3.1. Phase Composition (General Remarks)
The main features of the PEO layers are summarized in Table 2
. In several layers, the main phases detected by XRD (which analyses the coating full-thickness composition, as demonstrated by the presence of peaks from the substrate) are the two crystalline forms of TiO2
: anatase and rutile (Table 2
). Representative XRD spectra, recorded on single layers from silicate- or phosphate-rich baths (samples S and P, respectively) are reported in Figure 2
XRD patterns for sample S in Figure 2
also show the influence of abrasive blasting on phase composition but this point will be discussed subsequently (Section 3.2
Anatase was the dominating phase in XRD patterns of all samples (Table 2
) and it was the only phase detected by micro-Raman (Figure 3
), using spectra reported in Bouchard and Smith [21
] as references for Raman peaks indexing. Further reference spectra were found in Friedemann et al. [22
The inability of Raman (characterized by a shallower penetration through the PEO layer than XRD) to detect rutile may be due to the higher concentration of the anatase in the outermost portion of the PEO layers. In fact, prevalence of the metastable oxide in the outer portion of PEO layer, where a high cooling rate and hence rapid quenching is likely to predominate on annealing effects, has been observed also in the case of aluminium alloys by Xue et al. [23
The prevalence of anatase over rutile both in single layers (P, PN) and in double layers (DA) obtained in pulsed DC is most likely due to the use of phosphate-based baths (either for the single layer or for the external layer in double-layered architectures). In fact Khan [24
] showed that phosphate baths in DC mode are known to produce anatase-rich PEO layers, where phosphorus, incorporated during the growth, tends to limit the anatase to rutile transformation (which is the final transformation in the sequence from amorphous titania to metastable anatase (T > 550 °C) and then to thermodynamically stable rutile (T > 850 °C) under the action of micro-arc discharges.
By comparing samples DA and DB (treated in the same conditions and differing only in the abrasive blasting post-treatment carried out on DA, Table 1
), it is possible to notice that a minor contribution of rutile was detectable by XRD only in the abrasive-blasted sample DA, probably due to the smoother surface that improves the signal-to-noise ratio.
In all the layers produced by phosphate-based electrolyte (with the exception of the DC mode single layer P), also some orthorhombic aluminium phosphate (AlPO4
) was detected (Table 2
), as observed by Martini et al. [25
] in a previous work on PEO treatment of Ti-6Al-4V. According to Wang et al. [26
], crystalline AlPO4
is supposed to form as a consequence of high-temperature thermolysis of hydrated aluminium polyphosphates inside discharge channels. In single layers S and P (DC mode), also traces of α-Al2
were detected by XRD (Figure 2
), as previously observed by Yerokhin [27
] for PEO treatment of Ti-6Al-4V. In this work, silicate-based electrolyte in DC (sample S) did not produce a rutile-dominated layer, as the one observed by Yerokhin as well as by Wang [27
] in AC-treated Ti-6Al-4V (phosphate-based bath). In our case, the prevalence of anatase may be due to less intensive arcing and lower temperatures inside discharge channels.
As regards the influence of the incorporation of P-based compounds in the PEO layers on biological response (which has not been investigated yet for these layers), literature data indicate that P-containing PEO coatings on Ti-6Al-4V induce a homogeneous distribution of growing MG-63s cells as well as a higher collagen deposition per cell than plasma-sprayed hydroxyapatite [29
]. Also, the incorporation of Si (as amorphous silicate) is not expected to have a negative impact on biological response (SiO2
-based bioactive materials are known for their excellent bioactivity [30
]). However, the actual biological response to these PEO layer will require specific investigations in a further step of the work.
The influence of treatment conditions and surface finishing on microstructure, phase composition and micro-mechanical properties will be discussed in the following Section 3.2
, Section 3.3
and Section 3.4
). Subsequently, the tribological behaviour of all the PEO layers will be discussed and compared to selected reference materials in the final Section 3.5
3.2. Influence of Abrasive Blasting
The influence of abrasive blasting on surface morphology and roughness is shown in Figure 1
for PEO single layers obtained in DC mode (P and S in Table 1
The comparison of images and Ra values in Figure 1
, before and after abrasive blasting, shows that this mechanical surface finishing treatment effectively removes the brittle and porous external layer of the anodic oxides, most notably in the case of samples from silicate-based baths (S). For the S samples, the free surface displays a rough appearance, with nodular features (ranging from about 5 to 20 μm) and large pores between the nodules, typical of PEO layers grown in silicate-rich electrolyte as shown by Aliasghari et al. [31
]. In the case of samples treated in the phosphate-based solution (P), which already showed a rather smooth surface morphology in the as-treated condition, the decrease of surface roughness as well as the morphological modification is less remarkable. In fact, also after abrasive blasting, typical PEO defects due to stochastic discharge events and gas evolution, such as cavities and volcano-like features, are still visible in P samples.
