deconvoluted bands are shown in Table 1. The degree of hydrogen bonding corresponding to the urethane carbonyl was calculated using Equation (2) [47]:

$$X\_o = \frac{A\_o}{\left(A\_o + A\_{dio} + A\_{frc}\right)}$$

$$X\_{dio} = A\_{dio} / \left(A\_o + A\_{dio} + A\_{frc}\right) \tag{2}$$

$$X\_b = X\_o + X\_{dio}$$

where Ao, Adiso, and Afree denote the area of ordered, disordered, and free hydrogen bonded urea carbonyl, respectively. Whereas, Xo, Xdiso, and Xb showed the percentage of ordered, disordered, and free hydrogen bonding. A comparative analysis of the relative trends in the hydrogen bonding degree of PCB and PCA needed to be considered separately. According to the data in Table 1, major differences between the two sets of samples towards the hydrogen bonding behavior were the degree of ordered, disordered, and total urea hydrogen bonding.


**Table 1.** Degree of hydrogen bonding in the polyurea obtained from the deconvolution of the carbonyl stretching region.

The Xo value of PCB and PCA were 57.66% and 46.13%, respectively, whereas the Xb values were 83.41% and 74.56%. In contrast to PCB, a significant decrease in the Xo and Xb values of PCA was observed. The Xo value showed a slight increase tendency compared with PCB. Based on the observations above, the ordered hydrogenated urea carbonyl content in PCA was less than PCB, which was related to the intermolecular force, for a reduction in the degree of hydrogen bonding can result in the weakening of the aggregation between hard segments and a decrease in the order of the hard segment structure [49]. Thus, results clearly demonstrated that the intermolecular force of the polyurea after aging was reduced. Literature pointed out [46], the hydrogen bonding in the region of carbonyl can limit the phase separation, as compared to the cases with less hydrogen bonding. The results showed the reduction of hydrogen bonding urea carbonyl content, which improved the degree of microphase separation in the soft and hard sections, in good agreement with the results of AFM.

In the XPS spectrum shown as Figure 12, three obvious signals were detected at 285.3 eV (C 1s), 399.8 eV (N 1s), and 532.7eV (O 1s). It can be seen that relative content of N, C, and O elements was slightly changed and no new elements were produced. The contents of N, C, and O elements and the ratio of elements of PCB and PCA were shown in Table 2. After aging, the relative content of the N element increased more obviously, while the O element and the C element decreased slightly. The oxygen–to–carbon (O/C) and the nitrogen–to–carbon (N/C) ratio of PCB was 0.2598 and 0.0565, respectively, whereas that of PCA was 0.2605 and 0.0765, respectively. The increase of O/C and N/C ratio indicated that the coating degraded during exposure in the marine atmosphere, and the degradation caused the inefficiency of coating.

**Figure 12.** XPS full spectrum analysis before and after polyurea aging.


**Table 2.** Elemental composition and atomic ratios of polyurea before and after aging.

In order to elucidate the changes in chemical groups, the C1S spectrum was chosen and deconvolved with the XPSPEAK software as shown in Figure 11. The C1S spectrum of PCB can be resolved into four contributions (Figure 13a). The peak with lowest BE, 284.8 eV, was assigned to carbon atoms which were linked to hydrogen (C–H) or carbon (C–C). The peak observed at 285.7 was assigned to the carbon atom bonded to the nitrogen atom (C–N). The peak located at 286.3 eV corresponded to carbon atoms single-bonded to oxygen atoms (C–O–C) [50]. According to the data in Table 3, the relative content of C–O–C before aging was estimated to be 39.860%, while C=O was about 1.859%. After 150 days of aging, the percentage of C–O–C and C=O content changed greatly, the C–O–C content fell to 24.626%, while the C=O content increased to 15.963%. This phenomenon can be explained as the bond breakage degree of the C–O–C in the soft segment after aging, which was much higher than that of the hard segment urea bond –NHCONH–, and the soft segment phase was more damaged than the hard segment. This indicated the different sensitiveness of the hard and soft segment to the attack of the marine atmosphere exposure, and that the soft segments were more susceptible to erosion, which explained the reason for a greater degree of microphase separation of PCA.

**Figure 13.** Component-fitted XPS spectra of the C1s peaks of polyurea coating (**a**) before and (**b**) after aging.

**Table 3.** Component-Fitted C1s Spectra of polyurea before and after aging.


Through the above analysis results, we could draw up the partial aging mechanism shown in Scheme 1. As observed, aging had a great influence on the hydrogen bonding of the hard segment, especially the carbonyl region. After aging, the degree of hydrogen bonding in the carbonyl region was significantly reduced. The decrease of hydrogen bonding degree could weaken the aggregation between hard segments and reduce the order of the hard segment structure, which resulted in the decrease in the mechanical properties. Further studies found that the sensitivity of polyurea between soft and hard segment were significantly different, and the degree of C–O–C bond breakage in the soft segment was higher than that in the hard segment, which caused the soft segment to be etched, and

then resulted in the increase of roughness and the microphase separation degree. Macroscopically, the gloss of the material was decreased, and the micro-cracks were generated on the surface of the coating on the microscopic surface, which was one of the important reasons for the decline in the mechanical properties.

**Scheme 1.** Schematic illustration of the aging mechanism of polyurea coating.

#### **4. Conclusions**


**Author Contributions:** Conceptualization, K.C., P.L., and F.W.; Data curation, K.C. and P.L.; Formal analysis, K.C., P.L., and F.W.; Funding acquisition, P.L. and M.M.; Investigation, K.C., P.L., F.W., and M.M.; Methodology, K.C., P.L., and F.W.; Project administration, K.C., P.L., and F.W.; Resources, K.C. and P.L.; Supervision, K.C., P.L. and F.W.; Validation, K.C., P.L., and F.W.; Writing—original draft, K.C. and P.L.; Writing—review & editing, K.C., P.L., and F.W.

**Funding:** National Natural Science Foundation of China (No. 51503116 and No.51578298).

**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* **Material Analysis and Molecular Dynamics Simulation for Cavitation Erosion and Corrosion Suppression in Water Hydraulic Valves**

#### **Masoud Kamoleka Mlela 1, He Xu 1,\*, Feng Sun 1, Haihang Wang <sup>1</sup> and Gabriel Donald Madenge <sup>2</sup>**


Received: 29 December 2019; Accepted: 13 January 2020; Published: 17 January 2020

**Abstract:** In the milestone of straggling to make water hydraulics more advantageous, the choice of coating polymer for water hydraulics valves plays an essential role in alleviating the impact of cavitation erosion and corrosion, and this is a critical task for designers. Fulfilling the appropriate selection, we conflicted properties that are vital for erosion and corrosion inhibitors, as well as the tribology in the sense of coefficient of friction. This article aimed to choose the best alternative polymer for coating on the selected substrate, that is, Cr2O3, Al2O3, Ti2O3. By applying PROMETHEE (Preference Ranking Organization Method for Enrichment Evaluations), the best polymer obtained with an analyzed performance attribute is Polytetrafluoroethylene (PTFE) that comes up with higher outranking (0.5932052). A Molecular Dynamics (MD) simulation was conducted to identify the stronger bonding with the regards of the better cleave plane between Polytetrafluoroethylene (PTFE) and the selected substrate. Polytetrafluoroethylene (PTFE)/Al2O3 cleaved in (010) plane was observed to be the strongest bond in terms of binding energy (3188 kJ/mol) suitable for further studies.

