**Lotus-Inspired Multiscale Superhydrophobic AA5083 Resisting Surface Contamination and Marine Corrosion Attack**

#### **Binbin Zhang 1,2,3,\*, Weichen Xu 1,2,3, Qingjun Zhu 1,2,3, Shuai Yuan 1,2,4 and Yantao Li 1,2,3,\***


Received: 22 April 2019; Accepted: 11 May 2019; Published: 15 May 2019

**Abstract:** The massive and long-term service of 5083 aluminum alloy (AA5083) is restricted by several shortcomings in marine and industrial environments, such as proneness to localized corrosion attack, surface contamination, etc. Herein, we report a facile and cost-effective strategy to transform intrinsic hydrophilicity into water-repellent superhydrophobicity, combining fluorine-free chemisorption of a hydrophobic agent with etching texture. Dual-scale hierarchical structure, surface height relief and surface chemical elements were studied by field emission scanning electron microscopy (FE-SEM), atomic force microscopy (AFM), energy dispersive X-ray spectroscopy (EDS) and X-ray photoelectron spectroscopy (XPS), successively. Detailed investigations of the wetting property, self-cleaning effect, NaCl-particle self-propelling, corrosion and long-term behavior of the consequent superhydrophobic AA5083 surface were carried out, demonstrating extremely low adhesivity and outstanding water-repellent, self-cleaning and corrosion-resisting performance with long-term stability. We believe that the low cost, scalable and fluorine-free transforming of metallic surface wettability into waterproof superhydrophobicity is a possible strategy towards anti-contamination and marine anti-corrosion.

**Keywords:** fluorine free; silanization; superhydrophobic; corrosion protection; self-cleaning

#### **1. Introduction**

Metals and their alloys are central engineering materials in numerous industrial fields. As a typical representative, 5083 aluminum alloys (AA5083), featuring a high strength-to-weight ratio and good weldability, are widely employed in military equipment, marine constructions, automobile manufacture and aerospace applications. However, proneness to localized corrosion attacks [1–4] in corrosive environments restricts the large-scale and long-term application of AA5083 materials. Corrosion attacks in aggressive environments can produce a premature failure of AA5083 structural materials, resulting in environmental disruption, enormous economic loss, as well as catastrophic safety accident. Thus, how to endow AA5083 materials with superior corrosion resistance is an extensively concerning issue.

In recent years, inspired from the unique water-repellent property of natural organisms [5–9], the artificial fabrication of bionic superhydrophobic surfaces attracted intensive attention of scientists and engineers owing to their multi-functional applications, such as self-cleaning [10,11], oil–water

separation [12,13], drag reduction [14,15], anti-icing/frosting [16,17], microdroplet transportation [18,19], water collection [20,21], marine anti-corrosion [22,23], anti-biofoulings [24,25], etc. It is believed and demonstrated that superhydrophobic surfaces could effectively reduce the solid/liquid interfacial contacts and provide a functional corrosion-resistant barrier. Thus, transforming surface wettability from intrinsic hydrophilicity to water-repellent superhydrophobicity is considered to be a possible strategy resisting marine corrosion attacks [26,27].

At present, much efforts [28–31] have been devoted to exploring the fabrication and investigating the consequent corrosion-resistant behavior of a superhydrophobic surface on aluminum/aluminum alloy substrates. For instance, Boinovich et al. [32,33] reported a combination of nanosecond laser texturing and fluorinated hydrophobic agent chemisorption, achieving a superhydrophobic aluminum-magnesium alloy with an extremely low corrosion current. Wang et al. [34] used a hydrothermal in situ growth method to fabricate superhydrophobic Mg-Al-layered double hydroxide films on 6061 aluminum alloy substrates, presenting a highly improved corrosion resistance. In our previous reports, we developed an ammonia etching approach [35] and an anodization method [36] followed by 1H,1H,2H,2H-Perfluorodecyltriethoxysilane chemisorption to achieve superhydrophobic surfaces on aluminum/aluminum alloy substrates with multiscale hierarchical topography, and greatly enhanced corrosion inhibition performance.

However, most of the fabrication methods above are hindered by several shortcomings, such as being time-consuming, high-cost and fluorine reagents employment, etc., restricting their large area usage and severely threatening ecological environments and human security. Until now, only limited attempts have been achieved to develop scalable and fluorine-free superhydrophobic aluminum surfaces for marine corrosion protection. So, these issues push us to explore a cost effective, facile and non-fluorinated approach, transforming intrinsically hydrophilic AA5083 with a Wenzel contact [37,38] to superhydrophobic AA5083 with a Cassie–Baxter contact [39,40].

Herein, we report a facile, cost-effective and non-fluorinated fabrication strategy to prepare a superhydrophobic surface on an AA5083 substrate by combining etching texture and a hexadecyltrimethoxysilane (HDTMS) hydrophobic molecules assembly. Detailed studies about the surface morphology and chemical composition were carried out successively. In addition, wetting property, self-cleaning ability, corrosion-resisting behavior and long-term stability were investigated to display the typical characteristics and promising functional applications.

#### **2. Experimental Section**

#### *2.1. Materials and Reagents*

A 5083 aluminum alloy (AA5083) plate with a 0.3 mm thickness was obtained from Dongguan Wanxing Metal Co., Ltd. (Dongguan, China), and tailored into 25 mm × 20 mm specimens. The main composition of the pristine AA5083 substrate was 4.0–4.9% Mg, 0.4–1.0% Mn, 0.25% Zn, 0.4% Si, 0.15% Ti, 0.1% Cu, 0.05–0.25% Cr, 0.1–0.4% Fe and the balance was Al. Sodium hydroxide (NaOH), sodium chloride (NaCl), methylene blue trihydrate (C16H18ClN3S·3H2O), ethanol absolute and graphite powder were purchased from Sinopharm Chemical Reagent Co., Ltd. (Beijing, China). Manganese monoxide (MnO) was bought from Aladdin Industrial Corporation. Hexadecyltrimethoxysilane (C19H42O3Si, HDTMS) was received from J&K Scientific Ltd. (Shanghai, China). All experimental reagents mentioned above were used as received without further purification.

#### *2.2. Preparation of Superhydrophobic AA5083*

The brief schematic illustration of the fabrication process is presented in Figure 1, involving an etching-texture process and HDTMS assembly. Firstly, the pristine AA5083 substrates were sanded through SiC sandpaper with different grades (400, 800, 1200, etc.) and cleaned by sonication in ethanol and deionized water for more than 5 min, respectively. The cleaned and pristine AA5083 specimens were rinsed with deionized water and dried under an air blower before etching. Subsequently, the pristine AA5083 samples were immersed in 15 g/L sodium hydroxide aqueous solution for 30 min to accomplish the etching texture. After deionized water cleaning, ethanol cleaning and an 80 ◦C oven drying treatment, the etching-textured AA5083 specimens were immersed in a 3 vol.% HDTMS/ethanol solution for 1 h to chemically assemble HDTMS molecules. After the modification process, the modified AA5083 specimens were heated at 120 ◦C in a drying oven for 20 min.

**Figure 1.** Schematic illustration of the fabrication process.

#### *2.3. Characterization*

The surface topography of pristine and as-fabricated superhydrophobic AA5083 surfaces were characterized by field emission scanning electron microscopy (FE-SEM, FEI Nova Nano SEM450, Hillsboro, OR, USA) and atomic force microscopy (AFM, Bruker Multimode 8, Karlsruhe, Germany). The AFM images were obtained under tapping mode. Energy dispersive X-ray spectroscopy (EDS, Oxford X-MaxN50, Hillsboro, OR, USA) and X-ray photoelectron spectroscopy (XPS, Thermo Scientific Escalab 250Xi, Massachusetts, MA, USA) were performed to determine the existence of the key elements upon the specimens. The XPS measurements were performed using a monochromated Al K<sup>α</sup> irradiation and the chamber pressure was 3 <sup>×</sup> 10−<sup>8</sup> Torr during the test. The binding energy of adventitious carbon C1s (284.8 eV) was used as a basic reference. A Dataphysics OCA25 instrument (Stuttgart, Germany) was used to measure the static water contact angles and sliding angles of the as-prepared superhydrophobic AA5083 samples. For each measurement, at least three different positions were performed to obtain average data.

#### *2.4. Electrochemical Test*

The electrochemical tests were all carried out in a 3.5 wt.% NaCl aqueous solution through an Ametek Parstat 4000+ electrochemical workstation. A typical three-electrode measure system, including counter electrode (Pt sheet), reference electrode (saturated silver/silver chloride) and working electrode (pristine/superhydrophobic AA5083), was employed to proceed with the electrochemical tests. Prior to the test, the working electrode was exposed to a 3.5 wt.% NaCl aqueous solution for more than 1 h, achieving a stable measuring system. Electrochemical impedance spectroscopy (EIS) was measured under OCP (open circuit potential) at a frequency range of 100 kHz–10 MHz. *ZsimpWin* software was utilized to analyze and fit the EIS data for anti-corrosion evaluation.

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

#### *3.1. Surface Morphology and Wettability Behavior*

The surface topography of pristine AA5083 and superhydrophobic AA5083 surfaces were characterized by FE-SEM. Figure 2 displays the SEM images of the pristine AA5083 and superhydrophobic AA5083 surface. As for pristine AA5083 shown in Figure 2a, some micro-grooves/scratches can be seen, which is attributed to the pre-treatment of the pristine AA5083 specimen. The surface of the pristine AA5083 is relatively smooth. For the etching-textured superhydrophobic AA5083 surface, displayed in Figure 2b, some micro-textured rough structure can be found. Figure 2c shows the higher magnification image of Figure 2b and presents an obviously nano-scale petaloid surface architecture of the as-prepared superhydrophobic AA5083 sample. These micro-nano hierarchical structures contribute to the increase of surface roughness and provide sufficient structural clearance for the formation of the trapped air cushion between the solid surface and the water droplet, which is beneficial for the final Cassie–Baxter contact state. Figure 2d shows the static water contact angle of pristine AA5083 and superhydrophobic AA5083 surface. The contact angle of the pristine AA5083 surface is about 81.6◦ ± 1◦. After the etching texture and HDTMS assembly, however, the contact angle of the as-prepared superhydrophobic AA5083 surface is about 156.3◦ ± 1◦ with a sliding angle lower than 1◦. The etching-textured, dual-scale surface rough structure and the low surface energy of the HDTMS molecules endowed the AA5083 substrate with a high static water contact angle and a low sliding angle.

**Figure 2.** FE-SEM images (**a**–**c**) and static contact angles (**d**) of pristine 5083 aluminum alloy (AA5083) and as-fabricated superhydrophobic AA5083 surfaces.

As is well known, wetting behavior of a solid surface is mainly determined by roughness and surface energy. The microscale roughness and height relief of the pristine and as-fabricated superhydrophobic AA5083 surfaces were revealed by AFM, as shown in Figure 3. Figure 3a,b display the topographic fluctuation of pristine AA5083 and superhydrophobic AA5083 specimens. It was found that the height relief of the pristine AA5083 substrate was about 160 nm. On the contrary, the height relief of the as-prepared superhydrophobic AA5083 was approximately 970 nm, presenting a significant improvement. Figure 3c,d show the 3D AFM images of pristine AA5083 and superhydrophobic AA5083 surfaces. The surface roughness can be clearly observed and contrasted. Generally, *R*<sup>a</sup> (average roughness), *R*q (root mean square roughness) and *R*max (maximum roughness) were utilized to present surface roughness. In this case, the surface of the pristine AA5083 was relatively smooth with *R*a, *R*<sup>q</sup> and *R*max values (scanning area 10 μm × 10 μm) being 33.3 nm, 42.8 nm and 444 nm, respectively. While for the as-prepared superhydrophobic AA5083 surface, the *R*a, *R*q and *R*max values (scanning area 20 μm × 20 μm) were apparently increased to 107 nm, 142 nm and 1229 nm, respectively. As discussed above, the resultant superhydrophobic AA5083 surface features an obviously improved surface roughness. The etching-textured process contributes to this roughness enhancement, which is in favor of the air cushion formation and Cassie–Baxter contact of the water/solid/air interface.

**Figure 3.** Surface relief and 3D atomic force microscopy (AFM) images of pristine AA5083 (**a**,**c**) and superhydrophobic AA5083 (**b**,**d**).

#### *3.2. Low Surface Adhesivity*

The larger etching-textured roughness and lower HDTMS surface energy played key roles in the resultant water-repellent superhydrophobicity. Figure 4a displays the optical image of spherical water droplets, illustrating a waterproof property. Figure 4b,c show the surface response with a dynamic jet of water flow and faucet water impact. It can be clearly seen that the water flow jet and faucet water cannot remain on the superhydrophobic AA5083 surface. The water flow reflects, bounces and finally reflecting/rolling away easily from the specimen. Figure 4d–f presents the optical images of platform movement to contact and departing of the water droplet using the Dataphysics OCA25 instrument. With the gradually approaching, contacting and departing, the as-prepared superhydrophobic AA5083

surface could completely depart from the water droplet after tight contact, suggesting an extremely low adhesivity.

**Figure 4.** Optical images of the as-prepared superhydrophobic AA5083 surface (**a**) with spherical water droplets, (**b**) with a jet of water flow, (**c**) with faucet water impact, and (**d**–**f**) with platform movement to contact and departing of the water droplet.

The above wetting property and low adhesivity of the as-prepared superhydrophobic AA5083 can be explained by the combination action of the two-tier hierarchical rough structure and the HDTMS modification. Hierarchical petaloid rough structures benefit from the formation of the trapped air cushion, which significantly restrain the interfacial contact of the water–solid phase. Furthermore, the introduction of HDTMS molecules decrease the surface energy and further suppresses the penetration of water droplets into the surface structure. Thus, it can be concluded from the interaction between water and the superhydrophobic AA5083 surface that the water-repellent superhydrophobicity is attributed to the microscale roughness of the etching-textured surface and the low surface energy of HDTMS molecules.

#### *3.3. Chemical Composition*

The chemical composition of the pristine and HDTMS assembly superhydrophobic AA5083 surfaces were characterized by EDS (Figure 5a) and XPS (Figure 5b). The surface is rich in C, O, Mg, Al and Si elements as evidenced by the EDS spectrum, preliminarily verifying the assembly of the long carbon-chain tail on the as-prepared superhydrophobic AA5083 surface. As is well known, the XPS spectra present a surface chemical composition with a detecting depth of a few nanometers. The detailed elemental composition was further investigated through an XPS spectrum, as shown in Figure 5b. It is obvious that C, O, Si and Al elements were detected and demonstrated in the XPS spectrum of the as-prepared superhydrophobic AA5083 surface. Strong binding energy located at 284.7 eV, 532.4 eV, 102.3 eV and 74.6 eV were confirmed and ascribed to C 1s, O 1s, Si 2p and Al 2p, respectively, further confirming the existence of HDTMS species in the superhydrophobic AA5083 substrate. The C 1s, O 1s and Si 2p peaks in XPS spectra of the as-prepared superhydrophobic AA5083 sample were contributed to the HDTMS molecule assembly. This is in accordance with the EDS spectrum. It can be concluded, based on the analyses of surface topography, microscale roughness,

wetting property and chemical composition, that the AA5083 substrates were endowed with excellent water repellent superhydrophobicity.

**Figure 5.** Energy dispersive X-ray spectroscopy (EDS) and X-ray photoelectron spectroscopy (XPS) spectra of the fabricated superhydrophobic AA5083 sample. (**a**) EDS spectrum, and (**b**) XPS spectrum.

#### *3.4. Self-Cleaning Ability and NaCl Self-Propelling*

As for real-world applications, the self-cleaning ability of superhydrophobic material is an essential and promising function. In this work, MnO powder and graphite powder were successively applied as surface contaminations of the as-prepared superhydrophobic AA5083 samples. The superhydrophobic AA5083 specimen was firstly inclined at an angle lower than 10◦. Water droplets were subsequently dropped from above to evaluate the self-cleaning effect of the specimen. Figure 6a,b display the optical photos of the self-cleaning process. Given the excellent water repellence of the superhydrophobic AA5083 specimen, the water droplets can instantaneously roll off the sample surface, effectively picking up and taking away the MnO powder and graphite powder without difficulty. The water droplets rolled down the superhydrophobic AA5083 sample without moistening the solid surface, leaving several traces of rolling upon the surface. After washing, the contaminated AA5083 samples were totally clean and had no differences compared to the original uncontaminated specimen.

