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Corrosion and tribocorrosion behaviour of thermally sprayed ceramic coatings on steel
Surface & Coatings Technology 205 (2011) 3683–3691
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Corrosion and tribocorrosion behaviour of thermally sprayed ceramic coatings on steel C. Monticelli ⁎, A. Balbo 1, F. Zucchi 2 Centro di Studi sulla Corrosione “A. Daccò”, Università degli Studi di Ferrara, Via Saragat 4A, 44122 Ferrara, Italy
a r t i c l e
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Article history: Received 29 November 2010 Accepted in revised form 10 January 2011 Available online 15 January 2011 Keywords: Plasma spraying Ceramic coatings Corrosion Tribocorrosion Wear rate Chloride
a b s t r a c t This research aims at investigating the corrosion and tribocorrosion behaviour of thermally sprayed ceramic coatings deposited on steel specimens and exposed to a 3.5% NaCl solution. The coatings have been prepared by plasma spraying Cr2O3 and Al2O3/13% TiO2 powders on a Ni/20% Cr bond coating. Combined wear– corrosion conditions have been achieved by sliding an alumina antagonist on the lateral surface of coated steel cylinders, during their exposure to the aggressive solution. Polarization resistance values monitored during 3 days exposures and polarization curves recorded at the end of the immersion period show that both coatings only partially protect steel substrate from corrosion. Sliding conditions (under 2 N load and 20 rpm or 10 N and 100 rpm) induce a limited increase of the substrate corrosion rates, likely as a consequence of an increase in the defect population of the ceramic coatings. On Cr2O3-coated specimens, tribocorrosion is more severe at 10 N and 100 rpm, while on Al2O3/13% TiO2coated specimens, a stronger corrosion attack is achieved at 2 N and 20 rpm. Profilometer analysis and wear track observations by optical and scanning electron microscopes evidence that on both coatings abrasion of the surface asperities produce both a surface polishing effect and, at high loads, the formation of a tribofilm, more continuous on Al2O3/13% TiO2. On this coating the tribofilm reduces the amount of surface defects and limits the corrosion attack to a certain extent. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Wear is the major cause of material wastage and loss of mechanical performances and friction is the principal cause of wear and energy dissipation [1]. Thermally sprayed ceramic coatings on steel represent an efficient and economic way to improve the wear resistance of mechanical components, when common thermal treatments (such as quenching and tempering,…) or thermo-chemical processes (such as carburizing, nitriding…) result inadequate. In particular, plasma sprayed Cr2O3 and Al2O3–TiO2 coatings can be an excellent choice, providing protection against abrasive wear and resistance to galvanic and high temperature corrosion [2]. They often result to be superior to traditional wear resistant hard chromium [3,4] and molybdenum [5,6] coatings. A low as-sprayed surface roughness characterizes Cr2O3 coatings and represents a very important technological feature, because it reduces the number of post-deposition mechanical treatments necessary in many applications [3]. Cr2O3 coatings have low friction coefficients and can be conveniently deposited on piston
⁎ Corresponding author. Tel.: + 39 0532 455136; fax: + 39 0532 455011. E-mail addresses: [email protected] (C. Monticelli), [email protected] (A. Balbo), [email protected] (F. Zucchi). 1 Tel.: + 39 0532 455134. 2 Tel.: + 39 0532 455135. 0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2011.01.023
ring and cylinder liners in the automotive industry [7–9], where they reduce fuel and oil consumption and increase the engine life [10]. If compared to Cr2O3 coatings, Al2O3 and Al2O3–TiO2 ones are less expensive and more biocompatible and can be applied in the food and medicine packaging industry because they ensure the absence of heavy metal contamination [3]. Many papers concern the wear behaviour of plasma sprayed oxide coatings and investigate the corresponding wear mechanisms. Under dry sliding conditions, both Cr2O3 and Al2O3–TiO2 coatings are reported to form a tribofilm. On the former coating, this film is formed by plastically deformed and compacted wear debris, responsible of the low measured friction coefficients [3]. XPS analysis reveals that the film formed at room temperature is constituted by CrO3 and Cr2O3 [11]. On the contrary, on Al2O3–TiO2 coatings the tribofilm has a rather loose structure which does not adequately protect the underlying material from wear [3]. On these ceramic coatings different wear mechanisms are detected, involving abrasive wear [12,13] delamination of weakly adherent successive lamellae [14], crack nucleation and delamination [15] and adhesive wear [14]. Only a few studies involve the corrosion behaviour of thermally sprayed oxide coatings on steel [16–19]. They show the clear dependence of the corrosion resistance of the coated materials from the coating porosity [20,21], which is reduced by a proper choice of the spraying parameters [22,23]. In some cases, interconnected porosity decreases at increasing coating thickness [17]. However,
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the coating porosity is often found to increase with the deposition time (and coating thickness) [22] and higher corrosion resistance can be achieved with thin coatings [19,21]. Porosity also depends on the composition of the sprayed powders. As an example, ZrO2 addition to the Al2O3 powder tends to increase the coating porosity [17], while contrasting results are obtained by TiO2 addition [3,24]. To the authors' knowledge no literature information is available concerning tribocorrosion of ceramic coatings. As the concomitant presence of wear and corrosion processes usually induces a synergistic stimulation of material degradation [25], it has been reputed interesting a characterization of both the corrosion and tribocorrosion behaviour of Cr2O3- and Al2O3–13%TiO2-coated steel specimens in the presence of widespread aggressive species, such as chlorides.
2. Materials and methods Onto cylindrical AISI-SAE 1040 steel specimens (10 mm in height; 20 mm in diameter; nominal composition: 0.37–0.44% C; 0.5–0.8% Mn; 0.15–0.4% Si; balance Fe), two commercial 250 μm-thick ceramic coatings were prepared, that is a Cr2O3 coating (indicated as C1) and a Al2O3/13% TiO2 coating (indicated as C2). As part of a standard commercial deposition process, both coatings (obtained by air plasma spray (APS) technology) were deposited onto a 20 μm-thick Ni/20% Cr bond coat (also obtained by APS), which aimed at improving the ceramic material adhesion onto the steel substrate. Details of the feedstock powders are reported in Table 1, while the deposition parameters are confidential. The phase transformations occurring during the thermal spray of the ceramic powders were investigated by comparing the X-ray diffractograms of the powders to those of the coatings. Moreover, properly polished cross sections of the powders and coatings were observed by a Scanning Electron Microscope equipped by Variable Pressure technology (VPSEM) and Energy Dispersion Spectroscopy (EDS) microprobe, to investigate their microstructures. On the polished cross-sections of the coatings, the porosity (through an image analysis software), and the Vickers microhardness (at 0.3 kg load) were also measured. The coated steel specimens had a central threaded hole permitting them to be mounted on a steel shaft which afforded the electrical contact for electrochemical tests and let the specimen rotate at 20 or 100 rpm. Before testing, the coated specimens underwent a commercial grinding process, inducing a surface roughness, Ra, of about 0.1 and 0.2 μm on C1 and C2 coatings, respectively. A grinding process is usually applied to industrial components to reduce surface roughness and increase the wear resistance. As a reference, bare steel and steel specimens coated by a thick Ni/20% Cr bond coat (250 μm, C3 coating) were also tested. They were ground by emery papers down to no. 600 and carefully degreased before exposure to the aggressive solutions. This solution was a naturally aerated 3.5% NaCl aqueous solution (pH 6–6.5), kept at room temperature (22 ± 1 °C).