**Keywords:** polymer; water hydraulic valve; cavitation; erosion; corrosion; coating selection; molecular dynamics

#### **1. Introduction**

Today, water hydraulics still face some major adversity to expand their application. Generally, the initial cost of hydraulic fluid components is higher than the hydraulic oil components. This property can be reduced using a deficient cost pressure medium (water), much lower insurance, and disposal costs, if the time used is long enough [1]. Specifically, water hydraulics have a more rapid response and better efficiency compared to oil hydraulics. They are also more stable (in terms of flow velocity and efficiency) over a wide range of operating temperatures due to the higher volume module of water (about 50% higher than that of mineral oil). They also have a lower viscosity (less than 1/30 at 50 ◦C mineral oil) and higher specific heat capacity (multiple of 2.2 higher than mineral oil). All of the above benefits make water hydraulics particularly interesting for robotics for high-performance actuation techniques, in addition to improving people's awareness of environmental protection and sustainable development requirements [2–5]. Hydraulic systems that utilize water as a pressure medium could be the correct solution for the environmental and safety problems of most oil hydraulic systems [6]. Water hydraulic systems have been widely used in the fields of steel and glass production, ocean exploration, food and medicine processing, and coal mining [3,7]. Water-based systems are highly profoundly vulnerable to cavitation. The imploding of vapour cavities inside the flow motivates

pressure pulsations, which may conjointly cause intense noise, part vibration, energy loss, erosion and corrosion of internal surfaces and eventually, scale back the performance of the system or failure of the part [8,9]. In rivers and seawater, several works studied the mechanism of cavitation, erosion abrasion, and corrosion. Cavitation bubbles are created when the fluid's static pressure drops below the vapour pressure at a specific temperature. [3,8]. The formation and collapse of bubbles of vapor would result in high local temperatures and high pressure [10]. The valve is one of the critical hydraulic components and due to the sharp drop in pressure caused by throttling, the occurrence of cavitation is normal [10]. Significantly, fluid valves can be classified as non-continuous valves (e.g., shut-off and change valves) and continuous valves (e.g., servo and proportional valves) [11]. There are already numerous water hydraulic power control systems on the market. Usually, their components are made of stainless steel to ensure satisfactory performance in harsh, conventional operating conditions. They do not, however, provide the necessary quality and long-term, low-friction, and low-wear efficiency [6]. Tensile properties are made up of the materials' reaction to resist when tension forces are applied. It is vital to establish tensile properties because it provides information on the elasticity module, elastic limit, elongation, proportional limit, area reduction, tensile strength, yield point, yield strength, and other tensile properties [12]. Fairfield in [13] declares that the studied erosion resistance was a substitute for a primarily unknown combination of other properties, including fracture resistance, strength, impact resistance, surface roughness, hardness, and service temperature limitation. He additionally clarifies that material properties with the largest contribution to jets resistance are compressive, tensile, and flexural, without forgetting surface roughness, thickness, peak service temperature, and thermal conductivity. Due to polymer significance in many applications, the characterization and study of superhydrophobic surfaces are of significant interest. Many material science studies have focused on surfaces with thrilling behaviors of wettability, such as superhydrophobic and superhydrophilic. The surface must have hydrophobic chemistry for the coating of polymer with other metal oxides of choice. The angle of contact is the angle that a liquid and solid surface create when they come into contact. This angle depends on both the material properties and the attractive and repulsive interaction between the materials. When the liquid spreads over the surface, it forms a small contact angle, whereas the contact angles are higher when the contact area between the surface and the liquid is smaller. The surface properties were analyzed and showed a good anti-corrosion ability [14]. The estimation of polymer solubility in solvents is one of the most important applications of solubility parameters [15]. Joel H. Hildebrand in [16] (who performed definitive work on the solubility of non-electrolytes in 1916 and lay the groundwork for solubility theory) recommended the square root of the cohesive energy density as a numerical value showing a particular solvent's solvency behavior [15]. The cohesive energy density is a numerical value describing the vaporizing energy in calories per cubic centimeter and a direct representation of the severity of van der Waals forces binding the substance's well-organized molecules [17]. Nevertheless, water's low viscosity, strong corrosion, and inadequate lubricating capacity have created significant challenges for this component's development and application. It is, therefore, an urgent task to screen materials suitable for the hydraulic water valve, which is the purpose of this research [18]. Hardness has long been considered a key criterion for the measurement of wear resistance. The erosion rate is correlated with these material properties, such as dynamic hardness, critical failure strain, and toughness [19].

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

#### *2.1. Polymer Selection for Water Hydraulics Valves*

In the engineering design and manufacturing process, the material plays a vital role. The appropriate selection of materials for a specific function is one of the designers' vital tasks. To meet the end specifications of the consumer, developers need to evaluate, with specific functionalities, the quality of various materials and find appropriate materials. The selection of materials is a challenging and time-consuming task because of the presence of a large number of materials with different

attributes. [20]. The advancement of MCDA methods was not only inspired by a variety of real-life issues that need to address several considerations but also by experts' ability to incorporate stronger decision-making methods using recent developments in mathematical modeling, statistical analysis, and computer technology. To date, many computational approaches have been developed and applied to solve problems of material selection arising from various fields of engineering [20]. The PROMETHEE (Preference Ranking Organization Method for Enrichment Evaluations) approach is one of the latest MCDA methods developed by Brans [21] and further extended by Vincke and Brans [22].

PROMETHEE is an outranking method for a finite set of alternative actions to be ranked and selected among the often widely divergent criteria. PROMETHEE is, indeed, a straightforward design and application ranking method compared to other multi-criteria analysis methods [23]. This part will deal with the polymer coating selection and particularly utilizing the PROMETHEE (Preference Ranking Organization Method for Enrichment Evaluations) method. Figure 1 shows a flowchart for coating selection that uses the PROMETHEE technique, and Table 1 shows the used criteria used for polymer selection.

**Figure 1.** Flowchart for polymer coating selection.

#### 2.1.1. Polyvinylchloride (PVC)

PVC was examined as amongst the most chemically antagonistic polymers and hydrophilicity [24]. Generally, PVC as seen in Figure 2a dissolves into polar solvents but is very resistant to hydrocarbons,

alcohols, esters, acids, bases, and salts. Its chemical defiance can be predicted using solubility parameters [25]. When warmed exceedingly its Tg (about 87 ◦C), unstabilized PVC encounters dehydrochlorination [26]. PVC is the third most widely produced polymer in the world, with its outstanding mechanical strength, flame choke off, thermal stability, and insulation peculiarity, and has been mostly used as a contentedness resin in many products [27]. Table 1 gives some useful PVC properties for water hydraulics valves.