**Figure 6.** Self-cleaning ability with different surface contaminations: (**a**) MnO powder, (**b**) graphite powder, and (**c**) NaCl self-propelling property.

Figure 6c presents the NaCl self-propelling property of the superhydrophobic AA5083 surface. The NaCl particle was put on the surface of the sample. After the water droplet fell to the surface, the NaCl particle was fused with the water droplet. The outstanding water-repellence of the superhydrophobic AA5083 surface overcame the gravity of the NaCl particle, exhibiting a typical droplet/NaCl bouncing and deformation. An extremely low sliding angle propelled the droplet/NaCl coalition.The NaCl self-propelling process only takes 360 ms from the water droplet falling to the final sliding away. The self-cleaning and NaCl self-propelling ability mentioned above can be mainly ascribed to the extremely low surface energy, low adhesivity and water-repellence property of the surface. Therefore, it can be concluded that the fabricated superhydrophobic AA5083 substrate possesses an excellent self-cleaning ability resisting different surface contaminations.

#### *3.5. Marine Corrosion Protection*

In order to quantitatively characterize the corrosion behavior of the as-prepared superhydrophobic AA5083 film, we carried out and analyzed the electrochemical impedance spectroscopy (EIS) in an open circuit condition of a 3.5 wt.% NaCl aqueous corrosive solution. Figure 7a presents the EIS plots and fittings of pristine AA5083 and superhydrophobic AA5083 substrates. It can be seen from superhydrophobic AA5083 EIS plots that the diameter of the capacitive loop is much larger than that of the pristine AA5083 specimen. Figure 7b displays Bode plots of log |Z| vs. frequency and fittings of the pristine AA5083 and superhydrophobic AA5083 surfaces, presenting a three orders of magnitude higher impedance modulus value than that of the pristine AA5083 substrate. The anti-wetting and water-repellent property of the as-fabricated superhydrophobic AA5083 surface contributes to the remarkable improvement of impedance modulus.

**Figure 7.** (**a**) Electrochemical impedance spectroscopy (EIS) plots, and (**b**) Bode plots of log |Z| vs. frequency and fittings of the pristine AA5083 and superhydrophobic AA5083 surfaces.

Different equivalent circuits were employed to analyze the EIS results through *ZsimpWin* software, as shown in Figure 8. Figure 8a,b display the equivalent circuit of pristine AA5083 and superhydrophobic AA5083 specimens. For the convenience of EIS data analyzing, *R*s, *R*film, *R*oxide and *R*ct in Figure 8 represent solution resistance, superhydrophobic film resistance, oxide resistance and charge transfer resistance, respectively. *Q*film, *Q*oxide and *Q*dl represent the constant phase elements (CPE) modelling capacitance of the as-prepared superhydrophobic film, the oxide layer and the double-layer, respectively. Wherein, the impedance of the CPE could be defined as 1/*Y*0(*j*ω) <sup>n</sup> [41,42], where *Y*0, *j*, ω and n represent the modulus, imaginary number, angular frequency and the phase, respectively.

**Figure 8.** Equivalent circuit of (**a**) pristine AA5083 and (**b**) superhydrophobic AA5083.

The fitted electrochemical parameters are shown in Table 1. In General, *R*ct is utilized to calculate and evaluate the corrosion resisting property of the fabricated protective layers. From Table 1, it can be clearly seen that the *R*ct of the as-prepared superhydrophobic AA5083 specimen is 1.14 <sup>×</sup> 106 <sup>Ω</sup>·cm2, while the *R*ct of the pristine AA5083 sample is only 4.44 <sup>×</sup> 104 <sup>Ω</sup>·cm2. The *R*ct of superhydrophobic AA5083 is two orders of magnitude higher than that of pristine AA5083, demonstrating a remarkable enhanced corrosion-resisting performance. In addition, the *Q*dl of pristine AA5083 and as-prepared superhydrophobic AA5083 was 6.77 <sup>×</sup> 10−<sup>6</sup> <sup>Ω</sup>−1·cm−2·sn and 3.54 <sup>×</sup> 10−<sup>10</sup> <sup>Ω</sup>−1·cm−2·sn, respectively. The lower *Q*dl value and higher *R*ct value of the as-prepared superhydrophobic AA5083 substrate suggest that the charge transfer process of corrosive ions occurs with difficulty.


**Table 1.** The electrochemical parameters of simulated pristine AA5083 and superhydrophobic AA5083 surfaces in 3.5 wt.% NaCl aqueous solution.

In general, the *R*ct value is utilized to calculate the inhibition efficiency (η) of the protective film using the formula <sup>η</sup> <sup>=</sup> (*R*ct <sup>−</sup> *<sup>R</sup>*ct0)/*R*ct [43], in which *<sup>R</sup>*ct and *<sup>R</sup>*ct<sup>0</sup> represent the charge transfer resistance of the superhydrophobic AA5083 specimen and the pristine AA5083 specimen. It was calculated using the *R*ct and *R*ct<sup>0</sup> values discussed above that the inhibition efficiency in this case was approximately 96.11%, indicating an impressive performance resisting marine corrosion attack.

Figure 9 shows the anticorrosion mechanism of the as-fabricated superhydrophobic AA5083 surface. When the pristine AA5083 is exposed to environments containing aggressive corrosive ions, the natural oxide film breaks down at specific points leading to the formation of localized corrosion on AA5083 surface, viz. pitting corrosion or intergranular corrosion. On the contrary, the as-prepared superhydrophobic AA5083 surface could trap air within the micro-nano petaloid hierarchical structure, presenting a greatly decreasing fraction of the water/solid interface. The superior corrosion resistance is mainly attributed to the air cushion trapped in the petaloid rough structure, which suppresses the penetration and diffusion of aggressive corrosion species crossing the superhydrophobic protective film to the underlying AA5083 substrate. So, it can be implied that the as-fabricated superhydrophobic AA5083 specimen could largely improve the corrosion resistance of the substrate, displaying an outstanding protection ability towards marine corrosion attack.

**Figure 9.** Anticorrosion mechanism of the as-prepared superhydrophobic AA5083.

#### *3.6. Long-Term Stability*

For practical applications, it is significant to characterize the stability of the superhydrophobic AA5083 specimen subjected to air exposure and a corrosive solution immersion for prolonged period. In this case, we carried out air exposure and a 3.5 wt.% NaCl aqueous solution immersion experiments to evaluate the durability of the as-produced superhydrophobic AA5083 substrate, as shown in Figure 10a,b. After 90 days of air exposure and 12 days of a 3.5 wt.% NaCl immersion/contact of the as-fabricated superhydrophobic AA5083 surface, it can be found from the contact angle and sliding angle variation that the as-prepared surface causes nearly no degradation of water-repellent superhydrophobicity with a contact angle higher than 155◦, indicating outstanding long-term stability. The multiscale hierarchical structure and chemisorbed HDTMS hydrophobic molecules facilitate the formation of an air cushion, significantly contributing to the eventual water-repellence, air-exposure stability and corrosive-medium immersion stability.

**Figure 10.** The variation of the contact angle and sliding angle under different exposure time in (**a**) air and (**b**) 3.5 wt.% NaCl aqueous solution.

#### **4. Conclusions**

In conclusion, to successfully design a lotus-inspired superhydrophobic AA5083 surface with a facile, cost-effective and fluorine-free strategy, we provide an approach combining etching texture followed by chemisorption of HDTMS hydrophobic molecules. It was shown that the surface topography of the as-prepared superhydrophobic AA5083 possesses a micro-nano hierarchical petaloid structure with a contact angle of 156.3◦ ± 1◦ and a sliding angle lower than 1◦. The wetting behavior investigations presented water-repellence, extremely low adhesivity and a self-cleaning ability. The EIS and fitting results showed that the *R*ct of the as-prepared superhydrophobic AA5083 specimen is two orders of magnitude higher than that of the pristine AA5083. In addition, the *Q*dl of the as-prepared superhydrophobic AA5083 was four orders of magnitude lower than the pristine AA5083. The inhibition efficiency in this case was approximately 96.11%. In addition, after 90 days of air exposure and 12 days of a 3.5 wt.% NaCl immersion/contact, the as-fabricated superhydrophobic AA5083 surface sustained a durable and stable superhydrophobicity. Therefore, the resultant superhydrophobic AA5083 surface possesses superior corrosion-resisting performance and long-term stability. We greatly anticipate that this research work has important significance for the large-scale manufacturing of water-repellent superhydrophobic surfaces for long-term marine and industrial applications.

**Author Contributions:** B.Z. and Y.L. conceived the concept of this work; B.Z. conducted the fabrication and characterization and wrote the manuscript; and W.X., Q.Z. and S.Y. performed the electrochemical experiments.

**Funding:** This research work was financially supported by the National Natural Science Foundation of China (Grant No. 41806089), the Strategic Priority Program of Chinese Academy of Sciences (Grant No. XDA13040401).

**Conflicts of Interest:** The authors declare no conflicts 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* **Optimization of Cathodic Protection Design for Pre-Insulated Pipeline in District Heating System Using Computational Simulation**

#### **Min-Sung Hong, Yoon-Sik So and Jung-Gu Kim \***

School of Advanced Materials Engineering, Sungkyunkwan University, 300 Chunchun-Dong, Jangan-Gu, Suwon 440-746, Korea; smith803@skku.edu (M.-S.H.); soy2871@naver.com (Y.-S.S.)

**\*** Correspondence: kimjg@skku.edu

Received: 13 May 2019; Accepted: 28 May 2019; Published: 30 May 2019

**Abstract:** Cathodic protection (CP) has been used as a primary method in the control of corrosion, therefore it is regarded as the most effective way for protecting buried pipelines. However, it is difficult to apply CP to a pipeline for district heating distribution systems, because the pipeline has thermally insulated coatings which could disturb the CP. Theoretical calculation and field tests alone are not enough for a reliable CP design, and therefore additional CP design methods such as computational analysis should be used. In this study, the CP design for pre-insulated pipelines is tested considering several environmental factors, such as temperature and coating defect ratio. Additionally, computational analysis is performed to verify and optimize the CP design. The simulation results based on theoretical methods alone failed to satisfy the CP criteria. Then, a re-design is conducted considering the IR drop. Consequently, all of the simulation results of defective pipelines satisfied the CP criteria after adding the proper CP current.

**Keywords:** cathodic protection; corrosion mitigation method; potentiodynamic polarization test; simulation; pre-insulated pipeline

#### **1. Introduction**

In district heating (DH) systems, heated water is distributed through a double-pipe network and transferred to buildings for use in space heating, hot water generation, and process heating [1]. DH systems have three main elements: the heat source, the distribution system, and the customer interface. The distribution system supplies hot water from the heat source to the heat consumer and returns with temperatures in the range of 40 ◦C to 120 ◦C [2]. Generally, pipelines in DH distribution systems use a thermally insulated coating to minimize heat loss during transfer. As shown in Figure 1, the coating consists of two layers: an inner layer of polyurethane foam (PUR) to reduce heat loss, and an outer layer of high-density polyethylene (HDPE) to protect the PUR [3]. The coatings effectively mitigate corrosion by blocking the outer environment, which contains corrosive elements such as water, oxygen, and chloride ions, when the coating is maintained perfectly. However, the HDPE is susceptible to unpredictable mechanical damage, and the PUR can be vanished by heat, humidity, and oxygen during its long operational life [4,5]. Several studies have reported that the main source of corrosion is groundwater introduced through failure of the HDPE and PUR [5–7].

Cathodic protection (CP) has been used as the primary method in the control of corrosion, in conjunction with protective coatings. CP can reduce the corrosion rate, and a properly maintained system will provide protection in accordance with the designed structural life [8]. The impressed current CP (ICCP) system has a power supply (rectifier) that generates larger potential differences between the anode and the structure [9]. For this reason, ICCP is applied to many industrial pipelines. However, despite the availability of CP, there are still several limitations in applying CP to pre-insulated

pipelines [10]. The National Association of Corrosion Engineers (NACE) reported that the CP design for pre-insulated pipelines is ineffective because the protection current cannot reach the corroded area through the insulating layer [11]. Additionally, according to previous studies, the corrosion rate of PUR-insulated carbon steel is much lower than that of uninsulated (bare) steel, even when the PUR is fully immersed in groundwater [12]. Therefore, CP for pre-insulated pipelines may be unnecessary when the PUR layer is intact. However, the immersed PUR layer can deteriorate and vanish during long operating periods, causing exposure of the bare carbon steel to the corrosive environment. For this reason, it is important to apply CP to operating pipelines with external coating defects, as a precaution against sudden fracture. Nevertheless, it is difficult to design CP systems for operating pipelines using only theoretical methods and a limited number of standards. It is also difficult to verify the appropriate protecting current required to reach the external surface of the pipeline with proper CP potential. For this reason, additional CP design methods, such as computational analysis, should be applied to optimize the design [13–15]. To improve the reliability of simulation results, several essential factors should be considered, such as polarization data for real materials and appropriate environmental information.

**Figure 1.** Schematic diagram of the pre-insulated coated pipe (600 A).

In this study, a CP system was designed for an existing pipeline with damaged insulation, taking into consideration environmental factors, such as corrosion properties of real materials, operating temperatures, and structural effects. Additionally, electrochemical tests were performed in synthetic groundwater to obtain input data for the computer simulation. Finally, a computational analysis was performed to verify and optimize the CP design of pre-insulated pipelines.

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

#### *2.1. Materials and Test Conditions*

The corrosion environment used was synthetic groundwater. Table 1 gives the chemical composition of the synthetic groundwater, and HNO3 was used to control the pH of the solution. A welded carbon steel specimen consisting of a base metal, heat affected zone, and a weld metal was used during testing to calculate the required CP current. Table 2 shows the chemical composition of the SPW400 (carbon steel), and Table 3 shows the welding methods used in all experiments. The surface of the specimen was polished with 600-grit silicon carbide (SiC) paper, degreased with ethanol, and dried with N2.


**Table 1.** Chemical composition of synthetic groundwater.



#### *2.2. Electrochemical Test Methods*

All electrochemical experiments were performed using a three-electrode system, in a 1000 mL Pyrex glass corrosion cell connected to an electrochemical apparatus. The test specimens were connected to a working electrode, a graphite rod was used as the counter electrode, and a saturated calomel electrode (SCE) was used as the reference electrode. The area of the test specimen exposed to the electrolyte was 2.25 cm2 (1.5 cm <sup>×</sup> 1.5 cm). An open-circuit potential (OCP) was established within three hours to carry out the electrochemical test. Potentiodynamic polarization tests were carried out in accordance with ASTM G5-14 (Standard Reference Test Method for Making Potentiodynamic Anodic Polarization Measurements), using a VMP2 (Bio-Logic Science Instruments, Seyssinet-Pariset, France) with a potential sweep of 0.166 mV/sec, from an initial potential of −2000 mV versus the reference to a final potential of 200 mV versus the OCP. The electrochemical tests were performed at 80 ◦C, because a previous study found that the highest protection current for carbon steel was required at this temperature [7].

#### *2.3. CP Design and Computational Analysis Method*

The computational analysis tool BEASY S/W (BEASY Ltd., Southampton, England), which is based on the boundary element method (BEM), was used to conduct 3D modeling and computational analysis of the pre-insulated pipeline. The required CP current (Ireq) for the pipeline was calculated, taking into consideration the current density of real material measured by electrochemical tests. The cathodic polarization curve, which was used as input data for the simulation, was obtained from the potentiodynamic polarization test, which incorporated the environmental information.

