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Wear 267 (2009) 1452–1461

Contents lists available at ScienceDirect

Wear

journa l homepage: www.e lsev ier .com/ locate /wear

ardness properties and high-temperature wear behavior ofitrided AISI D2 tool steel, prior and after PAPVD coating

.H. Staiaa, Y. Pérez-Delgadoa, C. Sancheza, A. Castroa,. Le Bourhisb, E.S. Puchi-Cabreraa,c,∗

School of Metallurgical Engineering and Materials Science, Faculty of Engineering, Universidad Central de Venezuela,aracas, VenezuelaUniversité de Poitiers, Laboratoire de Physique des Materiaux, UMR 6630 CNRS, SP2MI-Téléport 2-Bd M&P Curie, BP 30179,6962 Futuroscope-Chasseneuil Cedex, FranceVenezuelan National Academy for Engineering and Habitat, Postal Address 1723, Caracas 1010, Venezuela

r t i c l e i n f o

rticle history:eceived 26 February 2008eceived in revised form 31 January 2009ccepted 23 March 2009vailable online 6 April 2009

eywords:igh-temperature wearISI D2 tool steelitridingiN and TiAlN PAPVD coatings

a b s t r a c t

Wear experiments in the range of 25–600 ◦C have been conducted on samples of D2 tool steel in differentconditions involving unnitrided, nitrided and nitrided and coated with Balinit® A (TiN) and Balinit®

Futura (TiAlN) deposited industrially at Balzers (Amherst, NY, USA), by means of PAPVD. The resultsindicate that coating the nitrided D2 tool steel substrate with these two films gives rise to an improvementof ∼97% (TiN) and 99% (TiAlN) in the wear behavior at the test temperature of 300 ◦C, in comparisonwith the uncoated substrate. However, at a temperature of 600 ◦C, besides oxidation of the coatings, themechanical strength of the substrate decreases giving rise to fracture and delamination of the films. At thistemperature the uncoated substrate exhibited the highest resistance to sliding wear, presumably due tothe formation of a well bonded surface glazed layer which gives rise to a significant reduction in the frictioncoefficient. The indentation experiments that were conducted with the nitrided steel substrate and the

ndentationardness profile

coated systems indicates that the nitriding process applied to the D2 steel prior to PAPVD coating providesa satisfactory load support which contributes to the improvement of the coated systems capability towithstand indentation loads at room temperature. In this regard, the coated system with a TiAlN coatingdisplayed a better behavior than that shown by the system with a TiN coating. An experimental procedureis proposed in order to predict the hardness profile of the nitrided tool steel, along the cross section ofthe material, just from hardness measurements taken on the surface of the sample, employing different

indentation loads.

. Introduction

The elevated temperature behavior of materials has consti-uted a main concern in the last decades as a consequence of theevelopment of industries such as metal extraction, alloy manu-

acture, chemical processing, power engineering, etc., where theifferent parts and components have to be able to perform in aarsh environment, maintaining a high mechanical strength and

esistance against fatigue, oxidation, corrosion and wear. There-ore, understanding the materials properties at high temperaturesould provide the basis for the production of new materials,s well as improvements of old ones, thereby contributing to a

∗ Corresponding author at: School of Metallurgical Engineering and Materials Sci-nce, Faculty of Engineering, Universidad Central de Venezuela, Caracas, Venezuela.el.: +58 212 7539017; fax: +58 212 7539017.

E-mail address: epuchi@reacciun.ve (E.S. Puchi-Cabrera).

043-1648/$ – see front matter © 2009 Elsevier B.V. All rights reserved.oi:10.1016/j.wear.2009.03.045

© 2009 Elsevier B.V. All rights reserved.

greater efficiency, safety, lower costs, as well as to a longer life-time.

