Corrosion potentials and corrosion current densities obtained by polarization curve measurements in 3.5 mass% NaCl aqueous solution for films prepared by various methods.
This chapter introduces a novel, chemical-free “steam coating” method for preparing films on magnesium (Mg) alloys and assesses their effectiveness in improving the corrosion resistance of two different Mg alloys. A film composed of crystalline Mg(OH)2 and Mg-Al layered double hydroxide (LDH) was successfully formed on AZ31 Mg alloy, and its corrosion resistance was evaluated through electrochemical measurements and immersion tests in an aqueous solution containing 5 wt.% NaCl. An anticorrosive film was also formed on Ca-added flame-resistant AM60 (AMCa602) Mg alloy via the same steam coating method and found to be composed of crystalline Mg(OH)2 and Mg-Al layered double hydroxide (LDH). Its corrosion resistance was also investigated, and the effectiveness of the steam coating method for improving the corrosion resistance of Mg was fully explored.
- magnesium alloy
- steam coating
- corrosion resistance
- surface treatment
- composite film
Magnesium (Mg) alloys have excellent physical and mechanical properties, including a high specific strength, excellent formability, good vibration adsorption, and high damping capacity [1–5]. However, the greatest advantage of Mg alloys is their lightweight, which has made it possible to achieve energy savings in automobiles, ships, trains, and airplanes through the use of steel-based hybrid materials. In such applications, Mg is superior to aluminum and CFRP (carbon fiber reinforced plastics) in terms of its density and specific strength but offers a much lower corrosion resistance. Various surface treatments have therefore been developed to produce protective coatings on Mg alloys through anodizing, chemical conversion, plasma electrolytic oxidation, or polymer films [6–12].
Chemical conversion has already been used with automotive Mg components because of its ease of operation and low cost, with chromate-based systems the most popular and effective option for producing a protective layer on various metal surfaces [13–15]. The use of chromate, however, is becoming increasingly more regulated due to the high toxicity of hexavalent chromium compounds . Other chemical conversion methods for imparting Mg alloys with a high corrosion resistance have been developed around using tin, rare earth salts, phosphates/permanganates [17–22], vanadium, or composite films of molybdenum-lanthanum or zinc-phosphate-calcium [23, 24].
Films produced by anodizing have also been used to improve the corrosion resistance of Mg alloys, with good results being achieved through the use of micro-arc oxidation (MAO) [25–27]. Using this method, a high anodic voltage (>200 V) is used to locally melt the surface of the Mg alloy and create a ceramic coating . Alkaline electrolytic solutions containing aluminate [28–31], silicate [32–38], or phosphate [39–45] are often used during the MAO of Mg alloys, but the need to treat the liquid waste before disposal creates a risk of environmental pollution due to heavy metal ions. Furthermore, their use may inhibit the recycling of post-use Mg product scraps into Mg ingot. There is therefore a need for an environmentally friendly surface treatment method that is capable of improving the low corrosion resistance of Mg alloys.
A chemical conversion method that creates a protective film of Mg-Al layered double hydroxide (LDH) has recently been used to significantly improve the corrosion resistance of Mg alloys [46–50]. The layered structure of the LDH consists of positively and negatively charged substances with a general molecular formula of [M1−x2+Mx3+(OH)2]x+[Ax/nn−] · mH2O , where M2+ and M3+ represent divalent and trivalent metal cations, respectively; An− is an anion such as CO32−, SO42−, or OH−; and x has a value of between 0.2 and 0.33 . This structure serves to store and release corrosion inhibitors on demand, thereby creating anion exchange between these inhibitors and chloride ions. This, in turn, results in a self-healing and corrosion-resistant film [52–54]. Tedim et al. reported that LDH nanocontainers loaded with different corrosion inhibitors such as vanadate, phosphate, and 2-mercaptobenzothiazolate (MBT) can improve active corrosion protection , with the release of MBT from LDH nanocontainers in a sol-gel film preventing degradation at the metal/film interface. They have also since developed a novel self-healing protective coating using LDH nanocontainers with corrosion inhibitors  formed directly on 2024 aluminum alloy via a conversion reaction. Montemor et al. demonstrated that using LDHs loaded with mercaptobenzothiazole in combination with cerium molybdate hollow nanospheres filled with 2-mercaptobenzothiazole results in a synergistic corrosion inhibition effect that has self-healing potential .
Though it is clear that LDHs loaded with corrosion inhibitors are effective in substantially improving corrosion resistance, it has also been reported that LDHs devoid of any corrosion inhibitor could also be effective in delaying corrosion by trapping anions such as chloride ions in LDHs [58, 59]. Lin et al., for example, have shown that the CO32− in carbonate Mg-Al LDH has Cl− anion-exchangeability in a corrosive environment and so can protect Mg alloy against corrosion . Tedim et al. have also reported that Cl− ions in a corrosive environment can be trapped in the interlayer of Zn-Al LDH intercalated with nitrate anions and that the addition of this LDH to a polymer layer drastically reduces the permeability of corrosive Cl− ions through the protective coating. This confirms that LDHs are a promising material for improving corrosion resistance , but anticorrosive films containing LDHs are usually prepared by a multistep process in which LDHs are first synthesized in powder form and then incorporated into a film for corrosion protection . For industrial applications, it is important to develop an environmentally friendly surface treatment method for achieving an anticorrosive film containing LDHs in a single step.