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

#### *3.1. Potentiodynamic Polarization Tests*

The applied current density (iapp) for the pre-insulated pipeline was calculated using the Evans diagram, as shown in Figure 2. According to the diagram, anodic current density is under activation control (activation polarization), and cathodic current density is limited at a higher current density (concentration polarization). As the applied current density for CP is increased, the potential and the

corrosion current density are reduced simultaneously [16,17]. According to the previous study, since the pre-insulated pipeline has a high corrosion rate at 80 ◦C, the reasonable maximum CP potential is −1350 mVSCE [7]. Figure 3 shows the results of the potentiodynamic polarization test in synthetic groundwater at 80 ◦C. The corrosion current density was determined using the Tafel extrapolation method. Table 4 shows the calculated CP current density, which will apply to the CP design. The applied current density was calculated as the difference between the anodic polarization curve and cathodic polarization curve at −1350 mVSCE, as shown in Figure 2. The cathodic polarization curve, which contains the corrosion properties of real material, was used as the input data for computational analysis.

**Figure 2.** Evan's diagram, indicating the relationship between the applied current density and protection potential [16].

**Figure 3.** Potentiodynamic polarization curves in the synthetic groundwater at 80 ◦C.

**Table 4.** Results of Potentiodynamic Polarization Test at 80 ◦C in the Synthetic Groundwater.


#### *3.2. Cathodic Protection Design and Computational Analysis*

The pre-insulated pipeline was connected every 6 m by welds, therefore, the CP design was preformed to 6 meters of 600 A pipe (Figure 1). In addition, ICCP anodes were installed at both edges of the pipeline, which are the parts most sensitive to corrosion because it will connect using welding. In this study, the CP design was applied to operating pipelines with slight defects. Therefore, the CP design was tested at a range of defect ratios (1, 5, 10, 20%), and it is assumed that the insulating part of the pipeline has no defect. Figure 4 shows the 3D modeling of the pipeline according to defect ratio. For modeling and calculations, the approach was based on the assumption that the crevice between the coating and pipeline was not effective as a CP [18]. Table 5 shows the basic design parameters related to the structural factors. The surface area of the pipe used in the CP design was 12.62 m2, which included an additional 10% safety factor. The resistivity of soil was assumed to be 1000 Ω·cm, corresponding to a highly corrosive environment. The required current (Ireq) for CP was calculated from the following equation [19,20]:

$$\mathbf{I}\_{\rm req} = \mathbf{C}\_{\rm defect} \cdot \mathbf{i}\_{\rm app} \cdot \mathbf{A}\_{\rm pipe},\tag{1}$$

where Cdefect is the defect ratio of the pipeline, iapp is the applied current density of the pipe material calculated from the electrochemical test, Apipe is the surface area of the pipe. Ireq is calculated with the defect ratio, as listed in Table 6.

**Figure 4.** 3D modeling of pipeline according to the defect ratio: (**a**) 1%, (**b**) 5%, (**c**) 10%, (**d**) 20%.


**Table 5.** Basic design parameters related to the structural factors.

**Table 6.** Required current calculation for cathodic protection (CP).


The computational analysis was performed using the cathodic polarization curve data, obtained from the electrochemical tests. Figure 5 shows the simulation results for CP. All of the simulation results failed to satisfy the CP criteria for pre-insulated pipelines (under −1350 mVSCE) because the IR drop caused by soil and structural factors was not considered in the CP design.

**Figure 5.** Simulation results (averaged protection potential, mVSCE) according to the defect ratio: (**a**) 1%, (**b**) 5%, (**c**) 10%, (**d**) 20%.

The additional CP current required to satisfy the CP criteria should be calculated taking into consideration the polarization curve, as shown in Figure 6. The maximum CP potentials were defined based on the simulation results according to the defect ratio. Then, the applied current densities were calculated at the maximum CP potential from the simulation results, using the same method as above. To obtain the additional CP current densities, the difference was calculated between the calculated applied current densities, according to the defect ratio and applied current density at −1350 mVSCE. The additional CP currents were then calculated using Equation (1). The calculated values are listed in Table 7.

**Figure 6.** Calculation of additional current caused by IR drop in the polarization curve: potential difference of (**a**) 1%, (**b**) 5%, (**c**) 10%, (**d**) 20% defected pipelines.


**Table 7.** Results of calculated additional CP current and optimized current for CP.

Then, the entire simulation was re-conducted. Figure 7 shows the optimized simulation results, and it was verified that all of the pipelines with different defect ratios satisfied the CP criteria. Another important point is over-protection due to the low CP criteria of district pipelines. The simulation results show that the minimum CP potentials have a range from −1.7 VSCE to −2.6 VSCE. This is quite a low potential value, which could cause hydrogen embrittlement risk. However, according to the international standards, such as NACE (RP0169-96), ARAMCO (SAES-X-400), and BSI (BS 7361-1), the over protection range of the steel pipeline ranges from −2.5 VSCE to −5 VSCE. Therefore, the simulation results can apply up to 10% of the defect ratio, which has a minimum potential of about −2.49 VSCE. When the CP applies over 10% of the defect ratio, the site of defect should be previously investigated. Then, the anode should be installed as close as possible to the defect area, to avoid over protection and reduce CP current requirement. Therefore, the investigation of the defect area is one of the significant design parameters in practical CP installation.

**Figure 7.** Optimized simulation results (averaged protection potential, mVSCE) according to the defect ratio: (**a**) 1%, (**b**) 5%, (**c**) 10%, (**d**) 20%.

#### **4. Conclusions**

In this study, a credible CP design method for existing pre-insulated pipelines was conducted, taking into consideration the environmental factors, and computational analysis was performed to verify and optimize the CP design. According to the results, the following conclusions were drawn:


**Author Contributions:** Conceptualization, M.-S.H.; methodology, Y.-S.S.; software, M.-S.H.; validation, M.-S.H. and J.-G.K.; formal analysis, M.-S.H. and Y.-S.S.; investigation, M.-S.H.; resources, M.-S.H.; data curation, M.-S.H. and Y.-S.S.; writing—original draft preparation, M.-S.H.; writing—review and editing, M.-S.H. and J.-G.K.; visualization, M.-S.H.; supervision, J.-G.K.; project administration, M.-S.H. and J.-G.K.

**Funding:** This research was supported by the program for fostering next-generation researchers in engineering of National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT (2017H1D8A2031628).

**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* **Microstructure and Corrosion Resistance of Zn-Al Di**ff**usion Layer on 45 Steel Aided by Mechanical Energy**

#### **Jianbin Tong 1, Yi Liang 2,\*, Shicheng Wei 2, Hongyi Su 2, Bo Wang 2, Yuzhong Ren 3, Yunlong Zhou <sup>1</sup> and Zhongqi Sheng <sup>1</sup>**


Received: 30 August 2019; Accepted: 17 September 2019; Published: 18 September 2019

**Abstract:** In harsh environments, the corrosion damage of steel structures and equipment is a serious threat to the operational safety of service. In this paper, a Zn-Al diffusion layer was fabricated on 45 steel by the Mechanical Energy Aided Diffusion Method (MEADM) at 450 ◦C. The microstructure and composition, the surface topography, and the electrochemical performance of the Zn-Al diffusion layer were analyzed before and after corrosion. The results show that the Zn-Al diffusion layer are composed of Al2O3 and Γ<sup>1</sup> phase (Fe11Zn40) and δ<sup>1</sup> phase (FeZn6.67, FeZn8.87, and FeZn10.98) Zn-Fe alloy. There is a transition zone with the thickness of about 5 μm at the interface between the Zn-Al diffusion layer and the substrate, and a carbon-rich layer exists in this zone. The full immersion test and electrochemical test show that the compact corrosion products produced by the initial corrosion of the Zn-Al diffusion layer will firmly bond to the Zn-Al diffusion layer surface and fill the crack, which plays a role in preventing corrosion of the corrosive medium and reducing the corrosion rate of the Zn-Al diffusion layer. The salt spray test reveals that the initial corrosion products of the Zn-Al diffusion layer are mainly ZnO and Zn5(OH)8Cl2H2O. New corrosion products such as ZnAl2O4, FeOCl appear at the middle corrosion stage. The corrosion product ZnAl2O4 disappears, and the corrosion products Zn(OH)2 and Al(OH)3 appear at the later corrosion stage.

**Keywords:** Zn-Al diffusion layer; mechanical energy aided diffusion; microstructure; corrosion resistance; electrochemistry

#### **1. Introduction**

Metal materials have long been exposed to environments with high temperatures, high humidity, high salt spray, and intense sunlight, making their corrosion levels several times or even dozens of times higher than that in other environments at the same time [1–10]. In the process of metal corrosion, the mechanical properties and internal microstructure of the metal have changed. The corrosion harm includes not only the damage of internal metal structures but also the destruction of overall metal structures. Metal corrosion causes tremendous economic loss and becomes a severe threat to the development of various fields. Therefore, corrosion protection of metal materials is essential in many industrial applications [11–14]. In various environments, steel is the most used type of material in various facilities and equipment. Therefore, the corrosion protection of steel structures and equipment have flourished over recent decades [15–20].

Chemical heat treatment is usually used to improve the corrosion resistance, high-temperature oxidation resistance and hardness of metal parts [21–24]. However, due to the high temperature and long production time, it not only consumes a lot of energy, but also affects the mechanical properties of parts. Therefore, the Mechanical Energy Aided Diffusion Method (MEADM) that emerged in recent years has become an attractive metal materials anti-corrosion technology in the field of material surface strengthening. The MEADM, developed by the addition of mechanical energy (rotation, vibration, friction, etc.) in the traditional solid pack-cementation process, achieves the purpose of reducing the preparation temperature required and shortening the preparation time. As such, the MEADM develops rapidly. There are three steps in the MEADM as shown in Figure 1: (a) with the assistance of mechanical energy, the active powder particles in the diffusion agent rub and impact the surface of the heated substrate, resulting in micro vacancies and plastic deformation on the substrate surface; (b) the active particles enter the generated vacancies or adsorb into the deformed surface to form a surface solid solution or an intermetallic compound, forming an initial diffusion layer; (c) a dense protective diffusion layer is achieved by adsorbing the active particles in a continuous fashion.

**Figure 1.** Schematic illustration of the mechanical energy aided diffusion processes.

Yuan et al. [25] have employed a pack-cementation process to produce aluminide coatings on both ferritic-martensitic (RAFM) and austenitic (316L) alloys at the temperature of 600–800 ◦C. Lee [26] has applied a pack-cementation process to form aluminide coatings at 850 ◦C for 2 h. However, Wang et al. [27,28] studied the MEADM and found that mechanical energy can improve the adhesion ability of Al powder on the surface of the substrate, increase the chemical activity and adhesion strength of Al, and thereby significantly reduce the aluminizing temperature, which can be reduced to 500 ◦C [27]. The MEADM attributes to the temperature reduction for the following reasons: (a) the powder, the workpiece, and the vessel wall collide with each other in the mechanical energy aided diffusion process to form flow and heat transfer, enhance the heat conduction and thereby accelerating the diffusion rate; (b) the moving powder particles hit the surface of the workpiece to cause activation and misalignment of the surface lattice atoms, forming supersaturated vacancies and thereby reducing the diffusion temperature; (c) the movement of the powder particles accelerates the chemical reaction of the diffusion and increases the concentration of the active Al atoms; (d) the movement of the powder particles purifies the surface of the workpiece.

Zhang et al. [29] have used the MEADM to prepare the Zn diffusion layer on Q235 steel, and found that the diffusion layer, composed of FeZn15, FeZn11, FeZn9, and FeZn7 phases, has excellent resistance to high temperature oxidation and corrosion. He et al. [30] obtained Al-Zn-Cr diffusion layer on 20 steel at 600 ◦C by the MEADM and concluded that the Al-Zn-Cr diffusion layer, a multi-layered structure, is resistant to high temperature oxidation and corrosion. In addition to excellent corrosion resistance and high temperature oxidation resistance, the Zn-Al diffusion layer (based on the MEADM) has the characteristics of uniform thickness, high hardness, good scratch resistance, high powder recyclability, and low environmental pollution. At present, Zn layers, Al layers, and Zn-Al layers are studied widely [27,28,31–33].

In this paper, in order to analyze the severe corrosion problem of steel structure and equipment in harsh corrosive environments, 45 steel was used as the substrate material, and the Zn-Al diffusion layer was prepared by the MEADM at 450 ◦C without changing the properties of the substrate. In this paper, we investigated the microstructures and element distribution of the Zn-Al diffusion layer. The 3D topography and corrosion mechanism of the Zn-Al diffusion layer surface were studied by full immersion test with 3.5 wt. % NaCl. The surface topography and corrosion product composition of the Zn-Al diffusion layer in different salt spray corrosion periods were analyzed. The results can provide a technical basis for long-term corrosion protection of steel structures and equipment in harsh environments.

#### **2. Experimental**

#### *2.1. Experimental Materials and Sample Preparation Process*

The 45 steel was used as the substrate material, and its specific components are shown in Table 1. The sample size was 10(L) × 10(W) × 4(H) mm.

The preparation process of the Zn-Al diffusion layer is mainly divided into three parts: pre-treatment, mechanical energy aided diffusion process, and post-treatment. The oil and rust on the surface of the substrate were removed by sanding, ultrasonic cleaning, and shot blasting (0.2 mm steel shot) during the pre-treatment. The mechanical energy aided diffusion processes are as follows:

(1) Stirring: The diffusion agent was weighed with a certain pre-calculated percentage (0.02% NH4Cl, 0.05% Rare Earth, 49.93% Al2O3, 35% Zn, and 15% Al) and mixed uniformly with a mixer.

(2) Loading furnace: A part of the mixed agent was taken out into the rotary furnace (a special mechanical energy aided diffusion device), the workpiece was put into the furnace, then the remaining agent was put into the furnace, and the furnace lid was tightened finally.

(3) Heating and holding: The furnace was heated in the electric furnace and rotated at a constant rotational speed of 7 r/min. The timing started when the temperature reaches 450 ◦C. The heating was stopped when the set holding time (4 h) was reached.

(4) Cooling: The furnace was kept rotating at a constant speed (can be increased to 10 r/min). The furnace was cooled naturally to room temperature in the air.

(5) Separation: The workpiece and the powder in the furnace were separated by filtration. The post-treatment included cleaning, alcohol wiping, drying, and testing.


**Table 1.** Elemental composition of 45 steel (wt. %).

#### *2.2. Experimental Equipment and Parameters*

(1) The cross-sectional topography, composition and element distribution of the Zn-Al diffusion layer before and after corrosion were analyzed by Scanning Electron Microscope (SEM) (Nova Nano SEM50, FEI, Hillsboro, OR, US).

(2) The phase composition of the Zn-Al diffusion layer before and after corrosion was analyzed by X-ray Diffractometer (XRD) (Smartlab, Rigaku, Tokyo, Japan). The specific test conditions were Cu

target (9 KW), accelerating voltage 40 kV, tube current 40 mA, scanning rate 5◦/min, test angle 10◦–90◦, and scanning step length 0.02◦.

(3) The 3D surface topography of the sample in different immersion stages was measured by 3D Laser Scanning Microscope (LEXT OLS4100, OLYMPUS, Tokyo, Japan).

(4) The samples of different corrosion stages were obtained on Salt Spray Testing Chamber (YWX-010, SHUANGKE, Beijing, China). The specific test conditions were that the corrosive medium was 5 wt. % NaCl solution, the temperature was 35 ◦C, and the spray method was continuous.

(5) Electrochemical impedance spectroscopy (EIS) (amplitude 5 mV, scanning frequency range 10 mHz–100 kHz) and potentiodynamic polarization (scanning speed 1 mV/s, scanning range ±250 mV of electrode potential) were tested on an Electrochemical Workstation (IM6, Zahner, Kronach, Germany).

