Surface & Coatings Technology 205 (2011) 3683–3691

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Surface & Coatings Technology

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Corrosion and tribocorrosion behaviour of thermally sprayed ceramic coatingson 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

⁎ Corresponding author. Tel.: +39 0532 455136; fax:E-mail addresses: cecilia.monticelli@unife.it (C. Mon

(A. Balbo), zhf@unife.it (F. Zucchi).1 Tel.: +39 0532 455134.2 Tel.: +39 0532 455135.

0257-8972/$ – see front matter © 2011 Elsevier B.V. Aldoi:10.1016/j.surfcoat.2011.01.023

a b s t r a c t

a r t i c l e i n f o

Article history:Received 29 November 2010Accepted in revised form 10 January 2011Available online 15 January 2011

Keywords:Plasma sprayingCeramic coatingsCorrosionTribocorrosionWear rateChloride

This research aims at investigating the corrosion and tribocorrosion behaviour of thermally sprayed ceramiccoatings deposited on steel specimens and exposed to a 3.5% NaCl solution. The coatings have been preparedby 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 steelcylinders, during their exposure to the aggressive solution.Polarization resistance values monitored during 3 days exposures and polarization curves recorded at the endof the immersion period show that both coatings only partially protect steel substrate from corrosion. Slidingconditions (under 2 N load and 20 rpm or 10 N and 100 rpm) induce a limited increase of the substratecorrosion 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% TiO2-coated specimens, a stronger corrosion attack is achieved at 2 N and 20 rpm. Profilometer analysis and weartrack observations by optical and scanning electron microscopes evidence that on both coatings abrasion ofthe 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 andlimits the corrosion attack to a certain extent.

+39 0532 455011.ticelli), andrea.balbo@unife.it

l rights reserved.

© 2011 Elsevier B.V. All rights reserved.

1. Introduction

Wear is themajor cause of material wastage and loss ofmechanicalperformances and friction is the principal cause of wear and energydissipation [1]. Thermally sprayed ceramic coatings on steel representan efficient and economic way to improve the wear resistance ofmechanical components, when common thermal treatments (such asquenching and tempering,…) or thermo-chemical processes (such ascarburizing, nitriding…) result inadequate. In particular, plasmasprayed Cr2O3 and Al2O3–TiO2 coatings can be an excellent choice,providing protection against abrasive wear and resistance to galvanicand high temperature corrosion [2]. They often result to be superior totraditional wear resistant hard chromium [3,4] andmolybdenum [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 mechanicaltreatments necessary in many applications [3]. Cr2O3 coatings havelow friction coefficients and can be conveniently deposited on piston

ring and cylinder liners in the automotive industry [7–9], where theyreduce fuel and oil consumption and increase the engine life [10].

If compared to Cr2O3 coatings, Al2O3 and Al2O3–TiO2 ones are lessexpensive andmore biocompatible and can be applied in the food andmedicine packaging industry because they ensure the absence ofheavy metal contamination [3].

Many papers concern the wear behaviour of plasma sprayed oxidecoatings and investigate the corresponding wear mechanisms. Underdry sliding conditions, both Cr2O3 and Al2O3–TiO2 coatings arereported to form a tribofilm. On the former coating, this film isformed by plastically deformed and compacted wear debris, respon-sible of the lowmeasured friction coefficients [3]. XPS analysis revealsthat the film formed at room temperature is constituted by CrO3 andCr2O3 [11]. On the contrary, on Al2O3–TiO2 coatings the tribofilm has arather loose structure which does not adequately protect theunderlying material from wear [3]. On these ceramic coatingsdifferent 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 thermallysprayed oxide coatings on steel [16–19]. They show the cleardependence of the corrosion resistance of the coated materials fromthe coating porosity [20,21], which is reduced by a proper choice ofthe spraying parameters [22,23]. In some cases, interconnectedporosity decreases at increasing coating thickness [17]. However,

3684 C. Monticelli et al. / Surface & Coatings Technology 205 (2011) 3683–3691

the coating porosity is often found to increase with the depositiontime (and coating thickness) [22] and higher corrosion resistance canbe achieved with thin coatings [19,21]. Porosity also depends on thecomposition of the sprayed powders. As an example, ZrO2 addition tothe Al2O3 powder tends to increase the coating porosity [17], whilecontrasting results are obtained by TiO2 addition [3,24].