For samples obtained in silicate baths (S), the influence of abrasive blasting on phase composition is shown in Figure 2
, where XRD patterns recorded before and after mechanical surface finishing are compared. In this case, the most evident effect of abrasive blasting is the removal of the amorphous contribution (wide band at low diffraction angle), likely due to amorphous silica as suggested by Wang et al. [28
]. Also, Yerokhin et al. [27
] detected amorphous silica in PEO layers produced on Ti-6Al-4V in silicate-based baths. In fact, also large-area EDS analysis (Table 3
) displays a remarkable decrease in Si concentration after blasting (from about 28 to 15 wt.%).
The influence of abrasive blasting on surface composition was also evaluated for single layers deposited from phosphate baths in pulsed DC mode (PP, PM, PN, Table 1
), in order to check for accumulation of Si from the blasting medium also in this set of samples. Results of large-area EDS analysis on the free surface demonstrated that sample PP (obtained at lowest duty cycle, that is, at lowest pulse-on times) shows a significantly higher Si concentration (visible also in cross-section X-ray maps discussed in Figure 4
) than the others.
The shorter pulse-on times for the treatment of sample PP may be responsible for a lower density of the layer, as shown by cross-sections in Figure 4
, hence for the increased tendency to abrasive incorporation by comparison to other PEO layers obtained in pulsed DC. In this current regime, also Dehnavi et al. [32
] observed that layer density and microstructure improved with increasing pulse-on time. This Si enrichment is likely to be responsible for the relatively high microhardness of sample PP (Table 2
), thus beneficially influencing its tribological behaviour.
These results show that, also for the S single layers discussed above (Table 3
), surface contamination due to Si from the blasting medium cannot be ruled out, because it can be masked by the decrease of Si% as a consequence of the removal of the outermost amorphous silica layer. Therefore, in the case of silicate-phosphate double layers (DA and DB, Table 1
), the possible surface enrichment of Si due to abrasive blasting was assessed by measuring GD-OES depth profiles (Figure 5
), in order to take into appropriate consideration also the layered structure of the coatings.
Glow Discharge-Optical Emission Spectroscopy (GD-OES) depth profiles in Figure 5
show the trend of Ti, Si and P signal intensity as a function of depth for the same double layer coating, both before (as-treated) and after abrasive-blasting. Oxygen was not included in this graph because its profile typically has a lower S/N ratio than the others and it would affect readability without yielding further useful information. Also, Al and V, which showed similar trends as Ti but with proportionally lower intensity due to their lower alloy concentration, were not added in the graph so as to preserve its readability. Based on the comparison of Si profiles, abrasive blasting decreases the total layer thickness of about 10 μm (from 15 to 5 μm). Such an estimate is probably more accurate than the one obtainable by polished cross-sections (Table 2
), because GD-OES data are averaged over a relatively large analysed area whilst the intrinsic brittleness of as-treated PEO layers makes them prone to damage during metallographic preparation.
In the as-treated double layer (DB), P can be detected through the whole thickness of the PEO layer and its signal shows a slight intensity increase at around 10 μm from the surface. In the abrasive-blasted layer (DA), thinned and compacted by the finishing procedure, the P signal is more intense than in the previous case and rather constant throughout the layer thickness, indicating in both cases that the electrolyte was able to penetrate the inner regions of the coating, probably through breakdown channels, as observed also by Galvis et al. [33
] for single layers obtained in DC mode from phosphate baths. Even though the immersion in the phosphate solution was the second step of the treatment, after the first step in silicate bath, it induced P enrichment of the whole layer (as shown also by EDS X-ray maps in Figure 4
). GD-OES depth profiles also show a remarkable increase of Si in the outermost portion of the abrasive-blasted layer (DA), by comparison to the same coating in the as-treated condition (DB). The different trend and the higher concentration of Si in the abrasive-blasted layer (DA) is most likely due to embedding of silicate glass fragments during abrasive blasting.
3.3. Influence of Current Mode (DC Versus Pulsed DC)
In the case of phosphate-based single layers, the influence of current mode can be estimated by comparing sample P (DC mode) with PP, PM, PN (pulsed DC). Cross-section images in Figure 1
(sample P) and 4 (PP, PM, PN), as well as average thickness values in Table 2
, show that the use of pulsed DC current induces a slight densification of the PEO layer, accompanied by a thickness decrease. This is probably due to the beneficial influence of pulse-off time which, in the pulsed DC mode, contribute to interrupt spark discharges, decreasing the growth rate but also limiting disruptive discharge events. In terms of phase composition, the use of pulsed DC slightly reduced the tendency towards the formation of stable rutile (Table 2
), probably due to the attainment of lower temperatures than in DC mode. In terms of micromechanical properties (Table 2
), there is no remarkable difference between microhardness and critical loads for full delamination (Lc3) among samples obtained in DC or pulsed DC mode. However, the PN layer, obtained in pulsed DC at the highest duty cycle (Table 1
), makes an exception, with its highest Lc3 value due to improved microstructure (discussed in Section 3.4
In the case of double layers, the influence of current mode on microstructure can be assessed by comparing sample D5 (DC mode) with DA (pulsed DC). The cross-section images in Figure 6
show that both samples do not display significant porosity.