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

#### *3.1. Cross-Sectional Topography and Composition Analysis of the Zn-Al Di*ff*usion Layer*

Figure 2 shows the SEM images and corresponding Energy Dispersive Spectroscopy (EDS) spectra of the Zn-Al diffusion layer in cross section. Figure 2b is an enlarged image of the rectangular area in Figure 2a,c,d are EDS spectra of the Zn-Al diffusion layer. It can be found that the main elements of the diffusion layer are Zn, Al, and Fe. After the Al powder is diffused by the MEADM at 450 ◦C, a high Al content is only detected near the diffusion layer surface, which indicates that the Al layer with a thickness of 2–4 μm is mainly present in the superficial layer. In the vertical direction of the Zn-Al diffusion layer, the content of Zn changes little, and the overall content shows a slight downward trend, whereas the content of Fe increases slowly.

On observation of Figure 2b, there is a transition zone with a thickness of 5 μm at the interface between the Zn-Al diffusion layer and the substrate. The content of Zn and Fe in the transition zone changes abruptly. The content of Zn reduces to almost 0. The Fe element content remains constant after a sharp increase. The transition zone is divided into two parts. The first part (Part A) is close to the Zn-Al diffusion layer and has a thickness of 3 μm. The other part (Part B) is close to the substrate and has a thickness of 2 μm. This part is formed by the diffusion of Zn into the substrate. A small number of pores (average diameter 0.6 μm) are found at the boundary of the transition zone by the SEM image. These pores are caused by a small amount of entrained air that cannot be discharged in time when the diffusion element permeates into the substrate. This is attributed to the fact that the ambient preparation temperature is not stable and the substrate temperature is low at the initial stage of the Zn-Al diffusion layer growth.

**Figure 2.** *Cont*.

**Figure 2.** Scanning Electron Microscope (SEM) images and corresponding Energy Dispersive Spectroscopy (EDS) spectra of the Zn-Al diffusion layer in cross section: (**a**) sectional topography of the Zn-Al diffusion layer; (**b**) enlarged image of the rectangular area; (**c**,**d**) EDS spectra of the Zn-Al diffusion layer.

Figure 3 shows the XRD patterns of the Zn-Al diffusion layer. It is found that Al2O3 and Zn-Fe alloys of Γ<sup>1</sup> phase (Fe11Zn40) and δ<sup>1</sup> phase (FeZn6.67, FeZn8.87, FeZn10.98) are mainly formed in the Zn-Al diffusion layer. Combined with SEM images and EDS spectra, the element ratios at positions 1, 2, 3, and 4 in Figure 4, are shown in Table 2. The Zn-Fe content ratios (WZn:WFe) are 10, 8.2, 8.4, and 6.4, respectively. According to XRD patterns, the Zn-Fe alloy near the surface of the Zn-Al diffusion layer is mainly FeZn10.98, the Zn-Fe alloy near the boundary between the diffusion layer and the substrate is mainly FeZn6.67, and the Zn-Fe alloy in the middle area of the diffusion layer is mainly FeZn8.87.

**Figure 3.** XRD patterns of the Zn-Al diffusion layer.

**Figure 4.** EDS spot scanning position distribution of the Zn-Al diffusion layer in cross section.


**Table 2.** Percentage of constituent elements of the Zn-Al diffusion layer in different positions.

The content and distribution of elements in the region near the interface of the Zn-Al layer were analyzed. The elements and ratios at the positions numbered by 1, 2, 3, 4, and 5 in Figure 5 are shown in Table 3. The Zn-Fe content ratios (WZn:WFe) in the positions numbered by 1, 2, 3, 4, and 5 are 6.8, 6.2, 5, 5.87, and 3, respectively, further indicating that the Zn-Fe alloy near the substrate is FeZn6.67. Analyses of SEM images and EDS spectra reveal that the closer to the substrate the Zn-Al diffusion layer is, the higher the C content is. At the positions numbered by 3, 4, and 5, the C content reaches 51.85%, 47.12%, and 47.78%, respectively, indicating that the region of the Zn-Al diffusion layer forms a carbon-rich layer near the substrate. The C content of the carbon-rich layer is much higher than that of the 45 steel substrate. This is because Zn and Al are non-carbide forming elements, which cause the C atom crowding-out effect during the formation of the diffusion layer [34]. Therefore, the diffusion layer will have a carbon-rich layer near the substrate.

**Figure 5.** Position distribution of EDS spot scanning in the region near the interface of the Zn-Al layer.

**Table 3.** Percentage of constituent elements in the region near the interface of the Zn-Al layer.


#### *3.2. Corrosion Resistance*

3.2.1. 3D Surface Topography Analysis of the Zn-Al Diffusion Layer during Full Immersion

In order to study the corrosion resistance of the Zn-Al diffusion layer, the samples were immersed fully in 3.5 wt. % NaCl solution, and the 3D topography of the samples in different immersion stages was characterized at room temperature. The topographical images shown in Figure 6 are composed of a real graph and a color graph, and the color graph is the color mark of the height of the Zn-Al diffusion layer in different regions of the real one.

Observing the surface topography, it is found that the surface of the Zn-Al diffusion layer is uniform and has a low roughness before the full immersion test, and the drop between the high and low points is within 54 μm. The reason is that the substrate was subjected to shot blasting during pre-treatment, which caused a certain roughness on the surface of the substrate. Therefore, the Zn-Al diffusion layer is uniformly distributed on the surface of the substrate, which causes a certain degree of surface drop.

After 240 h full immersion, the surface color of the sample changed. Compared with the surface before the full immersion, more pronounced peak and pit features are exhibited, and the maximum drop between the high and low points increase to 74 μm. After full immersion, some corrosion products accumulate on the Zn-Al diffusion layer surface and form some corrosion pits, which aggravates the surface roughness. With the immersion time extended to 600 h, a large number of white corrosion products appear on the Zn-Al diffusion layer surface. As the accumulated corrosion products further increase, the area and depth of corrosion pits further increase, and the maximum drop between the high and low points reaches 110 μm. When the sample is immersed for 1000 h, the corrosion products on the Zn-Al diffusion layer surface further increase, and the surface corrosion pits are obvious. The maximum drop on the surface is increased to 131 μm and some red rust spots are observed in the real image. At this moment, a small number of corrosion pits have penetrated the entire Zn-Al diffusion layer to the substrate.

**Figure 6.** 3D topography of the Zn-Al diffusion layer during full immersion: (**a**) before immersion; (**b**) immersed for 240 h; (**c**) immersed for 600 h; (**d**) immersed for 1000 h.

Figure 7 depicts the microscopic corrosion topography and EDS spectra of the Zn-Al diffusion layer after immersing for 360 h and ultrasonic cleaning for 10 min. The corrosion topography show that the metal powder on the Zn-Al diffusion layer surface is actively dissolved, and the flocculent corrosion products deposit on the surface, covering the entire surface. The energy spectrum analysis of the filler in the surface crack of the diffusion layer is shown in Table 4. In addition to the elements of O, Al, Fe, and Zn, the Cl element which is the main element causing corrosion is detected. It indicates that the filler in the surface crack is corrosion products produced by Cl− corroding in solution. After ultrasonic cleaning, the corrosion products are still present in the crack, indicating that the corrosion products are firmly bonded to the Zn-Al diffusion layer. The firmly combined corrosion products fill the crack to help block the intrusion tunnel of the corrosive medium, which can slow down the corrosion rate of the Zn-Al diffusion layer and improve the protection ability for the substrate [35,36].

**Figure 7.** (**a**) Microstructure of the Zn-Al layer for 360 h immersion; (**b**) EDS spectra of the Zn-Al layer for 360 h immersion.


**Table 4.** Main elemental weight ratio and atomic ratio of the Zn-Al diffusion layer for 360 h immersion.

#### 3.2.2. Corrosion Behavior Analysis of the Zn-Al Diffusion Layer

In order to further study the corrosion resistance of the Zn-Al diffusion layer, a neutral salt spray test was developed. The sample was placed in a salt spray test chamber. The surface corrosion topography and the corrosion product changes of the Zn-Al diffusion layer in different corrosion stages were studied.

The surface corrosion topography and corrosion product XRD results of the Zn-Al diffusion layer in different salt spray corrosion stages are shown in Figures 8 and 9. At the initial stage of salt spray corrosion (within 168 h), a layer of flocculent corrosion products uniformly forms on the Zn-Al diffusion layer surface. The corrosion products cover the surface, so that the tunnels (the corrosion solution can invade the substrate through these tunnels) are reduced, thereby the corrosion resistance of the Zn-Al diffusion layer is improved. XRD results show that the corrosion products on the Zn-Al diffusion layer surface are mainly composed of ZnO, Al2O3, and Zn5(OH)8Cl2H2O. From the corrosion products formed, with the electrochemical reaction proceeding, Na<sup>+</sup> moves toward the cathodic region, and Cl<sup>−</sup> moves toward the anodic region. Zinc hydroxychloride (Zn5(OH)8Cl2H2O) and Zinc oxide (ZnO) gradually form in the anodic dissolution region.

When corroding to the middle stage of corrosion (480 h), the flocculent corrosion products on the surface have become the needle-like corrosion products that are shown by a network-like structure on the surface. According to the XRD results, it is obvious that the main corrosion products are comprised by ZnO, Al2O3, Zn5(OH)8Cl2H2O, ZnAl2O4, and FeOCl. Compared with the initial corrosion stage, the number of Zn5(OH)8Cl2H2O on the surface increases, and the density of corrosion product layer increases, which helps slow down the corrosion from corrosive medium and reduce the corrosion rate of the Zn-Al diffusion layer. In addition, the newly formed corrosion product, iron oxychloride (FeOCl), is a structurally unstable corrosion intermediate that can accelerate corrosion. However, when FeOCl releases Cl−, it can form FeO(OH), and the migrated OH− can react with the metallic ions in the corrosive medium to develop new products. These products cover the Zn-Al diffusion layer surface, which further suppress the corrosion of the Zn-Al diffusion layer to some extent.

In the later stage of corrosion (1000 h), more agglomerated products appear on the surface of the diffusion layer. Compared with the initial stage and the middle stage of corrosion, a small number of flocculent and needle-like corrosion products distribute on the surface. The agglomerated corrosion product layer is easy to fall off. Therefore, the corrosion solution easily passes through the pores between the corrosion products and penetrates the diffusion layer. At this time, the corrosion resistance of the Zn-Al diffusion layer is weakened. XRD analyses show that ZnAl2O4 disappear on the surface and Zn(OH)2 and Al(OH)3 appear in comparison with the middle corrosion stage. The disappearance of ZnAl2O4 is due to the decrease of Al content in the diffusion layer and ZnAl2O4 formed and the shedding of ZnAl2O4 of the surface with the prolongation of corrosion time.

**Figure 8.** Surface topography of the Zn-Al diffusion layer during salt spray corrosion: (**a**) 168 h; (**b**) 480 h; (**c**) 1000 h.

**Figure 9.** XRD patterns of the Zn-Al diffusion layer corrosion products during salt spray corrosion: (**a**) 168 h; (**b**) 480 h; (**c**) 1000 h.

#### *3.3. Electrochemical Performance Analysis of the Zn-Al Di*ff*usion Layer*

Figure 10 depicts the potentiodynamic polarization curves of the Zn-Al diffusion layer in 3.5 wt. % NaCl solution for different immersion times, and Table 5 shows the corresponding polarization curve fitting data. Through the polarization curve analysis, it is found that as the immersion time is prolonged, the self-corrosion potential of the Zn-Al diffusion layer increases significantly, but the self-corrosion current density decreases by an order of magnitude. In other words, the corrosion rate of the Zn-Al diffusion layer decreases with the prolongation of immersion time in a certain time range.

In the whole process of full immersion corrosion, the anode Tafel slope βa of the Zn-Al diffusion layer is less than the cathode Tafel slope βc, indicating that the corrosion reaction of the Zn-Al diffusion layer is mainly controlled by the cathodic reaction. The specific reaction is as follows: the anodic reaction Zn <sup>−</sup> 2e<sup>−</sup> <sup>→</sup> Zn2<sup>+</sup>, Al <sup>−</sup> 3e<sup>−</sup> <sup>→</sup> Al3<sup>+</sup>; the cathodic reaction O2 <sup>+</sup> 2H2O <sup>+</sup> 4e<sup>−</sup> <sup>→</sup> 4OH−; the total reaction 2Zn + O2 + 2H2O → 2Zn(OH)2, 4Al + 3O2 + 6H2O → 4Al(OH)3 [37]. It is found that the cathode Tafel slopes βc and the anode Tafel slopes βa are not much different during the initial immersion stage. At this time, the anodic reaction is that Zn and Al in the Zn-Al diffusion layer dissolve to produce Zn2<sup>+</sup> and Al3<sup>+</sup> in the corrosive medium, and the cathode absorbs oxygen to form OH−. With the prolongation of immersion time, the anode Tafel slope βa remains unchanged, but the cathode Tafel slope βc increases gradually. It indicates that the cathodic oxygen-absorbing reaction (the formation of the corrosion products layer) controls the corrosion reaction of the Zn-Al diffusion layer with the βc value increases. The accumulating rate of electrons in the cathode region accelerates, resulting in a decrease of the self-corrosion potential difference and the corrosion current density between anode and cathode. In addition, the polarization resistance Rp also increases greatly with the prolongation of immersion time, indicating that the corrosion products formed on the Zn-Al diffusion layer surface accumulate gradually and the corrosion rate of the Zn-Al diffusion layer decreases. The compact corrosion products formed adhere to the surface, which acts as a protective layer and

slows the corrosion rate. The results of the polarization potential test are also consistent with the surface analysis results of the full immersion test.

**Figure 10.** Potentiodynamic polarization curve of the Zn-Al diffusion layer.


**Table 5.** Polarization curve fitting data of the Zn-Al diffusion layer.

Figure 11 presents the EIS of the Zn-Al diffusion layer in different immersion time, and Table 6 shows the impedance modulus in different immersion times. The impedance modulus diagram (Figure 11a) shows that the impedance modulus of the low-frequency region decreases within 24 h of immersion. When the immersion time is 72–360 h, the impedance modulus of the low-frequency region increases sharply, indicating that the corrosion rate of the Zn-Al diffusion layer decreases at this time. The phase angle diagram (Figure 11b) shows that only a time constant characteristic appears in the Zn-Al diffusion layer, and the peaks in the phase angle diagram gradually become higher as the immersion time is prolonged. It is concluded that the prolongation of immersion time (more than 24 h) makes the Zn-Al diffusion layer to be an isolating layer with high resistance and low capacitance, which plays a protective role for the Zn-Al diffusion layer. In the Nyquist diagram (Figure 11c), there is only one capacitive reactance arc in the Zn-Al diffusion layer, and the radius of the capacitive reactance arc decreases firstly and then increases with the prolongation of the immersion time, which is consistent with the change law of the corrosion resistance of the Zn-Al diffusion layer in the impedance modulus diagram and phase angle diagram.

**Figure 11.** EIS spectra plots of different immersion time of the Zn-Al diffusion layer: (**a**) Bode impedance modulus; (**b**) Bode phase angle; (**c**) Nyquist.

**Table 6.** The impedance modulus and corrosion rate of different immersion time of the Zn-Al diffusion layer.


The corrosion of the Zn-Al diffusion layer is a controlled process in which the electrochemical reaction gradually changes to the diffusion of corrosive medium or corrosion products. During the initial immersion stage, the impedance modulus of the low-frequency region in the 3.5 wt. % NaCl solution is low. The reactions on the Zn-Al diffusion layer surface are mainly zinc oxide, aluminum oxide, and zinc aluminum active dissolution. When the immersion time is 10 and 24 h, the impedance modulus of the low-frequency region decreases. The prolongation of the immersion time makes the electrolyte solution continuously penetrate the Zn-Al diffusion layer, causing continuous corrosion damage of the Zn-Al diffusion layer and the decreases of the impedance modulus. Therefore, as the immersion time is prolonged, the impedance modulus of the low-frequency region decreases. However, as the immersion time continues to increase, the corrosion products continuously deposit on the surface and fill into the crack. The tunnel from the corrosive medium to the substrate reduces, slowing the penetration rate of the corrosive electrolyte. Therefore, the impedance modulus of the Zn-Al diffusion layer rises rapidly in the 3.5 wt. % NaCl solution, and the corrosion rate decreases remarkably.