To the authors' knowledge no literature information is availableconcerning tribocorrosion of ceramic coatings. As the concomitantpresence of wear and corrosion processes usually induces asynergistic stimulation of material degradation [25], it has beenreputed interesting a characterization of both the corrosion andtribocorrosion behaviour of Cr2O3- and Al2O3–13%TiO2-coated steelspecimens in the presence of widespread aggressive species, such aschlorides.

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 ceramiccoatings were prepared, that is a Cr2O3 coating (indicated as C1) and aAl2O3/13% TiO2 coating (indicated as C2). As part of a standardcommercial deposition process, both coatings (obtained by air plasmaspray (APS) technology) were deposited onto a 20 μm-thick Ni/20% Crbond coat (also obtained by APS), which aimed at improving theceramic material adhesion onto the steel substrate. Details of thefeedstock powders are reported in Table 1, while the depositionparameters are confidential.

The phase transformations occurring during the thermal spray ofthe ceramic powders were investigated by comparing the X-raydiffractograms of the powders to those of the coatings. Moreover,properly polished cross sections of the powders and coatings wereobserved by a Scanning Electron Microscope equipped by VariablePressure technology (VPSEM) and Energy Dispersion Spectroscopy(EDS) microprobe, to investigate their microstructures. On thepolished cross-sections of the coatings, the porosity (through animage analysis software), and the Vickers microhardness (at 0.3 kgload) were also measured.

The coated steel specimens had a central threaded hole permittingthem to be mounted on a steel shaft which afforded the electricalcontact for electrochemical tests and let the specimen rotate at 20 or100 rpm.

Before testing, the coated specimens underwent a commercialgrinding process, inducing a surface roughness, Ra, of about 0.1 and0.2 μm on C1 and C2 coatings, respectively. A grinding process isusually applied to industrial components to reduce surface roughnessand increase the wear resistance.

As a reference, bare steel and steel specimens coated by a thickNi/20%Cr bond coat (250 μm, C3 coating)were also tested. Theywere ground byemerypapers down tono. 600 and carefully degreasedbefore exposure tothe 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 1Feedstock powders used to produce top and bond coatings.

Cr2O3 Al2O3/13% TiO2 Ni/20% Cr

Type Amperit 707.001by Starck

FSTC-335.23 by FlameSprayTechnologies

Metco 43CNSby Sulzer

Composition 0.06% SiO2, 0.03%Fe2O3, b0.02%TiO2, and balanceCr2O3

13.12% TiO2, 0.22% ZrO2,0.10% SiO2, 0.09% MgO,0.07% CaO, 0.39% otheroxides, and balance Al2O3

19.07% Cr, 1.1%Si, 0.4% Fe,0.02% C, andbalance Ni

Particle dimensions −45+22.5 μm −45+15 μm −106+45 μm

For tribocorrosion tests (Fig. 1), a non-commercial tribometerinduced a sliding-type wear on the lateral surface of the rotatingcylinders during exposure to the aggressive solution (rotation speeds:20 rpm or 100 rpm, involving sliding velocities of 21.5 mm s−1 or107 mm s−1, respectively). The normal loads (L), applied throughthe flat surface of an alumina cylinder with a 2 mm diameter, were2–10 N. The nominal hardness of the α-alumina counterbody was1950 HV.

During corrosion tests, the surface exposed to the aggressivesolution was the whole lateral surface of the cylindrical specimens(6.4 cm2), while during the tribocorrosion tests the lateral surfacewaspartially screened by Lacomit varnish, in order to expose only therubbed surface (1.6 cm2) to the solution.

The evolution of the corrosion process was monitored during 3 daysof exposure to the aggressive solution, by measuring the linearpolarization resistance (RP) values (in the potential range from −7 mVup to+7mVvs ECOR,with a potential scan rate of 0.1 mV s−1). At the endof the 3 days immersion, under both corrosion and tribocorrosionconditions, potentiodynamic polarization curves were recorded at ascan rate of 1 mV s−1, always starting from the ECOR value.

All the potentials in the text are referred to the Saturated CalomelElectrode (SCE).