Wear at elevated temperatures has been considered in detailfor different metals and alloys. Pauschitz et al. [1] have reportedrecently a systematic survey of the current status and future trendsof protection against high-temperature wear, with special emphasison the mechanisms that explain the formation of glazed layers. Theinfluence of oxide-forming alloying elements, such as Si and Al, onthe tribological behavior of Fe alloys has been pointed out in theliterature [1–3]. It has been indicated that such elements tend tosegregate at the surface of the material forming a thin oxide layer,which promotes further oxidation of the steel. Thus, a thin ironoxide layer is formed on top of the existing aluminum and silicon

oxide layer, whose mechanical stability is insufficient to protect thesteel against wear.

Surface engineering of metallic and ceramic materials, withthe aim of improving surface mechanical properties such as wearresistance at high temperatures, has constituted another way of

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Wear tests were carried out at temperatures of 25, 300 and600 ◦C employing a ball on disk configuration high-temperature tri-bometer (CSEM Switzerland,). The tests were conducted accordingto the ASTM G99-95a standard [13], under the conditions presentedin Table 1. Topographic profiles of the tested wear samples were

Table 1Conditions employed during the wear tests.

Normal load (N) 5

Radius (mm) 8.5Speed (m/s) 0.1

M.H. Staia et al. / We

ncreasing the components lifetime. An extensive survey has beeneported by PalDey and Deevi [4], where the wear resistant proper-ies of (Ti, Al)N coatings for various machining applications haveeen compared to other coatings such as TiN, Ti(C, N) and (Ti,r)N, indicating also the benefit of using the multicomponentystems, based on different metallic and nonmetallic elements.hese authors have mentioned that only when a hard wear resis-ant coating is complemented with a hard and tough substrate,he wear resistance of the coated system is increased, since aard coating deposited on a soft substrate leads to poor proper-ies.

One common approach to overcome this problem has beenhe use of duplex treatments, which combine the hardening ofhe substrate by nitriding prior to the PVD deposition processf the hard coatings. Such treatments not only reduce the hard-ess gradient between the coated surface and substrate but also

mprove the coating adhesion and increase the durability of theools [5–8].

Zukerman et al. [8] have indicated that there is still some uncer-ainty regarding the choice of the underlying nitrided layer thatmparts the best tribological performance of the uppermost hardoating and some contradictory results have been reported in theiterature [9,10]. In general, it has been found that the tribologicalerformance is related mainly to the control of the plasma nitrid-

ng processing conditions in relation to the quality of the nitridedayer, i.e. porosity, roughness, thickness, nature and amount of theron nitrides formed. When a uniform and dense compound layers formed, favorable sliding wear properties of the composite arebtained [7,10]. Recently, the authors [11] have also presented thexperimental wear results of a gas nitrided H13 steel/(Ti0.7Al0.3)NAPVD duplex coating at 600 ◦C, typical temperature that could bechieved during aluminum extrusion processes. It was found that aatisfactory wear resistance was obtained for this duplex system asconsequence to the presence of a very dense and thin compound

ayer.However, many industrial nitriding processes based on thermo-

hemical treatments do not allow either the strict control of theirrocessing parameters, or the control of the chemical compositionf the compound layer, as in the case of ion nitriding. Therefore, dueo the intrinsic porosity of the compound layer, additional polishingf the nitrided surface prior to coating deposition is always neces-ary. In this sense, Hernández et al. [12] have recently conducted annvestigation in order to develop a rational approach able to eval-ate quite accurately the depth of a nitrided tool steel that shoulde removed previous to hard PAPVD coating deposition, and theependence of such depth on the nitriding conditions.

The method proposed is based on the use of the hardnessata corresponding to the hardness depth profile of the nitridedaterial, in conjunction with Ficks’s second law. The analytical pro-

edure was described employing the experimental data obtainedrom a gas nitrided AISI H11 steel (X38CrMoV5.1), and was vali-ated by computing a rough estimate of the diffusivity parametersf nitrogen in the �-Fe matrix. According to these authors, nitridingt a temperature of 580 ◦C for 8 h seems to have important advan-ages over the process conducted at 510 ◦C and 48 h, particularlyn terms of the characteristics of the hardness profile obtained andhe ability of the nitrided material to withstand indentation loads,measure of its load-carrying capacity, as a feasible substrate forard PVD coating deposition.