Two equivalent electrical circuits shown in Figure 12 were utilized to fit the EIS data and account for the corrosion behavior of the Zn-Al diffusion layer. The first equivalent circuit (Rs(Qc(Rc(QctRct)))) was used to fit the EIS data displaying the Zn-Al diffusion layer within 24 h of immersion, whereas the second one (Rs(QcRc)(QctRct)) was used for the EIS data displaying the impedance after immersing for 24 h. In Figure 12, Rs is the solution resistance, Rc is the Zn-Al diffusion layer resistance, Rct is the

charge transfer resistance, Qc is the Zn-Al diffusion layer capacitance, and Qct is the electric double layer equivalent capacitance between the Zn-Al diffusion layer and the substrate.

**Figure 12.** EIS equivalent circuit models of the Zn-Al diffusion layer in different immersion times: (**a**) for fitting the diffusion layer impedance data within 24 h of immersion(Rs(Qc(Rc(Qct Rct)))); (**b**) for fitting the diffusion layer impedance data after immersing for 24 h (Rs(QcRc)(Qct Rct)).

The equivalent circuit fitting data of the Zn-Al diffusion layer is shown in Table 7. When the immersion time is 0–360 h, the Zn-Al diffusion layer capacitance value Qc in the corrosion solution increases gradually. At this stage, the Zn-Al diffusion layer is invaded by a corrosive medium, increasing the value of Qc. The microscopic fluctuation of the Zn-Al diffusion layer surface leads to a large or small deviation of the electric double layer capacitance Qct, so that the Qct value changes irregularly, that is, the dispersion phenomenon. When immersed for 0–24 h, the Rc value decreases gradually, and the Rct value decreases slightly, which is consistent with the radius variation of the capacitive reactance arc in the Nyquist diagram. Analyses of the diagram data reveal that the diffusion layer has a much larger resistance Rc than the solution resistance Rs. In other words, the penetration of the corrosion solution makes the Zn-Al diffusion layer resistance Rc decrease. As the immersion time continues to increase, the Rc and Rct values increase gradually. The Zn-Al diffusion layer reacts with the corrosive solution, which generates a large number of corrosion products on the surface or inside and accumulates continuously. The phenomenon of the above reaction reduces the porosity of the Zn-Al diffusion layer, and strengthens the self-sealing effect, so that it slows down the progress of the electrochemical reaction.


**Table 7.** Fitting data of the equivalent circuit of the Zn-Al diffusion layer.

#### **4. Conclusions**

(1) The Zn-Al diffusion layer aided by mechanical energy is composed of mainly Al2O3 and Γ<sup>1</sup> phase (Fe11Zn40) and δ<sup>1</sup> phase (FeZn6.67, FeZn8.87, FeZn10.98) Zn-Fe alloy. The FeZn10.98 alloy is mainly located in the region near the Zn-Al diffusion layer surface, and the FeZn6.67 alloy is mainly located in the interface between the Zn-Al diffusion layer and the substrate. There is a transition zone between the substrate and the Zn-Al diffusion layer with a thickness of about 5 μm, and a carbon-rich layer exists in this zone. In the direction of perpendicular to the surface, the content of Zn reduces gradually, the content of Fe increases slowly, and abrupt changes occur in this transition zone.

(2) As the immersion time is prolonged, the corrosion products accumulate on the surface and form corrosion pits, which increase the surface roughness. At the same time, the corrosion products of the Zn-Al diffusion layer accumulate on the surface to fill the surface crack and prevent the erosion of the corrosive medium, thereby reducing the corrosion rate of the Zn-Al diffusion layer and exerting the effect of enhancing the corrosion resistance.

(3) The results of salt spray corrosion test show that the corrosion products of the Zn-Al diffusion layer surface are flocculent in the initial corrosion stage, and the corrosion products are mainly ZnO and Zn5(OH)8Cl2H2O. In the middle corrosion stage, the needle-like corrosion products appear on the Zn-Al diffusion layer surface, and new corrosion products are ZnAl2O4 and FeOCl. In the later corrosion stage, the agglomerated corrosion products appear, the ZnAl2O4 disappears in the corrosion products, and Zn(OH)2 and Al(OH)3 appear. The appearance of Zn5(OH)8Cl2H2O increases the compactness of the corrosion products layer and slows down the corrosion rate of the Zn-Al diffusion layer. FeOCl is unstable and reacts with OH− to form FeO(OH) that accumulates on the surface, which further slows down the corrosion rate.

(4) The results of the Zn-Al diffusion layer electrochemical test show that as the immersion time is prolonged, the self-corrosion potential of the Zn-Al diffusion layer shifts gradually to the positive electrode after moving to the negative electrode. The self-corrosion current density first rises and then falls, the polarization resistance Rp increases, and the corrosion resistance of the Zn-Al diffusion layer increases gradually.

(5) The impedance modulus of low-frequency region decreases, and the radius of the capacitive reactance arc becomes smaller within 0–24 h. At this time, the corrosion solution invades the Zn-Al diffusion layer. The corrosion diffusion stage is 72–360 h, and the impedance modulus and the capacitive reactance arc radius of the low-frequency region increases gradually. In this stage, the corrosion rate of the Zn-Al diffusion layer gradually decreases. After immersing for 360 h, the Zn-Al diffusion layer still has excellent corrosion resistance.

**Author Contributions:** Y.L., S.W. and Z.S. designed the experiments, J.T., Y.Z. and Y.R. performed the experiments, J.T., H.S. and B.W. analyzed the data, J.T., Y.L. and B.W. wrote the paper.

**Funding:** This research was funded by The National Natural Science Foundation of China (51675533 and 51701238), and Equipment Pre-research Sharing Technology Project of '13th five-year' (41404010205).

**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* **Mechanical and Corrosion Resistance Enhancement of Closed-Cell Aluminum Foams through Nano-Electrodeposited Composite Coatings**

#### **Yiku Xu 1,\*, Shuang Ma 1, Mingyuan Fan 1, Hongbang Zheng 1, Yongnan Chen 1,\*, Xuding Song <sup>2</sup> and Jianmin Hao <sup>1</sup>**


Received: 26 July 2019; Accepted: 26 September 2019; Published: 29 September 2019

**Abstract:** This work aims to improve the properties of aluminum foams including the mechanical properties and corrosion resistance by electrodepositing a SiC/TiN nanoparticles reinforced Ni–Mo coating on the substrate. The coatings were electrodeposited at different voltages, and the morphologies of the coating were detected by SEM (scanning electron microscope) to determine the most suitable voltage. We used XRD (x-ray diffraction) and TEM (transmission electron microscope) to analyze the structure of the coatings. The aluminum foams and the substrates on which the coatings were electrodeposited at a voltage of 6.0 V for different electrodeposition times were compressed on an MTS (an Electro-mechanical Universal Testing Machine) to detect the mechanical properties. The corrosion resistance before and after the electrodeposition experiment was also examined. The results showed that the coating effectively improved the mechanical properties. When the electrodeposition time was changed from 10 min to 40 min, the Wv of the aluminum foams increased from 0.852 J to 2.520 J and the σ<sup>s</sup> increased from 1.06 MPa to 2.99 MPa. The corrosion resistance of the aluminum foams was significantly improved after being coated with the Ni–Mo–SiC–TiN nanocomposite coating. The self-corrosion potential, pitting potential, and potential for primary passivation were positively shifted by 294 mV, 99 mV, and 301 mV, respectively. The effect of nanoparticles on the corrosion resistance of the coatings is significant.

**Keywords:** aluminum foam; electrodeposition; compression test; corrosion resistance

#### **1. Introduction**

Aluminum foams have the characteristics of both metal materials and foam materials due to their special structure. They are functional materials with the properties of both structural materials and functional materials [1–4]. Aluminum foams have a unique stress–strain curve including a linear elastic region, a plastic collapse region, and a densification region, which makes aluminum foam materials suitable for use as an energy absorber [5]. Due to its excellent properties including light weight, high sound absorption and insulation performance, heat resistance, and high cushioning performance, it is widely used in sound absorption and sound insulation structures such as sound barriers and sound insulation boards, and for energy absorption and collision protection in automobiles [6–9]. However, the high porosity of aluminum foam significantly lowers its mechanical strength. When applied in an engineering field, it often fails prematurely, which greatly limits its potential range of applications. For example, when an aluminum foam is used for sound insulation and heat transfer in an environment that requires a certain load, the aluminum foam will fail and be crushed after the load acts on it for a long period of time [10]. The seawater, micro-organisms, and salt spray in a marine environment can corrode aluminum foams when they are used in marine transportation applications. In order to expand the range of uses of aluminum foams and make them better for practical applications, surface modification methods are used to simultaneously improve their mechanical properties and corrosion resistance. An aluminum foam with a high energy absorption capacity should have a longer and higher stress platform. At present, using the same material to thicken the foam pillar is a common method for improving the energy absorption capacity [11–14]. However, this method has certain limitations. When the pillar of the aluminum foam is thickened, the platform stress will be improved, but the aluminum foam will be dense [15–17]. Densification may limit the ability of aluminum foams to absorb energy. At present, the commonly used surface modification methods for aluminum foams include micro-arc oxidation, anodization, electro-less plating, sol–gel deposition, and electrodeposition [18–20]. Of these methods, electrodeposition is widely used because it is simple, low cost, and easy to control.

At present, there are a number of reports on the use of deposited layers to enhance the properties of aluminum foams. Yuttanant Boonyongmaneerat et al. [21] electrodeposited a nanocrystalline Ni–W coating on open-cell aluminum foams to improve their properties including compressive strength and energy abortion. Zhendong Li et al. [22] confirmed that a thermally evaporating Zn film could significantly enhance open-cell aluminum foams and increase their yield strength. Liu Huan et al. [10] studied the enhancements that a Ni coating could provide to closed-cell aluminum foams. It demonstrated that a Ni coating could improve the properties of aluminum foams including both the mechanical and corrosion resistance properties. Jiaan Liu et al. [23] showed that the corrosion resistance of closed-cell aluminum foams could be improved by an electro-less Ni–P coating. Due to the excellent mechanical properties, corrosion resistance, and wear resistance of the Ni–Mo coating, it often used as an protective coating [24–26]. SiC nanoparticles have a high degree of hardness, wear resistance, and thermal stability, and TiN nanoparticles have a high degree of hardness, high strength, and corrosion resistance [27–29]. As far as we know, no research has been done on the use of a Ni–Mo coating and a duplex nanoparticles reinforced Ni–Mo coating to enhance the properties of closed-cell aluminum foams.

In this work, the influence of a Ni–Mo coating and a duplex nanoparticles reinforced Ni–Mo coating on the mechanical properties and corrosion resistance of closed-cell aluminum foams was studied. The effects of electrodeposition voltage and electrodeposition time on the morphology, mechanical properties, and corrosion resistance of the closed-cell aluminum foams were investigated. The deposition mechanism of the duplex nanoparticles reinforced Ni–Mo coating is also discussed.

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

#### *2.1. Samples and Solution*

We used the method of melt foaming to prepare the closed-cell aluminum foams in this experiment. The density of the samples was 0.2 g/cm3. The pore diameter was 4 mm. We used an electrical discharge machine to reduce the sample's dimensions to 20 mm × 20 mm × 9 mm.

To create a good bond between the substrate and the coating, the aluminum foam was pretreated before the electrodeposition experiment. The aluminum foam sample was immersed in a 10~15% H2SO4 solution at 60 ◦C for 1–3 min. After immersion, the oil was removed. Then, a 5% NaOH solution was used to remove the Al2O3 film from the surface of the samples. The sample was immersed for 2 min. Finally, the aluminum foam was immersed for 5 min in a 10% HNO3 solution. The corrosion products were removed and activated. After each of the steps was completed, the sample was washed with distilled water to prevent the pretreatment liquid from being contaminated. After the pretreatment steps were completed, the aluminum foam sample was placed in the electrolyte immediately to prevent it from being oxidized in the air.

Table 1 shows the electrolyte components that were used in this experiment. The electrolyte was composed of analytically pure reagent and distilled water. The added SiC and TiN nanoparticles (Shanghai Chaowei Nanotechnology Co. Ltd., Nanxiang Hi-Tech Industrial Park, Jiading District, Shanghai) both had a mean particle diameter of 20 nm and purity of 99 wt.%. Since nanoparticles tend to agglomerate in the electrolyte, SDS was chosen as the dispersing agent. The electrolyte was subjected to ultrasonic treatment for 2 h and the electrolyte was stirred using a magnetic stirrer with a speed of 300 rpm during the electrodeposition experiment. Electrolyte (200 mL) was placed in a bath, pure nickel plate (99.99 wt.%) was used as the anode, and the aluminum foam sample was used as the cathode. The anode and the cathode had a distance of 30 mm between them.


**Table 1.** Components of the electrolyte for the electrodeposition of the Ni–Mo–SiC–TiN coating.

#### *2.2. Morphology Investigation*

The surface morphology and a cross-section of the coating were observed using a Hitachi S4800 field scanning electron microscope (SEM, Hitachi, Ltd., Tokyo, Japan). The elements were analyzed by energy dispersive x-ray spectroscopy. The structure of the coating was examined by D8 ADVANCE x-ray diffraction (XRD, Bruker, Karlsruhe, Germany). Cu-k<sup>α</sup> radiation was selected, and the 2θ range was 20~80◦. In order to further analyze the specific structure of the nanocomposite coating and the distribution of nanoparticles, the coating was examined by an FEI Talos F200X transmission electron microscope (TEM, FEI™, Hillsboro, OR, USA) including high-resolution TEM (HR-TEM) and selected area electron diffraction (SAED).

#### *2.3. Properties Investigation*

An electrodeposited aluminum foam with the dimensions of 20 mm × 30 mm × 40 mm was subjected to a quasi-static compression test on an MTS (an Electro-mechanical Universal Testing Machine, American MTS Corporation, MN, USA) with a selected load of 10 KN, a compression speed of 5 mm/min, and a compression amount greater than 70%.

The corrosion resistance of the sample at room temperature was measured by the three-electrode working system. In this experiment, a 3.5 wt% NaCl solution was used as the etching solution. The working electrode was the aluminum foam sample, the reference electrode was the saturated calomel electrode, and the counter electrode was the platinum plate electrode. The selected voltage range was −2 to 1 V and the scan rate was 2 mV/s.

The samples were placed in an immersion test for 120 h to measure the corrosion rate at 25 ◦C. The immersion solution was a 3.5 wt% NaCl. The samples were weighed to calculate the mass loss every 24 h. Distilled water was used to rinse the samples, and they were dried thoroughly before each weighing. The weight of a sample was expressed as the average of three measurements. The analytical balance that was used to weigh the samples had an accuracy of 0.01 mg.

#### **3. Theoretical Models**

Electrodeposition of metals and alloys refers to the reduction of metal ions from an electrolyte, where electrons (*e*) are provided by an external power supply. The reaction time and the current can optimize the thickness of a coating. Molybdenum cannot be electrodeposited from the electrolyte

solution, but the co-deposition of nickel and molybdenum can be achieved using sodium citrate as an inducer. During the electrodeposition of Ni–Mo composite coatings on an aluminum foam, the following chemical reactions occur at the cathode and anode [30]:

Anode:

$$\text{Ni} - 2\text{e} \to \text{Ni}^{2+} \tag{1}$$

Cathode:

$$2\text{ Ni}^{2+} + 2\text{ e} \to \text{Ni} \tag{2}$$

$$\text{MoO}\_4\text{}^{2-} + 2\text{H}\_2\text{O} + 2\varepsilon \rightarrow \text{MoO}\_2 + 4\text{OH}^-\tag{3}$$

$$\text{NiCl}^- + \text{MoO}\_2 \rightarrow \left[ \text{NiClMoO}\_2 \right]^-\_{\text{ads}} \tag{4}$$

$$\left[\mathrm{NiClMoO\_2}\right]^-\_{\mathrm{ads}} + 2\mathrm{H\_2O} + 4e \rightarrow \mathrm{Mo} + \mathrm{NiCl}^- + 4\mathrm{OH}^- \tag{5}$$

With respect to the co-deposition of nanoparticles with a Ni–Mo matrix, the processes include three main steps, as illustrated in Figure 1. According to Gugliemi's absorption model, Ni ions and Mo ions in the electrolyte solution are first adsorbed on the nanoparticles to form Ni/Mo ionic clouds. Under the electric field force, metal ions and ionic clouds move toward the cathode and are tightly adsorbed on the aluminum foams. Then, the Ni and Mo ions adsorbed on the surface of the nanoparticles are reduced partially at the surface of the foam. Simultaneously, nanoparticles are trapped by the metal matrix and embedded in the Ni–Mo plating layer.