Some specimens exposed to tribocorrosion conditions were notsubjected to destructive electrochemical tests. On these specimens,wear rates were calculated as WR=VW/L D, where VW is the wearvolume, L is the normal load and D corresponds to the sliding distance.VW was obtained by multiplying the area of the wear track crosssection (measured as an average from the wear track profiles at fivepositions along the specimen circumference by a Hommelwerk T2000profilometer with a TK300 piezoelectrical tip) by the circumference ofthe cylindrical specimens. Each wear rate value was the average ofthree measurements.

The morphology of the corrosion and tribocorrosion attack wascharacterized 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 C1coating (Fig. 2a). SEM observation of polished cross sections of theparticles indicates that they are monophasic (Fig. 2b) and XRDanalysis reveals that they are constituted by eskolaite (hexagonalCr2O3, Fig. 3a). C2 coating has been produced by a fused and crushedAl2O3–13% 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 andAl2TiO5 (Fig. 3b, upper diffractogram), as predicted by the Al2O3–TiO2

phase diagram [26], which indicates the stability of two structuralcomponents, that is α-Al2O3 and the eutectic α-Al2O3–Al2TiO5.

3.2. Coating characterization

SEM micrographs shown in Figs. 5a and 6 evidence the lamellarstructure of C1 and C2 coatings, typical of thermally sprayed coatings.C1 shows a fine homogeneous structure, still constituted by hexagonalCr2O3 (Fig. 3a, lower diffractogram) and characterized by small inter-and intralamellar cracks and finely distributed pores (Fig. 5a). Anirregular NiCr bond coat facilitates the adhesion of the ceramic coatingto the substrate (Fig. 5b), in spite of the presence of rare pores andsandblast residues, visible at the steel/bond coat interface. Withrespect to C1, C2 coating exhibits a coarser structure, with thickerlamellae and longer inter- and intralamellar cracks. Moreover, itshows isolated light grey titanium-rich splats (Fig. 6), suggesting the

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.

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inhomogeneous interdissolution of α-Al2O3 and Al2TiO5 phases,during the spraying process. In agreement with the literatureinformation [3], after the spraying process only metastable γ-Al2O3

and traces of α-Al2O3 are detected as crystalline phases, whilealuminium titanate remains amorphous (Fig. 4 b, lower curve). The

Fig. 2. VPSE SEM micrographs showing a) as-received Cr2O3 powder and b) Cr2O3

powder cross section.

bond coat morphology under C2 coating is equal to that exhibited inFig. 5b.

Coating porosity and Vickers microhardness values are reported inTable 2. The results obtained indicate that C1 exhibits a slightly loweraverage porosity with a lower standard deviation than those of C2coating. Moreover, the average microhardness of C1 is much higherthan that of C2, because of its higher bulk hardness value and lowerporosity. The hardness of Al2O3–13% TiO2 coating is lower than that ofthe alumina counterbody, because in the coatingmany factors, such asporosity, α- to γ-alumina conversion during thermal spraying andTiO2 addition, contribute to reduce the material hardness [3].

Table 2 also shows the porosity percentage of the bond coatingevaluated under SEM observations. It is lower than that of the ceramiccoatings, in agreement with literature information [27–29].

3.3. Linear polarization resistance measurements

Fig. 7 shows the RP and ECOR values collected on steel duringexposures to the aggressive solution under a rotation speed of 20 rpm.Low RP values are recorded after 1 h of immersion (540 Ω cm2) whichremainmore or less constant throughout the 3 days of exposure. After1 h of immersion, ECOR is about −0.560 VSCE, that is typical of steelunder active corrosion in neutral chloride solutions [30]. Then, itundergoes a moderate reactivation, with ECOR values decreasing downto −0.670 VSCE. During the immersion, steel is slowly covered byreddish iron hydroxide corrosion products.

Under the same exposure conditions, a quite different behaviour isexhibited by Ni/20% Cr bond coat material (Fig. 7). In fact, specimenscarrying a thick bond coating (C3-coated specimens) show RP (higherthan 105 Ω cm2) and ECOR (−0.11/−0.06 VSCE) values typical of apassive metal.

Tribocorrosion (TC) at 2 N load and 20 rpm stimulates thecorrosion attack on steel (Fig. 7). In fact, on this material it inducesRP values (about 200 Ω cm2, at the end of the immersion period)lower than those measured under pure corrosion conditions, withslightly ennobled ECOR values (−0.61 VSCE, after 3 days of immersion).