Thus, given the relatively lack of information regarding theehavior and mechanical properties of duplex systems involving D2

ool steels, the present investigation has been conducted in ordero study the hardness properties at room temperature and the wearehavior, in a wide range of temperatures (25–600 ◦C), of this steel

n four different conditions: as-received, nitrided and nitrided andoated with TiN or TiAlN films deposited by PAPVD.

(2009) 1452–1461 1453

2. Experimental techniques

The present investigation has been conducted with samples ofan AISI D2 steel, with the following chemical composition (wt.%):1.55 C, 11.80 Cr, 0.95 V, 0.80 Mo, 0.35 Mn, 0.25 Si. The material wasprovided as a bar of 27 mm in diameter, from which samples of4 mm in height were machined and subsequently heat treated inorder to achieve the required hardness. The specimens were firstlyheat treated at a temperature in the range of 550◦–580 ◦C for 20 min,austenitized at 1050 ◦C during 4 min. and, subsequently, quenchedand tempered at 200 ◦C for 30 min. After heat treatment, the sam-ples were prepared metallographically by grinding with differentSiC papers, varying from grit no. 80 to 1000, polished with aluminaof 1 �m and finally cleaned ultrasonically with acetone and iso-propyl alcohol. The mean roughness of the specimens after sucha preparation was of 0.020 ± 0.004 �m, and this value was deter-mined by employing a Zygo New View 200 profilometer by meansof optical interference techniques.

Gas nitriding of the samples was carried out in an industrialfacility (Promotremp S. L, Barcelona, Spain) at a temperature of520 ◦C for 3–4 h. The nitrided layer formed was of ∼60 �m, whosehardness at about 10 �m from the surface was of ∼1200 VHN(∼11.8 GPa). The characterization of such a layer was carried outon the cross sections of the nitrided samples which were preparedmetallographically and attacked with both a mixture of picric andnitric acids, and also Marble’s reagent (cupric sulphate), in orderto get the best possible definition of such a layer under the opticalmicroscope. The nitrided layer depth was measured by means ofimage analysis.

After nitriding, the specimens were machined and polishedagain in order to eliminate the white layer, which had been esti-mated from optical microscopy observations, as well as from thehardness profile determined on the cross section of the nitridedspecimens, employing the methodology suggested by Hernándezet al. [12]. Polishing was conducted on a rotating wheel at 200 rpm.The final mean roughness of the samples was of 0.114 ± 0.009 �m.Part of the nitrided samples was coated in a Balzers Inc. industrialfacility. A group was coated with Balinit® A (TiN) and another groupwith Balinit® Futura (TiAlN). The coatings thicknesses were deter-mined by means of Calotest measurements (CSEM, Switzerland)and image analysis.

Hardness tests were carried out employing a microhardnesstester (Leco). The measurements were conducted on the cross sec-tion of the nitrided samples up to a depth of ∼100 �m at steps of7 �m. Five Vickers indentations were performed at each selectedlocation from the surface, employing a load of 10 g. In the case ofthe coated samples, the composite hardness was determined byconducting Vickers indentations on the plane of the coating, nor-mal to the coating–substrate interface, applying loads of 25, 50, 100,200, 300, 500 and 1000 g.

Test temperatures (◦C) 25; 300; 600Room relative humidity (%) 65 ± 6Test distance (m) 1000Static counterpart Alumina ball (6 mm diameter)Contact pressure (GPa) 1.34 (DTiN, DTiAlN)–1.50 (SA)

1454 M.H. Staia et al. / Wear 267

Ft

dpdStoef

hardness of the nitrided material at the surface, k1 and k2, material

Fco

ig. 1. Optical micrograph showing the typical quenched and tempered microstruc-ure of the AISI D2 tool steel.

etermined by means of optical profilometry employing the samerofilometer indicated above. For each wear scar, four profiles wereetermined, which allowed the computation of the wear volume.EM techniques were used in order to analyze the morphology of

he wear scars and determine the operative wear mechanisms. Thebservations were conducted under both primary and secondarylectron modes With a Phillips XL 30 microscope coupled with EDSacilities.

ig. 2. (a) Optical micrograph showing a nitrided layer of ∼60 �m in depth. (b) Change inonducted both on the surface and along the cross section of the specimen. (c and d) Crobserved after fracture.