This chapter presents experimental results pertaining to the preparation and corrosion resistance of magnesium hydroxide films containing Mg-Al LDH on two types of Mg alloy: AZ31 and a flame-resistant Ca-added AM60 (AMCa602). A comparison is also made of the corrosion potentials and corrosion current densities obtained by polarization curve measurements in 3.5 mass% NaCl aqueous solution of films prepared on AZ31 alloy by various methods such as physical vapor deposition, chemical vapor deposition, anodization, chemical conversion, plasma electrolytic oxidization (PEO), spin coating, and a newly developed steam coating method for producing a hydroxide film containing Mg-Al LDH [10–12, 23, 24, 60–65]. As shown in Table 1, steam coating produces a relatively positive corrosion potential and the lowest corrosion current density of all the films tested, which indicates that this method imparts superior corrosion resistance. Furthermore, the direct growth of such films from the Mg alloy itself can greatly improve their adherence and mechanical stability when compared to other methods such as spin coating and dip coating .
|Treatment method||Main component in each film||Corrosion potential(V vs. Ag/AgCl)||Corrosion current density (A/cm2)||Ref.|
|PVD||Al2O3||−1.51||9.20 × 10−5|||
|PVD||TiO2||−1.50||1.29 × 10−4|||
|CVD||SiOx||−1.37||1.58 × 10−7|||
|CVD||SiOx||−1.59||4.3 × 10−7|||
|Anodization||Mg-Al LDH + MgO||−1.19||3.82 × 10−7|||
|Chemical conversion||MgO + Mg2SiO4 + Al2O3||−0.06||9.47 × 10−6|||
|Chemical conversion||Vanadium||−1.05||8.60 × 10−6|||
|Chemical conversion||Mo/La||−1.54||1.15 × 10−7|||
|PEO||Zn3(PO4)2||−1.76||1.04 × 10−7|||
|PEO||MgO + Mg3(PO4)2||−1.50||4.00 × 10−8|||
|PEO||Mg2SiO4 + CeO2||−1.20||2.00 × 10−8|||
|Spin coating||TiO2 + MgAl2O4||−1.26||4.90 × 10−6|||
2. Preparation of Mg-Al layered double hydroxide films on AZ31 Mg alloy
In this section, we introduce a method for preparing a protective magnesium hydroxide film containing Mg-Al LDH on specimens of AZ31Mg alloy (composition: 2.98% Al, 0.88% Zn, 0.38% Mn, 0.0135% Si, 0.001% Cu, 0.002% Ni, 0.0027% Fe, remainder Mg) measuring 20 × 20 × 1.5 mm. These AZ31 substrates were prepared by ultrasonically cleaning them in absolute ethanol for 10 min and then drying with inert Ar gas. Once clean, the substrates were set on a substrate stage in a 100 ml-capacity, Teflon-lined autoclave, which had 20 mL of ultrapure (resistance of 18.2 MΩ cm) located at the bottom to produce steam. The distance between the water surface and substrate stage was c.a. 3 cm. The autoclave was heated to a temperature of 423–453 K and then held at this temperature for 1–8 h before being allowed to cool naturally to room temperature. After this steam coating treatment, the samples were ultrasonically cleaned in ethanol for 10 min and dried with Ar gas.
The appearance of the samples after steam coating (Figure 1) reveals that all the substrates treated at 423–443 K were uniformly covered with film, regardless of the treatment time. However, as partial detachment of the film was observed in the case of the substrates treated at 433 K for 8 h and at 453 K for 3 and 4 h, these treatment conditions are not considered suitable for preparing an anticorrosive film. The detachment of the film was also found to become more prominent with an increase in the treatment time.
The GAXRD patterns of the films formed on AZ31 with different treatment times (423, 433, 443, and 453 K) shown in Figure 2 reveal clear peaks attributable to the Mg alloy substrate at treatment times less than 2 h with all treatment temperatures. An additional peak at around 2θ = 18° assigned to the 0 0 1 diffraction peak of brucite-type Mg(OH)2 is present in samples treated at 423 and 433 K for 2 h and at 443 and 453 K for 1 h. A diffraction peak consistent with [1 0 1] reflection at around 2θ = 38°, which has the highest intensity according to the JCPDS file number 44-1482, becomes stronger with increasing treatment time at all treatment temperatures, which suggests the formation of hexagonal Mg(OH)2. With an increase in treatment time, several reflection lines at 2θ angles of approximately 18, 33, 38, 51, 58, 62, 68, and 72° are clearly observed in all the GAXRD patterns, and these can be assigned to the 0 0 1, 1 0 0, 1 0 1, 1 0 2, 1 1 0, 1 1 1, 2 0 0, and 2 0 1 reflections of brucite-type Mg(OH)2, respectively. When the treatment time is increased to more than 3 h at any of the treatment temperatures, then one or two peaks attributable to a hydrotalcite (HT)-like structure are also observed. These are attributed to the formation of Mg1−xAlx(OH)2(CO3)x/2·nH2O (Mg-Al LDH) intercalated with carbonate anions. The two peaks at 2θ angles of around 11° and 22° were assigned to the [0 0 3] and [0 0 6] reflections of Mg1−xAlx(OH)2(CO3)x/2·nH2O, respectively (JCPDS No. 41-1428). Based on these XRD results, it is concluded that the most suitable treatment conditions for preparing a film composed of Mg(OH)2 and Mg-Al LDH are 423 K for 6–8 h, 433 K for 4–6 h, 443 K for 3–4 h, and 453 K for 2 h.
The most dominant feature of the FT-IR spectra of the samples treated at 423, 433, 443, and 453 K in Figure 3 is an intense sharp peak at 3696.9 cm−1. It has been reported that pure Mg(OH)2 exhibits a single band at 3698 cm−1 due to the high basicity of its O—H groups , and this peak was clearly presented in those samples treated at 423 K for more than 4 h, 433 K for 2–6 h, 443 K for over 2 h, and 453 K for 2 h. These results are in agreement well with their respective XRD profiles. The peak and shoulder bands observed at 950, 781, and 558 cm−1 can be attributed to Al—OH translation modes , while the bands at 1370–1520 cm−1 are attributed to symmetric and asymmetric stretching modes of CO32− in the interlayer [68, 69]. These peaks and bands could only be detected in the spectra of samples which exhibited peaks attributable to Mg-Al LDH in their XRD pattern. An additional band at 1630–1650 cm−1 was assigned to the bending mode of water molecules in the interlayer. A shoulder band at 3080 cm−1 revealed the presence of hydrogen bonding between water and CO32− in the interlayer . These results indicate that the films treated at 423 K for more than 4 h; at 433 K for 2, 4, and 6 h; at 443 K for more than 2 h; and at 453 K for 2 h were composed of crystalline Mg(OH)2 and carbonate Mg-Al LDH.