Fig. 8 shows the time dependence of RP and ECOR values collectedon 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), butthey are definitely lower than those measured on C3-coated speci-mens. The corresponding ECOR values (about −0.64 VSCE) are close tothose exhibited by bare steel specimens. Chromia, the ceramic

Fig. 4. VPSE SEM micrographs showing a) as-received Al2O3–13% TiO2 powder andb) Al2O3–13% TiO2 powder cross section.

20 40 60 80 100 1202-Theta / degree

20 40 60 80 100 1202-Theta / degree

Cou

nts

/ Arb

itra

ry u

nits

Hex-Cr2O3 All peaks

Cou

nts

/ Arb

itra

ry u

nits

α - Al2O3

γ - Al2O3

Al2TiO5

a)

b)

Fig. 3. XRD diffractograms recorded on Cr2O3 powder (a, upper curve) and coating (a, lowercurve) and on Al2O3–13% TiO2 powder (b, upper curve) and coating (b, lower curve).

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material of C1 coating, is an electrically insulating material whichcannot be involved in any electrochemical process. So on coated steelspecimen, corrosion may affect either the bond coat or the metallicsubstrate or both. However, at the measured ECOR values, only steelmay undergo a corrosion attack, while the bond coat material resultsin being galvanically protected. This suggests that the ceramic coatingand the underlying thin bond coating cannot hinder the aggressivesolution penetration through the coating pores and the onset of steelcorrosion at the pore bottom. The limited substrate area in contactwith the solution justifies the relatively high RP values. Fig. 8 alsoevidences that the corrosion of C1-coated steel specimens is notaffected by a variation in the rotation speed from 20 to 100 rpm.

During 3 days of exposure of C1-coated specimens to tribocorro-sion conditions at 2 N and 20 rpm, a diminution of RP values by only afactor of 2 or less is detected and no further RP reduction is observedby increasing the load up to 10 N and the rotation speed up to100 rpm (Fig. 8). This slight stimulation of the corrosion process islikely connected to an increase in the population of the coatingdefects, which permits an easier access of the solution to the steelsubstrate.

The behaviour of C2-coated specimens is similar to that exhibitedin the presence of C1 coatings (Fig. 9). Under pure corrosionconditions, 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 ofthe tests. The ECOR values remain close to −0.64 VSCE, throughout theimmersion period. This again indicates corrosion of the steelsubstrate. Under tribocorrosion conditions, a RP decrease is measured

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 at10 N and 100 rpm).

Under tribocorrosion conditions the slow formation of reddishiron corrosion products on both C1- and C2-coated specimens and inthe solution is clearly visible.

3.4. Polarization curves

The polarization curves recorded on the studied materials at theend of the immersion period in the neutral chloride solution arecollected in Fig. 10. Table 3 reports the corresponding ECOR and iCORvalues, as estimated by the Tafel method by extrapolation from thecathodic polarization curve.

As anticipated by the results of the linear polarization resistancemeasurements, steel exhibits an active corrosion behaviour, indicatedby its active ECOR value and low anodic overvoltage. The anodic Tafelslope (ba=0.080 V decade−1, Table 3) is slightly higher than thatexpected for active dissolution of iron (0.060 V decade−1), because ofthe formation of insoluble surface corrosion products which moder-ately hinder the anodic process.

Under free corrosion conditions, C3-coated specimens (with athick Ni–20% Cr coating) are under passive corrosion conditions (highba and low iCOR values, Table 3) and can suffer pitting corrosion atpotentials nobler than +0.45 VSCE (Fig. 10).

As shown in the previous section, C1- and C2-coated specimenshave ECOR values almost equal to that of steel (Fig. 10). However, theyexhibit higher anodic overvoltages (ba values in Table 3) and loweranodic and cathodic currents than those recorded on bare steel

1000000 0Rp - Steel - 20rpm

Fig. 5. VPSE SEM micrographs showing cross section microstructure of C1 coating(a) and underlaying bond coat (b).

Table 2Porosity and Vickers microhardness values for the studied coatings.