Table 1 Feedstock powders used to produce top and bond coatings. Cr2O3 Type
Amperit 707.001 by Starck Composition 0.06% SiO2, 0.03% Fe2O3, b 0.02% TiO2, and balance Cr2O3 Particle dimensions −45 + 22.5 μm
Al2O3/13% TiO2
Ni/20% Cr
FST C-335.23 by Flame Spray Technologies 13.12% TiO2, 0.22% ZrO2, 0.10% SiO2, 0.09% MgO, 0.07% CaO, 0.39% other oxides, and balance Al2O3 −45 + 15 μm
Metco 43CNS by Sulzer 19.07% Cr, 1.1% Si, 0.4% Fe, 0.02% C, and balance Ni −106 + 45 μm
For tribocorrosion tests (Fig. 1), a non-commercial tribometer induced a sliding-type wear on the lateral surface of the rotating cylinders during exposure to the aggressive solution (rotation speeds: 20 rpm or 100 rpm, involving sliding velocities of 21.5 mm s− 1 or 107 mm s−1, respectively). The normal loads (L), applied through the flat surface of an alumina cylinder with a 2 mm diameter, were 2–10 N. The nominal hardness of the α-alumina counterbody was 1950 HV. During corrosion tests, the surface exposed to the aggressive solution was the whole lateral surface of the cylindrical specimens (6.4 cm2), while during the tribocorrosion tests the lateral surface was partially screened by Lacomit varnish, in order to expose only the rubbed surface (1.6 cm2) to the solution. The evolution of the corrosion process was monitored during 3 days of exposure to the aggressive solution, by measuring the linear polarization resistance (RP) values (in the potential range from −7 mV up to+7 mV vs ECOR, with a potential scan rate of 0.1 mV s−1). At the end of the 3 days immersion, under both corrosion and tribocorrosion conditions, potentiodynamic polarization curves were recorded at a scan rate of 1 mV s−1, always starting from the ECOR value. All the potentials in the text are referred to the Saturated Calomel Electrode (SCE). Some specimens exposed to tribocorrosion conditions were not subjected to destructive electrochemical tests. On these specimens, wear rates were calculated as WR = VW / L D, where VW is the wear volume, L is the normal load and D corresponds to the sliding distance. VW was obtained by multiplying the area of the wear track cross section (measured as an average from the wear track profiles at five positions along the specimen circumference by a Hommelwerk T2000 profilometer with a TK300 piezoelectrical tip) by the circumference of the cylindrical specimens. Each wear rate value was the average of three measurements. The morphology of the corrosion and tribocorrosion attack was characterized by optical and scanning electron microscope (SEM) observations. 3. Results 3.1. Powder characterization A fused and crushed Cr2O3 powder is the raw material for C1 coating (Fig. 2a). SEM observation of polished cross sections of the particles indicates that they are monophasic (Fig. 2b) and XRD analysis reveals that they are constituted by eskolaite (hexagonal Cr2O3, Fig. 3a). C2 coating has been produced by a fused and crushed Al2O3–13% TiO2 powder, composed of irregular and angular particles (Fig. 4a), biphasic in nature, as evidenced by SEM observation (Fig. 4b). XRD analysis detects the presence of crystalline α-Al2O3 and Al2TiO5 (Fig. 3b, upper diffractogram), as predicted by the Al2O3–TiO2 phase diagram [26], which indicates the stability of two structural components, that is α-Al2O3 and the eutectic α-Al2O3–Al2TiO5. 3.2. Coating characterization SEM micrographs shown in Figs. 5a and 6 evidence the lamellar structure of C1 and C2 coatings, typical of thermally sprayed coatings. C1 shows a fine homogeneous structure, still constituted by hexagonal Cr2O3 (Fig. 3a, lower diffractogram) and characterized by small interand intralamellar cracks and finely distributed pores (Fig. 5a). An irregular NiCr bond coat facilitates the adhesion of the ceramic coating to the substrate (Fig. 5b), in spite of the presence of rare pores and sandblast residues, visible at the steel/bond coat interface. With respect to C1, C2 coating exhibits a coarser structure, with thicker lamellae and longer inter- and intralamellar cracks. Moreover, it shows isolated light grey titanium-rich splats (Fig. 6), suggesting the
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Fig. 1. Schematic view of the experimental arrangement for wear corrosion tests. W.E., C.E. and R.E. respectively correspond to the working, counter and reference electrode; Hg indicates the mercury pool which electrically connects the rotating shafts (and the W.E.) to the apparatus for electrochemical measurements.
inhomogeneous interdissolution of α-Al2O3 and Al2TiO5 phases, during the spraying process. In agreement with the literature information [3], after the spraying process only metastable γ-Al2O3 and traces of α-Al2O3 are detected as crystalline phases, while aluminium titanate remains amorphous (Fig. 4 b, lower curve). The
bond coat morphology under C2 coating is equal to that exhibited in Fig. 5b. Coating porosity and Vickers microhardness values are reported in Table 2. The results obtained indicate that C1 exhibits a slightly lower average porosity with a lower standard deviation than those of C2 coating. Moreover, the average microhardness of C1 is much higher than that of C2, because of its higher bulk hardness value and lower porosity. The hardness of Al2O3–13% TiO2 coating is lower than that of the alumina counterbody, because in the coating many factors, such as porosity, α- to γ-alumina conversion during thermal spraying and TiO2 addition, contribute to reduce the material hardness [3]. Table 2 also shows the porosity percentage of the bond coating evaluated under SEM observations. It is lower than that of the ceramic coatings, in agreement with literature information [27–29]. 3.3. Linear polarization resistance measurements
Fig. 2. VPSE SEM micrographs showing a) as-received Cr2O3 powder and b) Cr2O3 powder cross section.
Fig. 7 shows the RP and ECOR values collected on steel during exposures to the aggressive solution under a rotation speed of 20 rpm. Low RP values are recorded after 1 h of immersion (540 Ω cm2) which remain more or less constant throughout the 3 days of exposure. After 1 h of immersion, ECOR is about −0.560 VSCE, that is typical of steel under active corrosion in neutral chloride solutions [30]. Then, it undergoes a moderate reactivation, with ECOR values decreasing down to −0.670 VSCE. During the immersion, steel is slowly covered by reddish iron hydroxide corrosion products. Under the same exposure conditions, a quite different behaviour is exhibited by Ni/20% Cr bond coat material (Fig. 7). In fact, specimens carrying a thick bond coating (C3-coated specimens) show RP (higher than 105 Ω cm2) and ECOR (−0.11/−0.06 VSCE) values typical of a passive metal. Tribocorrosion (TC) at 2 N load and 20 rpm stimulates the corrosion attack on steel (Fig. 7). In fact, on this material it induces RP values (about 200 Ω cm2, at the end of the immersion period) lower than those measured under pure corrosion conditions, with slightly ennobled ECOR values (−0.61 VSCE, after 3 days of immersion). Fig. 8 shows the time dependence of RP and ECOR values collected on C1-coated specimens under free corrosion conditions at 20 rpm. RP values are higher than those recorded on steel (5.5 ÷ 3.5 kΩ cm2), but they are definitely lower than those measured on C3-coated specimens. The corresponding ECOR values (about −0.64 VSCE) are close to those exhibited by bare steel specimens. Chromia, the ceramic
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a) Counts / Arbitrary units
Hex-Cr2O3 All peaks
20
40
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2-Theta / degree
b)
α - Al2O3
Counts / Arbitrary units
Al2TiO5 γ - Al2O3
20
40
60
80
100
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2-Theta / degree Fig. 3. XRD diffractograms recorded on Cr2O3 powder (a, upper curve) and coating (a, lower curve) and on Al2O3–13% TiO2 powder (b, upper curve) and coating (b, lower curve).
material of C1 coating, is an electrically insulating material which cannot be involved in any electrochemical process. So on coated steel specimen, corrosion may affect either the bond coat or the metallic substrate or both. However, at the measured ECOR values, only steel may undergo a corrosion attack, while the bond coat material results in being galvanically protected. This suggests that the ceramic coating and the underlying thin bond coating cannot hinder the aggressive solution penetration through the coating pores and the onset of steel corrosion at the pore bottom. The limited substrate area in contact with the solution justifies the relatively high RP values. Fig. 8 also evidences that the corrosion of C1-coated steel specimens is not affected by a variation in the rotation speed from 20 to 100 rpm. During 3 days of exposure of C1-coated specimens to tribocorrosion conditions at 2 N and 20 rpm, a diminution of RP values by only a factor of 2 or less is detected and no further RP reduction is observed by increasing the load up to 10 N and the rotation speed up to 100 rpm (Fig. 8). This slight stimulation of the corrosion process is likely connected to an increase in the population of the coating defects, which permits an easier access of the solution to the steel substrate. The behaviour of C2-coated specimens is similar to that exhibited in the presence of C1 coatings (Fig. 9). Under pure corrosion conditions, at both 20 and 100 rpm, RP values of about 7.8 kΩ cm2 are measured which decrease down to about 4.3 kΩ cm2, at the end of the tests. The ECOR values remain close to −0.64 VSCE, throughout the immersion period. This again indicates corrosion of the steel substrate. Under tribocorrosion conditions, a RP decrease is measured
Fig. 4. VPSE SEM micrographs showing a) as-received Al2O3–13% TiO2 powder and b) Al2O3–13% TiO2 powder cross section.
which is more marked at 2 N and 20 rpm (1.4 kΩ cm2, after 3 days) than at higher load and rotation speed (2.8 kΩ cm2, after 3 days at 10 N and 100 rpm). Under tribocorrosion conditions the slow formation of reddish iron corrosion products on both C1- and C2-coated specimens and in the solution is clearly visible. 3.4. Polarization curves The polarization curves recorded on the studied materials at the end of the immersion period in the neutral chloride solution are collected in Fig. 10. Table 3 reports the corresponding ECOR and iCOR values, as estimated by the Tafel method by extrapolation from the cathodic polarization curve. As anticipated by the results of the linear polarization resistance measurements, steel exhibits an active corrosion behaviour, indicated by its active ECOR value and low anodic overvoltage. The anodic Tafel slope (ba = 0.080 V decade−1, Table 3) is slightly higher than that expected for active dissolution of iron (0.060 V decade−1), because of the formation of insoluble surface corrosion products which moderately hinder the anodic process. Under free corrosion conditions, C3-coated specimens (with a thick Ni–20% Cr coating) are under passive corrosion conditions (high ba and low iCOR values, Table 3) and can suffer pitting corrosion at potentials nobler than +0.45 VSCE (Fig. 10). As shown in the previous section, C1- and C2-coated specimens have ECOR values almost equal to that of steel (Fig. 10). However, they exhibit higher anodic overvoltages (ba values in Table 3) and lower anodic and cathodic currents than those recorded on bare steel
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Table 2 Porosity and Vickers microhardness values for the studied coatings.