The SEM images in Figure 4 show the surfaces of samples treated at 433 K, in which we see that steam coating for 2 h led to a rougher and more nodular surface than the untreated AZ31. Several pores of up to several hundred nanometers in diameter were also observed on the surface. When the preparation time was prolonged to 4 h, granular structures were formed and the film was found to be denser than after 2 h. However, a number of pores measuring several hundred nanometers in diameter were still observed on the surface. With the sample treated for 6 h, the film was found to be denser than after 4 h of treatment, with no pores evident on the surface.
Figure 5 presents SEM images showing the sample surfaces treated at 443 K. Note that when the deposition time was less than 1 h, there was only a very low localized surface coverage of the film. However, the surface became fully covered with film when the treatment time was extended to more than 2 h, and this coverage became denser with increasing treatment time. As such, the sample treated for 4 h had a denser film than the sample treated for 3 h, which means that surface topography could be effective way of improving the corrosion resistance.
Figure 6 shows a cross-sectional SEM image and EDX spectrum of the sample treated at 443 K for 4 h, as well as elemental mapping images for Mg, O, and Al. As shown in Figure 6(a), the film produced on this sample had a thickness of about 52 μm, whereas treatment at 423 K for 4 and 6 h, at 433 K for 4 and 6 h, at 443 K for 3 h, and at 453 K for 2 h produced film thicknesses of about 5, 35, 20, 68, 19, and 7 μm, respectively, based on their cross-sectional SEM images. The elemental mapping images in Figure 6(b) revealed that the film mainly consisted of Mg and O, which indicates that it was comprised of magnesium hydroxide. This is in agreement well with the XRD results and FT-IR spectrum for the sample. A trace amount of Al was also detected, which provides evidence that Mg-Al LDH exists slightly and locally in the film. The XRD, FT-IR, and SEM-EDX results reveal that steam coating method can produce Mg-Al LDH in the resulting surface film.
Based on these results presented thus far, a model for the formation of the film is proposed as follows. In the closed reaction vessel, steam is produced from water by thermal energy until the vessel becomes saturated. The high pressure and temperature mean that this water vapor has a high kinetic energy and reactivity, which leads to the dissolution of Mg and Al metal into Mg2+ and Al3+ ions. It is considered likely that Mg(OH)2 crystallites grow via a dissolution-precipitation mechanism, as Nordlien et al. have suggested that amorphous platelets of Mg(OH)2 are initially formed by the precipitation of Mg2+ or other soluble Mg species . In addition, Mg2+ and Al3+ ions might react with the steam to form carbonate Mg-Al LDH, with carbonate ions being sourced from CO2 present in the enclosed autoclave atmosphere prior to steam treatment. The water vapor can also penetrate the composite film due to the difference in pressure and its high kinetic energy, where it can then react with the Mg alloy substrate. As a result, a film composed of magnesium oxide/hydroxide and carbonate Mg-Al LDH is prepared directly on the AZ31 substrate. This direct growth of a film from a substrate is a strong advantage over other methods, as this can greatly improve the mechanical stability and adherence of the film. It has also been reported that the existence of Mg-Al LDH in a film prevents metals such as Al and Mg from occurring the corrosion reaction, and so the films produced here are expected to provide a high corrosion resistance.
The corrosion resistance of the films formed on AZ31 Mg alloy was investigated via potentiodynamic polarization curve measurements, i.e., the lower the polarization current, the better the corrosion resistance. Figure 7 shows the potentiodynamic polarization curves for the samples treated at 423, 433, 443, and 453 K and untreated AZ31. The corrosion potentials (Ecorr) and corrosion current densities (icorr) for samples treated under different conditions are listed in Table 2. The Ecorr and icorr values for the bare AZ31 were estimated to be ca. −1.43 V vs. Ag/AgCl and 5.39 × 10−5 A/cm2, respectively. The anodic region of the curve for the bare AZ31 can be divided into three fields: (i) first region beginning from Ecorr (low anodic over potential) that exhibits a linear increase in current density with potential; (ii) a range of abrupt increase in current density from −1.4 V, where the dissolution of Mg to Mg+ or Mg2+ ions predominantly occurred ; and (iii) a current plateau at more positive potentials, where the electrode surface is covered with a film identified as Mg(OH)2 . In the cathodic branch, hydrogen evolution is more dominant at negative potentials than Ecorr, resulting in an increase in the cathodic current density. The current densities of the anodic and cathodic branches of all samples decreased with film formation, which indicates that the films are effective in improving the corrosion resistance of AZ31 alloy.