**Figure 1.** A schematic diagram representing the electrodeposition process of a duplex nanoparticles reinforced Ni–Mo coating.

Based on a theoretical model of Cu electrodeposition, the model of the electrodeposited Ni–Mo alloy coating in this experiment is now described [31].

The plating deposition rate is expressed by P%, and its expression is:

$$\mathbf{P}\% = \left[ (\mathbf{P}\_2)\_i - \mathbf{P}\_{1i} \right] / \mathbf{P}\_{1i} \tag{6}$$

where P1 indicates the mass of the substrate before the electrodeposition experiment; P2 indicates the mass of the aluminum foam covered with a coating; and *i* indicates the sample number. P% is the ratio of the mass of the aluminum foam covered with a Ni–Mo coating to the mass of the aluminum foam before electrodeposition.

$$\mathbf{P}^{\rm o}\_{\rm o} = \frac{\mathbf{M}\_{\rm NiMo}}{\rho \ast V\_i} = (\mathbf{M} \mathbf{M}\_{\rm NiMo} \ast \mathbf{n}\_{\rm NiMo}) / (\rho \ast \mathbf{V}\_i) \tag{7}$$

where MNiMo is the mass of the deposited Ni–Mo alloy coating; MMNi–Mo is the molar mass of Ni–Mo alloy; and ρ and *Vi* indicate the density and volume of the aluminum foam before the electrodeposition experiment, respectively.

The Ni–Mo alloy that was formed in this experiment is a Ni–Mo solid solution. When 1 mole of Ni–Mo alloy coating is deposited, 14 moles of electron are required. Then, P% also can be expressed as:

$$P\% = (\text{MM}\_{\text{NiMo}} \* \text{n}\_{\text{c}}) / \ (14 \* \rho \* V\_{\text{i}}) \tag{8}$$

where *e* is the electric charge of an electron.

It is known that *ne* = q/(Na ∗ e), q = *i* ∗ *t*. Then,

$$\mathbf{P\%} = \mathbf{M} \mathbf{M}\_{\text{NiMo}} / \left( \mathbf{14} \ast \mathbf{N}\_{\text{a}} \ast \mathbf{e} \right) \ast \left[ \left( i \ast t \right) / \left( \rho \ast V\_{i} \right) \right] \tag{9}$$

where *<sup>t</sup>* is the electrodeposition time (in minutes); *Na* is Avogadro's number with a value of 6.02 <sup>×</sup> 1023; MMNiMo is 331 g/mol; and *<sup>e</sup>* is 1.6 <sup>×</sup> <sup>10</sup>−<sup>19</sup> C. Then,

$$P\% = 2.45 \times 10^{-4} \ast [(i \ast t) / (\rho \ast V\_i)]. \tag{10}$$

As this experiment was carried out under a certain voltage, the expression is written as

$$P\% = 2.45 \times 10^{-4} \ast \left[ (u \ast t) / (\rho \ast V\_i \ast r) \right] \tag{11}$$

where *u* is the electrodeposition voltage and *r* is the total resistance.

The relationship between the deposition rate of a Ni–Mo coating, the electrodeposition voltage *u*, and the time *t* can be obtained by Equation (11), and the P% that is obtained by experiments can be verified using theoretical calculations.

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

#### *4.1. Coating Characterization*

Figure 2b show the SEM images of the two kinds of nanoparticles with an original size of approximately 20 nm. As can be seen, both kinds of nanoparticles were agglomerated due to the surface effect.

Figure 2d–f show the morphologies of the electrodeposited duplex nanoparticles reinforced Ni–Mo coatings, applying electrodeposition voltages ranging from 2.5 V to 6.0 V, respectively. It has been reported that nanoparticles can make a coating have a finer grain and a higher microhardness. In accordance with this, SiC- and TiN-reinforced coatings have structures with finer grain sizes than Ni–Mo composite coatings. As the voltage increased, coating particles were gradually formed and completely covered the substrate. When the voltage was increased to 6.0 V, a uniform coating was prepared on the aluminum foam. The SEM image shown in Figure 2i revealed that the surface of the coating had nanoparticles dispersed upon it. At the voltage of 6.0 V, a nodular and homogenous Ni–Mo coating was also obtained. It is known that a larger electrodeposition voltage can increase the nucleation driving force, so plating particles are formed. The metal ion deposition rate was sufficiently high to form a uniform and dense coating on the substrate at the voltage of 6.0 V. Figure 2h shows the morphology of a cross-section of aluminum foam that was subject to electrodeposition for 10 min at 6.0 V. The coating had a thickness of about 25 μm. The bond between the plating layer and the substrate was good, the thickness of the plating layer was relatively uniform, and there were no cracks or discontinuities.

**Figure 2.** SEM (scanning electron microscope) images of (**a**) SiC nanoparticles, (**b**) TiN nanoparticles, (**c**) the Ni–Mo coating, and the duplex nanoparticles reinforced Ni–Mo coating electrodeposited (**d**) at 2.3 V, (**e**) 4.5 V, and (**f**) 6.0 V. The enlarged SEM images of the duplex nanoparticles reinforced Ni–Mo coating (**g**) at 2.3 V, (**i**) 6.0 V; and (**h**) morphology of a cross-section of the duplex nanoparticles reinforced Ni–Mo coating.

Figure 3 shows the XRD patterns of the coatings. The body-centered cubic structure that corresponds to nickel's (111), (200), and (220) diffraction peaks. No diffraction peak of molybdenum was detected, indicating that the nickel atom and the molybdenum atom existed in the form of a Ni–Mo solid solution. The nanoparticles did not change the structure of the Ni–Mo coating. In the XRD patterns, there were no diffraction peaks related to nanoparticles. This is mainly because the size of the nanoparticles was too small, their content too low, and the distribution was uniform [32]. The intensity of the peaks of the XRD patterns of the coatings electrodeposited at 6.0 V for different times were different. We used the Scherrer formula to calculate the crystallite size:

$$\mathbf{D} = \mathbf{K}\lambda / (\beta \mathbf{C} \mathbf{O} \mathbf{S} \theta) \tag{12}$$

where λ represents the wavelength of the x-ray (0.15406 nm); K is the Scherrer constant (0.9); β is the full width of the reflection line at half maxima; and θ is a Bragg diffraction angle.

**Figure 3.** XRD (x-ray diffraction) patterns of coatings electrodeposited by different times.

Table 2 shows the results. As the electrodeposition time increased, the grains of the coatings accumulated and the crystallite size increased. Comparing the crystallite size of the Ni–Mo coating to that of the duplex nanoparticles reinforced Ni–Mo coating, it can be seen that the nanoparticles decreased the crystallite size of the coating. Nanoparticles can inhibit the grain growth because they provide nucleation dots.



Figure 4 shows the EDS (energy dispersive spectrometer) elements mapping of the duplex nanoparticles reinforced Ni–Mo coating. Ni, Mo, Si, C, Ti, and N elements were detected. The existence of Si, C, Ti, and N elements indicates that duplex nanoparticles were successfully electrodeposited in the Ni–Mo composite coating. The EDS element mapping demonstrates the specific distribution of duplex nanoparticles. Nanoparticles were uniformly dispersed in the coating, but partial agglomeration occurred. Since the coating used for EDS (energy dispersive spectrometer) detection was a 100 nm thin layer, the distribution of nanoparticles inside the coating can be known.

**Figure 4.** EDS (energy dispersive spectrometer) element mapping of the Ni–Mo–SiC–TiN nanocomposite coating.

Figure 5a presents the TEM images under a bright field. From the images, it was found that the nanoparticles were tightly embedded in the Ni–Mo metal matrix and there were no voids between them. The interface between the nanoparticles and the Ni–Mo metal matrix was clear and there were no harmful interfacial reaction products. The selected area electron diffraction rings in Figure 5b correspond to the (111), (200), (220), and (311) crystal faces of the nickel–molybdenum solid solution, respectively. The fast Fourier transform (FFT) and inverse FFT of the nanoparticles in Figure 5a are shown in Figure 5c. The nanoparticles were proven to be 6H–SiC, which have a hexagonal closed-packed (HCP) structure with a Lattice constant of 3.08 Å.

**Figure 5.** (**a**) Bright field image of the duplex nanoparticles reinforced Ni–Mo coating; (**b**) diffraction ring of the metal matrix; and (**c**) HR-TEM image of the nanoparticles.

#### *4.2. Mechanical Behavior*

Figure 6a shows the stress–strain curves of the aluminum foam and the aluminum foams subjected to electrodeposition for different times. The enlarged elastic region of the curve is shown in Figure 6b. The stress–strain curve of the aluminum foam includes three parts: an elastic deformation stage, a yield stage, and a densification stage. In the initial stage of the compression experiment, the stress increased as the strain increased. The relationship between stress and strain was linear. When the curve entered the yield stage, the stress appeared to be small or substantially constant as the strain increased. In the densification stage, since the pores inside the aluminum foam burst and collapsed, the stress increased sharply at this stage and the strain remained substantially unchanged.

**Figure 6.** (**a**) Stress–strain curves of aluminum foam and aluminum foam subject to electrodeposition at the voltage of 6.0 V; (**b**) the enlarged elastic region of (**a**). (**c**) the yield strength and unit volume energy in function of the final density of samples.

Compared with the aluminum foam matrix, the elastic modulus and the platform stress of the aluminum foams after the coating was deposited were improved. When the strain remained the same, the stress of the aluminum foam after electrodeposition was larger than that of the aluminum foam substrate. There are two main reasons for the increase in strength and elastic modulus of aluminum foams with electrodeposited coatings. The first reason is that, due to the particularity of the structure of the aluminum foam, the deformation of the aluminum alloy during the compression experiment was not synchronized, resulting in separation of the coating from the substrate. The second reason is the friction and extrusion between the coating and the substrate. The stress–strain curve also showed that the coating increased the energy absorption of the aluminum foam.

The density of an aluminum foam affects its mechanical strength. An increase in density will increase the compressive properties. The density is related to the electrodeposition time. The electrodeposition time determines the quality of the coating that is deposited on the substrate, so the quality affects the mechanical strength of the substrate. Table 3 lists the coefficient of variation P% of different electrodeposition times. The resistance, which includes the external contact resistance *r*1, the solution resistance *r*2, and the resistance of the cathode film r3 during the electrodeposition

process, are all uncertain. The resistance at 10 min of electrodeposition was used as the resistance in this experiment.


**Table 3.** Characteristics of different samples.

The coefficient of variation of 20 min, 30 min, and 40 min of electrodeposition, as calculated by the established electrodeposition theoretical model, was 21.35%, 33.76%, and 42.16%, respectively, while the P% obtained from the experiment was 23.5%, 38.8%, and 53.4%, respectively. Density of Ni–Mo–SiC–TiN coatings does not obviously change with deposition time. As the deposition time increased, the error between the theoretical model and the results obtained from deposition rate increased. The main reason for this is that an increase in the electrodeposition time will cause a large change in resistance.

Table 3 also lists the density, yield strength, densification strain, and Wv of the aluminum foam after the electrodeposition experiments. Wv represents the energy absorbed per unit volume when the aluminum foam is deformed. When comparing the compression properties of the samples after different electrodeposition times, it was found that the stress–strain curves of aluminum foams moved upward with the increase of electrodeposition time. This is mainly because an increase in electrodeposition time will increase the quality of the coating on the aluminum foam. The quality of the coating on the surface increases the strength and stiffness of the aluminum foam. From the stress–strain curve, it can be seen that the curve appeared to fluctuate in the stress platform stage, which is due to the instability of the aluminum foam. This can be attributed to the non-uniformity of the aluminum foam's cell structure and its rough surface. When the stress–strain curve passes the linear elastic phase, the stress tends to decrease; the reasons for this are discussed in the literature [33,34].

The stress remained almost constant as the strain increased in the stress platform stage, which allowed the sample to absorb a large amount of energy during the compression process. Figure 6 shows the absorbed energy per unit volume of the aluminum foam during the quasi-static compression experiment in the stress–strain curve. Its calculation expression is [35]

$$\mathcal{W}\_{\rm V} = \int\_0^{\varepsilon\_{\rm D}} \sigma(\varepsilon) \mathrm{d}\varepsilon \tag{13}$$

where ε<sup>D</sup> represents the densification strain, which corresponds to a sharp rise in stress during compression because the aluminum foam is crushed and deformed and the cell structure completely collapses, and σ(ε) represents the stress.

Gibson et al. proposed the following relationship between the densification strain of closed-cell aluminum foam, ε*<sup>D</sup>* [36], and the relative density ρ:

$$
\varepsilon\_{\rm D} = 1 - 1.4 \,\overline{\rho} \tag{14}
$$

where ρ is the ratio of the apparent density ρ of the aluminum foam to the density ρ<sup>s</sup> (2.70 g/cm3) of the aluminum foam matrix.

The relationship between the Wv and the apparent density ρ of aluminum foams is obtained from the above two formulas:

$$\mathcal{W}\_{\mathbf{V}} = \int\_0^{1-0.518\rho} \mathbf{\sigma}(\varepsilon) \, \mathrm{d}\varepsilon. \tag{15}$$

The relationship indicates that Wv is related to the density of aluminum foams.

The specific relationship between the density and mechanical properties of samples was explored after the electrodeposition experiments, the density of aluminum foams after the electrodeposition of a coating between the Wv, and yield strength σ**<sup>s</sup>** were respectively fitted. The fitting results are shown in Figure 6c, d. The relationship between the unit volume energy absorption Wv and the density ρ is

$$\mathbf{W}\_{\rm V} = \mathbf{a} + \mathbf{b}\_1 \boldsymbol{\rho} + \mathbf{b}\_2 \boldsymbol{\rho}^2. \tag{16}$$

From the fitting results, the value of a, b1, and b2 is −12.05942 ± 1.72093, 40.51736 ± 5.70059, and <sup>−</sup>28.13555 <sup>±</sup> 4.65988, respectively. The value of the correlation coefficient R2 is 0.99376.

The relationship between the yield strength σ<sup>s</sup> and the density ρ is

$$
\mathfrak{o}\_{\mathfrak{k}} = \mathfrak{d} + \mathfrak{c}\rho.\tag{17}
$$

From the fitting results, the value of d, c, and the correlation coefficient R<sup>2</sup> is <sup>−</sup>0.5350 <sup>±</sup> 0.44692, 5.0664 ± 0.74711, and 0.93748, respectively.

Due to the particular structure of each cell of the aluminum foams and the differences in the deposition rate, the data shown in Figure 6c are relatively discrete.

When the aluminum foams were subjected to electrodeposition for 10 min, 20 min, 30 min, and 40 min, Wv was quadratic with ρ, and σ<sup>s</sup> increased linearly with ρ. Comparing the mechanical properties of the substrates, which were coated with a Ni–Mo coating and a duplex nanoparticles reinforced Ni–Mo coating, the addition of nanoparticles only slightly increased Wv and σs. This limited enhancement of the compressive properties is due to the small amount of nanoparticles in the Ni–Mo coatings.