C1 C2 Bond coat
Fig. 5. VPSE SEM micrographs showing cross section microstructure of C1 coating (a) and underlaying bond coat (b).
specimens. These observations are in agreement with the hypothesis that the anodic process is connected to steel dissolution through the double coating. As on C1- and C2-coated specimens, both steel dissolution and oxygen reduction reaction occur at the bottom of the coating pores, the rates of these reactions are limited by the restricted access of the aggressive solution to the underlying metals. At the bottom of the ceramic coating pores, the noble interlaying bond coat forms efficient cathodic regions (as shown by the cathodic polarization curve on C3-coated specimens in Fig. 10) of relatively high area (depending on the ceramic coating porosity). Then the bond
Porosity/%
HV0.3/kg mm−2
7.5 ± 1.2 9.7 ± 2.8 1.1 ± 1.0
1556 ± 104 1099 ± 96
coat can induce galvanic corrosion on the small substrate regions in contact with the aggressive solution through the porosity of the bond coat itself. The relatively high cathodic to anodic metal area ratio likely stimulates the galvanic attack. With respect to C1-coated specimens, C2-coated ones exhibit lower anodic slopes and higher anodic current densities, while the cathodic polarization curves and the iCOR values are quite close to each other (Table 3). The differences in the anodic curve suggest that the slightly higher porosity of C2 coatings facilitates the migration of iron corrosion products through the coatings driven by the anodic polarization. The presence of quite similar cathodic characteristics and iCOR values suggests that on these specimens the corrosion process is mainly under cathodic control. Fig. 10 also shows the polarization curves recorded on steel under tribocorrosion at 2 N load and 20 rpm. The analysis of these curves evidences that sliding wear speeds up the cathodic reaction on steel and causes the slight ECOR ennoblement (to −0.61 VSCE), already noticed during the discussion of the polarization resistance measurements. This is due to the stimulation of oxygen diffusion towards the metallic surface, caused by the continuous removal of corrosion products from the steel surface, which induces a fivefold increase of iCOR (Table 3). Wear application on C1- and C2-coated specimens slightly stimulates both the anodic and the cathodic reactions of steel corrosion (Figs. 11 and 12). In the case of C1 coating (Fig. 11), the stimulation is quite limited and iCOR reaches its maximum value, at the highest load and sliding velocity applied (10 N, 100 rpm, Table 3). Under this condition, iCOR is twice the value measured under pure corrosion conditions. On the contrary, in the case of C2 coating (Fig. 12), tribocorrosion under 2 N and 20 rpm is more severe than that experienced at 10 N load and 100 rpm. In fact, under the former condition iCOR is about fivefold higher than that obtained in pure corrosion, while under the latter it is about twice.
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Time / h Fig. 6. VPSE SEM micrograph showing cross section microstructure of C2 coating.
Fig. 7. RP and ECOR values collected on steel and bond-coated (C3) specimens (solid lines refer to pure corrosion conditions, while dotted lines to tribocorrosion conditions).
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3.5. Wear track observations The surface micrographs of the ceramic coatings tested under different tribocorrosion conditions are shown in Fig. 13. Fig. 13a and b exhibits the surface aspect of both coatings at the wear track borderline after tests at 10 N and 100 rpm: inside the track (lower part of the pictures) the surface is much smoother than outside, as a consequence of the removal of surface asperities, left by the grinding process. Fig. 13a shows that on the hardest C1 coating the surface is rougher than that on C2 coating (Fig. 13b) and the holes in the track are due to both the coating porosity and the troughs left by the commercial grinding process. On C2 coating (Fig. 13b), the polishing effect is more pronounced and large surface regions of the track are apparently poreless. However, surface observations at higher magnification reveal that the surface pores are partially obstructed by a tribofilm produced by the smearing of the wear debris on the surface
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Fig. 10. Polarization curves recorded on steel and coated specimens, after 3 days of immersion (after 1 day of immersion, in the case of the C3-coated specimens).
Rp - C1 - 100rpm Rp - C1 - 100rpm - 10N Ecor - C1 - 100rpm Ecor - C1 - 100rpm - 10N
Fig. 8. RP and ECOR values collected on C1-coated specimens, under both corrosion (solid lines) and tribocorrosion (dotted lines) conditions.
RP/ ohm cm2
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3
Current Density (A cm-2)
Time / h Rp - C1 - 20rpm Rp - C1 - 20rpm - 2N Ecor - C1 - 20rpm Ecor - C1 - 20rpm - 2N
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(Fig. 13c, d, where the arrows indicate some pores filled by the wear debris). On C1 coating (Fig. 13e) or on C2 at 2 N and 20 rpm (Fig. 13f) this phenomenon is definitely less evident, as many large craters are there still present. Then, on C1 coating under tribocorrosion conditions incomplete or negligible pore obstruction occurs, while on C2 coating at sufficiently high load and rotation speed a more continuous tribofilm forms. It partially obstructs the access of the aggressive solution towards the steel substrate, so explaining the lower iCOR recorded on C2-coated electrodes under heavier wear conditions (10 N, 100 rpm) with respect to those measured under milder ones (2 N, 20 rpm).
3.6. Profilometry After 3 day exposures to tribocorrosion conditions, wear rates were calculated and collected in Table 4. The table shows that, at constant load, the wear rates always diminish on going from 20 to 100 rpm. A fivefold higher rotation speed corresponds to a fivefold longer sliding distance, as the tests always lasted 3 days. At low rotation speed (and short sliding distance travelled), the wear track is relatively rough and it is reasonable to obtain high wear volume per meter by the removal of surface asperities. On the contrary, at high rotation speed (that is after long sliding distances) the wear track becomes smoother and smoother, producing lower wear rates. However, it must also be mentioned that, according to the Stribeck curve [31,32], during sliding of liquid-lubricated surfaces under mixed lubrication (ML) regime the friction coefficient decreases at increasing sliding velocities. Then, it cannot be excluded that a decreasing friction coefficient can contribute to the observed decrease in wear rates, going from 20 to 100 rpm.
-0.6 -0.7
Table 3 Electrochemical parameters calculated from the polarization curves recorded after 3 days of immersion, under pure corrosion and tribocorrosion conditions.
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Time / h Rp - C2 - 20rpm Rp - C2 - 100rpm - 10N Ecor - C2 - 20rpm - 2N
Rp - C2 - 20rpm - 2N Ecor - C2 - 20rpm Ecor - C2 - 100rpm - 10N
Fig. 9. RP and ECOR values collected on C2-coated specimens, under both corrosion (solid lines) and tribocorrosion (dotted lines) conditions.
20 rpm – 2 N
100 rpm – 10 N
ECOR/ VSCE
iCOR/ A cm−2
ba/ Vdec−1
ECOR/ VSCE
iCOR/ A cm−2
ECOR/ VSCE
iCOR/ A cm−2
−0.66 −0.61 −0.63 −0.06*
5·10−5 1.0·10−5 9·10−6 2.5·10−6*
0.080 0.240 0.100 0.460*
−0.61 −0.61 −0.71 –
2.3·10−4 1.5·10−5 4.6·10−5 –
– −0.60 −0.68 –
– 2.0·10−5 2.2·10−5 –
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80
Steel C1 C2 C3
*measured after 1 day of exposure.
C. Monticelli et al. / Surface & Coatings Technology 205 (2011) 3683–3691
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Current Density (A cm-2) Fig. 11. Polarization curves recorded on C1-coated specimens, after 3 days of immersion. Steel under tribocorrosion conditions is also reported, as a reference.