(2009) 1452–1461

3. Experimental results

3.1. Microstructural and mechanical characterization of the steelsubstrate prior and after gas nitriding

Fig. 1 illustrates the typical quenched and tempered microstruc-ture of the AISI D2 tool steel employed in the present work, whichmainly consists of carbide particles within a fine tempered marten-site structure. The hardness of this material was found to be in therange of ∼4 GPa. Gas nitriding of the steel substrate gave rise to thedevelopment of a nitrided layer of ∼60 �m in depth, as shown inFig. 2a. The nitrided layer depth was also characterized by means ofhardness measurements conducted both on the surface, employinga range of loads between 25 and 1000 g as indicated above, as wellas on the cross section of the sample, employing a constant load of10 g.

Fig. 2b illustrates the change in hardness with indentation depthfor both orientations. As suggested by Hernández et al. [12], theexperimental hardness data presented in Fig. 2b has been describedby means of a simple parametric relationship of the form:

HNS = H0 + HS − H0

1 + (ı/k1)k2(1)

where HNS represents the hardness of the nitrided steel substrate,H0 = 4 GPa is the hardness of the steel prior to nitriding, HS the

constants that are computed from the experimental data and ı thedepth in �m measured from the surface of the sample.

Parametric relationships of the form of Eq. (1) are commonlyused for the description of data points which display an “inverted-

hardness with depth for the nitrided D2 tool steel. Hardness measurements weress section views of the TiN and TiAlN films deposited onto the nitrided substrate,

M.H. Staia et al. / Wear 267 (2009) 1452–1461 1455

Table 2Parameters involved in Eq. (1) for the hardness curves determined for the nitridedsteel substrate along different orientations, shown in Fig. 2b.

O

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Table 3Parameters involved in Eq. (2) for the duplex systems coated with TiN and TiAlNfilms.

rientation H0 (GPa) HS (GPa) k1 (�m) k2

urface 4 16.3 5.9 1.52ross section 4 16.3 26.2 1.04

” shape [14] and, as shown in Fig. 2b, it provides a satisfactory fitf the experimental hardness values. The hardness measurementsonducted on the surface of the specimen employing low appliedoads (25–100 g), allowed the determination of the parameter HS,

hich was found to be in the range of ∼16.3 GPa. Table 2 summa-izes the values of all the parameters involved in Eq. (1) for bothurves, which were determined by means of non-linear least squarenalysis, employing a conjugate gradient method.

By means of the methodology also developed by Hernández et al.12] and employing the hardness data obtained along the cross sec-ion of the nitrided specimen, it was determined that not less than14 �m had to be machined off from the nitrided layer surface prior

o the PVD deposition of the TiN and TiAlN films, in order to ensurehe proper adhesion of the later to the nitrided steel substrate.

Finally, in relation to this section, Fig. 2c and d illustrate a crossection view of the fracture surface of the nitrided substrate coatedith the TiN and TiAlN films, respectively. From these micrographs

t can be appreciated that both coatings are apparently quite dense,ypical of the deposition process employed in the present work.

.2. Composite hardness of the coating–substrate duplex systems

The Calotest measurements indicated that the deposited TiNoating had a mean thickness of ∼4.5 �m, whereas the TiAlN coat-ng of ∼4 �m. Such values are very important in order to conduct

description of the composite hardness of the coated system asfunction of the relative indentation depth (RID or ˇ), defined

s ˇ = ı/t, where t represents the coating thickness. The compos-te hardness involves both the hardness of the coating and that ofhe substrate. At low indentation loads the composite hardness is

ainly dominated by the coating hardness, whereas at high inden-ation loads it is dominated by the substrate hardness.