Double-layered sample DA is slightly thinner and less compact that D5, mostly in the outer zone. The only difference in terms of phase composition between these samples was a higher amount of crystalline AlPO4
in the layer obtained in DC mode (D5), where high-temperature thermolysis of hydrated aluminium polyphosphates inside discharge channels was probably more likely than in pulsed DC (where pulse-off time may allow cooling during coating growth in a more effective way). In pulsed DC mode (sample DA), a lower current density was employed (Table 1
), further contributing to the achievement of lower temperatures during discharge events. The use of pulsed DC lead to lower thickness and microhardness in sample DA but it induced a higher practical adhesion (Lc3) by comparison to D5 (Table 2
3.4. Influence of Duty Cycle (Pulsed DC)
For phosphate-based single layers in pulsed DC (PP, PM and PN, Table 1
), thickness decreased whilst compactness increased with increasing duty cycle (Figure 4
), due to the beneficial microstructural effect of increasing pulse-on time, as previously discussed in Section 3.2
. Accordingly, in terms of phase composition, rutile was detected only at the highest duty cycle values, in sample PN (Table 2
). Correspondingly, a relatively high microhardness was detected in the same sample. The highest microhardness recorded for phosphate-based single layers was recorded in sample PP, obtained at the lowest duty cycle value: however, in this case the measured value in probably affected by the abrasive residues embedded in the surface layer, as previously discussed in Section 3.2
In general, the low hardness of these PEO layers, which are only slightly higher than the substrate, may be due to the predominance of soft anatase (Table 2
), as well as to residual non-oxidised titanium, as suggested also by Yerokhin et al. [27
]. Also, Diamanti et al. [34
] obtained a similar result, that is, thin anatase-based PEO layers on Ti-6Al-4V, produced in calcium glycerophosphate bath, displayed a hardness lower than the untreated substrate.
The microstructural modifications induced by the increase of duty cycle also showed a beneficial influence on practical adhesion (Lc3, Table 2
), which can be ascribed to the denser microstructure [35
3.5. Dry Sliding Tests
Average values of coefficient of friction (COF) as well as of maximum wear depth are plotted as a function of normal load in Figure 7
Each PEO layer is characterised by a critical normal load (Table 2
) at which failure of the coating occurs during the test, hence the load range for COF values in Figure 7
a is wider for the best-performing coatings than for the worst ones.
shows the typical graph recorded during the tests which induced coating failure: when the substrate starts to be involved in the contact (after about 100 m in this case), COF decreases whilst system wear (i.e., material removal from both the block and the ring) increases with increasing sliding distance. The friction transition after coating failure lead to COF values comparable to that of the untreated substrate.
The observation of wear scar morphology after the above described friction and wear transitions (Figure 9
) shows that, after coating failure, the underlying substrate is deeply ploughed (Figure 9
b) due to abrasion.
Adhesion damage is also noticeable between the grooves, showing the same typical morphologies observed also for the untreated substrate (Figure 9
a). The occurrence of these friction/wear transitions and the morphology of worn surfaces is completely comparable to the case of other PEO-treated Ti-6Al-4V samples tested in dry sliding conditions against bearing steel, discussed in a previous work by Martini et al. [25
For this specific set of samples, the highest COF values were recorded for the as-treated (not abrasive-blasted) double layer DB (Figure 7
) at 5 N. The COF of all abrasive-blasted PEO layers (with the exception of PM, discussed here below) is lower than for DB, demonstrating that the surface finishing procedure improved the frictional behaviour. Fei et al. [9
] reported similar results on this phenomenon, which is due to the decrease of the abrasive component of friction, brought about by the decrease of surface roughness induced by abrasive blasting. Also, the single layer PM coating, with high roughness (Table 2
) and low compactness (Figure 4
), showed high COF values at 5 N. Similar to DB, also this coating failed already at 10 N, due to its detrimental combination of high roughness, low compactness and low hardness (the latter two parameters being related, since pores and cracks negatively affect the hardness of PEO layers as shown by Curran and Clyne [37
For PEO layers that survived in a wider load range (namely D5 and PN), COF slightly increases on going from 30 to 40 N, due to destabilization of the iron oxide transfer layer that covers all the treated surfaces as a consequence of mild tribo-oxidation of the steel counterface. The presence of these iron oxide transfer layers, which is typically observed in the PEO-steel contact [25
], is documented by the images of wear scars taken at 5 N in Figure 10
and Figure 11
(the latter also reporting EDS and micro-Raman data).