|Temp.||1 h||2 h||3 h||4 h||6 h||8 h|
|423 K||Ecorr V vs. Ag/AgCl||–||−1.427||–||−1.446||−1.389||−0.316|
|icorr A/cm2||–||1.23 × 10−5||–||3.22 × 10−5||5.29 × 10−5||1.29 × 10−7|
|433 K||Ecorr V vs. Ag/AgCl||–||−1.425||–||−0.376||−0.161||–|
|icorr A/cm2||–||1.28 × 10−5||–||1.79 × 10−7||4.82 × 10−11||–|
|443 K||Ecorr V vs. Ag/AgCl||−1.427||−1.421||−0.899||−0.265||–||–|
|icorr A/cm2||3.33 × 10−6||2.48 × 10−7||1.90 × 10−9||9.84 × 10−8||–||–|
|453 K||Ecorr V vs. Ag/AgCl||−1.405||−1.104||–||–||–||–|
|icorr A/cm2||1.21 × 10−5||7.71 × 10−9||–||–||–||–|
At all treatment temperatures, the Ecorr tends to shift in a positive direction with increasing process time. Furthermore, at potentials more positive than Ecorr, clear passive ranges can be presented in the curves of the samples treated at 423 K for over 6 h, at 433 K for more than 4 h, at 443 K for more than 2 h, and at 453 K for 2 h. This existence of a long passive range means that the film coatings on AZ31 exhibit highly corrosive-resistant properties in solutions containing Cl− ions. An abrupt increase in current density at more positive potentials than Ecorr could be seen in the polarization curves of the samples prepared at 423 K for more than 6 h, at 433 K for 4 h, at 443 K for more than 2 h, and at 453 K for 2 h. For example, the samples treated at 423 K for more than 6 h show an increase in current density at approximately −0.7 and 0.3 V, respectively. This increase of current density could be ascribed to pitting corrosion, which would imply that the electrolytic solution permeated through the film due to the presence of cracks or pores. Similar behavior has been reported before, even in aluminum alloys . However, there was no evidence of pitting corrosion in the curve of the sample treated at 433 K for 6 h, which had Ecorr and icorr values estimated to be −0.161 V and 4.824 × 10−11 A/cm2, respectively (Table 2). These values represent the most positive potential and lowest current density of all the samples, indicating that these conditions produced the best corrosion resistance. We therefore subsequently investigated the durability of the corrosion resistance of a sample treated at 433 K for 6 h when immersed in a 5 wt.% aqueous NaCl solution.
Figure 8(a) shows the XRD patterns of samples treated at 433 K for 6 h before (black line) and after (dashed red line) immersion in 5 wt.% aqueous NaCl solution for 240 h. The only clear difference in peak intensity is attributable to 0 0 1 reflection of the Mg(OH)2 crystallite, which is much stronger after immersion. This indicates that immersion in a 5 wt.% NaCl solution changes the orientation of the film. The SEM images in Figure 8(b and c) show this change in the orientation of the Mg(OH)2 film, with immersion making the film aligned at fairly steeply inclined angles with respect to the surface.
The SEM images in Figure 9 show the sample treated at 433 K for 6 h before and after its immersion in 5 wt.% NaCl solution for 240 h. The visible appearance of these samples is shown in the inset images and reveals that although no physical damage occurred as a result of immersion, there was a slight change in the color of the film. The SEM images show that there was a slight microscale change following immersion, with submicron pores being formed on the surface. A morphological change in the film surface was also observed, with small amounts of nanosheets becoming aligned at fairly inclined angles relative to the surface. This nanosheet formation is in agreement with the XRD patterns and is likely the result of dissolution caused by the film being in contact with the NaCl solution. Figure 9(c) presents potentiodynamic polarization curves for the sample treated at 433 K for 6 h before and after immersion in 5 wt.% NaCl solution for 240 h. From this, we see that the Ecorr after immersion was approximately −0.65 V, which represents a considerable negative shift due to the change in the physicochemical properties of the film. The icorr value, on the other hand, was almost the same before and after immersion. A clear passive region at c.a. 650 mV was observed after immersion, and though there was a slight variation in the current density, this indicates that the film exhibits high corrosion resistance even when immersed in a 5 wt.% NaCl solution for 240 h, as shown by the SEM image in Figure 10(a).
Figure 10(b–e) shows the EDX spectra obtained at spots (i), (ii), (iii), and (iv) in Figure 10(a). At spots (i) and (ii), there are two peaks attributable to Cl and Al but no peak for Na, whereas no Cl was detected at spots (iii) and (iv). The presence of Al at spots (iii) and (iv) indicates that the Mg-Al LDH layer was still present, which suggests that CO32− ions in the Mg-Al LDH layer were partially exchanged for Cl− anions. From the XRD patterns and EDX spectra, it is evident that Mg-Al LDH is able to trap some Cl− ions within its structure, which has been reported to delay the degradation of protective coatings . It is also apparent that carbonate Mg-Al LDH has a weak anion exchangeability. This is significant given that Williams et al. have found LDHs loaded with carbonate and nitrates are more effective than a blank coating in preventing the corrosion of a metallic substrate during filiform corrosion tests .
This section has shown that a protective magnesium hydroxide film containing Mg-Al layered double hydroxide (LDH) can be produced on Mg alloy via a simple, environmentally friendly, inexpensive, and chemical-free steam coating method. Subsequent XRD and FT-IR studies revealed that these films are composed of Mg(OH)2 and carbonate Mg-Al LDH, and a formation mechanism was proposed based on the results of XRD, FT-IR, SEM, and EDX measurements. Potentiodynamic measurements and immersion tests in 5 wt.% NaCl aqueous solution revealed that a film prepared on AZ31 alloy at 433 K for 6 h provides the greatest corrosion resistance. Increasing the immersion time induced a change in the orientation of the crystalline Mg(OH)2 structure, but no change was observed in the appearance of the sample. With a growing demand for Mg alloys and an increasing focus on environmental issues, it is believed that this steam coating method can provide an effective means of improving the corrosion performance of large, complex-shape Mg alloy components.
3. Formation of Mg-Al layered double hydroxide films on flame-resistant AM60 Mg alloy
In the previous section, we characterized the properties of an anticorrosive film composed of Mg(OH)2 and Mg-Al LDH that is formed on commercial AZ31 Mg alloy by steam coating. In this section, we report on the formation and properties of the anticorrosive film that is produced in a similar manner on Ca-added AM60 Mg alloy (AMCa602), which is one of the most important lightweight materials for transportation components because of its flame-resistant nature. This unique trait is the result of the fact that Ca will oxide before Mg ignites , which increases the ignition temperature from 200 to 300 K and greatly restrains the inherently combustible nature of Mg alloys. However, as this addition of Ca does not improve the low corrosion resistance of Mg alloy, there is still a need to develop a suitable surface treatment technology.