#### *4.3. Corrosion Resistance*

The corrosion resistance of the aluminum foam and the aluminum foams with an electrodeposited Ni–Mo coating and a duplex nanoparticles reinforced Ni–Mo coating was detected. The obtained polarization curves are shown in Figure 7a. Table 4 lists the corrosion parameters extracted from the polarization curves. After the Ni–Mo coating was electrodeposited on the aluminum foam, the corrosion potential of the aluminum foam was positively shifted from −1160 mV to −937 mV and the corrosion current density decreased from 4.48 <sup>×</sup> 10−<sup>5</sup> A/cm<sup>2</sup> to 3.90 <sup>×</sup> 10−<sup>5</sup> A/cm2. The pitting potential and the potential for primary passivation were positively shifted by 48 mV and 225 mV, respectively. The positive shift of the Zero current potential was due to changes in the hydrogen evolution reduction process. Both the aluminum foam and the aluminum foam with an electrodeposited Ni–Mo coating formed passive films. The aluminum foam formed a passive film because of the oxide layer. The oxygen-rich surface reacted with the etching solution to form an adsorption layer. The adsorption layer prevented contact of the etching solution with the surface of the plating layer to prevent the hydration of nickel, which is the first step in the formation of a passive nickel film on the surface of the aluminum foam covered with the Ni–Mo coating.

**Figure 7.** The polarization curves of the samples (**a**) and weight loss versus time curves after the immersion test (**b**); SEM images of the substrate and coatings after the polarization test (**c**) and immersion test (**d**).

Compared with the substrate, the corrosion resistance of the samples after electrodeposition was greatly improved. The Ni–Mo coating was found to effectively protect the aluminum foam from corrosion. In order for a corrosive liquid to have a corrosive effect on the aluminum foam's substrate, the passivation film on the surface of the aluminum foam must first be destroyed. The *Cl*− in the etching solution was found to easily pass through the passivation film due to the small radius and adsorb on the samples to hinder the adsorption of oxygen. The cations in the passivation film combined with the *Cl*− to form a soluble chloride. The substrate was partially exposed due to the local corrosion. The aluminum foam had a galvanic effect with the oxide film to form a corrosive micro-battery. The matrix and the impurity elements Ca, Ti, and Si, which were contained in the aluminum foam, also formed a corrosive micro-battery. This resulted in an uneven accumulation and distribution of *Cl*−, which exacerbated the local corrosion.

The coating was able to effectively protect the aluminum foam matrix mainly because the electrodeposited coating could separate the aluminum foam matrix from the etching solution so the Ni-Mo coating had an initial corrosion. The amorphous Ni–Mo alloy coating had good corrosion resistance, and the Mo element could easily form an inert oxide with oxygen in solution to prevent further corrosion of the coating [37,38]. The plating layer was uniform and compact. The coating had a thickness of about 25 μm. The pinholes and cracks in the plating layer were reduced. These all made the distribution of *Cl*− become uniform. The corrosion on the surface of the coating was relatively uniform. Then, there was a better corrosion resistance.

From the polarization curves of the samples, corrosion parameters can be obtained. The corrosion potential was shifted from –0.937 V for the Ni–Mo coating to –0.866 V for the duplex nanoparticles reinforced Ni–Mo coating, and the corrosion current density was reduced from 3.90 <sup>×</sup> 10−<sup>5</sup> A/cm<sup>2</sup> to 2.72 <sup>×</sup> <sup>10</sup>−<sup>5</sup> <sup>A</sup>/cm2. This change illustrates that the SiC and TiN nanoparticles can have an improvement on the corrosion resistance of the Ni–Mo coating. The reason for this is that these two kinds of nanoparticles are inert nanoparticles that have a certain degree of corrosion resistance. Dispersed nanoparticles can enhance the corrosion resistance of the coating because the nanoparticles can block the etching solution and the coating from coming into contact.

Figure 7c shows the SEM of the sample after the polarization experiment in corrosive solution. The aluminum foam was severely corroded, and there were many corrosion products and corrosion pits on the aluminum foam. An EDS spectrum analysis was performed on the corrosion surface, and the oxygen content in the corrosion product was found to be 17.04 wt.%. Compared with the aluminum foam, the Ni–Mo coating provided better protection to the substrate. Corrosion occurred on the surface of the coating when corrosion occurred. Local corrosion cracks could be observed, and the oxygen content in the corrosion products decreased to 8.03 wt.%. After adding duplex nanoparticles, the number of corrosion products was significantly reduced. This is because the two inert types of nanoparticles protected the matrix coating. Nanoparticles filled the voids in the matrix coating and improved the compactness of the coating. The smaller contact area effectively reduced the corrosion rate. The low content (4.69 wt.%) of oxygen also indicated an improvement in corrosion resistance.

Figure 7d shows the corrosion morphologies of the aluminum foam, the Ni–Mo coating, and the duplex nanoparticles reinforced Ni–Mo coating deposited at the voltage of 6.0 V after the immersion test. There are many corrosion pits on the surface of the substrate. The aluminum foam was obviously corroded because there was no protection of the coatings. The Ni–Mo coating was slightly corroded, and only a few corrosion products existed. Almost no corrosion was observed on the surface of the duplex nanoparticles reinforced Ni–Mo coating due to the protection of the nanoparticles.

Table 4 lists the corrosion rates of different samples. Compared with the aluminum alloy matrix, the corrosion rates of the Ni–Mo coating and the duplex nanoparticles reinforced Ni–Mo coating were improved by 51.9% and 72.5%, respectively.


**Table 4.** The corrosion parameters extracted from the polarization curves and weight loss vs. time curves.

#### **5. Conclusions**


and the σ<sup>s</sup> increased from 1.06 MPa to 2.99 MPa. The addition of nanoparticles made a limited improvement to the mechanical properties.

4. The duplex nanoparticles-reinforced Ni–Mo coating was found to have better corrosion resistance. Compared to the aluminum foams, the self-corrosion potential, the pitting potential, and the potential for primary passivation were positively shifted by 294 mV, 99 mV, and 301 mV, respectively. The corrosion rate of the aluminum foam covered with a Ni–Mo coating was reduced by 51.9%. After adding nanoparticles, the corrosion rate was reduced by 72.5%. The nanoparticles obviously improved the corrosion resistance.

**Author Contributions:** Conceptualization, S.M., M.F., and H.Z.; Funding Acquisition, Y.X., Y.C., and J.H.; Investigation, Y.X., S.M., X.S., and M.F.; Methodology, Y.X. and S.M.; Project Administration, Y.X., Y.C., and J.H.; Data Curation, M.F. and H.Z.; Writing—Original Draft, S.M.; Writing—Review & Editing, Y.X., Y.C., J.H., and X.S.

**Funding:** This work was financially supported by the National Natural Science Foundation of China (No. 51,301,021), the China Postdoctoral Science Foundation (No. 2016M592730), the Fundamental Research Funds for the Central Universities (Nos. 300,102,318,205; 310,831,161,020; 310,831,163,401; 300,102,319,304), the Innovation and Entrepreneurship Training Program of Chang' an University (No. 201,910,710,144), the Key projects of Shaanxi Natural Science Foundation (2019JZ-27), and the Shaanxi Natural Science Basic Research Program-Shaanxi Coal (2019JLM-47).

**Conflicts of Interest:** The authors declare no conflicts 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* **Electrochemical Corrosion Behavior of Ni-Fe-Co-P Alloy Coating Containing Nano-CeO2 Particles in NaCl Solution**

#### **Xiuqing Fu 1,2,\*, Wenke Ma 1, Shuanglu Duan 1, Qingqing Wang <sup>1</sup> and Jinran Lin 1,2**


Received: 29 July 2019; Accepted: 14 August 2019; Published: 16 August 2019

**Abstract:** In order to study the effect of nano-CeO2 particles doping on the electrochemical corrosion behavior of pure Ni-Fe-Co-P alloy coating, Ni-Fe-Co-P-CeO2 composite coating is prepared on the surface of 45 steel by scanning electrodeposition. The morphology, composition, and phase structure of the composite coating are analyzed by means of scanning electron microscope (SEM), energy dispersive spectroscopy (EDS), and X-ray diffraction (XRD). The corrosion behavior of the coatings with different concentrations of nano-CeO2 particles in 50 g/L NaCl solution is studied by Tafel polarization curve and electrochemical impedance spectroscopy. The corrosion mechanism is discussed. The experimental results show that the obtained Ni-Fe-Co-P-CeO2 composite coating is amorphous, and the addition of nano-CeO2 particles increases the mass fraction of P. With the increase of the concentration of nano-CeO2 particles in the plating solution, the surface flatness of the coating increases. The surface of Ni-Fe-Co-P-1 g/L CeO2 composite coating is uniform and dense, and its self-corrosion potential is the most positive; the corrosion current and corrosion rate are the smallest, and the charge transfer resistance is the largest, showing the best corrosion resistance.

**Keywords:** scanning electrodeposition; Ni-Fe-Co-P-CeO2 composite coating; electrochemical corrosion behavior; corrosion mechanism

#### **1. Introduction**

Corrosive environments are one of the most common service environments for metal components in engineering. Due to the high chemical activity of Fe in such environments, the engineering application of steel components therein is facing severe challenges due to insufficient corrosion resistance [1,2]. Surface modification is one of the most effective ways to solve this problem. The electroplating process for the preparation of nanocomposites is a process for the co-deposition of nanoparticles and metal ions on the surface of a cathode workpiece via the electrochemical principle, and a process to obtain nanocomposites that demonstrate superior performance [3,4]. Scanning electrodeposition technology, as an extension of electroplating technology, is widely used in machinery, aerospace, electronics industry, etc., due to its controllability, high efficiency, selectivity, and superior coating performance [5,6]. In recent years, many scholars have devoted to improving the performance of traditional nickel-based alloy coatings. Usually, tungsten [7], copper [8], iron [9], cobalt [10], and other metal ions [11] are introduced into the electrolyte, thereby processing a multi-component alloy. The multi-component overcomes the shortcomings of unary and binary alloy coatings, and has good wear resistance and corrosion resistance, which meet the varying performance requirements of composite materials [12,13].

It has been found in research that the application properties and functions of alloy coating can be further improved by co-depositing second phase nano-oxide particles in a nickel-based alloy coating [14,15]. The rare earth element cerium is the only stable tetravalent element. It has a unique oxidizing property and a large effective nuclear charge number, which can catalyze many reactions, and is widely used in various applications. Cerium oxide is a typical rare earth oxide with good wear resistance and corrosion resistance, and can be used as a nanoparticle reinforcement phase in various applications [16–18]. In order to further improve the corrosion resistance of traditional nickel-based coatings, Ni-Fe-Co-P-CeO2 composite coatings are prepared by scanning electrodeposition technology. The concentration of nano-CeO2 particles in the plating solution is applied to the coating of Ni-Fe-Co-P alloys. The influence of appearance and structure, and the electrochemical corrosion behavior of composite coatings, provide a reference for the development of new composite materials.

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

#### *2.1. Experimental Principle*

The scanning electrodeposition test apparatus is shown in Figure 1, wherein the anode nozzle is mounted on the machine tool spindle; the workpiece is mounted on the workpiece mounting platform by tightening the fixing screws. During the scanning electrodeposition process, the anode bed of the anode nozzle reciprocates in the Y direction, and the water pump presses the plating solution from the reservoir into the anode nozzle through the inlet tube and sprays it on the surface of the workpiece at high speed to spray the plating solution in the electrodeposition chamber. The liquid return tube flows back to the reservoir to realize the circulation of the plating liquid. After the power is turned on, the plating solution sprayed on the surface of the workpiece through the anode nozzle forms a closed loop, and under the action of the external electric field, a redox reaction occurs to realize deposition of metal ions. The scanning length during the test is 20 mm, and the scanning speed is 13.5 mm/s. The height between the bottom of the anode nozzle and the workpiece processing surface is 1.5 mm.

**Figure 1.** Scanning electrodeposition test device.

#### *2.2. Materials and Methods*

Fourty five steel with dimensions of 25 mm × 10 mm × 8 mm was used as substrate material, and its chemical composition is listed in Table 1. Table 2 shows the formulation of the plating solution used. The drugs used are of analytical grade and are prepared with deionized water. The particle size of the nano-CeO2 particles in the test was 100 nm, and the concentration of nano-CeO2 particles in the plating solution was 0, 0.5, 1, and 1.5 g/L, respectively. The cathode workpieces are polished with 800# and 1500# water sandpaper, respectively, and the workpiece is pretreated before scanning

electrodeposition process; the specific process is shown in the Table 3, after each step is rinsed with deionized water. The workpiece that has been subjected to the pre-treatment is placed in a spray electrodeposition test apparatus for a sputtering test. The current during the spray electrodeposition process is 0.6 A, the pH of the plating solution is 1.0–1.5, the bath temperature is 60 ◦C, and the plating time is 20 min. After the end of the scanning electrodeposition process, the workpiece was subjected to ultrasonic cleaning and drying treatment, and performance studies were performed.




#### **Table 2.** Composition of plating solution.


Nickel chloride hexahydrate

#### **Table 3.** The process of workpiece pretreatment.

#### *2.3. Characterization*

The morphology of the coating was observed by scanning electron microscopy (FEI-SEM, Quanta FEG250; FEI Instruments, Hillsboro, OR, USA), with an accelerating voltage of 15 kV and image type of secondary electron image (SEI); the chemical composition of the coating was determinated by energy dispersive spectroscopy (EDS, XFlash 5030 Bruker AXS, Inc., Berlin, Germany), with an accelerating voltage of 16 kV and the working distance of 11 mm; the phase structure of the coating was analyzed by X-ray diffraction (XRD, PANalytical X'pert; PANalytical Inc., Almelo, The Netherlands), with a radiation source of Cu Kα (λ = 0.15405 nm), operating voltage of 40 kV, scan rate of 5 ◦/min, and scanning range (2θ) of 10◦ ∼ 80◦, using HighScore Plus 3.0 to analyze the results.

(NiCl2·6H2O) <sup>3</sup>

The corrosion resistance of the coating was detected by electrochemical test of the three-electrode system (Figure 2). The working electrode is the workpiece, and the auxiliary electrode is Pt piece; the reference electrode is saturated calomel electrode (SCE), and the Tafel polarization curve measurement and electrochemical impedance spectroscopy (EIS) are completed by electrochemical workstation CS350 (Wuhan Corrtest Instruments Corp., Ltd., Wuhan, China). In the test, the workpiece to be tested was encapsulated with epoxy resin and immersed in a 50 g/L NaCl solution, and the Tafel polarization curve of the coating was obtained by a potentiodynamic scanning method and then obtained by polarization curve epitaxy. Corrosion potential, corrosion current, and other parameters were used to explore the corrosion resistance of the coating and the substrate. Under the open circuit potential, the impedance spectrum of the coating in NaCl solution was tested by the alternating current impedance method (EIS). The test frequency was 0.01–105 Hz, and the scanning direction was from high frequency to low frequency. The impedance fitting of different coatings was performed by Zview 2 software analysis.

**Figure 2.** Electrochemical detection device schematic.

#### **3. Results**

#### *3.1. Coating Morphology Analysis and Composition*

The SEM photographs in Figure 3 (image type is SEI) show the surface topography of the composite coating before corrosion. It can be seen that before the corrosion, the coating structure of different nano-CeO2 particles is composed of different sizes of cells, the arrangement is tight, and no obvious defects are found. When the concentration of nano-CeO2 particles is 0 g/L (Figure 3a), the cytoplasm is a spherical hillock-like structure, but the size difference is large, and there are also defects such as pores and protrusions. When a small number of nano-CeO2 particles is added to the plating solution (Figure 3b), the surface flatness of the coating is improved, but the cell structure has partial protrusions, the boundary is tortuous, and there are some defects such as pores. When the concentration is increased to 1 g/L (Figure 3c), the surface of the coating is dense and flat, the structure is compact, the cells are closely arranged, and the boundary is very blurred, and there are no obvious protrusions and impurity pores. When the concentration of nano-CeO2 particles in the plating solution reaches 1.5 g/L or more (Figure 3d), the surface morphology of the coating can be seen to have obvious agglomeration, and the surface of the coating is rough and uneven, with protrusions and defects generated. According to the analysis, the scanning jet of the plating solution accelerates the ion transport, increases the limiting current density, and strengthens the cathodic polarization, so that the deposition is performed at a high flow density [5]. The formation process of the compact nickel-based coating is similar to that of soil plant growth, and the nano-CeO2 particles dispersed in the plating solution are similar to the seeds, and are adsorbed on the surface of the substrate by tiny solid particles, because the rare earth element Ce is the third sub-group element. It has a large effective charge number and exhibits strong adsorption capacity. It can adsorb Ni2<sup>+</sup>, Fe2+, Co2+, and other ions [5]. As the deposition progresses, the seeds gradually grow, forming a cell structure with many different sizes. When the nano-CeO2 particles are excessive, they are excessively adsorbed on the surface of the metal substrate, causing the surface-active sites of the matrix to be masked and lose their activity, thereby greatly reducing or even inhibiting the nucleation sites, and uneven nanoparticle agglomerates are deposited on the surface

of the plating layer. The formation of larger protrusions affects the quality of the coating, and the advantage of nano-CeO2 particles is not obvious.