Table 4 also evidences that on both coatings at 20 rpm the wear rates markedly decrease on passing from 2 to 5 N, then they remain more or less constant at 10 N. Even at 100 rpm, the wear rates are constant or slightly decrease on going from 5 to 10 N. This is likely due to the formation of a tribofilm at loads higher than or equal to 5 N. As documented by surface observations (Fig. 13), this tribofilm is connected to the plastic deformation of the wear debris which fills up and partially obstructs the surface pores so reducing the rate of material removal. 4. Discussion C1 and C2 ceramic coatings are characterized by porosity values of 7.5 and 9.7% respectively, frequently encountered on these coating types [3,11,24], and by inter- and intralamellar cracks. Porosity is mainly due to splat stacking faults and gas entrapment, while interand intralamellar cracks are often caused by scarce intersplat cohesion and thermal contraction respectively [3]. The Cr2O3 coating shows a finer structure, a slightly lower porosity and a higher microstructural homogeneity than the Al2O3/13% TiO2 coating. However, when exposed to a neutral chloride solution, the former coating exhibits corrosion protection properties quite similar to those of the latter one. In fact, both coatings only partially protect steel substrate from corrosion, in spite of the presence of a rather compact interlaying bond layer, showing passive corrosion behaviour. This suggests that the aggressive solution penetrates through the top and bond coatings and reaches steel, where a corrosion process sets
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4
1E-3
up, stimulated by the galvanic coupling with the noble bond coating material. As a result, coated steel specimens undergo five times smaller corrosion rates, than those measured on bare steel specimens. Polarization curve analysis suggests that the corrosion of C1- and C2-coated specimens is mainly under cathodic control. This means that the substrate corrosion rate depends on the rate of oxygen diffusion towards the metallic surfaces through the coating pores, which in turn is affected by the coating thickness, the defect population and the interconnection degree and tortuosity of the defects themselves. On the contrary, oxygen diffusion rate inside the pores cannot be stimulated by higher electrode rotation speeds. This explains the invariance of the corrosion rates of coated specimens under free corrosion conditions, at 20 or 100 rpm. Fig. 10 shows that the porosity difference characterizing C1 and C2 coatings does not significantly affects oxygen diffusion. On both C1- and C2-coated specimens, sliding stimulates both the anodic and the cathodic process and induces a limited increase of the measured corrosion rates. On C1-coated specimens, a small iCOR increase is measured by passing from pure corrosion, to tribocorrosion at 2 N and 20 rpm and then to tribocorrosion at 10 N and 100 rpm. On the contrary, on C2-coated specimens under wear, the most severe corrosion attack is observed at 2 N and 20 rpm. As a consequence, C1 coating is the most protective one under mild wear conditions, while the coatings afford quite similar performances under more severe conditions. These findings can be explained by analysing how sliding wear modifies the coating surfaces in the track. It has been noticed that on both coatings abrasion of surface asperities occurs, with a consequent surface polishing. This smoothing effect is more evident at high load and rotation speed and particularly on C2 coating, because it is characterized by the lowest microhardness value. In this coating, the wear debris detached at 10 N and 100 rpm is found to partially fill up the surface porosity by plastic deformation and embedment in the coating. This obstruction justifies the lower iCOR value measured on C2-coated specimens under these experimental conditions, with respect to those evaluated at 2 N and 20 rpm. Microscale plastic deformation of Cr2O3 and Al2O3/13% TiO2 coatings under dry sliding wear has been often reported in the literature [3,11,33] and it has been attributed to plastic slipping at splat boundaries. The progressive polishing effect observed under sliding wear may justify the decrease of the wear rates, going from low to high rotation speeds, as the variation of this parameter implies an increase in the sliding distances which in turn involves smoother and smoother wear paths. A decrease in the friction coefficient going from 20 to 100 rpm could also occur and play a role in reducing the measured wear rates [31,32]. The observed formation of a tribofilm is the cause of the significant decrease of the wear rates at loads higher than 2 N. In fact when the load applied is sufficiently high the smearing of the wear debris on the coating surfaces limits the material removal. 5. Conclusions
2
2 - C2 20 rpm 2N
-0.4
3689
1E-2
Current Density (A cm-2) Fig. 12. Polarization curves recorded on C2-coated specimens, after 3 days of immersion. Steel under tribocorrosion conditions is also reported, as a reference.
- Cr2O3 and Al2O3/13% TiO2 coatings only partially protect steel substrate from corrosion. The corrosion process is mainly under cathodic control and is stimulated by the galvanic coupling with the interlaying noble Ni/20% Cr bond coat. - Tribocorrosion of Cr2O3- and Al2O3/13% TiO2-coated specimens stimulates both the anodic and the cathodic process so increasing the substrate corrosion rates. On the former coating, the highest corrosion rates are obtained under the most severe wear conditions (10 N and 100 rpm), while on the latter one the highest corrosion rates are measured under the mildest wear conditions tested (2 N and 20 rpm). - At the end of tribocorrosion tests, the wear mechanism on the ceramic coatings has been investigated. It involves polishing of the surface asperities and, at relatively high loads (higher than 2 N), formation of a tribofilm due to the smearing of the wear debris on the coating surface so inducing a partial fill up of the surface pores.
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Fig. 13. Top surface images of C1 (a,e) and C2 (b,c,d,f) coatings in the wear scars produced under tribocorrosion conditions at 10 N and 100 rpm (a,b,c,d,e) or at 2 N and 20 rpm (f). With the exception of Fig. 13d (VPSEM micrograph), all other images are optical micrographs. The sliding direction is always horizontal.
Table 4 Wear rates (10−6 mm3/N m) calculated at the end of 3 days exposures under tribocorrosion conditions. Cr2O3-coated steel
2N 5N 10 N
Al2O3/13% TiO2-coated steel
20 rpm
100 rpm
20 rpm
100 rpm
4.9 ± 0.1 1.8 ± 0.4 1.5 ± 0.1
– 0.4 ± 0.1 0.3 ± 0.1
7.5 ± 2.5 3.0 ± 0.6 3.0 ± 0.5
– 1.3 ± 0.3 0.7 ± 0.1
- A more continuous tribofilm is observed on Al2O3/13% TiO2 coating at 10 N and 100 rpm, which partially limits the corrosion of the steel substrate. As a consequence, under mild wear conditions, Cr2O3 coating is the most protective one, while under the most severe wear conditions, the coatings afford quite similar performances. - Wear rates of the coatings tend to decrease both at increasing rotation speed (from 20 to 100 rpm) and at increasing the load (particularly from 2 to 5 N).
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Acknowledgements The authors wish to thank Ing. Marco Caliari and Zocca Officine Meccaniche (Funo, BO), for the thermally sprayed coating manufacturing. References [1] G.W. Stachowiak, A.W. Batchelor, Engineering tribology, third ed, ButterworthHeinemann, Boston, 2005. [2] K.A. Habib, J.J. Saura, C. Ferrer, M.S. Damra, E. Giménez, L. Cabedo, Surf. Coat. Technol. 201 (2006) 1436. [3] G. Bolelli, V. Cannillo, L. Lusvarghi, T. Manfredini, Wear 261 (2006) 1298. [4] L. Fedrizzi, S. Rossi, F. Bellei, F. Deflorian, Wear 253 (2002) 1173. [5] M. Laribi, A.B. Vannes, D. Treheux, Wear 262 (2007) 1330. [6] T.A. Stolarski, S. Tobe, Wear 249 (2001) 1096. [7] F. Rasteger, A.E. Craft, Surf. Coat. Technol. 61 (1993) 36. [8] H.-S. Ahn, I.-W. Lyo, D.-S. Lim, Surf. Coat. Technol. 133–134 (2000) 351. [9] E. Celik, C. Tekmen, I. Ozdemir, H. Cetinel, Y. Karakas, S.C. Okumus, Surf. Coat. Technol. 174–175 (2003) 1074. [10] P. Ernst, G. Barbezat, Surf. Coat. Technol. 202 (2008) 4428. [11] H.-S. Ahn, O.-K. Kwon, Wear 225–229 (1999) 814. [12] F. Vargas, H. Ageorges, P. Fournier, P. Fauchais, M.E. López, Surf. Coat. Technol. 205 (2010) 1132.
