Investigation of the Protection Performance of Mg and Al Coated Copper in High Temperature or Marine Environments

Abstract

Copper is commonly used in many applications such as electric wires, industrial machinery, boat propellers and heat exchangers. Although Cu is resistant to corrosion in ambient conditions, it is highly susceptible when exposed to harsh environments such as high-temperature air or marine atmospheres. The application of magnesium or aluminum coatings on Cu can efficiently smooth this disadvantage without affecting its primer properties. In the herein investigation, the deposition of these coatings was accomplished by pack cementation, a simple electrochemical technique which can be easily incorporated in the industrial scale. The final coatings were tested under high temperature environments by non-isothermal (dynamic) and isothermal thermogravimetric measurements, and in artificial sea water (3.5 wt.%-NaCl) by electrochemical corrosion. The structure of both coatings was compact with elevated thickness. The oxidation and corrosion measurements revealed that the coated samples had much better performance compared to the uncoated copper coupons. Furthermore, Al coatings were found to have the best performance when exposed in high temperature environment.

1. Introduction

Copper, due to its characteristics, is one of the most common materials for a variety of applications such as electrical wires, industrial machinery, boat propellers and heat exchangers [1,2]. The wide use of copper can be attributed to its properties such as excellent thermal and electrical conductivity, ductility and malleability. Due to its importance, the corrosion of this metal has been studied for a long time. Copper’s corrosion resistance is attributed to the formation of a protective film of cuprous oxide (Cu₂O and CuO) [3]. This film retards both the anodic dissolution of the metal and the rate of O₂ reduction, meaning that any failure is associated with the deterioration of that film, depending on the environmental conditions. Particularly, copper is more susceptible to the presence of chlorides, sulfur and applied fields [4]. Furthermore, a typical copper corrosion could be caused by the presence of the organic decomposition products on the surface leading to the formation of carboxylic acid [5].

The application of protective coatings is an ideal method for enhancing the corrosion resistance of copper against different corrosion types—environments—(e.g., high temperature oxidation, corrosion in seawater). These coatings can be of a polymer type [4,6,7] or metallic made from zinc, chromium, magnesium and aluminum. Magnesium and aluminum are two good candidates that can be used for the corrosion and oxidation protection of copper [8]. Magnesium is preferred as coating material on copper components because it does not affect significantly the beneficial electric properties of Cu, whereas aluminum coatings have good oxidation resistance at high temperatures due to the formation of a thin Al₂O₃ layer which is an efficient diffusion barrier [9]. Up to now the formation of Cu-Mg and Cu-Al phases are investigated either in form of coating or in form of bulk alloy. Suwwan de Felipeet et al. [10] refer the co-deposition of both phases by magnetron sputtering, which improved the high temperature corrosion without affecting significantly the Cu ion mobility. The same conclusion is reported by Arcot et al. [11], who found that the Mg deposition with Cu had low electric resistivity, which makes it ideal for wire protection. Zhu et al. [12] examined the effect of Mg/Al alloying of copper by plasma arc melting and found that both elements improve the oxidation resistance up to 400 °C, with the Al alloyed samples having the best performance. Zhao et al. [13] investigated the Mg implantation on Cu by a high current ion source and found improvement of oxidation resistance at 300 °C without any noticeable influence on the electric conductivity. Aluminizing of copper is referred to have been accomplished by high temperature diffusion procedures [14,15,16,17], Cathodic arc plasma and sputtering [18,19] or hot pressing [20]. Al was found to provide significant long-term protection to copper up to 800–900 °C, depending on the Al concentration, mainly by the formation of Al₂O₃.

A promising, simple and low-cost technique for the formation of Cu-Mg and Cu-Al phases is chemical vapor deposition with pack cementation process, as the energy consumption during pack cementation is much lower than other techniques and it is environment-friendly due to the lack of polluting byproducts. Moreover, a major asset of this procedure is the formation of homogeneous coatings over irregularly shaped objects [21,22].