Fig. 3 illustrates the change in the composite hardness forhe two coating–substrate duplex systems investigated in theresent work, as well as for the uncoated nitrided steel substrate.he description of the experimental data corresponding to the

ig. 3. Change in the composite hardness for the two coating–substrate duplex sys-ems investigated in the present work, as well as for the uncoated nitrided steelubstrate.

Coating HF (GPa) ˇ0 n

TiN 24.0 0.34 0.53TiAlN 27.2 0.58 0.57

duplex systems can be conducted by means of the equation earlieradvanced by Puchi-Cabrera [15,16], according to which:

H = HSubst(ˇ) +[HF − HSubst(ˇ)

]exp

[−(

ˇ

ˇ0

)n](2)

In the above equation H represents the composite hardness,HSubst(ˇ) and HF the substrate and film hardness, respectivelyand ˇ the RID. ˇ0 and n represent constants characteristic of thecoating–substrate system. As can be seen in Fig. 3 and as expected,the curve that describes the change in hardness with RID for theuncoated nitrided substrate (identified as (a)) is not a constant, butalso presents an inverted “S” shape, as shown previously by curve(a) in Fig. 2b. Therefore, it is possible to describe such a changeby means of an expression similar to that given in Eq. (1), whichexpresses the hardness as a function of ˇ rather than ı:

HSubst(ˇ) = HS0 + HS1 − HS0

1 + (ˇ/kS1)kS2(3)

In this case, ˇ = ı/t, where t = 4.25 �m and represents a mean ofthe coating thickness of both films. From the experimental hard-ness data of this material, it was determined that HS0 = 4 GPa,HS1 = 16.3 GPa, kS1 = 1.39 and kS2 = 1.52. Thus, the evaluation of thechange in the composite hardness with the RID for the duplex sys-tems is accomplished by means of Eqs. (2) and (3), employing theabove values that were computed for the uncoated nitrided steel.Table 3 summarizes the values of the constants HF, ˇ0 and n for thesystems coated with TiN and TiAlN films.

3.3. Friction and wear behavior

On the other hand, Figs. 4–6 illustrate the change in the frictioncoefficient (�) with distance at the different test temperatures forthe steel substrate prior (SP) and after (SA) nitriding, as well as forthe duplex systems consisting of the nitrided substrate coated withTiN (DTiN) and TiAlN (DTiAlN) and the corresponding morphologicalobservations (SEM) of the wear track in each particular case.

Fig. 4a indicates that at 25 ◦C the change in the friction coefficientwith distance is similar for SP and SA, as well as for the DTiAlN sys-tem, remaining at a value of ∼0.83. In the case of the DTiAlN system,the oscillations observed in the friction curve indicate that debrisof both coating and counterpart are produced, which do not remainwithin the contact area. At a distance of 900 m, a decrease in thefriction coefficient can be seen, whose origin is not entirely under-stood at the present time. For the DTiN system, it is observed thatduring the first ∼200 m, � varies between 0.5 and 0.6, increasesup to ∼1.0 at 600 m and, finally, decreases to values between 0.4and 0.7 displaying significant fluctuations. Fig. 4b illustrates thatthe SP sample undergoes severe abrasive wear, in comparison withthe other materials, for which the depth and width of the weartrack are smaller. However, it has to be mentioned that in the weartrack a small amount of debris, mainly constituted of alumina parti-cles from the counterpart as determined by EDS analysis, were alsonoticed. The SA specimen (Fig. 4c) displays an abrasive wear mech-

anism, with accumulation of Fe oxides and alumina debris fromthe ball within the wear track, which agglomerate predominatelyat the center of it, impeding the ball-disk contact and hinderingsubsequent wear. Both the DTiN and DTiAlN systems (Fig. 4d and 4e,respectively) show a satisfactory behavior since the substrate has