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Corrosion and tribocorrosion behaviour of thermally sprayed ceramic coatings on steel C. Monticelli ⁎, A. Balbo 1, F. Zucchi 2 Centro di Studi sulla Corrosione “A. Daccò”, Università degli Studi di Ferrara, Via Saragat 4A, 44122 Ferrara, Italy
a r t i c l e
i n f o
Article history: Received 29 November 2010 Accepted in revised form 10 January 2011 Available online 15 January 2011 Keywords: Plasma spraying Ceramic coatings Corrosion Tribocorrosion Wear rate Chloride
a b s t r a c t This research aims at investigating the corrosion and tribocorrosion behaviour of thermally sprayed ceramic coatings deposited on steel specimens and exposed to a 3.5% NaCl solution. The coatings have been prepared by plasma spraying Cr2O3 and Al2O3/13% TiO2 powders on a Ni/20% Cr bond coating. Combined wear– corrosion conditions have been achieved by sliding an alumina antagonist on the lateral surface of coated steel cylinders, during their exposure to the aggressive solution. Polarization resistance values monitored during 3 days exposures and polarization curves recorded at the end of the immersion period show that both coatings only partially protect steel substrate from corrosion. Sliding conditions (under 2 N load and 20 rpm or 10 N and 100 rpm) induce a limited increase of the substrate corrosion rates, likely as a consequence of an increase in the defect population of the ceramic coatings. On Cr2O3-coated specimens, tribocorrosion is more severe at 10 N and 100 rpm, while on Al2O3/13% TiO2coated specimens, a stronger corrosion attack is achieved at 2 N and 20 rpm. Profilometer analysis and wear track observations by optical and scanning electron microscopes evidence that on both coatings abrasion of the surface asperities produce both a surface polishing effect and, at high loads, the formation of a tribofilm, more continuous on Al2O3/13% TiO2. On this coating the tribofilm reduces the amount of surface defects and limits the corrosion attack to a certain extent. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Wear is the major cause of material wastage and loss of mechanical performances and friction is the principal cause of wear and energy dissipation [1]. Thermally sprayed ceramic coatings on steel represent an efficient and economic way to improve the wear resistance of mechanical components, when common thermal treatments (such as quenching and tempering,…) or thermo-chemical processes (such as carburizing, nitriding…) result inadequate. In particular, plasma sprayed Cr2O3 and Al2O3–TiO2 coatings can be an excellent choice, providing protection against abrasive wear and resistance to galvanic and high temperature corrosion [2]. They often result to be superior to traditional wear resistant hard chromium [3,4] and molybdenum [5,6] coatings. A low as-sprayed surface roughness characterizes Cr2O3 coatings and represents a very important technological feature, because it reduces the number of post-deposition mechanical treatments necessary in many applications [3]. Cr2O3 coatings have low friction coefficients and can be conveniently deposited on piston
⁎ Corresponding author. Tel.: + 39 0532 455136; fax: + 39 0532 455011. E-mail addresses: [email protected] (C. Monticelli), [email protected] (A. Balbo), [email protected] (F. Zucchi). 1 Tel.: + 39 0532 455134. 2 Tel.: + 39 0532 455135. 0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2011.01.023
ring and cylinder liners in the automotive industry [7–9], where they reduce fuel and oil consumption and increase the engine life [10]. If compared to Cr2O3 coatings, Al2O3 and Al2O3–TiO2 ones are less expensive and more biocompatible and can be applied in the food and medicine packaging industry because they ensure the absence of heavy metal contamination [3]. Many papers concern the wear behaviour of plasma sprayed oxide coatings and investigate the corresponding wear mechanisms. Under dry sliding conditions, both Cr2O3 and Al2O3–TiO2 coatings are reported to form a tribofilm. On the former coating, this film is formed by plastically deformed and compacted wear debris, responsible of the low measured friction coefficients [3]. XPS analysis reveals that the film formed at room temperature is constituted by CrO3 and Cr2O3 [11]. On the contrary, on Al2O3–TiO2 coatings the tribofilm has a rather loose structure which does not adequately protect the underlying material from wear [3]. On these ceramic coatings different wear mechanisms are detected, involving abrasive wear [12,13] delamination of weakly adherent successive lamellae [14], crack nucleation and delamination [15] and adhesive wear [14]. Only a few studies involve the corrosion behaviour of thermally sprayed oxide coatings on steel [16–19]. They show the clear dependence of the corrosion resistance of the coated materials from the coating porosity [20,21], which is reduced by a proper choice of the spraying parameters [22,23]. In some cases, interconnected porosity decreases at increasing coating thickness [17]. However,
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the coating porosity is often found to increase with the deposition time (and coating thickness) [22] and higher corrosion resistance can be achieved with thin coatings [19,21]. Porosity also depends on the composition of the sprayed powders. As an example, ZrO2 addition to the Al2O3 powder tends to increase the coating porosity [17], while contrasting results are obtained by TiO2 addition [3,24]. To the authors' knowledge no literature information is available concerning tribocorrosion of ceramic coatings. As the concomitant presence of wear and corrosion processes usually induces a synergistic stimulation of material degradation [25], it has been reputed interesting a characterization of both the corrosion and tribocorrosion behaviour of Cr2O3- and Al2O3–13%TiO2-coated steel specimens in the presence of widespread aggressive species, such as chlorides.
2. Materials and methods Onto cylindrical AISI-SAE 1040 steel specimens (10 mm in height; 20 mm in diameter; nominal composition: 0.37–0.44% C; 0.5–0.8% Mn; 0.15–0.4% Si; balance Fe), two commercial 250 μm-thick ceramic coatings were prepared, that is a Cr2O3 coating (indicated as C1) and a Al2O3/13% TiO2 coating (indicated as C2). As part of a standard commercial deposition process, both coatings (obtained by air plasma spray (APS) technology) were deposited onto a 20 μm-thick Ni/20% Cr bond coat (also obtained by APS), which aimed at improving the ceramic material adhesion onto the steel substrate. Details of the feedstock powders are reported in Table 1, while the deposition parameters are confidential. The phase transformations occurring during the thermal spray of the ceramic powders were investigated by comparing the X-ray diffractograms of the powders to those of the coatings. Moreover, properly polished cross sections of the powders and coatings were observed by a Scanning Electron Microscope equipped by Variable Pressure technology (VPSEM) and Energy Dispersion Spectroscopy (EDS) microprobe, to investigate their microstructures. On the polished cross-sections of the coatings, the porosity (through an image analysis software), and the Vickers microhardness (at 0.3 kg load) were also measured. The coated steel specimens had a central threaded hole permitting them to be mounted on a steel shaft which afforded the electrical contact for electrochemical tests and let the specimen rotate at 20 or 100 rpm. Before testing, the coated specimens underwent a commercial grinding process, inducing a surface roughness, Ra, of about 0.1 and 0.2 μm on C1 and C2 coatings, respectively. A grinding process is usually applied to industrial components to reduce surface roughness and increase the wear resistance. As a reference, bare steel and steel specimens coated by a thick Ni/20% Cr bond coat (250 μm, C3 coating) were also tested. They were ground by emery papers down to no. 600 and carefully degreased before exposure to the aggressive solutions. This solution was a naturally aerated 3.5% NaCl aqueous solution (pH 6–6.5), kept at room temperature (22 ± 1 °C).
Table 1 Feedstock powders used to produce top and bond coatings. Cr2O3 Type
Amperit 707.001 by Starck Composition 0.06% SiO2, 0.03% Fe2O3, b 0.02% TiO2, and balance Cr2O3 Particle dimensions −45 + 22.5 μm
Al2O3/13% TiO2
Ni/20% Cr
FST C-335.23 by Flame Spray Technologies 13.12% TiO2, 0.22% ZrO2, 0.10% SiO2, 0.09% MgO, 0.07% CaO, 0.39% other oxides, and balance Al2O3 −45 + 15 μm
Metco 43CNS by Sulzer 19.07% Cr, 1.1% Si, 0.4% Fe, 0.02% C, and balance Ni −106 + 45 μm
For tribocorrosion tests (Fig. 1), a non-commercial tribometer induced a sliding-type wear on the lateral surface of the rotating cylinders during exposure to the aggressive solution (rotation speeds: 20 rpm or 100 rpm, involving sliding velocities of 21.5 mm s− 1 or 107 mm s−1, respectively). The normal loads (L), applied through the flat surface of an alumina cylinder with a 2 mm diameter, were 2–10 N. The nominal hardness of the α-alumina counterbody was 1950 HV. During corrosion tests, the surface exposed to the aggressive solution was the whole lateral surface of the cylindrical specimens (6.4 cm2), while during the tribocorrosion tests the lateral surface was partially screened by Lacomit varnish, in order to expose only the rubbed surface (1.6 cm2) to the solution. The evolution of the corrosion process was monitored during 3 days of exposure to the aggressive solution, by measuring the linear polarization resistance (RP) values (in the potential range from −7 mV up to+7 mV vs ECOR, with a potential scan rate of 0.1 mV s−1). At the end of the 3 days immersion, under both corrosion and tribocorrosion conditions, potentiodynamic polarization curves were recorded at a scan rate of 1 mV s−1, always starting from the ECOR value. All the potentials in the text are referred to the Saturated Calomel Electrode (SCE). Some specimens exposed to tribocorrosion conditions were not subjected to destructive electrochemical tests. On these specimens, wear rates were calculated as WR = VW / L D, where VW is the wear volume, L is the normal load and D corresponds to the sliding distance. VW was obtained by multiplying the area of the wear track cross section (measured as an average from the wear track profiles at five positions along the specimen circumference by a Hommelwerk T2000 profilometer with a TK300 piezoelectrical tip) by the circumference of the cylindrical specimens. Each wear rate value was the average of three measurements. The morphology of the corrosion and tribocorrosion attack was characterized by optical and scanning electron microscope (SEM) observations. 3. Results 3.1. Powder characterization A fused and crushed Cr2O3 powder is the raw material for C1 coating (Fig. 2a). SEM observation of polished cross sections of the particles indicates that they are monophasic (Fig. 2b) and XRD analysis reveals that they are constituted by eskolaite (hexagonal Cr2O3, Fig. 3a). C2 coating has been produced by a fused and crushed Al2O3–13% TiO2 powder, composed of irregular and angular particles (Fig. 4a), biphasic in nature, as evidenced by SEM observation (Fig. 4b). XRD analysis detects the presence of crystalline α-Al2O3 and Al2TiO5 (Fig. 3b, upper diffractogram), as predicted by the Al2O3–TiO2 phase diagram [26], which indicates the stability of two structural components, that is α-Al2O3 and the eutectic α-Al2O3–Al2TiO5. 3.2. Coating characterization SEM micrographs shown in Figs. 5a and 6 evidence the lamellar structure of C1 and C2 coatings, typical of thermally sprayed coatings. C1 shows a fine homogeneous structure, still constituted by hexagonal Cr2O3 (Fig. 3a, lower diffractogram) and characterized by small interand intralamellar cracks and finely distributed pores (Fig. 5a). An irregular NiCr bond coat facilitates the adhesion of the ceramic coating to the substrate (Fig. 5b), in spite of the presence of rare pores and sandblast residues, visible at the steel/bond coat interface. With respect to C1, C2 coating exhibits a coarser structure, with thicker lamellae and longer inter- and intralamellar cracks. Moreover, it shows isolated light grey titanium-rich splats (Fig. 6), suggesting the
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Fig. 1. Schematic view of the experimental arrangement for wear corrosion tests. W.E., C.E. and R.E. respectively correspond to the working, counter and reference electrode; Hg indicates the mercury pool which electrically connects the rotating shafts (and the W.E.) to the apparatus for electrochemical measurements.