**Table 1.** Selected polymers and properties that are suitable for water hydraulic valves.

Annotation: P1 is the polyvinylchloride, P2 is the polytetrafluoroethylene, P3 is the polydimethylsiloxane, P4 is the Polymethylmethacrylate, and P5 is the Polyaryletheretherketone. While non-beneficial criteria are water absorption or equilibrium in the water at 23◦ C (A1), Hildebrand Solubility(A2), and coefficient of friction(A3). The beneficial criteria are Contact angle (A4), Tensile strength (A5), Hardness Shore D (A6), Impact Strength (A7) and Chemical resistance (A8) [24–27].

#### 2.1.2. Polytetrafluoroethylene (PTFE)

Polytetrafluoroethylene (PTFE) in Figure 2b also recognized as Teflon. This material notably used in spacecraft design, automotive industries, and semiconductor design [38]. Disclosed utilization of the PTFE is divulging with the fact that it has meritorious chemical inactivity and high thermal stability. The coefficient of sliding friction between PTFE and many engineering materials is extremely low and has the self-lubrication capability [39,40]. It is insoluble in all common solvents and is resistant to almost all acidic and caustic substances [40]. PTFE is resistant to attack even by corrosive solutions and is practically unaffected by water (hydrophobicity) [26]. PTFE is subjected and considered as superplastic. Table 1 provides some useful PTFE properties suitable for water hydraulic valves. The essential characteristic of a superplastic material is its high strain rate sensitivity of flow stress that entertains high resistance [41]. Although PTFE has a high impact strength, its tensile strength and resistance to wear and creep are low compared to other polymers in the materials' database. Polytetrafluoroethylene is primarily used in applications needing extreme strength, exceptional chemical and heat tolerance, good electrical properties, low friction, or a blend of these properties [26].

#### 2.1.3. Polydimethylsiloxane (PDMS)

The Figure 2c shows a silicon-based elastomer Polydimethylsiloxane which has several interesting properties, along with biological and chemical inertness, extremely low Tg, high temperature, and oxidation ability to resist, and vapor permeability [42]. After preparing the PDMS solid, the hydrophilic uniqueness of the PDMS was lowered so that it can be covered with metals for electrode and microchannel applications [43]. Table 1 provides some vital water hydraulic valve PDMS properties.

#### 2.1.4. Polymethylmethacrylate (PMMA)

Poly (methyl methacrylate) as shown in Figure 2d, is the most significant member of acrylic polymers, distinguished by its thrilling weatherability, strong crystallinity, and extreme glassiness to visible light. It is mechanically robust with excellent insulating properties, making this thermoplastic polymer a choice for many mechanical, microelectronic and electrical uses, and good chemical and electrical and thermal tolerance [26,44]. In the arena of mechanical strength, PMMA has a low elongation at breakage and a high Young's Modulus. Therefore, it does not splinter upon rupture and happens to be one of the hardest thermoplastics with high scratch resistance [45]. Table 1 presents some useful PMMA properties for valves for water hydraulics.

#### 2.1.5. Polyaryletheretherketone (PEEK)

Polyetheretherketone (PEEK) is a tough semi-crystalline thermoplastic polymer with excellent mechanical and dielectric properties, as well as good chemical resistance [39,46]. Polyetheretherketone (PEEK) as seen in Figure 2e is a member of the Polyaryletherketone group that exhibit a strong-performance polymer (HPP) that is commonly used in relatively movement applications. It has extreme chemical resistance; it is suitable for tribological applications due to its high strength and wears resistance. Polyaryletheretherketone (PEEK) has outstanding properties due to its semi-crystalline structure and the molecular strength of its repeating units [46].

**Figure 2.** Shows (**a**) Polyvinylchloride; (**b**) Polytetrafluoroethylene; (**c**) Polydimethylsiloxane; (**d**) Polymethylmethacrylate; (**e**) Polyaryletheretherketone.

#### *2.2. PROMETHEE (Preference Ranking Organization Method for Enrichment Evaluations) II*

Material selection for engineering design is a multicriteria decision-making model (MCDM) problem that requires consideration of several available materials and conflicting evaluation criteria. The materials screened to be under selection are designated as P1 (Polyvinylchloride), P2 (Polytetrafluoroethylene), P3 (Polydimethylsiloxane), P4 (Polymethylmethacrylate) and P5 (Polyaryletheretherketone), as seen in Table 2. In addition, the beneficial and non-beneficial attributes or criteria to perform well in the water-logged condition of water hydraulic valves under the impact of water cavitation were chosen.

Figure 3 shows the hierarchy of the Material Selection based on the non-beneficial and beneficial criteria. Among the non-beneficial criteria are water absorption or equilibrium in the water at 23 ◦C (A1), Hildebrand Solubility (A2), and Coefficient of friction (A3). The beneficial criteria are Contact angle (A4), Tensile strength (A5), Hardness Shore D (A6), Impact Strength (A7) and Chemical resistance (A8). The criterion must fit the requirement of the system to survive from the prone to hydrodynamic cavitation surge pressure and erosion wears due to massive pressurized fluid flow in the valve chamber. Corrosion of the valve's internal wall must be prevented to sustain the longer life of the valves and to resist the erosion notch. Water absorption or equilibrium in the water at 23 ◦C is based on the superhydrophobic surface. However, as a result, the low viscosity and poor lubrication properties of water lead to potential risks of higher wear and friction, especially in the proportional spool valves [6].

For some materials, hardness (macro and micro) is a good cavitation erosion resistance indicator [47]. In terms of hydrophobicity, water absorption is commonly inferred as a measure of a solute's relative tendency to prefer a non-aqueous environment rather than an aqueous one. Hydrophobicity plays an essential role in the biological and physicochemical behaviour of numerous types of organic compounds [48]. Simply, hydrophobicity composite material, remarkable mechanochemical robustness, stain repellency, oil-water separation. Figure 4 presents a stepwise procedure for implementing PROMETHEE II [49,50]. Based on the requirement of the valve for water hydraulics and considering the flow of most significant selected attributes, engineering and material fields are integrated by aggregating the significant weights of criteria and the ratings of alternatives. Table 6 applies the weight distribution of the selected attribute.

**Figure 3.** The hierarchy of the Material Selection.


**Table 2.** Polymer properties and introducing the 5-point scale.

Annotation: Apply the 5-point scale for Chemical Resistance attribute (A8) in Table 1. Poor = 1, Satisfactory = 2, Good = 3, Very Good = 4 and Excellent = 5.

**Step 1:** Determine the objective and the evaluation criteria or attributes. Also determine / Select the weights forȱȱȱdifferent attributes by relative comparison of attributes.