Pack aluminization of copper is reported only by Chiang et al. and Shi et al., who deposited Al at high temperatures (over 800 °C) and for holding times over 4 h. Magnesium deposition is mainly referred by physical vapor procedures [10,11,12,23]. Furthermore, the oxidation and corrosion performance of the final products is not investigated in a wide range of temperatures or in marine environments. The aim of this study is to investigate the deposition and structure of magnesium and aluminum coatings on copper by pack cementation at low temperatures and holding times, and examine the high temperature oxidation and corrosion performance of the final products by thermal analysis and electrochemical corrosion, respectively. The characterization of the coatings was performed using Scanning Electron Microscopy (SEM) and X-ray diffraction (XRD) analysis, while high temperature oxidation and corrosion tests were examined by dynamic/isothermal oxidations and exposure in simulated seawater solution.

2. Materials and Methods

The substrates that were used for the experimental process had dimensions 6 mm × 4 mm × 2 mm and were cut from C10800 copper alloy. Before the coating process, the substrates were suitably prepared by ultrasonic cleaning and mechanical polishing using up to 1000-grid SiC paper. Finally, the substrates were thoroughly cleaned in alcohol solution. The as prepared substrates were buried in a powder mixture consisting of 25 and 28 wt.% magnesium (Mg) and aluminum (Al) powder, respectively, 3 wt.% ammonium chloride (NH₄Cl) and balance aluminum oxide (Al₂O₃). These mixtures were found to be optimum for such coatings as reported in previous published work [24,25]. Mg/Al was the donor material, NH₄Cl the halide activator and Al₂O₃ the inert filler. The mixture with the substrate was then packed and sealed in ceramic crucibles and inserted in an Ar purged furnace preheated at the deposition temperature. The experimental conditions for the deposition procedures were: temperature of heat treatment 550 °C and holding time 2 h for both Mg and Al deposition. After the 2 h coating the system was left to cool down to ambient temperature without interrupting the Ar flow. Afterwards, the coated samples were cleaned from the remaining powder on their surface and metallographically prepared for microscopic observation by mechanical polishing using up to 2000-grid SiC paper and Al₂O₃ paste (5 μm), and extensive cleaning in alcoholic solution.

The coated samples were examined by two different tests which corresponded to high temperature and marine corrosive environments. The examination of the high temperature oxidation resistance of the specimens was accomplished through dynamic thermogravimetric measurements (TG) from ambient temperature up to 1000 °C carried in a TG-DTA SetaramSetsys 16/18 system (SETSYS, Lyon, France), with heating rate 10 °C/min and air flow 50 mL per minute. Furthermore, isothermal thermogravimetric measurements were also performed at selective temperatures in order to evaluate the durability of the samples. Four different oxidation temperatures were implemented for each coating: 550, 600, 650 and 700 °C for Mg coatings and 800, 850, 900 and 950 °C for Al coatings, and the holding time was6 h. During these tests, the mass gain vs. temperature, for dynamic oxidations, and the mass gain vs. time, for isothermal oxidations, were recorded for the evaluation of the samples.

The performance in marine atmosphere was evaluated through electrochemical corrosion experiments using a conventional three-electrode cell in 3.5 wt.% NaCl aqueous (distilled water) solution at room temperature (~25 °C). The reproducibility of the electrochemical test was verified by carrying out three measurements for each sample. Prior to potentiodynamic polarization, the working electrode was immersed into the test solution at open circuit potential (OCP) for 60 min to attain a steady state condition. Potentiodynamic polarization measurements were performed from −300 mV vs. OCP to +500 mV vs. OCP with a scanning rate of 0.5 mV/s. The current density values were normalized to the exposed surface area. The corresponding mean values were used in the relative diagrams.

The microscopic examination of the samples was performed by Scanning Electron Microscopy (SEM) using a 20 kV JEOL 840A scanning electron microscope (SEM, Jeol Ltd., Akishima, Tokyo, Japan) equipped with an OXFORD INCA Energy Dispersive X-ray Spectroscopy (EDS) analyzer (Oxford Instruments Nanoanalysis, Oxford, UK) and the INCA 18D 4.15 software that enabled point microanalysis and linear microanalysis of the area of interest. The phase identification was accomplished by X-ray diffraction (XRD) using 2-circles Regaku Ultima+ diffractometer (CuKα radiation, 40 KV, 30 mA) with Bragg–Brentano (BB-XRD, Rigaku Corporation, Shibuya-Ku, Tokyo, Japan).

