3.1. Ni–P Coating Morphology and Chemical Analysis
Uniform Ni–P coating was successfully deposited on the previously pre-treated surface of the ZE10 magnesium alloy substrate. As shown in Figure 2
a, grain boundaries and intermetallic phase particles appeared on the surface of the material after the pre-treated, pickled substrate. The pickled surface also shows a “honeycomb-like microstructure.” Revealed intermetallic phase particles/α-Mg solid solution interface and the grain boundaries served as places where the initiation of the electroless deposition of the coating initiated due to the galvanic coupling [24
]. The higher roughness of pickled and activated surface increased the adhesion of deposited Ni–P coating to the magnesium substrate.
Usually, zinc immersion is used for the magnesium alloy pre-treatment before Ni–P coating deposition. El Mahallawy et al. [14
] studied the electroless Ni–P coating of different magnesium alloys, using zinc immersion as the pre-treatment of magnesium alloys. The zinc immersion was used to remove the residual oxides and hydroxides from the surface of the magnesium alloys, and a thin layer of zinc formed on the Mg surface preventing back oxidation.
Based on the obtained results, a partial re-oxidation of the ZE10 magnesium alloy surface occurred without the use of zinc immersion (Figure 2
a). During the following immersion of the activated sample into the electroless nickel bath, the substrate became catalytically active when the surface oxides were dissolved in the nickel bath, and the replacement reaction occurred between the substrate and nickel ions. Even though some contamination of the substrate surface was observed before Ni–P coating deposition, it seems that it did not negatively affect the coating process [9
The surface morphology of the deposited Ni–P coating with a nodular structure, formed by typical cauliflower-like shapes is shown in Figure 2
b. Wang et al. [25
] showed that deposited Ni–P coatings are formed by a columnar microstructure. However, the deposited Ni–P amorphous coatings improve the corrosion resistance of magnesium, the inherent columnar microstructure of the coating does not provide the best protection against the corrosion. The high concentration of inter-column defects, such as microvoids and micropores, [26
], form channels where the corrosion ions and environment can pass through the coating and react with the substrate. The presence of microcavities was not evident between nodular cusps of the deposited coating (Figure 2
b), which is in agreement with observations in [28
]. Based on an evaluation of SEM figures (Figure 2
b and Figure 3
a), no defects and cracks were observed in the deposited Ni–P coating at the ZE10 magnesium substrate/Ni–P coating interface.
The average thickness of the coating prepared for 60 min determined from the cross sections was approximately 10 µm. In the case of a longer deposition time (180 min), prepared with the aim to increase the coating thickness to obtain relevant microhardness values, thickness was 30 µm.
Deposited Ni–P coating with an average thickness of about 10 µm was chosen for EDS analysis (Figure 3
). Using EDS mapping analysis, it was determined that the distribution of Ni and P in deposited Ni–P coating was homogeneous in the entire cross section, as shown in Figure 3
b,c, respectively. The EDS analysis determined that the Ni content in the deposited Ni–P coating was 95.6 wt % and the P content was 4.4 wt %. Based on the literature [9
], it was determined that the deposited Ni–P coating is low-phosphorus, as in the cases of the AZ31 magnesium alloy presented in [30
] and the AZ61 magnesium alloy in [28
3.2. Ni–P Coating Microhardness Analysis
Based on the measured data, it was determined that the average value of the microhardness of the Ni–P coating was 690 ± 30 HV 0.025, measured in the cross section. The microhardness of the plain ZE10 magnesium substrate was 60 ± 4 HV 0.025.
The surface microhardness of the coated samples increased approximately 11-fold compared with the ZE10 magnesium alloy.