1456 M.H. Staia et al. / Wear 267 (2009) 1452–1461

F pondin with

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ig. 4. (a) Change in the friction coefficient (�) with distance at 25 ◦C. (b–e) Corresitrided substrate, nitrided substrate coated with TiN and nitrided substrate coated

ot been exposed after the wear test, indicating that it has providedn adequate load-carrying support for the coatings. An abrasiveear mechanism was identified for the TiN coated system. The wearebris produced was lost from the contact and a polished surface forhe TiN coated system can be noticed in Fig. 4d. As has been pointedut by Wilson and Alpas [17], at this temperature oxidation of theiN coating could occur due to the low values of the applied loadnd sliding velocity employed during the wear tests conducted inhe present study. However, in case of the TiAlN tribosystem, thelumina counterpart was observed to wear out and some material

ccumulated within the wear track, as shown in Fig. 4e.

At 300 ◦C (Fig. 5a), the change in the friction coefficient withistance indicates that for the SP material such a parameter fluc-uates between ∼0.70 and 1.0 with periods of ∼160 m, whichould be attributed to the formation and rupture of an unstable

ng SEM morphological observations of the wear track for the unnitrided substrate,TiAlN, respectively.

Fe and Cr oxide layer. On the contrary, for the SA material � isobserved to remain approximately constant at a value of ∼0.67.During the test, the SP specimen had a violet coloration whichis typical of the presence of chromium oxide. The wear track ofthe SP sample (Fig. 5c) is observed to be wider and deeper thanthat of the SA material and both systems present an abrasivewear mechanism. However, in the micrographs corresponding tothe SA specimen (Fig. 5b), a more irregular surface morphologyis observed with an appreciable amount of craters and consider-able accumulations of Ti-rich and Cr and Fe oxide particles that

were identified by EDS analysis. The duplex systems exhibiteda different behavior. For the DTiN system, � showed initial peri-odic oscillations and varied between ∼0.70 and 0.90 caused by thedebris produced due to the localized coating fracture indicated inFig. 5d. The EDS analyses that were conducted indicated the pres-

M.H. Staia et al. / Wear 267 (2009) 1452–1461 1457

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ig. 5. (a) Change in the friction coefficient (�) with distance at 300 ◦C. (b–e) Corresitrided substrate, nitrided substrate coated with TiN and nitrided substrate coated

nce of Ti-rich third-body particles and large amounts of Fe and Crxides, presumable from the substrate. After approximately 330 m,steady state friction coefficient was observed, with an average

alue of ∼0.9. In the case of the DTiAlN system, � showed severeuctuations up to ∼500 m, which lead to changes between ∼0.60nd 1.20. Such a behavior could be attributed to the presence oflumina third-body particles, which are kept in the contact andccumulated along the wear track, as a consequence of the peri-dic formation and rupture of small particles from the pin due

o the higher hardness of the coating. The DTiAlN system showedhe best behavior at this temperature. Although Fig. 5e shows aide wear track, the coating maintained its integrity and the onlyaterial accumulated within the wear track was that of the static

ounterpart.

ing SEM morphological observations of the wear track for the unnitrided substrate,TiAlN, respectively.

At 600 ◦C, the change in friction coefficient with distance (Fig. 6a)revealed that for both the SP and SA materials, up to ∼700 m, �remains approximately constant at a value between ∼0.50 and 0.60.