**Step 2**: Normalize the evaluation criteria to dissolve the difference in units For beneficial criteria, the properties that its small value is the prominent for consideration

$$R\_{\vec{ij}} = \frac{\{(\mathbf{X}\_{\vec{ij}}) \text{ - } \text{Min}(\mathbf{X}\_{\vec{ij}})\}}{\{\text{Max}(\mathbf{X}\_{\vec{ij}}) \text{ - } \text{Min}(\mathbf{X}\_{\vec{ij}})\}} \qquad \qquad (\text{i = 1, 2, ..., m; j = 1, 2, ...n)} \tag{1}$$

For nonȬbeneficial criteria, the properties that its large value is prominent to consideration

$$R\_{\vec{ij}} = \frac{[\text{Max}(\text{X}\_{\vec{ij}}) \cdot (\text{X}\_{\vec{ij}})]}{[\text{Max}(\text{X}\_{\vec{ij}}) \cdot \text{Min}(\text{X}\_{\vec{ij}})]} \qquad \qquad (i = 1, 2, \dots, m; \, j = 1, 2, \dots n) \tag{2}$$

Where *Xij* denote, the specific value of criteria, *Max(Xij)* denote maximum value of criteria, *Min(Xij)* denote the minimum vale of the criteria

**Step 3:** Calculate the evaluative differences of ith alternative with respect to other alternatives ȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱ *j jj d (a,b) = g (a)Ȭ g (b)* ȱȱ ȱ ȱ ȱ ȱ ȱ (3)

Where *<sup>j</sup> d (a,b)* denotes the difference between the evaluations of *a* and *b* on each criterion

**Step 4:** Calculate the evaluative differences of *ith* alternative with respect to other alternatives *P ab F d ab j k j jj* ( , ) [ ( , )]ȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱȱ( 1, 2,.. ) ȱȱ ȱ ȱ ȱ (4)

where *Pj(a, b)* denotes the reference of alternative *a* with regard to alternative *b* against the *jth* criterion, as a function of *dj(a, b)*. *Fj* is a preference function, which translates the difference between the evaluations of alternatives *a* and *b* on the *jth* criterion into *a* preference degree ranging from 0 to 1.

**Step 5:** Calculate the overall or global preference index

$$
\pi(a, b) = \sum\_{\vec{j}=1}^{M} P\_{\vec{j}}(a, b) w\_{\vec{j}} \tag{5}
$$

where, *wj* is the weight of the *jth* criterion, and *Δ(a, b)* of *a* over *b* is defined as the weighted sum of *P(a, b)* for each criterion.

**Step 6:** Calculated the positive and negative outranking flows

$$
\boldsymbol{\phi}^{+}(\boldsymbol{a}) - \frac{1}{m \cdot \mathbf{1}} \sum\_{X \in \mathcal{A}} \pi(\boldsymbol{a}, \mathbf{X}) \tag{6}
$$

$$
\boldsymbol{\phi}^{-}(\boldsymbol{a}) - \frac{1}{m \cdot \mathbf{1}} \sum\_{X \in \mathcal{A}} \pi(\mathbf{x}, \mathbf{a}) \tag{7}
$$

Where I *+(a)* ȱȱand I *Ȭ (a)* denote the positive outranking flow and negative outranking flow for each alternative respectively.


**Figure 4.** Stepwise complete ranking PROMETHEE II.

#### *2.3. Molecular Dynamics Simulation to Predict the Substrate and Suitable Coating Plane*

MD simulation is a modeling technique based on physics that has been commonly used to represent different material systems as it can provide detailed information on variations and conformational changes in the structure and behavior of material systems at the molecular level [51,52]. Figure 5 below shows the Molecular dynamics algorithms. Here, the MD Simulation was conducted to predict a polymeric coating on the selected substrates (Cr2O3, Al2O3, and Ti2O3) along the planes (100), (010) and (001). The simulation computation was created using the material science simulation software developed by Accelrys Software Inc. in San Diego, CA, USA, named Material Studio ® Software, after several phases of model design and calculation. The tasks were to create engineered Cr2O3, Al2O3, and Ti2O3, slabs with an equal thickness of 13.233 <sup>×</sup> 10−<sup>10</sup> m each in a cleaved plane along (100), (010) and (001) each. Using a smarter algorithm with a maximum iteration of 100,000 steps, the surfaces were relaxed at minimum energy using model typing for force-field and condensed-phase optimized molecular potentials (COMPASS) to ensure bonding between metal oxide (Cr2O3, Al2O3, and Ti2O3). The supercell was designed by growing the surface to the lattice vector of U × V = 6 × 6 and then constructing the slab without the vacuum in order to change the periodicity from 2D to 3D. The steps of making a slab are the same as all-metal oxides (Cr2O3, Al2O3, and Ti2O3) picked. The periodic layer of the amorphous polytetrafluoroethylene cell (PTFE) was constructed and the structure was geometrically refined from the atactic polymer used in this work. This was pre-constructed, manufactured from the 50-chain length repeat-unit tetrafluoroethylene with a number of configurations set to 1.0, the final design goal density agreed to 2 g/cm3. The next step was to build a layer with a vacuum of 70 <sup>×</sup> <sup>10</sup>−<sup>10</sup> m. Minimized polytetrafluoroethylene was deposited on the matched metal oxide slabs of Cr2O3, Al2O3, and Ti2O3.

A molecular dynamics simulation was conducted in each coating composite layer. The temperature control method of Andersen was used, ensembles of a constant number of particles, constant volume, and constant temperature (NVPT). Each arrangement of molecular dynamics (MD) simulation was done with an interval of 1 femtosecond (fs) and each frame of 5000. COMPASS forcefield was determined in all the arrangements. Simulation steps were set to 200,000 with a total dynamic time of 200 picoseconds (ps).

**Figure 5.** Molecular dynamics simulation algorithm.

#### **3. Results and Discussion**

PROMETHEE was used so solve the polymer selection problem based on the Equations (1)–(8) as shown in the Tables 3–8. The PROMETHEE analysis was done by following all the essential steps of the analysis. Finally, the PTFE was the best polymer based on the selected significant properties for coating in the valve in water hydraulics. Table 9 shows the net flow of polymer alternatives. The PTFE has a higher value of 0.5932052, as seen in Figure 6 of the PROMETHEE I-II partial and complete ranking that the (Φ*net* values) of alternative P2 is the better coating polymer for water hydraulics valves to resist the cavitation erosion and corrosion. Moreover, PDMS, PVC, PEEK, and PMMA, as shown in Figure 7 of the ranking sequence, follow PTFE.

**Table 3.** Applying normalization to the evaluation criteria to obtain Rij (beneficial and non-beneficial) using Equations (1) and (2).


**Table 4.** Applying the calculation of the evaluative differences of ith alternative with respect to other alternatives using Equation (3).



**Table 5.** Application of the preference function using Equation (4).

**Table 6.** Applying the calculation of the overall or global preference index, using Equation (5).



**Table 7.** Applying the calculation of the positive and negative outranking flows, using Equations (6) and (7).

**Table 8.** Applying the Equation (8) to calculate the net outranking flow and the complete ranking.


**Table 9.** The value of the net flow of polymer alternatives.


Water hydraulics valves are subjected to cavitation erosion, wear, and corrosion due to a high bubbles implosion impact; selected properties of PTFE fit well with the coating requirements. Furthermore, the molecular dynamics results show that the PTFE coated with Al2O3 in the cleaved plane of (010) gives more considerable binding energy compared to other composites in all the cleaved planes. This may be due to difference in the chemical reactivity functions within the frameworks of the density functional theory (DFT), the high rate of transfer of electrons to or from the reactants in the plane (010) of PTFE/Al2O3, or due to the change in the geometric structure of the molecules (equivalent in external potential). In addition, the chemisorption takes place due to the strong electrostatic (ionic) bond.

These results show the sign of promising the strength of the composite over the effect of the impairment of water hydraulics valves due to the sudden surge pressure caused by the imploding of cavitation bubbles. Figure 8 below shows the pictorial presentation and data obtained after simulation.