It is assumed that the measured hardness of low-phosphorus Ni–P coatings is higher compared to the high-phosphorus coatings [9
]. The addition of filler (SiC, Al2
) into the high-phosphorus Ni–P matrix led to a substantial increase in hardness [9
]. The microhardness (690 ± 30 HV 0.025) of the deposited low-phosphorus Ni–P coating on the ZE10 magnesium alloy reached a value higher than that of the Ni–P/SiC composite coating prepared on the AZ91 magnesium alloy presented in [31
] and [32
]. The microhardness of the Ni–P/SiC composite coating (7.33 wt % P) was 620 HV [31
], and that of the electroless Ni–P/SiC nanocomposite coating (10 wt % P) was 600 HV 0.025 [32
The microhardness of the deposited low-phosphorus Ni–P coating was higher compared with the values obtained for the high-phosphorus Ni–P coatings. The hardness of high-phosphorus Ni–P coatings ranges from 410 to 600 HV [29
]. As the content of phosphorus in Ni–P coatings increases, the microhardness of the coating decreased due to the microstructural changes (a decrease in crystallinity) [9
3.3. Analysis of the Physical Properties of the Ni–P Coating
The results of the scratch test performed on the Ni–P-coated ZE10 magnesium alloy sample are shown in Figure 4
. The measured values of the critical normal forces Lc1
and the corresponding friction forces Ft1
, respectively, are given in Table 2
. As indicated by Table 2
, the value of the critical normal force Lc1
was 7.9 N, and the formation of oblique and parallel cracks was observed on the coating surface (Figure 5
a). The value of the critical normal force Lc2
was 13.6 N, and the formation of transverse tensile arch cracks across the entire width of the track was observed on the coating surface (Figure 5
As stated in [22
], tensile and compressive stresses are generated during the scratch test and cause more complex mechanisms and damage. A crack can nucleate on a defect or at the coating/substrate interface. The crack is formed due to the localization of the stresses on the coating/substrate interface or in the coating (transverse crack). In the case of a layer, the tensile radial tension induced with the Rockwell tip can generate circular or transverse arch cracks that extend across the layer into the substrate. As the tip moves, several circular or transverse arch cracks can intersect. These cracks can also occur at the back of the contact as a response to tensile stresses during tip sliding. Cracks also occur on the back of the contact due to the friction-induced tensile stresses [34
As a result of the applied pressure load of the Rockwell diamond tip during the scratch test, ductile failure of the deposited Ni–P coating occurs due to the introduced internal stresses.
The character of the damage to the locating layer during the scratch test is dependent on many factors [22
]. In addition to the influence of the characteristics of the experimental device on the tested layer damage mechanism, there are geometric properties of the substrate-layer system (such as layer thickness, roughness, etc.), experimental parameters (tip and scratch rate), and properties of the substrate-layer system (thermal coefficients, microstructure and internal stresses, elasticity, and hardness modules). Figure 4
and Figure 5
show the scratch track morphology and the layer cracking character, which is similar to the case of Ni–P coatings on AZ31 and AZ61 magnesium alloys presented in [28
The formation of transverse tensile arch cracks [22
] across the entire width of the track was observed (Figure 4
). The adhesion strength of the experimental electroless deposited Ni–P coating on a wrought ZE10 magnesium alloy (Lc1
) was higher compared to the data presented in articles [28
], where the Ni–P coating was deposited on AZ31 and AZ61 magnesium alloys, respectively. The difference could be explained by the coated substrate pre-treatment process. The pre-treatment of AZ31 and AZ61 magnesium alloys before the deposition of the Ni–P coating included polishing to a roughness Ra
≈ 0.25 μm [28
]. However, the surface of the experimental ZE10 magnesium alloy was polished to a roughness Ra
≈ 2 μm, which is significantly rougher than AZ31 and AZ61. The higher roughness of the substrate surface can improve the adhesion strength between the deposited Ni–P coating and the ZE10 magnesium alloy due to the mechanical interlocking of the two components [34
This effect was also observed in [35
], where the adhesion strength between the deposited Ni–P coating and blasted or polished surface of the AZ91 magnesium alloy was studied. The scratch track morphology for the pre-blasted and pre-polished samples with the deposited Ni–P coating showed a similar trend, but it was observed that the scratch track width was narrower on the rougher surface when compared to the polished surface. The scratch track width was slightly narrower for coated samples after annealing for 1 h at 523 K. This effect can be contributed to the increase in the hardness of the Ni–P coating after annealing, which was demonstrated with the increase in hardness from ~600 to ~900 HV due to the coated sample annealing. As indicated by Table 2
, the critical load Lc
for the plain Ni–P coating was 14.0 N and 10.2 N for the blasted and polished surfaces, respectively. The increase in adhesion strength to 16.5 N was observed for the rough blasted AZ91 substrate after annealing for 1 h at 523 K. This increase was apparently linked to the hardness increase and the effect of the rough surface. The brittle cracking of the deposited coating was observed at substrates with the rough surface, and the wedge spallation was observed at substrates with the polished surface. The decrease in adhesion strength was observed for samples annealed at 673 K due to the embrittlement of the Ni–P coatings (Table 2
However, as indicated by Table 2
, resulting values of the critical loads of rough (blasted) samples of AZ91 [35
] are slightly higher when compared to the experimental Ni–P coating deposited on the ZE10 magnesium alloy. This can again be connected to the higher roughness of the coated substrate. The roughness of the blasted AZ91 magnesium alloy surface was Ra
≈ 4.5 μm [35
], and that of the experimental ZE10 magnesium alloy was Ra
≈ 2 μm. It is also possible to observe that the value of the critical load Lc
of experimental coating deposited on the ZE10 alloy is higher in comparison with the polished surface of the AZ91 alloy in [35
], where the roughness was Ra
≈ 0.05 μm. As is obvious, the roughness of the substrate surface has a significant effect on the coating adhesion strength due to the mechanical interlocking between Ni–P coating and the coated magnesium substrate.