3. Results and Discussion

3.1. Structure of Mg and Al Coatings

In Figure 1a, the cross section of the Mg coating deposited at 550 °C for 2 h is presented. The main remarks are that the coating has significant thickness of approximately 25 μm. While it is characterized by high homogeneity, as no cracks and porous areas are detected, the coating has two distinguishing layers. EDS line scanning analysis (Figure 1b) on these layers revealed that the inner layer contains 16 wt.% Mg, while the outer layer contains 44 wt.% Mg. Moreover, the thickness of the inner layer is 5 μm and the outer layer has an average thickness of 20 μm.

XRD analysis identified the existence of two distinct phases corresponding to the Cu₂Mg and to the CuMg₂ (Figure 1c). Combining these results with those of the EDS analysis, it is concluded that the inner layer is Cu₂Mg and the outmost layer is the CuMg₂ phase. This layer topography is also confirmed from the results of Ref. [24], in which it was found that for very short deposition time (30 min) the coating had one single layer corresponding to Cu₂Mg; for longer holding times, the CuMg₂ phase appeared over Cu₂Mg. The formation of these phases is the result of a two-step chemical procedure. In the beginning, Mg diffuses in the copper substrate which results to the formation of the MgCu₂ phase. The MgCu₂ increases as Mg diffusion continues through the Cu₂Mg. Moreover, some of the Mg ions do not reach the copper substrate but enrich the upper areas of the Cu₂Mg phase and transform it to CuMg₂. The rate of the formation of CuMg₂ increases as the thickness of the MgCu₂ grows, because the Mg ions have to traverse a bigger distance to reach Cu. After a particular point, the formation of extra MgCu₂ amounts is insignificant and that is the reason the CuMg₂ layer has bigger thickness than the inner Cu₂Mg layer.

The same characterization procedure was followed for the Al coated samples. In Figure 2a the coatings formed at 550 °C for 2 h is presented. In this case, a single phase 12 μm thick film is created, with similar homogeneity as the Mg coating as no cracks or pores were observed in its structure. EDS point microanalysis showed that the average Al concentration in the coating is 16 wt.% and the Cu is 84 wt.%. One single phase was traced in the XRD diffraction pattern corresponding to the Al₄Cu₉ (Figure 2b). Concluding from all the structural results, thick and homogeneous coatings were formed in every case. Furthermore, the phases that were detected in both cases are reported to have little effect on the electrical properties of Cu, thus, they can be effectively used in electrical applications [11].

In the following section, these coatings are investigated for their corrosion and oxidation performance in order to be evaluated for applications working under harsh high temperature or marine environments.

3.2. High Temperature Oxidation

3.2.1. Dynamic Oxidation

From the non-isothermal oxidation of copper, it is revealed that the uncoated copper begins to oxidize at 200 °C, as from this temperature and on there is a considerable increase of the mass gain due to the oxide formation (Figure 3). After 400 °C the oxidation rate is high and continues to increase up to the end of the test (approx. 1000 °C).

In the same mass gain diagram, the Mg-coated sample does not oxidize until 560 °C (Figure 3b). This protection is attributed to the spontaneous formation of a thin and dense MgO oxide film (20–50 nm) on the surface as a result of the atmospheric air exposure [26,27]. Moreover, Mg(OH)₂ is also formed, as a result of the high affinity of Mg to the atmospheric humidity, on the surface of the samples. At 350 °C a drop of mass gain occurs (Figure 3c), which is due to the decomposition of Mg(OH)₂. For temperatures over 380 °C the Mg(OH)₂ destabilizes [26] and is totally transformed to amorphous MgO following Equation (1):

$$\mathrm{Mg}{\left(\mathrm{OH}\right)}_2\to \mathrm{MgO}+{\mathrm{H}}_2\mathrm{O}$$

(1)