Although anodized, chemical conversion and polymer films have all been applied to Mg alloys [6–9, 76–81], there have been few reports pertaining to the surface treatment method of flame-resistant Mg alloys. Xia et al. did find that the corrosion resistance of Mg-4.0Zn-0.2Ca alloy is improved by micro-arc oxidization (MAO) treatment in basic aqueous solutions containing Na4SiO4, NaF, and NaOH due to the formation of a film composed mainly of MgO and MgF2 ; however, the presence of fluorine atoms in such coatings creates a risk of environmental pollution. Mori et al. reported on the preparation and corrosion resistance of a plasma electrolytic oxidation layer on nonflammable Ma-Al-Mn-Ca Mg alloy (AMCa602), which was mainly composed of Mg2SiO4 and MgO . The liquid waste produced during this anodization process, however, needs to be treated prior to disposal. There is therefore still a need for an environmentally friendly surface treatment capable of improving the low corrosion resistance of flame-resistant Mg alloys.
Specimens of AMCa602 Mg alloy (composition in wt.%: 5.92% Al, 0.36% Mn, 1.99% Ca, <0.01% Si, <0.01% Cu, <0.005% Ni, <0.005% Fe, balance Mg) measuring 20 × 20 × 1.5 mm in size were supplied by Fuji Light Metal Co., Ltd. These Mg alloy substrates were ultrasonically cleaned in absolute ethanol for 10 min and then dried with inert Ar gas. Once clean, they were placed on a substrate stage in a Teflon-lined, 100 ml-capacity autoclave, into which 20 mL of ultrapure water containing 0, 2, 20, or 200 mM Al(NO3)3 · 9H2O was added to the bottom to produce steam. The distance between the surface of the water and the substrate stage was around 3 cm. The autoclave was heated to a temperature of 433 K, where it was held for 6 h before being allowed to cool naturally to room temperature. Following this steam treatment, the samples were ultrasonically cleaned in ethanol for 10 min and dried with Ar gas.
Figure 11 shows the GAXRD patterns of the films formed on AMCa602 at 433 K over a 6 h period from aqueous solutions containing 0, 2, 20, or 200 mM of Al(NO3)3 · 9H2O in 20 mL of ultrapure water. A weak peak attributable to the Mg alloy substrate is visible at a 2θ angle of approximately 34° in all the patterns. A diffraction peak corresponding to 1 0 1 reflection at approximately 2θ = 38°, which has the highest intensity according to JCPDS No. 44-1482, appears slightly weaker with an increase in the Al(NO3)3 concentration. In addition, several peaks were observed in all the GAXRD patterns at 2θ angles of approximately 18, 33, 51, 58, 62, 68, and 72° that can be attributed to the 0 0 1, 1 0 0, 1 0 2, 1 1 0, 1 1 1, 2 0 0, and 2 0 1 diffraction peaks of brucite-type Mg(OH)2, indicating the formation of hexagonal Mg(OH)2. In addition to these peaks, one or two peaks attributable to a hydrotalcite (HT)-like structure can be distinguished that are assigned to Mg1−xAlx(OH)2(NO3)x·nH2O intercalated with nitrate anions . These diffraction peaks are indexed according to 3R symmetry, and so the two peaks at around 2θ = 11° and 22° are attributed to the 0 0 3 and 0 0 6 reflections, respectively, of Mg1−xAlx(OH)2(NO3)x·nH2O [51, 85, 86]. This suggests that the treated surfaces were crystallized with a preferential hexagonal orientation along the (0 0 3) plane, which is a characteristic feature of the spontaneous texture axis of Mg1−xAlx(OH)2(NO3)x·nH2O. Furthermore, these XRD profiles reveal that all films were composed of crystalline Mg-Al LDH and Mg(OH)2, regardless of the Al(NO3)3 concentration used.
The appearances of the samples after steam coating at 433 K for 6 h using aqueous solutions containing Al(NO3)3 at concentrations of 2, 20, and 200 mM are shown in Figure 12(a–c). Note that in all cases the substrate were uniformly covered with film, but the color of this film did change with Al(NO3)3 concentration. The SEM images in Figure 12(d–f) show the same samples, from which we see that an aqueous solution containing 2 mM Al(NO3)3 produces a relatively smooth surface with only a few minute pores measuring several hundred nanometers in diameter. The surface of the sample prepared using an aqueous solution containing 20 mM Al(NO3)3, on the other hand, was quite rough and had a small number of nanosheets that were slightly inclined with respect to the surface.
As shown in Figure 13, elemental mapping by EDX revealed traces of elemental Mg, Al, and O primarily in areas where nanosheets existed, which is consistent with the presence of Mg-Al LDH. Aggregated particles are also clearly discernible in these images, which likely formed during steam treatment and condensed on the substrate, and it is on these particles that the nanosheets formed. The sample prepared from an aqueous solution containing 200 mM Al(NO3)3 was smooth and dense compared to that produced from a 20 mM Al(NO3)3 solution, and although small nanosheets were formed that were slightly inclined with respect to the surface, no cracks were found in the film.