inhomogeneous interdissolution of α-Al2O3 and Al2TiO5 phases, during the spraying process. In agreement with the literature information [3], after the spraying process only metastable γ-Al2O3 and traces of α-Al2O3 are detected as crystalline phases, while aluminium titanate remains amorphous (Fig. 4 b, lower curve). The
bond coat morphology under C2 coating is equal to that exhibited in Fig. 5b. Coating porosity and Vickers microhardness values are reported in Table 2. The results obtained indicate that C1 exhibits a slightly lower average porosity with a lower standard deviation than those of C2 coating. Moreover, the average microhardness of C1 is much higher than that of C2, because of its higher bulk hardness value and lower porosity. The hardness of Al2O3–13% TiO2 coating is lower than that of the alumina counterbody, because in the coating many factors, such as porosity, α- to γ-alumina conversion during thermal spraying and TiO2 addition, contribute to reduce the material hardness [3]. Table 2 also shows the porosity percentage of the bond coating evaluated under SEM observations. It is lower than that of the ceramic coatings, in agreement with literature information [27–29]. 3.3. Linear polarization resistance measurements
Fig. 2. VPSE SEM micrographs showing a) as-received Cr2O3 powder and b) Cr2O3 powder cross section.
Fig. 7 shows the RP and ECOR values collected on steel during exposures to the aggressive solution under a rotation speed of 20 rpm. Low RP values are recorded after 1 h of immersion (540 Ω cm2) which remain more or less constant throughout the 3 days of exposure. After 1 h of immersion, ECOR is about −0.560 VSCE, that is typical of steel under active corrosion in neutral chloride solutions [30]. Then, it undergoes a moderate reactivation, with ECOR values decreasing down to −0.670 VSCE. During the immersion, steel is slowly covered by reddish iron hydroxide corrosion products. Under the same exposure conditions, a quite different behaviour is exhibited by Ni/20% Cr bond coat material (Fig. 7). In fact, specimens carrying a thick bond coating (C3-coated specimens) show RP (higher than 105 Ω cm2) and ECOR (−0.11/−0.06 VSCE) values typical of a passive metal. Tribocorrosion (TC) at 2 N load and 20 rpm stimulates the corrosion attack on steel (Fig. 7). In fact, on this material it induces RP values (about 200 Ω cm2, at the end of the immersion period) lower than those measured under pure corrosion conditions, with slightly ennobled ECOR values (−0.61 VSCE, after 3 days of immersion). Fig. 8 shows the time dependence of RP and ECOR values collected on C1-coated specimens under free corrosion conditions at 20 rpm. RP values are higher than those recorded on steel (5.5 ÷ 3.5 kΩ cm2), but they are definitely lower than those measured on C3-coated specimens. The corresponding ECOR values (about −0.64 VSCE) are close to those exhibited by bare steel specimens. Chromia, the ceramic
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a) Counts / Arbitrary units
Hex-Cr2O3 All peaks
20
40
60
80
100
120
2-Theta / degree
b)
α - Al2O3
Counts / Arbitrary units
Al2TiO5 γ - Al2O3
20
40
60
80
100
120
2-Theta / degree Fig. 3. XRD diffractograms recorded on Cr2O3 powder (a, upper curve) and coating (a, lower curve) and on Al2O3–13% TiO2 powder (b, upper curve) and coating (b, lower curve).
material of C1 coating, is an electrically insulating material which cannot be involved in any electrochemical process. So on coated steel specimen, corrosion may affect either the bond coat or the metallic substrate or both. However, at the measured ECOR values, only steel may undergo a corrosion attack, while the bond coat material results in being galvanically protected. This suggests that the ceramic coating and the underlying thin bond coating cannot hinder the aggressive solution penetration through the coating pores and the onset of steel corrosion at the pore bottom. The limited substrate area in contact with the solution justifies the relatively high RP values. Fig. 8 also evidences that the corrosion of C1-coated steel specimens is not affected by a variation in the rotation speed from 20 to 100 rpm. During 3 days of exposure of C1-coated specimens to tribocorrosion conditions at 2 N and 20 rpm, a diminution of RP values by only a factor of 2 or less is detected and no further RP reduction is observed by increasing the load up to 10 N and the rotation speed up to 100 rpm (Fig. 8). This slight stimulation of the corrosion process is likely connected to an increase in the population of the coating defects, which permits an easier access of the solution to the steel substrate. The behaviour of C2-coated specimens is similar to that exhibited in the presence of C1 coatings (Fig. 9). Under pure corrosion conditions, at both 20 and 100 rpm, RP values of about 7.8 kΩ cm2 are measured which decrease down to about 4.3 kΩ cm2, at the end of the tests. The ECOR values remain close to −0.64 VSCE, throughout the immersion period. This again indicates corrosion of the steel substrate. Under tribocorrosion conditions, a RP decrease is measured
Fig. 4. VPSE SEM micrographs showing a) as-received Al2O3–13% TiO2 powder and b) Al2O3–13% TiO2 powder cross section.
which is more marked at 2 N and 20 rpm (1.4 kΩ cm2, after 3 days) than at higher load and rotation speed (2.8 kΩ cm2, after 3 days at 10 N and 100 rpm). Under tribocorrosion conditions the slow formation of reddish iron corrosion products on both C1- and C2-coated specimens and in the solution is clearly visible. 3.4. Polarization curves The polarization curves recorded on the studied materials at the end of the immersion period in the neutral chloride solution are collected in Fig. 10. Table 3 reports the corresponding ECOR and iCOR values, as estimated by the Tafel method by extrapolation from the cathodic polarization curve. As anticipated by the results of the linear polarization resistance measurements, steel exhibits an active corrosion behaviour, indicated by its active ECOR value and low anodic overvoltage. The anodic Tafel slope (ba = 0.080 V decade−1, Table 3) is slightly higher than that expected for active dissolution of iron (0.060 V decade−1), because of the formation of insoluble surface corrosion products which moderately hinder the anodic process. Under free corrosion conditions, C3-coated specimens (with a thick Ni–20% Cr coating) are under passive corrosion conditions (high ba and low iCOR values, Table 3) and can suffer pitting corrosion at potentials nobler than +0.45 VSCE (Fig. 10). As shown in the previous section, C1- and C2-coated specimens have ECOR values almost equal to that of steel (Fig. 10). However, they exhibit higher anodic overvoltages (ba values in Table 3) and lower anodic and cathodic currents than those recorded on bare steel
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Table 2 Porosity and Vickers microhardness values for the studied coatings.