This amorphous scale gradually grows with the increase of temperature and acts as a barrier to the aggressive environment up to 460 °C. From that point it begins to be less protective, and a low-rate oxidation of the underlying coating begins. Over 560 °C, and especially over 600 °C, a dramatic mass gain is observed as a result of the high-rated oxidation of the sample. This change is mostly attributed to the residual stresses created in the oxide scale, which can provoke the creation of cracks in their structure. These stresses result from the difference between the crystal structure of the oxide and the coating material and from the growth mechanism of the scale based on the theory of Pilling-Bedworth ratio (PBR) for metals [28,29]. The PBR ratio in this case is:

$${\mathrm{PBR}}_{\mathrm{metal} \mathrm{alloy}} = \frac{\mathrm{Volume} \mathrm{of} \mathrm{a} \mathrm{mole} \mathrm{of} \mathrm{MgO}}{\mathrm{Volume} \mathrm{of} x \mathrm{atoms} \mathrm{of} \mathrm{Mg} \mathrm{in} \mathrm{the} \mathrm{alloy}}$$

(2)

When an oxide is formed on the surface of an alloy and not on pure metal, the metal with the most stable oxide is calculated for the factor V_oxide, here V_MgO. When the PBR ratio is higher than 1 then compressive stresses are created in the structure of the oxide scale, while when PBR < 1 then tensile stresses are created. The magnitude of these stresses is proportionate to the difference PBR−1.

In the herein case, the oxide formation is expressed from the chemical equation:

$$2\mathrm{Mg}+{\mathrm{O}}_2\to 2\mathrm{Mg}\mathrm{O}$$

(3)

The volume of 1 mole of MgO is:

$$V_{\mathrm{MgO}}=\frac{W}{d}=\frac{40.03 \mathrm{g}}{3.58 \raisebox{1ex}{$\mathrm{g}$}\!\left/ \!\raisebox{-1ex}{${\mathrm{cm}}^3$}\right.}=11.181 {\mathrm{cm}}^3$$

(4)

where, W is the weight and d the density of 1 mole of oxide, respectively. Considering that there is no Cu-rich area formed underneath the MgO oxide, the volume of Mg moles in the alloy which oxidize forming MgO can be calculated from the expression [28]:

$$\mathrm{Volume} \mathrm{of} x \mathrm{moles} \mathrm{Mg} \mathrm{in} \mathrm{alloy}=\left(\frac{x}{u+v}\right)\times N\times V_{{\mathrm{Mg}}_2\mathrm{Cu}}$$

(5)

where, x is the moles of Mg required for the formation of MgO (Equation (3)) which is 2, N is Avogadro′s number and u, v are the number of moles of every element in the Mg₂Cu alloy (surface layer of the as deposited coating), thus u = 2 and v = 1. From the JCPDS card #651116 and given that Mg₂Cu belongs to the orthorhombic Fddd system, the volume of the Mg₂Cu unit cell is:

$$\begin{gathered} V_{\mathrm{cell}}=\frac{a\times b\times c}{4}=\frac{0.904 \times {10}^{-7} \times 0.5275 \times {10}^{-7}\times 1.838 \times {10}^{-7}}{4} \\ {V}_{\mathrm{cell}}=\frac{0.876\times {10}^{-21}}{4}=0.219\times {10}^{-21} {\mathrm{cm}}^3 \end{gathered}$$

where, a, b, c: lattice constants of the unit cell.

The volume of Mg in the Mg₂Cu molecule is [28]:

$$V_{{\mathrm{Mg}}_2\mathrm{Cu}}=\frac{V_{\mathrm{cell}}}{n} = \frac{0.219\times {10}^{-21}}{4}=0.05475 \times {10}^{-21} {\mathrm{cm}}^3$$

n is the number of atoms per unit cell in Mg₂Cu (Fddd system).

