*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 *L*c1 and *L*c2 and the corresponding friction forces *F*t1 and *F*t2, respectively, are given in Table 2. As indicated by Table 2, the value of the critical normal force *L*c1 was 7.9 N, and the formation of oblique and parallel cracks was observed on the coating surface (Figure 5a). The value of the critical normal force *L*c2 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 5b).

**Figure 4.** Results of the scratch test for the Ni–P coating on the ZE10 magnesium alloy with a scratch pattern.

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].


**Table 2.** Values of critical normal forces and friction forces of Ni–P coating deposited on the ZE10 magnesium alloy and, for comparison, values found in other studies.

**Figure 5.** Detail of the damage of Ni–P coating during the scratch test: (**a**) *L*c1 and (**b**) *L*c2 (SEM).

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). Figures 4 and 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 [28,30].

The formation of transverse tensile arch cracks [22,34] 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 (*L*c1 and *L*c2) was higher compared to the data presented in articles [28,30], 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 *R*<sup>a</sup> ≈ 0.25 μm [28,30]. However, the surface of the experimental ZE10 magnesium alloy was polished to a roughness *R*<sup>a</sup> ≈ 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 *R*<sup>a</sup> ≈ 4.5 μm [35], and that of the experimental ZE10 magnesium alloy was *R*<sup>a</sup> ≈ 2 μm. It is also possible to observe that the value of the critical load *L*c 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 *R*<sup>a</sup> ≈ 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 [35,37]. 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 *R*<sup>a</sup> ≈ 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.