**Figure 3.** Surface morphology of the coatings before corrosion and EDS spectrum of coatings: (**a**) Ni-Fe-Co-P; (**b**) Ni-Fe-Co-P-0.5g/L CeO2; (**c**) Ni-Fe-Co-P-1g/L CeO2; and (**d**) Ni-Fe-Co-P-1.5g/L CeO2.

After cutting and inlaying the test piece, the cross-section of the test piece is observed by SEM, and the cross-sectional shape of the obtained coating is shown in Figure 4 (image type is SEI). It is obvious that the Ni-Fe-Co-P-CeO2 composite coating is uniform and dense, and there are no larger defects such as cracks and holes, effectively shielding the corrosion passage of the corrosive medium into the substrate and retarding the corrosion.

**Figure 4.** Cross-section morphology of the coatings: (**a**) Ni-Fe-Co-P; (**b**) Ni-Fe-Co-P-0.5 g/L CeO2; (**c**) Ni-Fe-Co-P-1 g/L CeO2; and (**d**) Ni-Fe-Co-P-1.5 g/L CeO2.

Using EDS technology, the EDS spectrum obtained by analyzing the composition of the surface of the coating is shown in Figure 3. Ni, Fe, Co, and P elements are present in all the energy spectra, and an appropriate number of nano-CeO2 particles are added to the plating solution. The energy spectrum of the surface of the coating shows a slight peak of Ce element (Figure 3b–d), which indicates that the prepared coating is a quaternary Ni-Fe-Co-P alloy coating and Ni-Fe-Co-P-CeO2 composite coating. Figure 5 shows the mass fraction of P element in the coatings of different nano-CeO2 particles obtained by EDS analysis. It can be seen that the mass fraction of P element increases first and then decreases with the increase of the concentration of nano-CeO2 particles, and when the concentration of nano-CeO2 is 1 g/L, the maximum value is 3.40%. Since P element will be enriched and hydrolyzed on the surface of the electrolyte to form hypophosphite, a phosphorus-rich film is formed between the coating and the interface of the corrosive medium to make the nickel-based coating exhibit high corrosion resistance. Adding an appropriate number of nano-CeO2 particles to the plating solution increases the P content in the coatings. The increase of the P element content shortens the film formation time of the phosphating film on the surface of the coatings, and also increases the thickness of the phosphating film, which contributes to the improvement of the corrosion resistance of the coatings [19,20].

Figure 6 shows an elemental view of the surface of the Ni-Fe-Co-P-1 g/L CeO2 composite coating, wherein the Ce element diagram (Figure 6f) represents nano-CeO2 particles, and it can be seen that the alloying elements and the nano-CeO2 particles are uniformly distributed on the surface of the plating layer. Studies have shown that the uniform distribution of elements and particles is due to the improved corrosion resistance of the coating.

**Figure 5.** P mass fraction of coatings with different concentration of nano-CeO2 particles.

**Figure 6.** Elemental surface mapping of Ni-Fe-Co-P-1 g/L CeO2 composite coating:(**a**) the SEM image of the analyzed surface (**b**) Ni content; (**c**) Fe content; (**d**) Co content; (**e**) P content; and (**f**) Ce content.

#### *3.2. Plating Phase Structure*

Figure 7 is an XRD pattern of the coating obtained by X-ray diffraction test. It can be seen that the coating is a typical amorphous structure, and there is a significant diffuse scattering broadening peak (Ni (110)) between 42◦ and 48◦ in 2θ. The peak width of the diffraction peak of the nanocrystalline alloy coating did not produce obvious changes, and peak intensity changes were not obvious, indicating that the nano-CeO2 particles did not obviously change in the phase structure of Ni-Fe-Co-P coating. For the nickel-phosphorus coating, the crystal structure depends mainly on the P element content in the coating. The authors of [21] have shown that when P content is lower than 5%, it is usually crystalline structure, and when P content is higher than 6.5%, it becomes amorphous structure. In this test, since the prepared plating layer is a quaternary alloy plating layer, and the atomic structure, size, and electronegativity of Ni, Fe, Co, and P elements are largely different, the amorphous forming ability is enhanced. Therefore, the plating layer is still amorphous when the P content is low. It is generally believed that the amorphous coating has better corrosion resistance due to the absence of local electrochemical potential difference between crystal grains and grain boundaries in the crystalline coating [18].

**Figure 7.** XRD patterns of coatings with different concentration of nanometer CeO2 particles.

#### *3.3. Tafel Polarization Curve*

Figure 8 shows the polarization curves of the composite coatings in the 50 g/L NaCl solution. The corrosion parameters obtained by Cview 2 software and polarization curve epitaxy are shown in Table 4. It can be seen from Figure 8 that the anodic polarization process of the composite coating is hindered and a significant passivation behavior occurs, and the composite coating is obtained when the concentration of nano-CeO2 particles in the plating solution is 1 g/L. The passivation zone is significantly larger than the remaining composite coating. It can be seen from Figure 8 and Table 4, compared with the polarization curve of pure Ni-Fe-Co-P alloy coating, that the polarization curve of composite coating prepared by co-deposition of a certain number of nano-CeO2 particles by scanning electrodeposition technology moves up and left as a whole. With the increase of the concentration of nano-CeO2 particles in the plating solution, the self-corrosion potential is continuously shifted, and the corrosion current density is gradually reduced. When the concentration of nano-CeO2 particles is 1 g/L, the prepared Ni-Fe-Co-P-CeO2 composite coating has the most positive self-corrosion potential (−0.19372 V) and the minimum corrosion current density (1.5375 <sup>×</sup> <sup>10</sup>−<sup>5</sup> <sup>A</sup>·cm<sup>−</sup>2). While continuing to increase the concentration of nano-CeO2 particles, the corrosion potential is negatively shifted, and the corrosion current density is significantly increased, indicating that corrosion resistance has begun to decline. According to the principle of corrosion electrochemistry, the larger the corrosion potential is, the smaller the corrosion current density is, the smaller the corrosion tendency of the material

is, and the better the corrosion resistance is. Therefore, the concentration of nano-CeO2 particles is 1 g/L. The Ni-Fe-Co-P-CeO2 composite coating has the best corrosion resistance. Studies have shown that anodic polarization can slow metal corrosion, and the degree of anodic polarization directly affects the speed of the anode process [22]. Compared with the pure Ni-Fe-Co-P coating, the addition of nano-CeO2 particles increases the hindrance of the corrosion process of the nickel-based coating. The Ba and Bc of the polarization curve of the composite coating are increased compared with the coating of the undoped nano-CeO2 particles; especially, the blocking effect (Ba) of the anode is more significant. When excessive nano-CeO2 particles are added to the plating solution, too much rare earth oxide adsorbs on the surface of the substrate, hindering the adsorption of Ni, Co, and Fe element on the surface of the substrate, which hinders the deposition of particles, which is not conducive to the plating. The formation of its corrosion resistance has been weakened.

**Figure 8.** Polarization curves of coatings with different concentrations of nanometer CeO2 particles.


**Table 4.** Composition of plating solution.

#### *3.4. Analysis of Electrochemical Impedance Spectroscopy*

In order to further explore the mechanism of electrochemical corrosion of Ni-Fe-Co-P-CeO2 composite coating, more corrosion kinetic information is obtained. The AC impedance analysis of the composite coating is performed under open circuit potential. The electrochemical impedance spectrum obtained in Figure 9 is shown. The Nyquist diagram of the composite coating (Figure 9a) shows a single capacitive reactance arc characteristic, and the Bode diagram (Figure 9c) has only one peak, indicating that the time constant is 1, and the electrode reaction process is mainly affected by the charge [23]. The transfer effect also indicates that the corrosive medium only contacts the interface of the coating and does not penetrate into the surface of the substrate due to diffusion. It can be seen from the phase angle curve of the Bode diagram (Figure 9c) that the maximum phase angle of the Ni-Fe-Co-P-CeO2 composite coating is higher than that of the pure nickel-based coating (56.379). From the impedance curve (Figure 9b), the impedance modulus of the composite coating doped

with nano-CeO2 particles is higher than that of the undoped nano-CeO2 particles throughout the scanning frequency interval. This shows that the corrosion resistance of the Ni-Fe-Co-P alloy coating is effectively improved by co-depositing nano-CeO2 particles. It can also be seen from the Nyquist diagram of Figure 9 (Figure 9a) that the radius of the capacitive reactance of the Ni-Fe-Co-P-CeO2 composite coating is much larger than that of the Ni-Fe-Co-P alloy coating. When the concentration of nano-CeO2 particles is 1 g/L, the radius of the capacitive anti-arc is the largest, and the radius of the capacitive anti-arc is used as the characterization of the corrosion resistance of the coating. The larger the radius, the greater the resistance of charge transfer and the harder the corrosion reaction. This result shows that the Ni-Fe-Co-P-CeO2 composite coating has better corrosion resistance. The AC impedance spectrum is modeled by the equivalent circuit diagram shown in Figure 10 and fitted by Zview software. The obtained fitting data is shown in Table 5. In the equivalent circuit diagram, Rs is the resistance in the solution. Rp is a charge transfer resistor, CPE is a constant phase angle element, and its impedance is

$$Z = 1/\text{Y}\_0(\text{j}\omega)^{-n}$$

its type has two parameters: constant Y0, its dimension is <sup>Ω</sup>−1·cm−2·s<sup>−</sup>n; parameter n, dimensionless index. When n = 1, the CPE component is the ideal capacitor. When n = 0, the CPE component is pure resistance, and in the actual solution, n is between 0 and 1 [7]. Obviously, with the addition of nano-CeO2 particles, the charge transfer resistance of the composite coating increases first and then decreases but the charge transfer resistor (Rp) of the doped nano-CeO2 particles is always larger than that of the pure nickel-based coating, and the corrosion resistance is extremely high great improvement. When the concentration of nano-CeO2 particles in the plating solution is too large, the nano-CeO2 particles are easily agglomerated, and the inclusions formed are increased, resulting in loose coating structure, and the strengthening effect of the nano-CeO2 particles is weakened.

**Figure 9.** *Cont*.

**Figure 9.** Alternating current impedance method EIS of coatings with of coatings with different concentrations of nanometer CeO2 particles: (**a**) Nyquist diagram, (**b**) Bode diagram—impedance curve, and (**c**) Bode diagram—phase Angle curve.

**Figure 10.** Equivalent circuit diagram.


**Table 5.** Equivalent circuit diagram parameter value.

#### *3.5. Surface Morphology after Corrosion of the Coating*

The SEM photographs in Figure 11 (image type is SEI) shows the surface topography of the composite coating after corrosion. It can be seen that after 5 days of etching in 50 g/L NaCl solution, many micro-protrusions of different sizes appear on the surface of the coating, and the surface appears more frequently black corrosion product. A large number of narrow and shallow microcracks extend along the boundaries of the cell structure. The degree of corrosion of the composite coating with different concentrations of nano-CeO2 particles in the plating solution is different. Among them, the coating of undoped nano-CeO2 particles (Figure 11a) is most corroded, and a large amount of corrosion product is deposited on the surface. When the concentration of nano-CeO2 particles is 1 g/L (Figure 11c), the coating is the least corroded and has a stronger retarding effect on corrosive media.

Generally speaking, the corrosion of metal in NaCl solution is mainly due to the presence of Cl−, the Cl− radius is small, the penetrating ability is very strong, and the adsorption is unevenly in the vicinity of the boundary and the impurity, so that the local dissolution is dominant, and pitting micropores are formed. Even if the surface has a passivation film formed by the metal, Cl− can form a soluble compound with the cation of the passivation film, destroying the dynamic balance of dissolution and repair of the passivation film, causing the passivation film to be gradually eliminated and continue to be corroded, resulting in etching. The deepening of the hole can quickly become a corrosion pit. When sprayed electrodeposition is used to prepare Ni-Fe-Co-P alloy coating, it is always accompanied by hydrogen evolution reaction, which retards the discharge deposition of Ni, Co, and Fe elements, which leads to the formation of pinholes or pits on the surface of the coating. In addition, the transition elements are sprayed. During the electrodeposition process, the amount of hydrogen absorption is large. The hydrogen atoms that penetrate into the cell by diffusion will cause distortion of the cytoplasm, forming a large internal stress, and stress corrosion occurs during the corrosion process. After the corrosion, the surface of the plating layer is easily cracked [24]. Since the deposition process of the sprayed electrodeposited Ni-Fe-Co-P alloy coating is a process of uneven reduction and accumulation of Ni, Co, Fe, and P, the atomic size is different and the arrangement is different, which can only be disorderly stacked and reflected to the plating layer [25]. On the top, the cell material is relatively dispersed, which provides conditions for the diffusion of corrosive media, but the corrosion condition after the addition of nano-CeO2 particles is improved. The reason is: first, the filling effect of the nano-CeO2 particles between the coating boundaries makes the structure of the composite coating is more uniform and dense, the porosity is greatly reduced, and the rare earth elements have strong affinity with impurity elements such as O and H, and these impurity elements can be wrapped to form a rare earth composite phase, and the surface of the coating is dispersed. The composite phase coverage causes the permeation channel of Cl− ions to be effectively intercepted, thereby enhancing the corrosion resistance of the composite coating [26]. Secondly, because the potential of nano-CeO2 particles is in Ni, Co, and Fe metals, it is easy to form microscopic galvanic cells at the interface between nano-CeO2 particles and nickel-based alloy. Nano-CeO2 particles are used as the cathode, and Ni, Co, and Fe are the anode. This galvanic reaction changes the coating from local spot corrosion to uniform corrosion, which helps to slow down the corrosion. When the concentration of nano-CeO2 particles is too large, the metal ions are precipitated in a large amount as a complex, and the amount of precipitation increases, but the actual deposition rate decreases, and the corrosion tendency of the coating increases [27].

**Figure 11.** Surface morphology after plating corrosion: (**a**) Ni-Fe-Co-P; (**b**) Ni-Fe-Co-P-0.5 g/L CeO2; (**c**) Ni-Fe-Co-P-1 g/L CeO2; and (**d**) Ni-Fe-Co-P-1.5 g/L CeO2.

#### **4. Conclusions**

In this paper, a Ni-Fe-Co-P-CeO2 composite coating was prepared using the scanning electrodeposition technique. To explore the impact of the concentration of nano-CeO2 particles in the plating solution on the micro morphology, structure, and composition of the coating, and to study the strengthening mechanism of nano-CeO2 particles on the electrochemical corrosion behavior of Ni-Fe-Co-P alloy coating, the following conclusions were drawn:


**Author Contributions:** X.F. and S.D. designed the experiments; W.M., Q.W., and S.D. performed the experiments and analyzed the data; X.F. contributed reagents and materials. S.D. and W.M. wrote the paper; and X.F. and J.L. provided corrections on the original draft.

**Funding:** This research was funded by the China Postdoctoral Science Foundation (Grant number 2017M621665), the Postdoctoral Science Foundation of Jiangsu Province of China (Grant number 2018K022A), and Postgraduate Research and Practice Innovate Program of Jiangsu Province (Grant number SJCX19\_0145).

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

#### **References**


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