**Figure 6.** PROMETHEE I-II partial and complete ranking.

**Figure 8.** Molecular dynamics Polytetrafluoroethylene coating on Metal Oxides (Ti2O3, Al2O3, and Cr2O3).

#### **4. Conclusions**

The selection and screening of polymer coating on metal oxides were studied for the design of liquid hydraulic valves. For this reason, PROMETHEE's decision-making methods were used. The results show the most suitable polymer materials being simulated. The improvement in the cavitation erosion resistance can mainly be linked to the increase in hardness and the elastic response of the PTFE/Al2O3 enhancement. Allison and Tong, in [53], using the density functional theory, depicted that binding energy is proportional to hardness. Moreover, simulation of molecular dynamics found the best substrate and cleaved plane in further studies should be considered by the characterization of a different PTFE/Al2O3 composition ratio.

**Author Contributions:** M.K.M. carried out the critical study and did all the comprehensive procedure and wrote the paper; H.X. provided instructions for this research and checked the findings of the analysis; F.S. and H.W. reviewed, revised the manuscript and looked for the literature; G.D.M. involved during the polymer attribute weighing. All authors have read and agreed to the published version of the manuscript.

**Funding:** The Natural Science Foundation of China supported this work under Grant 5187511, Natural Science Foundation of the Heilongjiang Province of China under Grant F2016003, 'Jinshan Talent' Zhenjiang Manufacture 2025 Leading Talent Project, "Jiangyan Planning" Project in Yangzhong City.

**Conflicts of Interest:** No conflict of interest declared by the authors.

#### **References**


© 2020 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* **Mechanically Robust and Repairable Superhydrophobic Zinc Coating via a Fast and Facile Method for Corrosion Resisting**

#### **Junfei Ou \*, Wenhui Zhu, Chan Xie and Mingshan Xue \***

School of Materials Science and Engineering, Nanchang Hangkong University, Nanchang 330063, China; 70299@nchu.edu.cn (W.Z.); 70362@nchu.edu.cn (C.X.)

**\*** Correspondence: oujunfei\_1982@163.com (J.O.); xuems04@mails.ucas.ac.cn (M.X.)

Received: 6 May 2019; Accepted: 28 May 2019; Published: 31 May 2019

**Abstract:** Zinc coatings and superhydrophobic surfaces have their own characteristics in terms of metal corrosion resistance. Herein, we have prepared a robust and repairable superhydrophobic zinc coating (SZC) based on a widely commercially available cold galvanized paint via a fast (within 10 min) and facile process for corrosion resistance. Specifically, the cold galvanized paint was sprayed onto the iron substrate, followed by acetic acid (HAc) etching and stearic acid (STA) hydrophobizing. The as-obtained sample was coded as Fe-Zn-HAc-STA and possessed an apparent contact angle of 168.4 ± 1.5◦ as well as a sliding angle of 3.5 ± 1.2◦. The Fe-Zn-HAc-STA sample was mechanically durable and easily repairable. After being ultrasonicated in ethanol for 100 min, the superhydrophobicity was still retained. The Fe-Zn-HAc-STA sample lost its superhydrophobicity after being abraded against sandpaper with a load of 100 g and regained its superhydrophobicity after HAc etching and subsequent STA hydrophobizing. The corrosion resistance of the SZC was investigated by immersing the Fe-Zn-HAc-STA sample into the static or dynamic aqueous solution of NaCl (3.5 wt.%) and the lasting life of the entrapped underwater air layer (EUAL) was roughly determined by the turning point at the variation curve of surface wettability against immersion time. The lasting life of the EUAL iwas 8 to 10 days for the SZC in the static NaCl solution and it decreased sharply to 12 h in a dynamic one with the flow rate of 2 and 4 m/s. This suggests that the superhydrophobic surface provided extra corrosion protection of 8 to 10 days or 12 h to the zinc coating. We hope that the SZC may find its practical application due to the facile and fast fabrication procedure, the good mechanical durability, the easy repairability, and the good corrosion protection.

**Keywords:** superhydrophobic; zinc-rich coating; cold galvanized coating; durability

#### **1. Introduction**

Iron and steel are the most common metals used in industry. However, the corrosion and damage of iron are serious; therefore, the anti-corrosion is an important issue in the iron industry [1]. People are using conversion coatings, such as chromate conversion coating and phosphate conversion coating, to improve its corrosion resistance. These conversion coatings have found wide practical application, however, also possess certain inherent disadvantages such as the carcinogenicity caused by the hexavalent chromium [2] and the high energy costs caused by the high treating temperature [3].

The cold zinc-spraying coating is an environment-friendly alternative with the standard electrode potential of −0.76 V, which is more active than iron (−0.44 V) [4]. When the coating is attacked, the zinc powder is corroded as the anode first, and the base iron is protected as the cathode, so the corrosion rate of iron can be slowed down significantly.

Superhydrophobic surface (SHS) with an apparent contact angle above 150◦ and a sliding angle below 10◦ is another newly-developed corrosion resistance strategy [5–15]. Once the superhydrophobic sample is immersed into water, a layer of air will be entrapped at the solid/liquid interface, which hinders the penetration of corrosive species (such as the Cl− ions dissolved in water) to reach the metallic substrate and improve the corrosion resistance greatly [5].

To combine the sacrificing anode effect of zinc coating with the air barrier effect of SHS, Zhang et al. [16] and Brassard et al. [17] have fabricated superhydrophobic zinc coatings (SZC) via electrochemical zinc deposition and subsequent surface passivation by polypropylene or silicone polymer. Such SZC samples improve the anti-corrosion performance of iron or steel significantly. However, the zinc deposits are obtained via an electrochemical deposition, which makes the fabrication of the SZC sample instrument-dependent.

Herein, we aimed to fabricate an SZC sample via a more facile method. The metallic zinc layer was obtained by spraying the cold galvanized paint onto the iron substrate, followed by acetic acid etching and subsequent stearic acid hydrophobizing. The procedure was time-saving (within 10 min) and easy to perform. The mechanical durability was assessed by ultrasonication in ethanol and abrasion against sandpaper. The corrosion resistance was investigated by immersing the Fe-Zn-HAc-STA sample into the static or dynamic aqueous solution of NaCl (3.5 wt.%), and the variation in surface wettability, surface morphology, and surface composition was monitored.

#### **2. Experimental**

#### *2.1. Sample Fabrication*

The cold galvanized coating (Shanghai Roval Zinc-rich Coating Co, Ltd., Shanghai, China) was shaken evenly and sprayed directly onto a clean iron substrate (2.5 mm × 2.5 mm × 1 mm). Then, the sample was dried at room temperature for 12 h to obtain a uniform cold galvanized coating. The sample was immersed into the aqueous solution of acetic acid (HAc, 18 wt.%) followed by rinsing with ultrapure water and blown dry with nitrogen. Finally, the sample was immersed into the ethanol solution of stearic acid (STA, 0.01 M, chemically pure, Shantou Xilong Chemical Co. Ltd., Shantou, China) for 30 s and blown dry with hot air. To obtain the optimum superhydrophobicity, the immersion in STA was repeated twice. For convenience, the sample after spraying, etching, and hydrophobizing was coded as Fe-Zn, Fe-Zn-HAc, and Fe-Zn-HAc-STA, respectively.