These transfer layers, when formed during dry sliding against AISI 52100 at room temperature and at relatively low sliding speed (0.3 m s−1
), consist of haematite (Fe2
) according to micro-Raman analyses (Figure 11
c). Tonelli et al. [38
] observed the same type of haematite-based transfer layers also for other PEO coatings in similar contact conditions.
The destabilization of the transfer layers starts to be appreciable at the load before friction and wear transition (e.g., in Figure 10
c,f for layers PM and DB respectively, already being at the load before failure at 5 N and in Figure 12
for the other layers). Also, the detachment of micro-fragments from the PEO layer, which mostly occurs before complete coating failure, may contribute to increase the abrasive component of friction at high load.
As for COF, also the trend of maximum wear depth (measured on PEO-treated and untreated stationary blocks at the end of the tests) versus normal load (Figure 7
b) shows the above described transitions, related to coating failure. After coating failure, in fact, wear depth of PEO-treated blocks noticeably increases with increasing normal load, achieving values comparable to the untreated substrate. However, before coating failure, all the PEO layers investigated in the present work performed better than the bare Ti-6Al-4V substrate, in terms of wear resistance.
As previously discussed for COF, the abrasive blasting process has a beneficial influence also on wear depth. By comparing the curves of samples DA and DB in Figure 7
b, it is possible to notice that the removal of the porous and brittle external layer increased the critical load to failure from 10 N (DB) to 40 N (DA), probably due to a decreased tendency towards micro-crack driven damage accumulation, in the case of the smoother coating (DA).
It is worth noting that the PEO treatment in phosphate-based bath after the production of the inner layer in silicate bath (i.e., comparing S (single layer) to D5 (double layer)) was beneficial, leading to an increased critical load to failure (from 30 to 40 N). This effect can be probably ascribed to the higher compactness of D5 and hence to its increased hardness.
The critical load to failure for single layers obtained in phosphate bath increased in the following order: PM < P < PP < PN. This indicates that most layers obtained in pulsed DC perform better than the one obtained in DC (sample P) and the best tribological behaviour can be achieved for pulsed DC treatment at highest duty cycle (PN), due to combination of dense microstructure, relatively high hardness and high adhesion. In the case of PP, the unexpected good wear resistance (despite its low adhesion and non-dense microstructure) may be due to the Si-rich top layer, formed as consequence of glass embedding during abrasive blasting (discussed in Section 3.2
), which may also contribute to local enhancement of surface hardness.
It is also worth noting that pulsed DC treatment at highest duty cycle (PN) attains the same critical load to failure in dry sliding as the best double layers. This is probably related to its dense microstructure and relatively high hardness. Moreover, the low roughness of PN by comparison to D5 and DA, induced by the absence of the intermediate silicate layer, is also likely to limit stress concentration at asperities, thereby limiting micro-crack driven damage, which is a typical wear mechanism associated with PEO layers as reported by Diamanti et al. [34
]. The promising behaviour of PN therefore suggests that single-layer high duty cycle pulsed DC can be considered as an alternative and simpler processing route than double-layer deposition.
As regards abrasive-blasted double layers, both D5 and DA achieved the highest values of critical load to failure (40 N, Figure 7
b). Their comparable performance indicates that microhardness alone is not the key parameter in influencing wear behaviour: both layers display a rather dense microstructure, hence practical adhesion plays a key role in determining the high critical load of DA, notwithstanding its lower hardness. It is also worth noting that embedding of glass residues was observed in the DA layer (as discussed in Section 3.2
). As previously discussed for PP, this Si-rich surface layer may have a non-negligible beneficial influence, also predominating over other features such as microstructure, adhesion and hardness.
As for the comparison carried out under the highest normal loads (30 and 40 N) between the best-performing PEO layers (PN, D5 and DA) and reference materials (i.e., uncoated or PVD (Ti,Nb)N coated Ti-6Al-4V and CoCrMo), the average maximum wear scar depth values are reported in Figure 13
As previously discussed, all PEO layers outperform the uncoated Ti-6Al-4V. At 30 N, all the selected PEO layers display wear depths comparable to CoCrMo (both uncoated and PVD-coated). The wear depth of PEO layers is also slightly lower than for PVD-coated Ti-6Al-4V, which is a promising result considering that in this case the thickness of PEO layers is only slightly higher than that of the PVD coating, which is 5 μm thick). At 40 N, both the PEO the PVD layers on Ti-6Al-4V are worn out and their substrate is markedly involved in the contact. Only the PVD-coated CoCrMo still shows a very low wear depth (lower than that of the uncoated CoCrMo), due to the high load-bearing capacity of this substrate.