Figure 14 shows the FT-IR spectra of samples steam coated at 433 K for 6 h using aqueous solutions containing 2, 20, or 200 mM of Al(NO3)3. The most dominant spectral feature is the intense sharp peak observed at 3696.9 cm−1, which is in good agreement with the XRD data in that pure Mg(OH)2 has been reported to exhibit a single band at 3698 cm−1 due to the high basicity of its O—H groups . The shoulder band at 781 cm−1 can be ascribed to the Al-OH translational mode , while the bands at 1370–1520 cm−1 are attributed to the symmetric and asymmetric stretching modes of CO32− in the interlayer [69, 84]. The peak at approximately 1385 cm−1 is attributed to the ν3 vibrational mode of NO3− with D3h symmetry in the interlayer . These peaks and bands were present in the spectra of all samples that had peaks attributable to Mg-Al LDH in the irrespective XRD patterns. An additional band at 1630–1650 cm−1 is assigned to the bending mode of water molecules in the interlayer, which are confirmed by the shoulder band at 3080 cm−1 and means that hydrogen bonding can occur between H2O and carbonate ions in the interlayer . These results demonstrate that all films were composed of crystalline Mg(OH)2, nitrate-based Mg-Al LDH (Mg1−xAlx(OH)2(NO3)x·nH2O), and carbonate-based Mg-Al LDH (Mg1−xAlx(OH)2(CO3)x/2·nH2O).
Figure 15(a) presents a cross-sectional SEM image of the film prepared at 433 K for 6 h from an aqueous solution containing 200 mM of Al(NO3)3, which reveals that it had a thickness of about 197 μm. In comparison, the films prepared under the same conditions from 2 and 20 mM Al(NO3)3 solutions were estimated to be about 308 and 193 μm, respectively. All of the films formed from an aqueous solution containing Al(NO3)3 were much thicker than the 53 μm achieved with only H2O, which implies that the existence of Al(NO3)3 in a closed autoclave accelerates film growth. The elemental mapping images in Figure 5(b–d) reveal that the films contained mostly Mg and O, implying that they were composed mainly of Mg(OH)2. This result is in agreement well with the XRD profiles and FT-IR spectra obtained from each of the samples. A small trace of Al was also detected, demonstrating that Mg-Al LDH is presented locally in the film. Based on these XRD, FT-IR, and SEM-EDX results, it is concluded that steam coating is capable of forming Mg(OH)2 and Mg-Al LDH on both AMCa602 and AZ31.
From the results presented in this section, the proposed mechanism by which films are formed by steam treatment can be described as follows. First, an aqueous solution containing Al(NO3)3 is heated in the electric oven to generate steam containing Al3+ and NO3− ions that then saturates a closed reaction vessel. When this steam makes contact with the AMCa602 Mg alloy, the Mg metal it contains is readily dissolved into Mg2+ ions due to the high temperature and pressure. It is likely that Mg(OH)2 crystallites are subsequently formed via a dissolution-precipitation mechanism, resulting in the formation of Mg(OH)2 on the surface of the AMCa602 substrate. This is predicated on a notion by Nordlien et al. that amorphous platelets of Mg(OH)2 are initially formed by the precipitation of Mg2+ or other soluble Mg species . In addition, Mg2+, Al3+, and NO3− ions have the potential to react with the steam and form nitrate-based Mg-Al LDH. Any carbonate ions generated from CO2 present in the enclosed autoclave atmosphere before steam treatment can result in the formation of carbonate-based Mg-Al LDH. Thus, along with Mg(OH)2, both nitrate- and carbonate-based Mg-Al LDHs can form on AMCa602. Even if these reactions only occur at the surface, they can still progress at the interface between AMCa602, and the overlying film as the steam is able to penetrate thanks to the difference in pressure and its high kinetic energy. The end result is that a film composed of Mg(OH)2 and Mg-Al LDH is directly formed on AMCa602. A porous film formed during the early stage of the process will allow more steam to pass through, resulting in a thicker final film. In contrast, it is difficult for steam to easily penetrate a dense initial film. Thus, the film thickness is considered to be dependent on the preparation conditions.
The corrosion resistance of the film formed on the AMCa602 Mg alloy at 433 K for 6 h from aqueous solutions containing 0, 2, 20, and 200 mM Al(NO3)3 was estimated from the potentiodynamic polarization curves shown in Figure 16. For reference, the potentiodynamic polarization curve of untreated AMCa602 is shown in Figure 6. The corrosion potentials (Ecorr) and corrosion current densities (jcorr) of the untreated AMCa602 and samples prepared under various conditions are listed in Table 3 along with the Tafel slopes (βa, anodic; βc, cathodic) obtained from the polarization curve for each sample. The Ecorr and jcorr values for the untreated AMCa602 were estimated to be −1.524 V vs. Ag/AgCl/sat.-KCl and 4.214 × 10−5 A cm−2, respectively. In the cathodic region of the bare AMCa602, hydrogen evolution dominated at potentials that were more negative than Ecorr (i.e., negative potentials with magnitudes above 1.524 V), resulting in an increase in the cathodic current density. The anodic region of the curve for the bare AMCa602 can be distinguished into two ranges. The first of these starts from Ecorr (low anodic overpotential, −1.524 V) and is defined by an abrupt linear increase in the logarithm of the current density, with a shift toward a positive voltage in the potential. It is in this region that dissolution reaction of Mg to Mg+ or Mg2+ ions is most likely to have occurred . As the potential moved toward values more positive than −1.4 V vs. Ag/AgCl/sat.-KCl (i.e., negative potentials with magnitudes below 1.4 V and positive potentials), a current plateau occurred as a result of the electrode surface being covered with Mg(OH)2 . It has been reported that the anodic region of the polarization curves for AZ31 can be distinguished into three ranges , which means that the overpotential for the dissolution of metallic magnesium to magnesium ions (Mg+ and/or Mg2+) for AMCa602 is lower than that of AZ31. This improved dissolution might have resulted from the addition of Ca, as this is more reactive than Mg.