From Equation (5), the volume of x = 2 moles of Mg that oxidize is:

$$\mathrm{Volume} \mathrm{of} 2 \mathrm{moles} \mathrm{of} \mathrm{Mg} \mathrm{in} \mathrm{alloy}= \frac{2\times 6.02\times {10}^{23}\times 0.05475\times {10}^{-21}}{3}=21.984 {\mathrm{cm}}^3$$

As a result, the PBR ratio (Equation (2)) is:

$$\mathrm{PBR}=\frac{11.181}{21.984}=0.5085$$

Since PBR < 1, it is concluded that tensile stresses are created at the interface between the alloyed coating (Mg₂Cu) and the oxide scale. Thus, with the increase of the oxidizing temperature, the PBR ratio decreases further and more intense tensile stresses are created. As a result, in the structure of the oxide scale more cracks are formed, providing easy paths for the aggressive agents to reach the coating and the copper substrate at longer exposure periods. The degradation of the coating occurs wherein a porous white oxide film grows over the surface, which stabilizes the oxidation rate at a lower value for the temperature range from 600 to 800 °C [12,28]. At 800 °C, the tensile stresses are so intense that it drives to the total collapse and delamination of the oxide scale, while a significant fraction of the Mg-Cu coating also oxidizes, leading to non-protective conditions of the Cu substrate.

The Al coated samples were found to be much more resistant under non-isotheral high temperature oxidation conditions. In Figure 4, the graph of mass gain vs. oxidation temperature is presented, in which it is revealed that, up to 650 °C, the sample does not oxidize as there is not any mass gain recorded. From this temperature, the oxidation rate slightly increases, until the temperature reaches 850 °C. After the latter temperature, the rate increases dramatically, which can result to the total deterioration of the coating and the substrate for long exposure periods. The stability of the coating up to 850 °C is attributed to the high affinity of Al to oxygen, which results in the formation of a very stable, thin and adhesive oxide on the top of the coating (Equation (6)). This scale acts as a barrier to the diffusion of oxygen in the coating. The grade of protection of this thin layer is relative to the Al concentration in the coating and of the oxidation temperature [16]. Thus, over 850 °C, this layer begins to destabilize; the rate of destabilization increases dramatically after 900 °C for 28 wt.% Al used in this case.

$$2\mathrm{Al}+\frac{3}{2}{\mathrm{O}}_2={\mathrm{Al}}_2{\mathrm{O}}_3$$

(6)

Comparing with the bare copper substrate, it is concluded that the Al coating provides superior protection to the substrate which can safely work under high-temperature aggressive environments.

Comparing the two types of coatings, the protection provided by the Al coatings is much better, as the samples with Mg coatings begin to oxidize 360 °C higher than Cu bare samples, while Al coated ones begin to oxidize 450 °C higher than Cu. As a conclusion, Al-coated components can be safely used up to 800–850 °C while Mg deposited components can be used in more moderate temperature, up to 450–500 °C. In any case, both limits can serve most of the application referred in the introduction.

3.2.2. Isothermal Oxidation

In order to examine the durability of both coatings under prolonged exposure periods and make a complete examination of the oxidation by thermal analysis, the samples were examined under constant temperatures for 6 h.

The Mg-coated sample was held under 550, 600, 650 and 700 °C dry air environment. As concluded from the dynamic oxidation of this sample in Figure 3, 550 °C is the temperature at which the oxidation was recorded at very low rate; 600 °C corresponds to the middle rate oxidation before the steep increase of the mass gain in the mass gain; and finally 650 and 700 °C are oxidation temperatures corresponding to oxidation conditions after this temperature point.

The results of the isothermal oxidations are illustrated in Figure 5 in the diagram of the mass gain vs. time (t). It is revealed that from the lowest oxidation temperatures (550, 600 °C) the samples have similar behavior; for the first minutes of exposure a linear increase of the mass gain is observed and afterwards the mass gain makes a plateau at which the oxidation rate is insignificant and the coating is highly protective to the substrate. For the higher oxidation temperatures (650, 700 °C) the rate of the mass gain is also linear in the first minute of exposure but afterwards this rate decreases but remains linear up to the end of the exposure but at a considerably lower rate. For further examination of the curves, the diagrams of mass gain vs. t^0.5 are presented in Figure 6a,b.