#### *2.2. Characterization*

The surface morphology of the sample was characterized by field emission scanning electron microscopy (FE-SEM, Nova, Nano-SEM 450, FEI, Hillsboro, OR, USA; at 1 kV) with attached energy dispersive spectroscopy (EDX, FEI Quanta 200, USA; at 20 kV) under a vacuum environment. The chemical compositions and valence states of the samples were measured using X-ray photoelectron spectroscopy (XPS, Physical Electronics, Chanhassen, MN, USA, PHI-5702) with the Al Kα X-ray source (hν = 1486.6 eV) at a base pressure of 10−<sup>7</sup> Pa. The surveyed (enlarged) spectra were recorded with a pass energy of 20 (5) eV and an energy step of 0.5 (0.05) eV. The binding energy of adventitious carbon (C1s: 284.8 eV) was used as a basic reference. The crystal orientation of the samples was characterized by X-ray diffraction (XRD, Bruker D8 ADVANCE, Karlsruhe, Germany) using CuKα radiation from 10 to 80◦ (2 theta) at 4 ◦/min. The powders scraped from the surface of the samples were recorded by Fourier Transform Infrared spectrometer (FTIR Nicolet IS10, Thermo Scientific, Waltham, MA, USA) by the KBr pellet methods with air as the reference. A resolution of 4 cm−<sup>1</sup> and 32 scans were chosen. The surface wettability was measured by a contact angle meter (CAM, DSA 20, Krüss, Hamburg, Germany) with a computer-controlled liquid dispensing system and a motorized tilting stage at ambient temperature (20 ◦C). For apparent contact-angle measurements, a 10 μL water droplet was used with the Laplace-Young fitting. Advancing and receding angles were measured by the addition and subtraction of 5 μL of water to/from the 10 μL droplets sitting on the surfaces. For the sliding angle measurement, the water droplet (10 μL) was dripped onto the sample, which was then rotated until the water drop slid away. The critical angle where a water droplet began to slide down

the inclined plate was measured to be the sliding angle. Each reported datum was an average of at least three measurements on the surface.

The adhesion of the coating to the iron substrate was assessed according to the standard of GB9286-98 by a crosscut tester (QFH-A, Dongguan Sanhe Instrument Manufacturing Factory, Dongguan, China). After the crosscut pattern at 90◦ was formed on the coating, the sample was brushed lightly with a soft brush to remove excess debris from the surface. A transparent tape was applied to the cut surface and rubbed with a rubber to ensure good contact with the coating and then removed after 90 s. Samples were evaluated under a magnifying glass (7×) and a bright light, and rated according to the American Society for Testing Materials (ASTM) rating scheme.

The variation of the surface wettability of the sample against ultrasonication (40 kHz, 100 W) in ethanol or against abrasion under a load of 100 g in contact with sandpaper (600 grits) was used to test the mechanical durability. The corrosion resistance was measured by immersing the sample into the static or dynamic aqueous solution of NaCl (3.5 wt.%) and the variation of surface morphology, chemical composition, and surface wettability was measured by FE-SEM, EDX, and CAM, respectively.

#### **3. Results and Discussion**

#### *3.1. Wettability, Surface Morphology, and Surface Chemistry*

The surface wettability of the Fe-Zn-HAc-STA sample varies against the etching time in the aqueous solution of HAc. The apparent contact angle or sliding angle levels off after rising or falling steeply for the first 6 min to a value of 161.8 ± 3.0◦ or 13.0 ± 1.5◦, respectively (Figure 1a). As shown in Figure 1(b1), the so-sprayed cold galvanized paint was composed of smooth zinc micro-particles (MPs). After for 6 min, the MPs became rough with nano-wrinkles (Figure 1(b2)). As the etching time increased to 12 min, the nano-wrinkles became much denser (Figure 1(b3)) and the superhydrophobicity was optimum (apparent contact angle of 168.4 ± 1.5◦ and sliding angle of 3.5 ± 1.2◦). The formation mechanism for these nano-wrinkles was due to the un-uniform etching between HAc and zinc as observed by Qian et al. [18]. It is explained that there are numerous dislocation defects in metals, which possess relatively higher energy and are prone to be attacked by chemical etchants such as HAc. Therefore, the etching rate in these dislocation sites was quicker and consequently, micro/nano-hierarchical surface structures were generated (Figure 1c).

**Figure 1.** Variation of the surface wettability for the Fe-Zn-HAc-STA sample against immersion time in the aqueous solution of HAc (**a**). Scanning electron microscopic images of the Fe-Zn sample (**b1**), the Fe-Zn-HAc with an etching time of 6 min (**b2**) and 12 min (**b3**). Representative for zinc particles in images (**b1**–**b3**) and the etching reaction (**c**). Compression of a water droplet on the Fe-Zn-HAc-STA sample with the etching time of 6 min in HAc by a Fe-Zn sample (**d1**–**d3**).

The wetting state of the Fe-Zn-HAc-STA sample with an etching time in HAc for 6 min was further assessed by measuring the contact angle hysteresis (CAH), which was approximately 4◦ and equivalent to the difference between the advancing angle (approximately 169◦) and the receding angle (approximately 165◦). The Fe-Zn sample possessed an advancing angle of approximately 143◦ and a receding angle of about approximately 35◦. A water droplet on the Fe-Zn-HAc-STA sample was compressed by the Fe-Zn sample and the typical images are shown in Figure 1d. After compression, the contact angle decreased slightly from 168.5◦ to 163.3◦; after relaxation, the water droplet adhered to the upper surface and there was no residual on the lower superhydrophobic surface. These analyses suggest that the water droplet on the Fe-Zn-HAc-STA sample with an etching time in HAc of 6 min was in the Cassie state. The hierarchical surface roughness was helpful to resist the destabilization from Cassie wetting state to Wenzel state [15].

The XRD pattern (Figure 2a) for the Fe-Zn sample shows the strongest diffraction peak at 2θ of 43.25◦, and weak ones at 36.34◦, 39.04◦, 54.41◦, 70.21◦, 70.72◦, and 77.19◦. These peaks are consistent with the standard PDF card of zinc (JCPDS 04-0831). Therefore, the main component of the Fe-Zn sample was zinc powder. Figure 2b shows the FTIR spectra of the Fe-Zn sample and the Fe-Zn-HAc-STA sample. The peaks at 1536 cm−<sup>1</sup> and 1455 cm−<sup>1</sup> in curve (i) for the Fe-Zn sample were attributed to the skeleton vibration of the benzene ring [19]. This was most likely attributed to the benzene ring in the bisphenol A epoxy resin, which was a routine ingredient for the cold galvanized paint to enhance the adhesion to the substrate. The peaks at 2919 cm−<sup>1</sup> and 2851 cm−<sup>1</sup> in curve (ii) were attributed to the C–H in epoxy resin [19], which became much stronger for the Fe-Zn-HAc-STA sample (curve ii). These stronger peaks were due to the anchored STA molecules abundant with C–H bonds. Compared with curve i for the Fe-Zn sample, two peaks emerged in curve ii for the Fe-Zn-HAc-STA sample, viz., a weak peak at 1752 cm−<sup>1</sup> attributed to the carboxyl groups (COOH) [20] and a strong peak at 1400 cm−<sup>1</sup> attributed to the deprotonated carboxyl groups (COO−) [17]. The stronger signal at 1400 cm−<sup>1</sup> suggests that most of the carboxyl groups are deprotonated to facilitate the anchoring of STA to the Zn MPs via coordination bonding (Figure 2(c1)) or ion bonding (Figure 2(c2)) [21].