C1 C2 Bond coat
Fig. 5. VPSE SEM micrographs showing cross section microstructure of C1 coating (a) and underlaying bond coat (b).
specimens. These observations are in agreement with the hypothesis that the anodic process is connected to steel dissolution through the double coating. As on C1- and C2-coated specimens, both steel dissolution and oxygen reduction reaction occur at the bottom of the coating pores, the rates of these reactions are limited by the restricted access of the aggressive solution to the underlying metals. At the bottom of the ceramic coating pores, the noble interlaying bond coat forms efficient cathodic regions (as shown by the cathodic polarization curve on C3-coated specimens in Fig. 10) of relatively high area (depending on the ceramic coating porosity). Then the bond
Porosity/%
HV0.3/kg mm−2
7.5 ± 1.2 9.7 ± 2.8 1.1 ± 1.0
1556 ± 104 1099 ± 96
coat can induce galvanic corrosion on the small substrate regions in contact with the aggressive solution through the porosity of the bond coat itself. The relatively high cathodic to anodic metal area ratio likely stimulates the galvanic attack. With respect to C1-coated specimens, C2-coated ones exhibit lower anodic slopes and higher anodic current densities, while the cathodic polarization curves and the iCOR values are quite close to each other (Table 3). The differences in the anodic curve suggest that the slightly higher porosity of C2 coatings facilitates the migration of iron corrosion products through the coatings driven by the anodic polarization. The presence of quite similar cathodic characteristics and iCOR values suggests that on these specimens the corrosion process is mainly under cathodic control. Fig. 10 also shows the polarization curves recorded on steel under tribocorrosion at 2 N load and 20 rpm. The analysis of these curves evidences that sliding wear speeds up the cathodic reaction on steel and causes the slight ECOR ennoblement (to −0.61 VSCE), already noticed during the discussion of the polarization resistance measurements. This is due to the stimulation of oxygen diffusion towards the metallic surface, caused by the continuous removal of corrosion products from the steel surface, which induces a fivefold increase of iCOR (Table 3). Wear application on C1- and C2-coated specimens slightly stimulates both the anodic and the cathodic reactions of steel corrosion (Figs. 11 and 12). In the case of C1 coating (Fig. 11), the stimulation is quite limited and iCOR reaches its maximum value, at the highest load and sliding velocity applied (10 N, 100 rpm, Table 3). Under this condition, iCOR is twice the value measured under pure corrosion conditions. On the contrary, in the case of C2 coating (Fig. 12), tribocorrosion under 2 N and 20 rpm is more severe than that experienced at 10 N load and 100 rpm. In fact, under the former condition iCOR is about fivefold higher than that obtained in pure corrosion, while under the latter it is about twice.
0
1000000 Rp - Steel - 20rpm Rp - Steel - 20rpm - 2N Rp - C3 - 20rpm Ecor - Steel - 20rpm Ecor - Steel - 20rpm - 2N Ecor - C3 - 20rpm
-0.2 -0.3 -0.4
10000
-0.5
ECOR / VSCE
RP/ ohm cm2
100000
-0.1
-0.6
1000
-0.7 -0.8
100 0
20
40
60
80
Time / h Fig. 6. VPSE SEM micrograph showing cross section microstructure of C2 coating.
Fig. 7. RP and ECOR values collected on steel and bond-coated (C3) specimens (solid lines refer to pure corrosion conditions, while dotted lines to tribocorrosion conditions).
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10000
1
0
1 - Steel 20 rpm 2 - Steel 20 rpm 2N 3 - C1 20 rpm 4 - C2 20 rpm 5 - C3 20 rpm
0.8
-0.1
RP/ ohm cm2
1000
-0.4 -0.5
ECOR / VSCE
-0.3
Potential (VSCE)
0.6 -0.2
0.4
5
3
0.2
4
0
1
-0.2
2
5
-0.4 -0.6
-0.6
-0.8 -0.7 100
-1 1E-7
-0.8 0
20
40
60
80
3.5. Wear track observations The surface micrographs of the ceramic coatings tested under different tribocorrosion conditions are shown in Fig. 13. Fig. 13a and b exhibits the surface aspect of both coatings at the wear track borderline after tests at 10 N and 100 rpm: inside the track (lower part of the pictures) the surface is much smoother than outside, as a consequence of the removal of surface asperities, left by the grinding process. Fig. 13a shows that on the hardest C1 coating the surface is rougher than that on C2 coating (Fig. 13b) and the holes in the track are due to both the coating porosity and the troughs left by the commercial grinding process. On C2 coating (Fig. 13b), the polishing effect is more pronounced and large surface regions of the track are apparently poreless. However, surface observations at higher magnification reveal that the surface pores are partially obstructed by a tribofilm produced by the smearing of the wear debris on the surface
10000
0 -0.1 -0.2
1000
-0.4 -0.5
ECOR / VSCE
-0.3
1E-5
1
1E-4
2
1E-3
1E-2
1E-1
Fig. 10. Polarization curves recorded on steel and coated specimens, after 3 days of immersion (after 1 day of immersion, in the case of the C3-coated specimens).
Rp - C1 - 100rpm Rp - C1 - 100rpm - 10N Ecor - C1 - 100rpm Ecor - C1 - 100rpm - 10N
Fig. 8. RP and ECOR values collected on C1-coated specimens, under both corrosion (solid lines) and tribocorrosion (dotted lines) conditions.
RP/ ohm cm2
1E-6
3
Current Density (A cm-2)
Time / h Rp - C1 - 20rpm Rp - C1 - 20rpm - 2N Ecor - C1 - 20rpm Ecor - C1 - 20rpm - 2N
4
(Fig. 13c, d, where the arrows indicate some pores filled by the wear debris). On C1 coating (Fig. 13e) or on C2 at 2 N and 20 rpm (Fig. 13f) this phenomenon is definitely less evident, as many large craters are there still present. Then, on C1 coating under tribocorrosion conditions incomplete or negligible pore obstruction occurs, while on C2 coating at sufficiently high load and rotation speed a more continuous tribofilm forms. It partially obstructs the access of the aggressive solution towards the steel substrate, so explaining the lower iCOR recorded on C2-coated electrodes under heavier wear conditions (10 N, 100 rpm) with respect to those measured under milder ones (2 N, 20 rpm).
3.6. Profilometry After 3 day exposures to tribocorrosion conditions, wear rates were calculated and collected in Table 4. The table shows that, at constant load, the wear rates always diminish on going from 20 to 100 rpm. A fivefold higher rotation speed corresponds to a fivefold longer sliding distance, as the tests always lasted 3 days. At low rotation speed (and short sliding distance travelled), the wear track is relatively rough and it is reasonable to obtain high wear volume per meter by the removal of surface asperities. On the contrary, at high rotation speed (that is after long sliding distances) the wear track becomes smoother and smoother, producing lower wear rates. However, it must also be mentioned that, according to the Stribeck curve [31,32], during sliding of liquid-lubricated surfaces under mixed lubrication (ML) regime the friction coefficient decreases at increasing sliding velocities. Then, it cannot be excluded that a decreasing friction coefficient can contribute to the observed decrease in wear rates, going from 20 to 100 rpm.
-0.6 -0.7
Table 3 Electrochemical parameters calculated from the polarization curves recorded after 3 days of immersion, under pure corrosion and tribocorrosion conditions.
-0.8
100 0
20
40
60
Time / h Rp - C2 - 20rpm Rp - C2 - 100rpm - 10N Ecor - C2 - 20rpm - 2N
Rp - C2 - 20rpm - 2N Ecor - C2 - 20rpm Ecor - C2 - 100rpm - 10N
Fig. 9. RP and ECOR values collected on C2-coated specimens, under both corrosion (solid lines) and tribocorrosion (dotted lines) conditions.
20 rpm – 2 N
100 rpm – 10 N
ECOR/ VSCE
iCOR/ A cm−2
ba/ Vdec−1
ECOR/ VSCE
iCOR/ A cm−2
ECOR/ VSCE
iCOR/ A cm−2
−0.66 −0.61 −0.63 −0.06*
5·10−5 1.0·10−5 9·10−6 2.5·10−6*
0.080 0.240 0.100 0.460*
−0.61 −0.61 −0.71 –
2.3·10−4 1.5·10−5 4.6·10−5 –
– −0.60 −0.68 –
– 2.0·10−5 2.2·10−5 –
20 rpm
80
Steel C1 C2 C3
*measured after 1 day of exposure.
C. Monticelli et al. / Surface & Coatings Technology 205 (2011) 3683–3691
0 1 - C1 20 rpm
1
3
2 - C1 20 rpm 2N
Potential (VSCE)
-0.2
3 - C1 100 rpm 10 4 - steel 20 rpm
2
-0.4 4
-0.6
-0.8 1 2 3
-1 1E-6
1E-5
4
1E-4
1E-3
1E-2
Current Density (A cm-2) Fig. 11. Polarization curves recorded on C1-coated specimens, after 3 days of immersion. Steel under tribocorrosion conditions is also reported, as a reference.