From these graphs, it is concluded that, ignoring the first minutes of exposure, the oxidation rate can be described by a linear fitting with different slope for 550/600 °C and for 650/700 °C. As a result, the mass gain (Dm) and so oxidation follow the parabolic law vs. time (t), which is [30,31]:

Dm² = k_p × t

(7)

k_p: is a rate constant which incorporates the rest of the factors affecting the phenomenon (see Equation (8)).

The first ignored minutes of oxidation correspond to the formation of a protecting oxide scale, which provides passivation on the surface. After this step, the oxidation proceeds by ion diffusion in the coating. Thus, the constant of the mass gain rate (k_p) can be calculated from this step of the test (linear part of Figure 6). The results are presented in

Table 1. Values of the constant of the mass gain rate (k_p) of Mg coated samples for 550, 600, 650, 700 °C.

T/°C	${\mathit{k}}_{\mathbf{p}}/\mathbf{\Delta }\mathit{m}\times {\mathit{S}}^{-1}\times {\mathit{t}}^{-0.5}$
550	41.3 × 10−4
600	82 × 10−4
650	205.2 × 10−4
700	246.9 × 10−4

.

The calculation of the activation energy of the reactions was realized by the use of Arrhenius law, which is expressed by the equation:

$$k_{\mathrm{p}}=k_0 \times e^{\left(-\frac{E}{RT}\right)}$$

(8)

where, k₀ is the pre-exponential factor, E is the apparent activation energy, R is the gas constant and T is the absolute temperature.

Thus, the activation energy E was calculated from the slope of the ln[k_p(T)] vs. 1000/T (Figure 7) using the correlations [32,33,34] Dm² vs. t^0.5. The activation was finally calculated 86.2 kJ/mol.

The Al coated sample was held under 800, 850, 900 and 950 °C in dry air environment. Following the same way of thinking with Mg coated samples and concluding from Figure 4, 800 °C is the temperature at which the oxidation was recorded at very low rate. Further, 850 °C corresponds to the middle rate oxidation. Finally, 900 °C corresponds to the critical temperature after which the oxidation rate increases dramatically, and at 950 °C the samples have significant oxidation rate. The graphs corresponding to isothermal measurements are presented in Figure 8.

From these graphs it is concluded that, at the lowest temperatures (800, 850 °C), the curves have approximately similar behavior; at the first steps of exposure the mass gain increases linearly, and afterwards it tends to a plateau and the rate of oxidation stabilizes to a very low constant rate. At higher oxidation temperatures (900, 950 °C) the mass gain rate is also linear and has much longer duration comparing with 800 and 850 °C curves. Afterwards, the rate decreases but still has a significant value. Furthermore, the curves do not seem to have the tension of making any plateau. The behavior of these last curves can be attributed to the too high oxidation temperature, which is detrimental for any protective scale formed on the surface of the samples.

In Figure 9 the diagrams of the mass gain vs. t^0.5 are presented and it is equally revealed that, ignoring the first time-limited step of oxidation, the remaining phenomenon follows the parabolic law of Equation (7) as a result of ion diffusion.

The constant of the mass gain rate (k_p) was calculated from the linear part of each graph of Figure 9 and is presented in

Table 2. Values of the constant of the mass gain rate (k_p) of Al coated samples for 800,850, 900, 950 °C.

T/°C	${\mathit{k}}_{\mathbf{p}}/\mathbf{\Delta }\mathit{m}\times {\mathit{S}}^{-1}\times {\mathit{t}}^{-0.5}$
800	185.2 × 10−4
850	261.2 × 10−4
900	300.6 × 10−4
950	385.7 × 10−4

.

Activation energy E was calculated from Equation (2) and from the slope of the ln[k_p(T)] versus 1000/T graph (Figure 10). The activation was finally calculated at 52.105 kJ/mol.

The low value of the activation energy of both coatings is mainly attributed to the high affinity of Mg and Al to oxygen. However, the chemical result of this reaction is the creation of protective scales, which hinders the diffusion of oxygen in the underlying coating and substrate, assuring safe, long-lasting service life compared to the bare Cu samples. Comparing the two coatings, it can be assessed that Al coatings have superior protection from Mg because, in dynamic oxidation, it begins to oxidize at much higher temperature and with significantly lower rate. Furthermore, in isothermal oxidations Al coatings have approximately same mass gain with Mg ones but at a temperature 250 °C higher.