As indicated in the literature [36
], applied surfactants in the nickel bath had a significant effect on the adhesion strength of deposited Ni–P/TiO2
composite coating on the AISI 1018 steel substrate (Table 2
). No cohesive or adhesive failure of the coating was observed up to ~13 N in the case of the Ni–P/TiO2
coating prepared on AISI 1018 without using the surfactant. The formation of the mild tensile cracks at ~19 N was evident for the Ni–P/TiO2
composite coating using sodium dodecyl sulfate (SDS) surfactant at 1.5× CMC (critical micelle concentration). In the case of the Ni–P/TiO2
composite coatings on AISI 1018 involving dodecyl trimethyl ammonium bromide (DTAB) at 1× CMC, the cohesive failure was observed at the applied load of ~29 N. Moreover, no linear or radial cracks were observed in the case of the coated steel substrate, nor of any of the analyzed coatings, which also indicates the importance of the surface of the substrate with respect to the adhesion of the coating.
The increase in the adhesion strength of Ni–P coatings to the magnesium substrate, along with a slight increase in the roughness of the substrate surface, was shown to be achieved by adding the proper surfactant into the nickel bath [35
]. This proves that a more effective adhesion of the coatings is caused by the excessive attractive forces between the Ni–P coatings and substrate [38
Based on the obtained result, a sufficient surface roughness of ZE10 reached via surface polishing on the roughness Ra ≈ 2 μm, in combination with the activation of the surface via acid pickling, seems to be reached during pre-treatment. Adequate pretreatment resulted in an adequate adhesion of the coating to the substrate and a considerably high resistivity against damage.
3.4. The Electrochemical Corrosion Test in 0.1 M NaCl
shows the potentiodynamic polarization curves of the ZE10 magnesium alloy and the ZE10 alloy with the deposited Ni–P coating in 0.1 M NaCl obtained at laboratory temperature. The polarization curve of the Ni–P-coated sample is significantly shifted to more electropositive values, which means better corrosion properties of the Ni–P-coated sample compared with the untreated ZE10 magnesium alloy.
Therefore, a deposited Ni–P coating appears to be suitable for the protection of magnesium alloys. Based on the Tafel extrapolation analysis [39
], the values of the corrosion potential, Ecorr
, and the corrosion current density, icorr
, for samples with the deposited Ni–P coating and the ZE10 magnesium alloy were determined. The average values of the Ecorr
, for the ZE10 magnesium alloy and the Ni–P-coated alloy, were −1701 and −505 mV, respectively, and the values of the icorr
, for the ZE10 magnesium alloy and the Ni–P-coated material, were 23.7 and 0.4 µA·cm−2
, respectively (Table 3
). It is generally known [25
] that the columnar structure of deposited Ni–P coatings contains a network of defects (grain boundaries, pores, and microcavities). These defects are precursors for a micropitting and may cause the corrosion attack of the material under the deposited Ni–P coating. However, no local corrosion attack (pitting) was observed in the anodic area of the polarization curves.