|Ecorr (V vs. Ag/AgCl)||jcorr (A/cm2)||Tafel slope (mV/decade)|
|Untreated AMCa602||−1.524||4.21 × 10−5 ± 4.82 × 10−5||ba: 21.2|
|Film coated AMCa602 prepared from ultrapure water||−0.385||1.85 × 10−9 ± 9.85 × 10−10||ba: 198.7|
|Film coated AMCa602 prepared from 2 mM Al(NO3)3||−0.322||7.86 × 10−8 ± 7.68 × 10−7||ba: 278.1|
|Film coated AMCa602 prepared from 20 mM Al(NO3)3||−0.316||1.16 × 10−9 ± 2.39 × 10−9||ba: 235.4|
|Film coated AMCa602 prepared from 200 mM Al(NO3)3||−1.007||9.56 × 10−11 ± 2.17 × 10−10||ba: 288.6|
Figure 17 shows the potentiodynamic polarization curves for the bare AMCa602 and the samples that were treated at 433 K for 6 h from aqueous solutions containing 0, 2, 20, and 200 mM Al(NO3)3. Note that the current densities in the anodic and cathodic regions for all film-coated samples were lowered compared to the bare AMCa602 (Figure 6) and reproducibly low (Figure 7), demonstrating that all films improved the corrosion resistance of AMCa602. The Ecorr also shifted toward a more positive value independent of the Al(NO3)3 concentration. At potentials more positive than Ecorr, extended passive regions were clearly observed in the curves of the samples prepared from aqueous solutions containing 0 and 200 mM Al(NO3)3, as shown in Figure 6. These perfect passive regions imply that the film effectively protected AMCa602 in a solution containing Cl− ions. The jcorr values decreased with an increase in the Al(NO3)3 concentration, and this was independent of the film thickness. Any increase in the thickness and/or density of a film is generally beneficial to improving corrosion resistance by inhibiting contact between the metal surface and a corrosive solution. In this case, the film prepared from an aqueous solution containing 2 mM Al(NO3)3 was the thickest of all of the samples, and so this is expected to provide the best corrosion resistance. However, the low density of this film produced the highest jcorr value, as explained earlier in the discussion regarding the formation mechanism. Thus, although the films prepared from aqueous solutions containing 20 and 200 mM Al(NO3)3 were of comparable thickness, their jcorr values suggest that they had very different densities.
An increase in current density at approximately 0.3 and 0 V vs. Ag/AgCl/sat.-KCl was observed in the polarization curves of the samples prepared from aqueous solutions containing 2 and 20 mM Al(NO3)3, which suggests that the corrosive solution permeated through cracks or pores in the film and initiated pitting corrosion. Similar behavior has been reported even in aluminum alloys , but in this case, the increase in current density was suppressed to produce a passive-like behavior with a shift in potential toward more positive values. This behavior suggests that the film might have some degree of self-healing functionality, which has been reported with Zn-Al LDH films formed on Al substrates that exhibited an increase in passive current density to initial values after immersion in a corrosive electrolyte . In contrast, there is no evidence of pitting corrosion in the curve of the sample prepared from an aqueous solution containing 200 mM Al(NO3)3, which had Ecorr and jcorr values estimated to be −1.007 V vs. Ag/AgCl/sat.-KCl and 9.560 × 10−11 A cm−2, respectively, as shown in Table 3. This jcorr value was the lowest of all the samples, and as this is closely related to the rate of corrosion, this would suggest that a film prepared from 200 mM Al(NO3)3 offers the best corrosion resistance.
The correlations between the Mg-Al LDH content of the films, Al(NO3)3 concentration in solution, and corrosion current density are shown in Figure 18. The Mg-Al LDH content of the films was roughly estimated from the XRD peak intensity ratio of the (0 0 3) plane of Mg-Al LDH, ILDH-003, to that of the (1 0 1) plane of Mg(OH)2, IMg(OH)2-101, i.e., ILDH content = ILDH-003/IMg(OH)2-101. The peak intensities for these planes correspond to the strongest intensities in the standard powder diffraction data for Mg(OH)2 and Mg-Al LDH. As the Al(NO3)3 concentration in the aqueous solution increased, the ILDH content increased and the corrosion current density decreased, suggesting a strong relationship between the two. It has been reported that carbonate-based Mg-Al LDH shows anion exchangeability, whereby Cl− ions are incorporated from NaCl solution and CO32− ions are exhausted . Moreover, nitrate-based Mg-Al LDH can abstract Cl− by ejecting NO3− and so has an anion exchangeability that is higher than that of carbonate-based Mg-Al LDH. The Mg-Al LDH content of the films is therefore considered to have a significant effect on the corrosion current density, which is closely related to the rate of corrosion.
The durability of the corrosion-resistance performance of the sample prepared from an aqueous solutions containing 200 mM Al(NO3)3 was investigated by immersing it in a 5 wt.% NaCl solution. Figure 19 shows the surface of the sample before immersion and after intervals of 5, 12, 24, 48, and 72 h. Note that the beige color of the film surface appeared to change slightly upon immersion, but no corrosion product or physical change in the appearance of the film was observed for up to 12 h. After 24 h of immersion, however, some spotty contrasts emerged that are likely the result of the film being damaged by the corrosive solution. An increase in the immersion time from 48 to 72 h led to a slight dissolution of the AMCa602 substrate at the corners of the sample, which is believed to be the progression of corrosion initiated by corrosive medium coming into contact with the substrate through a pit or minute pore in the film. Additionally, as corrosion resistance is known to be closely related to the film thickness, it is possible that the film was simply thinner at the corner than at other points on the sample. This nonuniformity in film thickness could have occurred through differences in crystal growth at the edge of the substrate.