Porosity/% HV0.3/kg mm−2

C1 7.5±1.2 1556±104C2 9.7±2.8 1099±96Bond coat 1.1±1.0

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specimens. These observations are in agreement with the hypothesisthat the anodic process is connected to steel dissolution through thedouble coating. As on C1- and C2-coated specimens, both steeldissolution and oxygen reduction reaction occur at the bottom of thecoating pores, the rates of these reactions are limited by the restrictedaccess of the aggressive solution to the underlying metals.

At the bottom of the ceramic coating pores, the noble interlayingbond coat forms efficient cathodic regions (as shown by the cathodicpolarization curve on C3-coated specimens in Fig. 10) of relativelyhigh area (depending on the ceramic coating porosity). Then the bond

Fig. 6. VPSE SEM micrograph showing cross section microstructure of C2 coating.

coat can induce galvanic corrosion on the small substrate regions incontact with the aggressive solution through the porosity of the bondcoat itself. The relatively high cathodic to anodicmetal area ratio likelystimulates the galvanic attack.

With respect to C1-coated specimens, C2-coated ones exhibitlower anodic slopes and higher anodic current densities, while thecathodic polarization curves and the iCOR values are quite close to eachother (Table 3). The differences in the anodic curve suggest that theslightly higher porosity of C2 coatings facilitates the migration of ironcorrosion products through the coatings driven by the anodicpolarization. The presence of quite similar cathodic characteristicsand iCOR values suggests that on these specimens the corrosionprocess is mainly under cathodic control.

Fig. 10 also shows the polarization curves recorded on steel undertribocorrosion at 2 N load and 20 rpm. The analysis of these curvesevidences that sliding wear speeds up the cathodic reaction on steeland causes the slight ECOR ennoblement (to −0.61 VSCE), alreadynoticed during the discussion of the polarization resistance measure-ments. This is due to the stimulation of oxygen diffusion towards themetallic surface, caused by the continuous removal of corrosionproducts from the steel surface, which induces a fivefold increaseof iCOR (Table 3).

Wear application on C1- and C2-coated specimens slightlystimulates both the anodic and the cathodic reactions of steelcorrosion (Figs. 11 and 12). In the case of C1 coating (Fig. 11), thestimulation is quite limited and iCOR reaches its maximumvalue, at thehighest load and sliding velocity applied (10 N, 100 rpm, Table 3).Under this condition, iCOR is twice the value measured under purecorrosion conditions. On the contrary, in the case of C2 coating(Fig. 12), tribocorrosion under 2 N and 20 rpm is more severe thanthat experienced at 10 N load and 100 rpm. In fact, under the formercondition iCOR is about fivefold higher than that obtained in purecorrosion, while under the latter it is about twice.

100

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OR / V

SCE

Rp - Steel - 20rpm - 2NRp - C3 - 20rpmEcor - Steel - 20rpmEcor - Steel - 20rpm - 2NEcor - C3 - 20rpm

RP/ o

hm c

m2

Fig. 7. RP and ECOR values collected on steel and bond-coated (C3) specimens (solid linesrefer to pure corrosion conditions, while dotted lines to tribocorrosion conditions).

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E) 4 - C2 20 rpm

1 - Steel 20 rpm2 - Steel 20 rpm 2N

5 - C3 20 rpm

3 - C1 20 rpm

1

1 2

2

3

3

4

4

5

5

Fig. 10. Polarization curves recorded on steel and coated specimens, after 3 days ofimmersion (after 1 day of immersion, in the case of the C3-coated specimens).

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EC

OR / V

SCER

P/ o

hm c

m2

Fig. 8. RP and ECOR values collected on C1-coated specimens, under both corrosion (solidlines) and tribocorrosion (dotted lines) conditions.

3688 C. Monticelli et al. / Surface & Coatings Technology 205 (2011) 3683–3691

3.5. Wear track observations

The surface micrographs of the ceramic coatings tested underdifferent tribocorrosion conditions are shown in Fig. 13. Fig. 13a and bexhibits the surface aspect of both coatings at the wear trackborderline after tests at 10 N and 100 rpm: inside the track (lowerpart of the pictures) the surface is much smoother than outside, as aconsequence of the removal of surface asperities, left by the grindingprocess. Fig. 13a shows that on the hardest C1 coating the surface isrougher than that on C2 coating (Fig. 13b) and the holes in the trackare due to both the coating porosity and the troughs left by thecommercial grinding process. On C2 coating (Fig. 13b), the polishingeffect is more pronounced and large surface regions of the track areapparently poreless. However, surface observations at higher magni-fication reveal that the surface pores are partially obstructed by atribofilm produced by the smearing of the wear debris on the surface