**Figure 2.** X-ray diffraction pattern of the Fe-Zn sample (**a**). Fourier transform infrared spectra of the Fe-Zn sample (curve i) and the Fe-Zn-HAc-STA sample (curve ii) (**b**). A possible mechanism for stearic acid anchored to the oxidized zinc: coordination (**c1**) or ionic (**c2**) bonding.

Figure 3 shows the XP spectra of the Fe-Zn-HAc-STA sample. In the survey spectrum (Figure 3a), peaks were attributed to the elements of oxygen, carbon, and zinc emerge. The Zn 2p peaks were located at 1021.9 eV and 1044.8 eV, which were higher than the binding energy of Zn (0) and ascribed to the Zn (II) species [22]. This suggests that, due to its high reactivity, the outermost surface of Zn MPs was oxidized. The as-formed oxide layer facilitated the anchoring of STA to the Fe-Zn-HAc sample just as shown in Figure 2c. The C 1s spectrum was deconvoluted to three components, viz., the C–H bonding at 284.8 eV, the C–O bonding at 286.2 eV, and the C=O bonding at 288.3 eV [23]. The deconvoluted peak at 531.8 eV for O 1s was due to the O=C group and the one at 530.6 eV was due to the O–C bond, both of which once again prove the presence of STA on the sample surface.

**Figure 3.** XP spectra for the Fe-Zn-HAc-STA sample: survey spectrum (**a**), Zn 2p spectrum (**b**), C 1s spectrum (**c**), and O 1s spectrum (**d**).

#### *3.2. Mechanical Durability and Repairability*

Ultrasonication has been widely used to evaluate the mechanical durability of the superhydrophobic samples [24–28]. Bubble implosions by quick forming and violent collapsing generate shock waves and exert mechanical forces on the surface. Herein, the Fe-Zn-HAc-STA sample was immersed in ethanol, and the variation of surface wettability was plotted (Figure 4a). After 120 min, the apparent contact angle decreased slightly from 168.4 ± 1.5◦ to 150.6 ± 1.1◦ and the sliding angle increased from 3.5 ± 1.2◦ to 13.9 ± 1.3◦. The most durable sample in the references was reported by Peng et al., which retained its superhydrophobicity after ultrasonication in ethanol for 14 h [21]. As reported by Liu et al. [25] and Huang et al. [26], the superhydrophobic samples can withstand ultrasonic treatment of 1 and 2 h, respectively. As reported by Liu et al. [27], after ultrasonication of 30 min, the apparent contact angle decreased from 167◦ to 141◦. As reported by He et al. [28], after ultrasonication of 10 min, the apparent contact angle decreased from approximately 153◦ to approximately 136◦. Comparing with these References [25–28], we can say that the Fe-Zn-HAc-STA sample was mechanically robust.

**Figure 4.** Variation of surface wettability for the Fe-Zn-HAc-STA sample against ultrasonication time in ethanol (**a**). The surface appearance after cross-cut testing (**b**).

The good mechanical durability was due to the strong interlayer interaction. To evaluate its adhesion strength to the iron substrate, the Fe-Zn-HAc-STA sample was crosscut and peeled off by Scotch tape according to the standardized coating adhesion test method (ASTM D 3359) [29]. As shown in Figure 4b, the edges of the cuts were completely smooth and none of the squares of the lattice were detached. Therefore, the Zn-HAc-STA coating adhered to the iron substrate in the level of 5B. The high adhesion strength was due to the epoxy resin, which was a routine ingredient in cold galvanized paint as confirmed by FTIR (Figure 2b). Moreover, as discussed earlier, the STA molecules were supposed to be anchored to the zinc coating firmly via coordination or ionic bonding (Figure 2c).

An abrading test was also performed to further evaluate the mechanical durability of the Fe-Zn-HAc-STA sample, which was subjected to a weight of 100 g while in contact with a 600 grit sandpaper at a length of 20 cm (Figure 5a). After 10 cycles, the micro/nanostructures on the Fe-Zn-HAc-STA sample were nearly destroyed (Figure 5(c1,c2)) and the apparent contact angle decreased to approximately 120◦. This suggests that the outermost hydrophobic STA layer and the rough structures on the Fe-Zn-HAc-STA sample were abraded away and consequently the sample becomes hydrophobic just as the Fe-Zn sample. However, this does not mean that the coating will fail. After re-impregnation in an aqueous solution of HAc and ethanol solution of STA, the micro/nanostructures appeared on the surface (Figure 5(c3,c4)) and the coating regained the superhydrophobicity (Figure 5b). The abrading not only caused the change in the surface morphology but also decreased the thickness. As shown in Figure 5d, the thickness of the Fe-Zn-HAc-STA sample was reduced by 49 μm after 100 cycles.

**Figure 5.** Schematic view of the setup for wearing test (**a**). Variation of apparent contact angle of the Fe-Zn-HAc-STA sample against wearing cycles (**b**). Scanning electron images for the Fe-Zn-HAc-STA sample after 10 cycles of wearing (**c1**,**c2**) and re-impregnation in acetic acid solution and ethanol solution of stearic acid (**c3**,**c4**). Scanning electron images for the cross-section of the Fe-Zn-HAc-STA sample before (**d1**) and after (**d2**) wearing for 100 cycles.

#### *3.3. Corrosion Resistance*

Figure 6a shows the variation of the surface wettability for the Fe-Zn-HAc-STA sample (curves i, iii) and the Fe-Zn sample (curve ii) in the aqueous solution of NaCl for 20 days. After 6 days, the apparent contact angle of the Fe-Zn sample decreased sharply to approximately 0◦; however, the Fe-Zn-HAc-STA sample still possessed good hydrophobicity (apparent contact angle of 150.7 ± 0.9◦ and sliding angle of 16.0 ± 0.3◦). The much slower deterioration rate of hydrophobicity for the Fe-Zn-HAc-STA sample was due to the entrapped underwater air layer (EUAL) at the solid/liquid interface, which was reflected by the "silver mirror" as shown in Figure 6(b1). In other words, once the superhydrophobic Fe-Zn-HAc-STA samplewas immersed into the aqueous solution of NaCl, most of the solid surface (up to 90%) was in contact with the EUAL rather than water [25]. The EUAL blocked the diffusion of corrosive ions to the solid substrate [5] and consequently, the surface hydrophobicity of the Fe-Zn-HAc-STA sample deteriorated much more slowly as compared with the hydrophobic Fe-Zn sample. We can speculate the lasting life of the EUAL from the curves in Figure 6a. Specifically, the apparent contact angle (curve i) and the sliding angle (curve iii) for the Fe-Zn-HAc-STA sample varied very gently before 8 days and 10 days, respectively; thereafter, the variation was much faster. Moreover, the underwater "silver mirror" after 8 days as shown in Figure 6(b2) was not so obvious and became discontinuous. These data hint that the lasting life of EUAL was about 8 to 10 days. After that, statistically speaking, the STA layer rather than the EUAL was in contact with water. The surface micro-morphology of the Fe-Zn sample after immersion is shown in Figure 6c. Some area was covered by the corrosion products. In contrast, the micro/nano- hierarchical surface structures (Figure 6d) for the Fe-Zn-HAc-STA sample varied a little but were still observable, which is the morphology requirement for the superhydrophobicity.