The SEM images in Figure 20 show the sample surface before and after immersion in 5 wt.% NaCl solution for 24 and 72 h. It is evident from this that with an increase in immersion time, there was a change in the surface morphology and nanosheets that were slightly inclined with respect to the surface were more likely to be noticed on the surface. The EDX spectrum in Figure 20(d) revealed that there were two peaks attributable to Cl and Al at spot (i) in Figure 20(c), but no evidence of a Na peak, which supports the earlier assertion that the Mg-Al LDH layer remained in the film and exhausted CO32− and NO3− in exchange for Cl− anions. Similarly, the absence of Cl and the presence of Al at spot (ii) also indicate that the Mg-Al LDH layer remained, which means that the nitrate or carbonate anions in the Mg-Al LDH layer were partially replaced with Cl− anions.
Figure 21(a) presents XRD patterns obtained from the film-coated AMCa602 before and after immersion in 5 wt.% NaCl solution for 24 to 72 h. This shows that the peak attributable to the (1 0 1) plane of the Mg(OH)2 crystallite decreased slightly in intensity with an increase in immersion time, whereas that attributable to the (0 0 3) plane of the Mg-Al LDH crystallite increased in intensity. This indicates that even though the Mg(OH)2 film was slightly dissolved in the corrosive solution, the Mg-Al LDH crystallite remained intact. The potentiodynamic polarization curves for the film-coated AMCa602 before and after immersion are shown in Figure 21(b), from which the Ecorr and jcorr values were determined to be −0.348 V vs. Ag/AgCl/sat.-KCl and 1.821 × 10−9 A cm−2, respectively, after 24 h immersion, and −0.065 V vs. Ag/AgCl/sat.-KCl and 1.326 × 10−9 A cm−2 after 72 h. The Ecorr shifted toward a more positive value with an increase in immersion time, which is possibly a result of the change in the ratio of anodic to cathodic sites on the surface that is caused by the dissolution of Mg(OH)2 in the film. The jcorr value of the film-coated AMCa602 increased after immersion, indicating an increase in the corrosion rate that agrees well with the results of the immersion tests. This increase, however, was independent of the immersion time, with the jcorr value almost comparable after 24 and 72 h of immersion.
From the XRD patterns, SEM observations, EDX spectra, and polarization curve measurements, the mechanism by which a film produced by steam coating provides corrosion resistance to AMCa602 alloy can be explained as follows. At first, Mg(OH)2 from the Mg(OH)2/Mg-Al LDH composite film formed earlier on AMCa602 is dissolved to some extent through contact with a corrosive solution. This dissolution of Mg(OH)2 causes the amount of Mg-Al LDH with anion exchangeability present on the film surface to increase gradually, leading to an increase in the intensity of the 0 0 3 reflection of Mg-Al LDH, as shown in Figure 21(a). Due to the presence of Mg-Al LDH on the film surface, Cl− ions can become trapped in the Mg-Al LDH, as shown in the EDS analysis, resulting in delaying corrosion. These results support the hypothesis that the Mg-Al content of a Mg(OH)2/Mg-Al LDH composite film is an key factor in improving the durability and corrosion resistance of the film. Moreover, steam coating produces a film that is thick and has a high density due to the presence of two phases of different sizes, namely, Mg(OH)2 and Mg-Al LDH, which further contributes to improving the corrosion resistance of the film.
In conclusion for this section, a film composed of Mg(OH)2 and Mg-Al LDH was successfully formed on combustion-resistant AMCa602 Mg alloy by using an environmentally friendly steam coating method. Subsequent XRD and FT-IR studies confirmed that the film was composed of Mg(OH)2 in combination with nitrate- and carbonate-based Mg-Al LDHs. The film thickness varied from 52 to 308 μm depending on the Al(NO3)3 concentration. A mechanism for the formation of such films was proposed based on the results of XRD, FT-IR, SEM, and EDX measurements, and the corrosion resistance was estimated by potentiodynamic measurements and immersion tests in 5 wt.% NaCl solution. A film prepared by steam coating at 433 K for 6 h using an aqueous solution of 200 mM Al(NO3)3 was found to have the highest corrosion resistance of all samples tested. The correlation between the Mg-Al LDH content of the film, Al(NO3)3, its corrosion current density, and the concentration of the aqueous solution used to create it was investigated, which revealed a strong correlation between the Mg-Al LDH content and corrosion current density. Based on the results achieved here with flame-resistant Mg alloy, we believe that our steam coating technique is an effective method for improving the corrosion resistance of both existing and future-generation Mg alloys.
4. Conclusion and outlook
This chapter has presented the results of experimental studies into the formation and corrosion resistance of an anticorrosive film composed of Mg(OH)2 and Mg-Al LDH produced on AZ31 and AMCa602 Mg alloys by steam treatment. The physicochemical properties of the films were investigated using XRD, SEM, EDX, and FT-IR, and a mechanism for their formation was proposed based on these results. The corrosion resistance of the films produced on AZ31 and AMCa602 was estimated based on potentiodynamic measurements and immersion tests in 5 wt.% NaCl solution, which revealed that corrosion resistance of both alloys was improved by steam treatment. The Mg-Al LDH content of the films was also found to be strongly related to the corrosion current density. As Mg alloys are one of the more promising materials for reducing vehicle weight, thereby lowering fuel consumption and reducing CO2 emission, any improvement in their inherently low corrosion resistance is of great value to increasing their wider-scale use. Thus, although there has already been a great deal of academic work and commercial interest in surface treatments for Mg alloys, we believe future work should focus on developing the steam coating process described here to greatly improve the corrosion resistance of a range of Mg alloys.
This research was partially supported by a grant for Advanced Industrial Technology Development from the New Energy and Industrial Technology Development Organization (NEDO) of Japan (no. 11B06024d), a grant-in-aid for young scientists (B) (no. 16 K18249) from the Japan Society for the Promotion of Science, and the Japan Science and Technology Agency (JST), Adaptable and Seamless Technology Transfer Program through Target-driven R&D (A-step: no. AS251Z02997K and AS2815047S).