Rp - C2 - 20rpm Rp - C2 - 20rpm - 2NRp - C2 - 100rpm - 10N Ecor - C2 - 20rpmEcor - C2 - 20rpm - 2N Ecor - C2 - 100rpm - 10N

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hm c

m2

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0

EC

OR / V

SCE

0 20 40 60 80

Time / h

Fig. 9. RP and ECOR values collected on C2-coated specimens, under both corrosion (solidlines) and tribocorrosion (dotted lines) conditions.

(Fig. 13c, d, where the arrows indicate some pores filled by the weardebris). 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 arethere still present.

Then, on C1 coating under tribocorrosion conditions incomplete ornegligible pore obstruction occurs, while on C2 coating at sufficientlyhigh load and rotation speed a more continuous tribofilm forms. Itpartially obstructs the access of the aggressive solution towards thesteel substrate, so explaining the lower iCOR recorded on C2-coatedelectrodes under heavier wear conditions (10 N, 100 rpm) withrespect to those measured under milder ones (2 N, 20 rpm).

3.6. Profilometry

After 3 day exposures to tribocorrosion conditions, wear rateswere calculated and collected in Table 4. The table shows that, atconstant load, the wear rates always diminish on going from 20 to100 rpm. A fivefold higher rotation speed corresponds to a fivefoldlonger sliding distance, as the tests always lasted 3 days. At lowrotation speed (and short sliding distance travelled), the wear track isrelatively rough and it is reasonable to obtain high wear volume permeter by the removal of surface asperities. On the contrary, at highrotation speed (that is after long sliding distances) the wear trackbecomes smoother and smoother, producing lower wear rates.However, it must also be mentioned that, according to the Stribeckcurve [31,32], during sliding of liquid-lubricated surfaces undermixedlubrication (ML) regime the friction coefficient decreases at increasingsliding velocities. Then, it cannot be excluded that a decreasingfriction coefficient can contribute to the observed decrease in wearrates, going from 20 to 100 rpm.

Table 3Electrochemical parameters calculated from the polarization curves recorded after3 days of immersion, under pure corrosion and tribocorrosion conditions.

20 rpm 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

Steel −0.66 5·10−5 0.080 −0.61 2.3·10−4 – –

C1 −0.61 1.0·10−5 0.240 −0.61 1.5·10−5 −0.60 2.0·10−5

C2 −0.63 9·10−6 0.100 −0.71 4.6·10−5 −0.68 2.2·10−5

C3 −0.06* 2.5·10−6* 0.460* – – – –

*measured after 1 day of exposure.

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1 - C1 20 rpm

2 - C1 20 rpm 2N

4 - steel 20 rpm

3 - C1 100 rpm 10

1

1

2

2

3

3

4

4

Fig. 11. Polarization curves recorded on C1-coated specimens, after 3 days ofimmersion. Steel under tribocorrosion conditions is also reported, as a reference.

3689C. Monticelli et al. / Surface & Coatings Technology 205 (2011) 3683–3691

Table 4 also evidences that on both coatings at 20 rpm the wear ratesmarkedly decrease on passing from 2 to 5 N, then they remain more orless constant at 10 N. Even at 100 rpm, the wear rates are constant orslightly decrease on going from 5 to 10 N. This is likely due to theformation of a tribofilm at loads higher than or equal to 5 N. Asdocumented by surface observations (Fig. 13), this tribofilm is connectedto the plastic deformation of the wear debris which fills up and partiallyobstructs the surface pores so reducing the rate of material removal.

4. Discussion

C1 and C2 ceramic coatings are characterized by porosity values of7.5 and 9.7% respectively, frequently encountered on these coatingtypes [3,11,24], and by inter- and intralamellar cracks. Porosity ismainly due to splat stacking faults and gas entrapment, while inter-and intralamellar cracks are often caused by scarce intersplat cohesionand thermal contraction respectively [3].