Beyond 700 m, for the SA sample � displays a slight increase,attaining a value of ∼0.70. The curve corresponding to the DTiNsystem shows periodic fluctuations due to the detachment of thecoatings and accumulation of debris in the wear track producing anincrease in � up to values of ∼0.90 at a distance somewhat less than200 m. Above such a distance the fluctuations are still observed, but

� changes between ∼0.40 and 0.60, due to the oxidation processwhich takes place on the exposed substrate. In the case of the DTiAlNsystem, during the first ∼200 m of run � attains a value between∼0.60 and 0.70, with the trend of decreasing progressively towardsthe value of the SP sample. It is observed that, whereas the TiN and

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458 M.H. Staia et al. / We

iAlN coatings maintained their integrity, the friction coefficientst high temperature in the first 200–300 m are higher than thoseor the uncoated systems, behavior which is characteristic for highardness mating surfaces. After 500 m sliding distance, the frictionurves indicate that in the DTiAlN system the coating has been totallyemoved and therefore, the friction coefficient achieves a value sim-lar to those of the SP and SA samples. Regarding the DTiN system,t is believed that partial delamination of the coating takes place,indering the formation of a uniform oxidized layer, leading to aon-uniform contact surface.

The significant reduction in the friction coefficient valuebserved for the SP sample (Fig. 6a) leads to the presumption thathe wear track of such a specimen presented in Fig. 6b, could exhibitn oxidized layer, as a consequence of the test temperature and

ig. 6. (a) Change in the friction coefficient (�) with distance at 600 ◦C. (b–e) Corresponditrided substrate, nitrided substrate coated with TiN and nitrided substrate coated with

(2009) 1452–1461

environment. As has been indicated by Pauschitz et al. [1], in thepresent case the formation of a glazed layer would also be favoredby the high hardness and the relatively elevated Cr content (∼12%)present in the D2 steel. Also, such a layer would be promoted by thehardness of the counterpart material and its capability for forminglarger grooves, able to retain the oxide debris formed in the earlystage of sliding. The results of the present research indicate that thesteel microstructure could play an important role in the formationof a protective glazed layer. Fontalvo and Mitterrer [2] have reportedsimilar results for a hot working steel alloy of the type X38CrMoV5-

1 and an alumina counterpart, tested at 500 ◦C. Stott and Jordan [18]observed that wear protective layers were able to develop for highchromium steel/carbon steel combination under 10–20 N load at550–600 ◦C, a behavior that was attributed to the high hardness of

ing SEM morphological observations of the wear track for the unnitrided substrate,TiAlN, respectively.

M.H. Staia et al. / Wear 267 (2009) 1452–1461 1459

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ume of SP and SA decreases in relation to that determined atroom temperature, whereas the wear volume of both DTiN andDTiAlN systems experiments a slight increase. In this case, the spec-imen SA displays the best behavior, with a wear volume of just0.01 mm3.

Fig. 7. Optical micrographs of the static alumina counterpa

he chromium steel. If a stable oxide layer were not present in theear track, the friction coefficient curve should exhibit significantuctuations associated to the formation and loss of the oxides thatre expected to form under these experimental conditions.

Fig. 7 illustrates an optical micrograph of the static counter-art corresponding to the different testing conditions. As can bebserved in Fig. 7a for the SP sample, the counterpart has beenorn off by a continuous Cr and Fe oxide layer produced on D2

teel substrate surface.In the case of the SA sample (Fig. 6c), a glazed oxide film was

ormed with the presence of agglomerated debris which is consti-uted by a mixture of alumina particles from the pin and Cr and Fexide particles from the surface layers of the nitrided steel insidehe wear track. Also, the presence of debris can be observed on bothides of the wear track. Fig. 7b shows that this new surface producedslightly increased worn volume of the counterpart as compared

o the SP tribosystem.The analysis of the wear track of the DTiN system showed the par-

ial delamination of the coating and a pronounced abrasive wearechanism. Considerable oxidation took place over the exposed

urface inside the track, which is unprotected by the coating,hereas the surface outside the track also showed significant oxi-ation at this temperature. In the case of the DTiAlN system, the wearrack is observed to be deeper than that of the TiN film, reaching intohe underlying substrate, which indicates the complete delamina-ion of the coating and the subsequent oxidation of the substrate,ith the formation of a glazed layer. The delamination of the TiAlN

oating produced a higher amount of debris during the test as com-ared to the TiN tribosystem, which explains the increased amountf the wear volume for the former. It is important to point out that

he EDS analyses carried out on the TiN coating outside the wearrack indicated the existence of ∼ 32% O2, which was approximatelytimes higher than that corresponding to TiAlN coated system.

