*3.2. Microhardness*

As a result of slide burnishing and grinding, the material was strengthened (Figure 12). Following grinding conducted close to the surface at a depth of 1 μm, the microhardness of the surface layer was approx. 15% higher than the microhardness of the core. From a depth of 5 μm, the microhardness of the surface layer of the ground sample was similar to the microhardness of the core. The distribution of microhardness for the slide burnished sample shows that the highest microhardness occurred at a depth of 3 μm from the surface. The obtained microhardness distribution is characteristic of this treatment method.

**Figure 12.** Distribution of microhardness on the surface layer of samples after grinding (pre-treatment) and slide burnishing.

Figure 13 shows the degree of strengthening *e* and hardened layer thickness *gh* as a function of burnishing feed (Figure 13a) and slide burnishing force (Figure 13b). The hardening degree was determined for the depth *gh* = 3 μm, for which the microhardness is the highest. The slide burnishing of X6CrNiTi18 steel specimens resulted in the degree of strengthening *e* ranging from 8.77% to 42.74%, which was higher than the degree of strengthening obtained for X19NiCrMo4 steel (*e* = 32%) [6]. The hardened layer thickness

*gh* after slide burnishing ranged from about 10 μm to 100 μm, while in [6], the reported changes in microhardness reached a depth of up to 18 μm.

**Figure 13.** Degree of strengthening and hardened layer thickness after slide burnishing as a function of: (**a**) feed *f* (*F* = 230 N, *vn* = 35 m/min, *i* = 1); (**b**) force *F* (*f* = 0.06 mm/rev., v*n* = 35 m/min, *i* = 1).

When burnishing is conducted with a higher feed (Figure 13a), the traces of the diamond tip passes are at a greater distance from each other. This causes a decrease in the structural homogeneity and, consequently, a lower degree of strengthening *e* and decrease in the hardened layer thickness *gh*.

The greatest differences between the values of the strengthening degree *e* occurred for the forces *F* = 230 N and *F* = 300 N. The application of a higher burnishing force caused an increase in the degree of strengthening *e*. This is most likely due to intense plastic deformation caused by friction, which leads to grain refinement of the microstructure. The use of a higher burnishing force causes the plastic deformation to take place deeper in the material, which results in increased microhardness of the surface layer extending further from the treated surface. The obtained results of the influence of *F* on *gh* are similar to the results described in [26], where the slide burnishing process was conducted on carbon steel with a hardness of 250 HV.

### *3.3. Residual Stress*

Figure 14 shows the results of the influence of the burnishing force on the distribution of stresses S11 occurring in the surface layer. As the burnishing force increased, the depth of compressive stresses increased. The maximum value of the compressive residual stress (about 400 MPa) was obtained for the burnishing forces *F* = 230 N and *F* = 300 N. The burnishing force significantly affected the depth of compressive residual stresses from *gσ* = 0.4 mm (for the force *F* = 90 N) to *g<sup>σ</sup>* = 1.1 mm (for the force *F* = 300 N). It should be assumed that the increase in the depth of residual stresses will allow the place of fatigue crack initiation to be shifted from the surface to the subsurface layers. This means that the phenomenon of nucleation and crack propagation will be delayed [51].

**Figure 14.** Distribution of residual stresses as a function of distance from the surface of X6CrNiTi18 steel samples after slide burnishing with a variable burnishing force (*f* = 0.06 mm/rev; *vn* = 35 m/min; *i* = 1; R = 3 mm).

Figure 15 shows the effect of burnishing feed on the distribution of S11 stresses. The depth of the compressive residual stresses was about *gσ* = 0.9 mm. No significant influence of the feed on the residual stress distribution was observed. The obtained depths of residual stresses are greater than those observed after the slide burnishing of C45 steel [52]. The compressive residual stress zone was much deeper than the plastically deformed zone. The results are similar to those presented in the paper [53].

**Figure 15.** Distribution of residual stresses as a function of distance from the surface of X6CrNiTi18 steel samples after slide burnishing with a variable feed (*F* = 230 N; *vn*= 35 m/min; *i* = 1; R = 3 mm).

Table 3 shows the visualization of the PEEQ equivalent plastic strains. For variable burnishing forces, the feed rate *f* = 0.06 mm/rev. was used; for variable feeds, the burnishing force was *F* = 230 N. A symmetrical cross-section was created to illustrate deformation on the surface as well as in the subsurface layers. The PEEQ color maps demonstrate that the plastic deformation wave resulting from the passages of successive burnishing

elements shifted the area of maximum plastic deformation concentration away from the axis of symmetry of the workpiece. This phenomenon was not observed for the feeds *f* = 0.16 mm/rev. and *f* = 0.20 mm/rev.

**Table 3.** Influence of burnishing force and feed on equivalent plastic strains.

#### *3.4. Positron Annihilation Lifetime Spectroscopy*

The obtained positron lifetimes and intensities of both components in individual samples slightly differed from each other. The short-lived component with an intensity of 80.5 ± 0.3% dominated in the spectra, and its lifetime of 168.7 ± 0.3 ps indicated its origin to mainly come from the positrons trapped at edge dislocations, where a lifetime of 162 was observed [54]. The lifetime of the second component was 360 ± 2 ps and corresponds to quite large vacancy clusters, i.e., approx. 15 vacancies [55]. However, it should be noted that this is an average value and that the vacancies can vary in size. Due to slight changes in lifetimes and intensities, the differences between the samples are best reflected by a change in the mean lifetime (*τmean*), defined as:

$$
\tau\_{\text{mean}} = \frac{\tau\_1 I\_1 + \tau\_2 I\_2}{I\_1 + I\_2} \tag{2}
$$

where τ1, τ<sup>2</sup> and *I*1, *I*<sup>2</sup> are the lifetimes and the intensities of the first and second components of the positron lifetime spectrum, respectively. An increase in *τmean* means an increase in the probability of positron trapping in vacancy clusters (most likely due to an increase in their concentration) and/or an increase in the average size of the clusters. It has been shown that this parameter is a good indicator of changes in the microhardness of the surface layer of samples [42].

The positron mean lifetime *τmean* increased slightly yet systematically with an increase in the burnishing force *F* (Figure 16). This dependence is consistent with the results of the microhardness measurements (Figures 12 and 13). On the other hand, the decreasing dependence of microhardness on burnishing feed *f* was not reproduced by the *τmean*. This suggests a different nature of the change in material defects when changing burnishing feeds than those induced by changing the burnishing force. The correlation of the dependence between *τmean* and burnishing feed with an analogous dependence for the *Sa* parameter suggests that the change in surface topography and roughness may also affect the positron implantation profile.

**Figure 16.** Mean positron lifetime *τmean* as a function of (**a**) burnishing force *F* (*f* = 0.06 mm/rev., *vn* = 35 m/min, *i* = 1), (**b**) burnishing feed *f* (*F* = 230 N, *vn* = 35 m/min, *i* = 1). The dashed lines are only an eye-guide.