Based on the determined values of icorr
, the corrosion rate, vcorr
, was calculated for a short-term experiment. As indicated by Table 3
, the corrosion rates of the ZE10 alloy and the Ni–P-coated samples in 0.1 M NaCl were 530.00 µmpy and 8.95 µmpy, respectively.
Experimental deposited Ni–P coating was ranked among the low-phosphorus Ni–P coatings, which are characterized by their lower corrosion resistance when compared to the medium- and high-phosphorus coatings, [9
]. As indicated in the work of S. Narayanan [12
], deposited Ni–P coatings with low phosphorus content (3.34 wt % P) showed an Ecorr
of −0.536 V and an icorr
of 4.22 µA·cm−2
, whereas medium- (6.70 wt % P) and high-phosphorus (13.30 wt % P) Ni–P coatings showed Ecorr
values of −0.434 and −0.411 V, respectively, and icorr
values of 1.17 and 0.60 µA·cm−2
, respectively. When comparing the experimentally obtained results and the results presented in [12
], it is evident that the value (Table 3
) of the corrosion potential, Ecorr
, is between Ecorr
values of 25-µm-thick low-phosphorus (3.34 wt % P) and medium-phosphorus (6.70 wt % P) Ni–P coatings presented in [12
Although the plating rate of high-phosphorus Ni–P coatings in [12
] was higher, high-phosphorus Ni–P coatings were deposited under more energy-intensive conditions (90 ± 1 °C) [12
]. It is evident that the obtained value of the corrosion current density, icorr
, of high-phosphorus Ni–P coating was worse compared to the presented low-phosphorus Ni–P coatings on the ZE10 magnesium alloy (Table 3
), and the electroless deposited coating on the ZE10 alloy exhibits greater corrosion resistance. This can be attributed to the fact that the experimentally prepared electroless deposited Ni–P coating was probably less defective or that the Ni–P coating contained only a small amount of microcavities when compared to the coatings analyzed in [12
3.5. The Immersion Test in 0.1 M NaCl
As indicated by Figure 7
a,b, after the exposition of the sample with deposited Ni–P coating in 0.1 M NaCl for 1 h, the degradation of the Ni–P coating occurred due to the corrosion of the magnesium substrate under the coating.
Corrosive agents may accumulate between the nodules (grains) creating the Ni–P coating and migrate to the substrate surface. According to the results in [25
], the inherent columnar porous microstructure of the coating (Figure 7
d) does not adequately protect the magnesium substrate against corrosive environments [41
]. Furthermore, electroless Ni–P coatings contain a certain amount of microcavities in their volume. These microcavities form nucleation sites for micropitting when the material is exposed to a corrosive environment. Microcavities are created as a result of the hydrogen evolution during the deposition process when small bubbles of hydrogen H2
, as a byproduct of the nickel reduction, are adsorbed on the surface of the growing Ni–P coating [40
]. Stirring of the nickel baths or the rotation of the deposited objects is a partial solution of this problem. However, the addition of surfactants into the nickel bath was shown to be a more effective solution, while the reduction of these microcavities resulted in an increase in the corrosion resistance of the coated substrate [42
As described above, due to the high concentration of inter-column defects, such as microvoids and micropores [27
], the interaction between the corrosive environment containing chloride ions and magnesium substrate under the deposited Ni–P coating resulted in the creation of oxides or chlorides of magnesium. This reaction has an effect on the increase in corrosion product volume when compared to the bulk material and the evolution of hydrogen is an accompanying process. The increase in the volume of the material under the Ni–P coating leads to the formation of cracks and the subsequent destruction of the coating and thus to the acceleration of the corrosion of the magnesium substrate (Figure 7
Even though the coating internal defects were not observed in SEM analysis, even the intercolumnar areas can provide a path for the transfer of corrosive medium elements into the coating and react with the substrate. Figure 7
d reveals the damaged structure of the Ni–P coating showing the columnar structure. The revealed columnar structure supports the theory that the grain boundaries and microdefects present on the grain boundaries acted as the paths for the corrosive medium transfer to the substrate.
Observed local corrosion attack did not correlate with the results obtained with potentiodynamic tests; however, the electrochemical test was limited to the range of polarization used, and a larger range for the measurement likely can reveal the pitting attack.