The Cr2O3 coating shows a finer structure, a slightly lower porosityand a higher microstructural homogeneity than the Al2O3/13% TiO2

coating. However, when exposed to a neutral chloride solution, theformer coating exhibits corrosion protection properties quite similarto those of the latter one. In fact, both coatings only partially protectsteel substrate from corrosion, in spite of the presence of a rathercompact interlaying bond layer, showing passive corrosion behaviour.This suggests that the aggressive solution penetrates through the topand bond coatings and reaches steel, where a corrosion process sets

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1 - C2 20 rpm

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4 - steel 20 rpm 2N

3 - C2 100 rpm 10 N

1

1

1

2

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23

34

4

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VSC

E)

Fig. 12. Polarization curves recorded on C2-coated specimens, after 3 days ofimmersion. Steel under tribocorrosion conditions is also reported, as a reference.

up, stimulated by the galvanic coupling with the noble bond coatingmaterial. As a result, coated steel specimens undergo five timessmaller corrosion rates, than those measured on bare steel specimens.

Polarization curve analysis suggests that the corrosion of C1- andC2-coated specimens is mainly under cathodic control. This means thatthe substrate corrosion rate depends on the rate of oxygen diffusiontowards themetallic surfaces through the coatingpores,which in turn isaffected by the coating thickness, the defect population and theinterconnection degree and tortuosity of the defects themselves. Onthecontrary, oxygendiffusion rate inside thepores cannotbe stimulatedby higher electrode rotation speeds. This explains the invariance of thecorrosion rates of coated specimens under free corrosion conditions, at20 or 100 rpm. Fig. 10 shows that the porosity difference characterizingC1 and C2 coatings does not significantly affects oxygen diffusion.

On both C1- and C2-coated specimens, sliding stimulates both theanodic and the cathodic process and induces a limited increase of themeasured corrosion rates. OnC1-coated specimens, a small iCOR increaseismeasuredbypassing frompure corrosion, to tribocorrosion at 2 N and20 rpmand then to tribocorrosion at10 Nand100 rpm.On the contrary,on C2-coated specimens underwear, themost severe corrosion attack isobserved at 2 N and 20 rpm. As a consequence, C1 coating is the mostprotective one under mild wear conditions, while the coatings affordquite similar performances under more severe conditions.

These findings can be explained by analysing how sliding wearmodifies the coating surfaces in the track. It has been noticed that onboth coatings abrasion of surface asperities occurs, with a consequentsurface polishing. This smoothing effect is more evident at high loadand rotation speed and particularly on C2 coating, because it ischaracterized by the lowest microhardness value. In this coating, thewear debris detached at 10 N and 100 rpm is found to partially fill upthe surface porosity by plastic deformation and embedment in thecoating. This obstruction justifies the lower iCOR value measured onC2-coated specimens under these experimental conditions, withrespect to those evaluated at 2 N and 20 rpm. Microscale plasticdeformation of Cr2O3 and Al2O3/13% TiO2 coatings under dry slidingwear has been often reported in the literature [3,11,33] and it hasbeen attributed to plastic slipping at splat boundaries.

The progressive polishing effect observed under sliding wear mayjustify the decrease of the wear rates, going from low to high rotationspeeds, as the variation of this parameter implies an increase in the slidingdistances which in turn involves smoother and smoother wear paths. Adecrease in the friction coefficient going from 20 to 100 rpm could alsooccur and play a role in reducing the measured wear rates [31,32].

The observed formation of a tribofilm is the cause of the significantdecrease of the wear rates at loads higher than 2 N. In fact when theload applied is sufficiently high the smearing of the wear debris on thecoating surfaces limits the material removal.

5. Conclusions

- Cr2O3 and Al2O3/13% TiO2 coatings only partially protect steelsubstrate from corrosion. The corrosion process is mainly undercathodic control and is stimulated by the galvanic coupling withthe interlaying noble Ni/20% Cr bond coat.

- Tribocorrosion of Cr2O3- and Al2O3/13% TiO2-coated specimensstimulates both the anodic and the cathodic process so increasingthe substrate corrosion rates. On the former coating, the highestcorrosion rates are obtained under the most severe wearconditions (10 N and 100 rpm), while on the latter one the highestcorrosion rates are measured under the mildest wear conditionstested (2 N and 20 rpm).

- At the end of tribocorrosion tests, the wear mechanism on theceramic coatings has been investigated. It involves polishing of thesurface asperities and, at relatively high loads (higher than 2 N),formation of a tribofilm due to the smearing of the wear debris onthe coating surface so inducing a partial fill up of the surface pores.