Table 4 also evidences that on both coatings at 20 rpm the wear rates markedly decrease on passing from 2 to 5 N, then they remain more or less constant at 10 N. Even at 100 rpm, the wear rates are constant or slightly decrease on going from 5 to 10 N. This is likely due to the formation of a tribofilm at loads higher than or equal to 5 N. As documented by surface observations (Fig. 13), this tribofilm is connected to the plastic deformation of the wear debris which fills up and partially obstructs the surface pores so reducing the rate of material removal. 4. Discussion C1 and C2 ceramic coatings are characterized by porosity values of 7.5 and 9.7% respectively, frequently encountered on these coating types [3,11,24], and by inter- and intralamellar cracks. Porosity is mainly due to splat stacking faults and gas entrapment, while interand intralamellar cracks are often caused by scarce intersplat cohesion and thermal contraction respectively [3]. The Cr2O3 coating shows a finer structure, a slightly lower porosity and a higher microstructural homogeneity than the Al2O3/13% TiO2 coating. However, when exposed to a neutral chloride solution, the former coating exhibits corrosion protection properties quite similar to those of the latter one. In fact, both coatings only partially protect steel substrate from corrosion, in spite of the presence of a rather compact interlaying bond layer, showing passive corrosion behaviour. This suggests that the aggressive solution penetrates through the top and bond coatings and reaches steel, where a corrosion process sets
0 1 - C2 20 rpm
Potential (VSCE)
-0.2
1
3 - C2 100 rpm 10 N 4 - steel 20 rpm 2N 3
4
1
-0.6 2
-0.8
-1 1E-6
1
1E-5
3
1E-4
2
4
1E-3
up, stimulated by the galvanic coupling with the noble bond coating material. As a result, coated steel specimens undergo five times smaller corrosion rates, than those measured on bare steel specimens. Polarization curve analysis suggests that the corrosion of C1- and C2-coated specimens is mainly under cathodic control. This means that the substrate corrosion rate depends on the rate of oxygen diffusion towards the metallic surfaces through the coating pores, which in turn is affected by the coating thickness, the defect population and the interconnection degree and tortuosity of the defects themselves. On the contrary, oxygen diffusion rate inside the pores cannot be stimulated by higher electrode rotation speeds. This explains the invariance of the corrosion rates of coated specimens under free corrosion conditions, at 20 or 100 rpm. Fig. 10 shows that the porosity difference characterizing C1 and C2 coatings does not significantly affects oxygen diffusion. On both C1- and C2-coated specimens, sliding stimulates both the anodic and the cathodic process and induces a limited increase of the measured corrosion rates. On C1-coated specimens, a small iCOR increase is measured by passing from pure corrosion, to tribocorrosion at 2 N and 20 rpm and then to tribocorrosion at 10 N and 100 rpm. On the contrary, on C2-coated specimens under wear, the most severe corrosion attack is observed at 2 N and 20 rpm. As a consequence, C1 coating is the most protective one under mild wear conditions, while the coatings afford quite similar performances under more severe conditions. These findings can be explained by analysing how sliding wear modifies the coating surfaces in the track. It has been noticed that on both coatings abrasion of surface asperities occurs, with a consequent surface polishing. This smoothing effect is more evident at high load and rotation speed and particularly on C2 coating, because it is characterized by the lowest microhardness value. In this coating, the wear debris detached at 10 N and 100 rpm is found to partially fill up the surface porosity by plastic deformation and embedment in the coating. This obstruction justifies the lower iCOR value measured on C2-coated specimens under these experimental conditions, with respect to those evaluated at 2 N and 20 rpm. Microscale plastic deformation of Cr2O3 and Al2O3/13% TiO2 coatings under dry sliding wear has been often reported in the literature [3,11,33] and it has been attributed to plastic slipping at splat boundaries. The progressive polishing effect observed under sliding wear may justify the decrease of the wear rates, going from low to high rotation speeds, as the variation of this parameter implies an increase in the sliding distances which in turn involves smoother and smoother wear paths. A decrease in the friction coefficient going from 20 to 100 rpm could also occur and play a role in reducing the measured wear rates [31,32]. The observed formation of a tribofilm is the cause of the significant decrease of the wear rates at loads higher than 2 N. In fact when the load applied is sufficiently high the smearing of the wear debris on the coating surfaces limits the material removal. 5. Conclusions
2
2 - C2 20 rpm 2N
-0.4
3689
1E-2
Current Density (A cm-2) Fig. 12. Polarization curves recorded on C2-coated specimens, after 3 days of immersion. Steel under tribocorrosion conditions is also reported, as a reference.
- Cr2O3 and Al2O3/13% TiO2 coatings only partially protect steel substrate from corrosion. The corrosion process is mainly under cathodic control and is stimulated by the galvanic coupling with the interlaying noble Ni/20% Cr bond coat. - Tribocorrosion of Cr2O3- and Al2O3/13% TiO2-coated specimens stimulates both the anodic and the cathodic process so increasing the substrate corrosion rates. On the former coating, the highest corrosion rates are obtained under the most severe wear conditions (10 N and 100 rpm), while on the latter one the highest corrosion rates are measured under the mildest wear conditions tested (2 N and 20 rpm). - At the end of tribocorrosion tests, the wear mechanism on the ceramic coatings has been investigated. It involves polishing of the surface asperities and, at relatively high loads (higher than 2 N), formation of a tribofilm due to the smearing of the wear debris on the coating surface so inducing a partial fill up of the surface pores.
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Fig. 13. Top surface images of C1 (a,e) and C2 (b,c,d,f) coatings in the wear scars produced under tribocorrosion conditions at 10 N and 100 rpm (a,b,c,d,e) or at 2 N and 20 rpm (f). With the exception of Fig. 13d (VPSEM micrograph), all other images are optical micrographs. The sliding direction is always horizontal.
Table 4 Wear rates (10−6 mm3/N m) calculated at the end of 3 days exposures under tribocorrosion conditions. Cr2O3-coated steel
2N 5N 10 N
Al2O3/13% TiO2-coated steel
20 rpm
100 rpm
20 rpm
100 rpm
4.9 ± 0.1 1.8 ± 0.4 1.5 ± 0.1
– 0.4 ± 0.1 0.3 ± 0.1
7.5 ± 2.5 3.0 ± 0.6 3.0 ± 0.5
– 1.3 ± 0.3 0.7 ± 0.1
- A more continuous tribofilm is observed on Al2O3/13% TiO2 coating at 10 N and 100 rpm, which partially limits the corrosion of the steel substrate. As a consequence, under mild wear conditions, Cr2O3 coating is the most protective one, while under the most severe wear conditions, the coatings afford quite similar performances. - Wear rates of the coatings tend to decrease both at increasing rotation speed (from 20 to 100 rpm) and at increasing the load (particularly from 2 to 5 N).
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Acknowledgements The authors wish to thank Ing. Marco Caliari and Zocca Officine Meccaniche (Funo, BO), for the thermally sprayed coating manufacturing. References [1] G.W. Stachowiak, A.W. Batchelor, Engineering tribology, third ed, ButterworthHeinemann, Boston, 2005. [2] K.A. Habib, J.J. Saura, C. Ferrer, M.S. Damra, E. Giménez, L. Cabedo, Surf. Coat. Technol. 201 (2006) 1436. [3] G. Bolelli, V. Cannillo, L. Lusvarghi, T. Manfredini, Wear 261 (2006) 1298. [4] L. Fedrizzi, S. Rossi, F. Bellei, F. Deflorian, Wear 253 (2002) 1173. [5] M. Laribi, A.B. Vannes, D. Treheux, Wear 262 (2007) 1330. [6] T.A. Stolarski, S. Tobe, Wear 249 (2001) 1096. [7] F. Rasteger, A.E. Craft, Surf. Coat. Technol. 61 (1993) 36. [8] H.-S. Ahn, I.-W. Lyo, D.-S. Lim, Surf. Coat. Technol. 133–134 (2000) 351. [9] E. Celik, C. Tekmen, I. Ozdemir, H. Cetinel, Y. Karakas, S.C. Okumus, Surf. Coat. Technol. 174–175 (2003) 1074. [10] P. Ernst, G. Barbezat, Surf. Coat. Technol. 202 (2008) 4428. [11] H.-S. Ahn, O.-K. Kwon, Wear 225–229 (1999) 814. [12] F. Vargas, H. Ageorges, P. Fournier, P. Fauchais, M.E. López, Surf. Coat. Technol. 205 (2010) 1132.
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