3.3. Electrochemical Corrosion

The samples were also examined under electrochemical corrosion in order to estimate their durability in marine atmospheres under the effect of oxygen and chlorine ions.

The open circuit potential (OCP), E_OCP, of all samples in 3.5 wt.% NaCl aqueous solution after a period of 60 min are shown in Figure 11. It can be seen that, compared to pure copper potential, the coated samples have less electronegative E_OCP as reported in

Table 3. OCP (E_OCP), corrosion potentials (E_corr), corrosion current densities (i_corr), Tafel slopes and corrosion rates of uncoated and Mg, Al coated samples in 3.5 wt.% NaCl aqueous solution at 25 °C (mean values with standard deviation for potentials <7% and for current densities <5%).

Type of Sample	EOCP (mV vs. SCE)	Ecorr (mV vs. SCE)	icorr (μΑ/cm2)	βα	−βc	Corrosion Rate (mpy)
Pure copper	−200	−170	3	46	53	1.38
Mg coat.	−210	−220	0.9	58	56	0.9
Al coat.	−311	−330	1	83	80	0.82

. These differences are related to the different nature of the materials (Cu, phases of Mg or Al) directly exposed to the solution. It should be noted that the high homogeneity of the coatings and the absence of cracks and porous areas in their structure is a first indication of the protection of copper from the corrosive environment.

The potentiodynamic polarization diagrams of the uncoated and coated samples are presented in Figure 12. In the cathodic part of the polarization curves, two regions can be clearly observed. The first one represents the reduction of dissolved oxygen in the solution and the second one the reduction of hydrogen. The corrosion potentials (E_corr), the corrosion current densities (i_corr) and the anodic/Cathodic Tafel slopes are reported in Table 3. The differences between E_OCP and E_corr values probably indicate a chemical change at the working electrodes surfaces as a result of alkalization from cathodic reactions occurring during the potentiodynamic scan [35]. These parameters were estimated by taking into consideration several zones from 50 to 200 mV to optimize the fitting of the Tafel slopes. Additionally, in Table 3 the corrosion rates of the samples are reported calculated [33,36] by Equation (9):

Corrosion rate = (0.13 × I_corr × E.W.)/d

(9)

where, E.W. is the equivalent weight of the oxidized element in g (63.5/2 Cu, 26.98/3 Al and 24.3/2 Mg) and d is the density of the corroded specimens in g/cm³.

From the results it is concluded that the corrosion rates follow the same trends as i_corr. The corrosion rates of the coated samples are lower than that of the copper uncoated samples. Apart from that, by examining the anodic regions of the potentiodynamic curves (Figure 12), it is revealed that the current of the uncoated samples increases rapidly after E_corr.

For the Mg coatings and pure copper samples, a passive-like region, where current densities are reduced, can be observed (Figure 12) due to the formation of any corrosion products like Mg(OH)₂ and Cu₂O or CuO on them, whereas the active dissolution for the Al coatings continues approaching a stable value of current density (around 10⁻⁴ A/cm²) due to the formation of Al₂O₃ on the surface [37,38]. Both coated samples are found to have similar corrosion resistance and relative significantly higher than the unprotected copper coupons.

4. Conclusions

From the results of this scientific examination, it is concluded that the structure morphology of Mg and Al coatings deposited by pack cementation, is compact with no cracks and areas with pores. For the case of Mg deposition, the coating has two distinct layers corresponding to MgCu₂(inner layer) and Mg₂Cu (outer layer). For the case of Al deposition, the coatings have one single layer corresponding to the Al₄Cu₉ phase. The Al coatings were found to be remarkably more resistant in high temperature air environments than Mg coatings. However, both coatings are much more durable than the unprotected copper substrate. From the corrosion tests corresponding to marine environment, it is concluded that both Al and Mg coatings have similar and elevated corrosion resistance as compared with copper substrate, and can offer long-term protection when working under similar conditions.