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 4Wear rates (10−6 mm3/N m) calculated at the end of 3 days exposures undertribocorrosion conditions.

Cr2O3-coated steel Al2O3/13% TiO2-coated steel

20 rpm 100 rpm 20 rpm 100 rpm

2 N 4.9±0.1 – 7.5±2.5 –

5 N 1.8±0.4 0.4±0.1 3.0±0.6 1.3±0.310 N 1.5±0.1 0.3±0.1 3.0±0.5 0.7±0.1

3690 C. Monticelli et al. / Surface & Coatings Technology 205 (2011) 3683–3691

- A more continuous tribofilm is observed on Al2O3/13% TiO2 coatingat 10 N and 100 rpm,which partially limits the corrosion of the steelsubstrate. As a consequence, under mild wear conditions, Cr2O3

coating is themost protective one,while under themost severewearconditions, the coatings afford quite similar performances.

- Wear rates of the coatings tend to decrease both at increasingrotation speed (from 20 to 100 rpm) and at increasing the load(particularly from 2 to 5 N).

3691C. Monticelli et al. / Surface & Coatings Technology 205 (2011) 3683–3691

Acknowledgements

The authors wish to thank Ing. Marco Caliari and Zocca OfficineMeccaniche (Funo, BO), for the thermally sprayed coatingmanufacturing.

References

[1] G.W. Stachowiak, A.W. Batchelor, Engineering tribology, third ed, Butterworth-Heinemann, 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.

[13] H. Liu, J. Tao, J. Xu, Z. Chen, Q. Gao, Surf. Coat. Technol. 204 (2009) 28.[14] P.P. Psyllaki, M. Jeandin, D.I. Pantelis, Mater. Lett. 47 (2001) 77.[15] B. Normand, V. Fervel, C. Coddet, V. Nikitine, Surf. Coat. Technol. 123 (2000) 278.[16] G. Bolelli, V. Cannillo, R. Giovanardi, L. Lusvarghi, T. Manfredini, E. Soragni, Proc.

Giornate Nazionali sulla Corrosione e Protezione, Senigallia (AN), Italy, 29 June – 1July, ISBN: 88-85298-53-2, 2005, p. 1.

[17] J. Zhang, A. Kobayashi, Vacuum 83 (2009) 92.[18] H. Liu, J. Tao, J. Xu, Z. Chen, Q. Gao, Surf. Coat. Technol. 204 (2009) 28.[19] Y. Zhijian, T. Shunyan, Z. Xiaming, Mater. Charact. 62 (2010) 90.[20] E. Celik, A.S. Demirkıran, E. Avcı, Surf. Coat. Technol. 116–119 (1999) 1061.[21] E. Celik, I.A. Sengil, E. Avcı, Surf. Coat. Technol. 97 (1997) 355.[22] O. Sarikaya, Surf. Coat. Technol. 190 (2005) 388.[23] A. Nusair Khan, J. Lu, J. Mat. Proc. Technol. 209 (2009) 2508.[24] R. Yılmaz, A.O. Kurt, A. Demir, Z. Tatlı, J. Eur. Ceram. Soc. 27 (2007) 1319.[25] D. Landolt, S. Mischler, M. Stemp, Electrochim. Acta 46 (2001) 3913.[26] S.M. Lang, C.L. Fillmore, L.H. Maxwell, J. Res. Nat. Bur. Stand. 48, 4 (1952) 2988

Research Paper 2316.[27] C.L. Li, H.X. Zhao, T. Takahashi, M. Matsumura, Mater. Sci. Eng. A308 (2001) 268.[28] N.F. Ak, C. Tekmen, I. Ozdemir, H.S. Soykan, E. Celik, Surf. Coat. Technol. 173–174

(2003) 1070.[29] C. Senderowski, Z. Bojar, J. Therm. Spray Technol. 18 (3) (2009) 435.[30] C. Monticelli, A. Frignani, F. Zucchi, Corros. Sci. 46 (2004) 1225.[31] M. Woydt, R. Wäsche, Wear 268 (2010) 1542.[32] L. Ma, W.M. Rainforth, Tribol. Int. 43 (2010) 1872.[33] V. Fervel, B. Normand, C. Coddet, Wear 230 (1999) 70.