**Figure 6.** Variation of surface wettability for the Fe-Zn-HAc-STA sample (i, iii) and the Fe-Zn sample (ii) against the immersion time in the aqueous solution of NaCl (3.5 wt.%) (**a**). Digital pictures to show the underwater mirror for the Fe-Zn-HAc-STA sample just immersed into the NaCl solution (**b1**) and after 8 days (**b2**). Surface micro-morphology for the Fe-Zn sample (**c1**,**c2**) and the Fe-Zn-HAc-STA sample (**d1**,**d2**) after immersed in the aqueous solution of NaCl for 8 days.

The element content of the samples before and after immersion was measured by EDX and is summarized in Table 1. As confirmed by FTIR, XPS, and XRD earlier, the Fe-Zn sample was mainly composed of Zn MPs and epoxy resin; for the sample of Fe-Zn-HAc-STA, STA molecules were attached. Therefore, the main elements for these two samples were carbon, oxygen, and zinc. After immersion, the signal attributed to the element of chlorine emerged, which was attributed to the adsorbed NaCl or the chlorine contained corrosion products. For the Fe-Zn sample, the element content of carbon decreased sharply from 61.83 ± 6.23% to 17.48 ± 4.45%, and the oxygen increased from 6.14 ± 0.75% to 53.77 ± 7.42%. This suggests that the carbon-rich organic molecules were detached severely from the Fe-Zn sample and lots of oxygen-rich corrosion products such as Zn(OH)2 were formed via the following electrochemical reactions.

$$\text{anodereaction}: \text{Zn} - \text{2e} \to \text{Zn}^{2+} \tag{1}$$

cathodereaction : O2 + 2H2O + 4e → 4OH<sup>−</sup> (2)

$$\text{Overall reaction}: 2\text{Zn} + \text{O}\_2 + 2\text{H}\_2\text{O} \rightarrow 2\text{Zn(OH)}\_2 \downarrow \tag{3}$$

However, for the Fe-Zn-HAc-STA sample, the content change of carbon and oxygen was not so obvious. This is due to the barrier effect of the EUAL with a lasting life of about 8 to 10 days as discussed earlier.


**Table 1.** Atomic element content of the Fe-Zn sample and the Fe-Zn-HAc-STA sample before and after immersing in the aqueous solution of NaCl for 8 days.

We monitored the variation of the surface wettability in the dynamic aqueous solution of NaCl for 48 h (Figure 7). The flow rate affected the variation of the surface wettability significantly. Curves in Figure 7a (static solution) and Figure 7b (1 m/s) vary smoothly. However, as the flow rate increased to 2 or 4 m/s, similar turning points as shown in Figure 6a at 8 to 10 days appeared after a short immersion of approximately 12 h. This suggests that the lasting life of the EUAL was shortened significantly under dynamic solution with the high flow rate of 2 to 4 m/s.

**Figure 7.** Variation of the surface wettability for the Fe-Zn-HAc-STA sample immersed in 3.5 wt.% NaCl solution for 48 h at a flow rate of 0 m/s (**a**), 1 m/s (**b**), 2 m/s (**c**) and 4 m/s (**d**).

Figure 8 shows the micro-morphological images for the Fe-Zn-HAc-STA sample after immersed in the aqueous solution of NaCl for 48 h. The sample the underwent static corrosion still possessed a hierarchical surface microstructure (Figure 8a). As the flow rate increased to 1 or 2 m/s, some area (the circle in Figure 8(b1)) of the surface became compact and the MPs were covered by the corrosion products (Figure 8(b2,c2)). At 4 m/s, the shape of Zn MPs was difficult to recognize.

**Figure 8.** FE-SEM images for the Fe-Zn-HAc-STA samples in the aqueous solution of NaCl (3.5 wt.%) for 48 h at a flow rate of 0 m/s (**a1**,**a2**), 1 m/s (**b1**,**b2**), 2 m/s (**c1**,**c2**) and 4 m/s (**d1**,**d2**).

The element content of the Fe-Zn-HAc-STA sample after dynamic corrosion at different flow rates was measured by EDX (Table 2). The oxygen content increased from 8.46 ± 3.20% to 13.63 ± 3.89% at 1 m/s or to 56.47 ± 5.89% at 4 m/s after 48 h. The carbon content remained almost unchanged at 1 m/s and decreased sharply to 10.50 ± 2.14% at 4 m/s. As discussed earlier, the decreasing carbon content was due to the detachment of STA molecules and the increasing of oxygen content is due to the formation of corrosion products. These data suggest that the flow rate affects the corrosion of the superhydrophobic sample significantly. The reason is that the flowing liquid made it difficult to maintain the EUAL at the submerged superhydrophobic surface.



#### **4. Conclusions**

A superhydrophobic zinc coating coded as Fe-Zn-HAc-STA was prepared on the iron substrate via a facial multi-step procedure including (a) the spraying of the widely commercially available cold galvanized paint, (b) the etching by acetic acid (HAc), and (c) the hydrophobizing by stearic acid (STA). The Fe-Zn-HAc-STA sample possessed good mechanical durability and its superhydrophobicity remained after being ultrasonicated in ethanol for 100 min. After being abraded against the sandpaper for 10 cycles, the surface superhydrophobicity was lost due to the destruction of the hierarchical surface structures and the detachment of the hydrophobic STA molecules. However, after quick HAc etching and STA hydrophobizing, the sample regained the superhydrophobicity easily. The Fe-Zn-HAc-STA sample was immersed into the aqueous solution of NaCl to evaluate the corrosion resistance, and an entrapped underwater air layer (EUAL) was formed at the solid/liquid interface. The lasting life of EUAL in the static solution was 8 to 10 days and decreased sharply to 12 h as the flow rate of the corrosive solution increased to 2 or 4 m/s. We can say that the superhydrophobic surface provided extra corrosion protection of 8 to 10 days or 12 h to the zinc coating. We hope that this research may be helpful for the design of anti-corrosion superhydrophobic coatings due to the facile and fast fabrication procedure, the good mechanical durability, the easy repairability, and the good corrosion protection.

**Author Contributions:** For research articles with several authors, a short paragraph specifying their individual contributions must be provided. The following statements should be used "conceptualization, J.O. and M.X.; methodology, W.Z.; validation, W.Z. and C.X.; formal analysis, W.Z.; investigation, W.Z.; resources, C.X.; writing—original draft preparation, J.O.; writing—review and editing, J.O.; visualization, C.X.; supervision, M.X.; project administration, J.O.; funding acquisition, J.O.

**Funding:** This research was funded by the National Natural Science Foundation of China (Grant No. 51563018).

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

#### **References**


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