## 4 Discussion
were able to detect significant differences between the zirconium-rich (about +170 mV) as well as the yttrium-rich precipitates (about +50 mV) and the \(\alpha\)-Mg matrix, whereas the second phase and \(\beta\)-phase showed negligible differences (+25 and +15 mV, respectively) [30]. Based on these potential differences, it can be assumed that micro-galvanic corrosion is mainly driven by zirconium-rich precipitates. In comparison, the hydrogen evolution of accelerated corrosion does not show any change in slope and thus no passivation tendency (Figure 5). This is associated with the increased number of corrosion pits with increasing polarization current density (Figure 6). It can be assumed that the increase in hydrogen evolution due to galvanostatic polarization is driven by the number of corrosion pits on the inhomogeneous microstructure. This theory is in accordance with the areal homogeneous corrosion morphology between thecorrosion pits (Figure 7c). Furthermore, anodically controlled corrosion causes a higher coefficient of determination of the regression lines (Table 3). The composition of the corrosion layer, as well as the deposit's morphology of the unaccelerated tests (Figure 9), is consistent with the current literature [11; 34]. In addition to the characteristic components Mg and oxygen (O), the layer contains the elements phosphorus (P), calcium (Ca), and carbon (C), whereby the proportion of the latter increases for the circular deposits. To determine the chemical compounds, the peaks of the Fourier-transform infrared spectroscopy (FTIR) scans are considered (Figure 9). The activity at 1030 cm\({}^{-1}\) can be attributed to phosphate ions (PO4\({}^{3-}\)), whereas the double peak at 1500 and 1401 cm\({}^{-1}\) and the peak at 851 cm\({}^{-1}\) are due to stretching and bending of carbonate ions (CO3\({}^{-}\)) [9; 22]. Consequently, the formation of a poorly crystalline carbonate-enriched apatite layer is assumed [22]. The classification of the activity at 556 cm\({}^{-1}\) is more difficult. Based on the current literature, it indicates the presence of hydroxide groups (OH\({}^{-}\)) [35] or also of phosphate compounds [22; 34]. Similarly, the broad activity between 3630 and 2460 cm\({}^{-1}\), occurring only for the accelerated immersion tests, indicates the presence of OH\({}^{-}\) [34; 35]. Furthermore, no circular or spherical deposits are observable for the accelerated tests, only volcano-like deposits can be found (Figure 8). The deposition of phosphate and carbonate compounds seems to be a time-dependent process [11; 22]; hence, they cannot be found due to the shortened duration of the accelerated tests. Accordingly, the corrosion layer consists mainly of the characteristic magnesium oxides. It is questionable to what extent the time-dependent deposition of phosphates and carbonates affects the actual corrosion process and the mechanical stability and whether application-oriented acceleration by galvanostatic polarization is possible. To quantify this, Figure 10 shows the calculated dissolution rates \(\dot{\mathrm{m}}_{\mathrm{diss}}\) (Table 3) plotted against the current density i in a double-logarithmic representation. In addition, the respective standard deviations are shown (Table 3), resulting for the unaccelerated corrosion rate from the three tests carried out and for the accelerated tests from the coefficient of determination (COD) R\({}^{2}\) of the regression lines. The data points lie in good approximation (R\({}^{2}\) = 99.886%) on a straight line, described by the following Equation (1)  
\[\dot{\mathrm{m}}_{\mathrm{diss}}\ =\ 10^{3.139}\cdot\mathrm{i}^{1.094} \tag{1}\]  
For this alloy, the dissolution rate is controllable by using galvanostatic anodic polarization under the test conditions described. The extent to which the corrosion morphology, deviating with increasing polarization, influences the mechanical stability in the functional phase of the implant remains open for the moment.