The wear volumes were determined from the mean wear trackadius and the surface topographic profiles through such a track

r wear tests at 600 ◦C. (a) SP, (b) SA, (c) DTiN and (d) DTiAlN.

obtained by means of optical profilometry, which allowed the com-putation of the wear cross section. Fig. 8 clearly shows that at25 ◦C, the largest wear volume corresponds to the sample SP, whichis in the range of ∼0.072 mm3, whereas for the other systemssuch a volume ranges between ∼0.014 and 0.021 mm3. However,as the test temperature increases to 300 ◦C, the wear volume ofthe SP and SA samples increases significantly, as consequence tothe formation and rupture of the oxide layers, whereas for theDTiN specimen it increases slightly and for the DTiAlN, it decreasesand displays the smallest volume of all. At 600 ◦C, the wear vol-

Fig. 8. Change in the wear volume as a function of the wear test temperature for thesteel substrate prior and after nitriding, as well as for the duplex systems consistingof the nitrided substrate coated with TiN and TiAlN.

1 ar 267 (2009) 1452–1461

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460 M.H. Staia et al. / We

. Discussion

Fig. 2 illustrates that Eq. (1) is quite satisfactory for the descrip-ion of the hardness data of the nitrided steel substrate, both onross section, where such a property is determined at constantoad (10 g), as well as on the surface, where hardness is esti-

ated employing a range of load values. As expected, it can belearly observed that both curves are significantly different. Hard-ess measurements conducted on the surface of the nitrided steelt low indentation loads (25–100 g) allow an accurate prediction ofhe surface hardness of the material to which both curves shouldxtrapolate at low indentation depths. Also, at large depths from theurface both curves should extrapolate to the core hardness of theaterial (4 GPa). The difference between both hardness curves can

e accounted for in terms of the nature of the hardness test itself, onhe following basis. It is well known that during the indentation testplastic zone develops under the indenter, whose volume increasess the indentation load increases. If it assumed that the volume oflastic zone under the indenter has a hemispherical shape and that

t could be divided into a number of discrete elements of volumei, whose corresponding hardness is Hi, the overall hardness of theaterial at a constant load could be expressed by means of a simple

inear law of mixtures as:

¯ =∑

iViHi∑iVi

(4)

f the microstructure of the material under the indent does nothange with depth, the hardness of all the elementary volumesncompassed within the plastic zone can be considered to be

constant and therefore, Hi = H = H. On the contrary, if theicrostructure of the material under the indent changes with

epth, as in the case of the surface hardness measurements con-ucted in a nitrided material, the hardness of the elementaryolumes will also change continuously with depth from the harderurface to the softer core and therefore, Hi = f(ıi).

Thus, for the hardness measurements conducted under a con-tant indentation load along the cross section of the nitridedample, it would be expected that H decreased with the distancerom the surface due to the approximation of the steel softer core.owever, for the hardness measurements conducted on the surfacef the nitrided sample with variable indentation loads, it would alsoe expected that the plastic zone volume increased with inden-ation load and that such a volume was composed of materiallements with different hardness. Elements near the surface woulde harder than those away from it. Also, as the plastic zone volume

ncreases with indentation load, the number of elements whoseardness is closer to that of the steel core increases significantly,

eading to a more pronounced decrease in H with depth from sur-ace to core, in comparison with the previous case.

In spite of their differences, the correlation of the surface andross section hardness curves, as shown in Fig. 9, could be veryseful from a practical point of view since it would be possibleo predict the cross section hardness profile of any nitrided D2ool steel only from surface hardness measurements, together withqs. (1) and (3). Thus, by conducting surface hardness tests withifferent indentation loads and knowing the core hardness of theaterial (HS0), the value of the parameter HS1 (Eq. (3)), equal to HS

Eq. (1)), can be determined.Fig. 9 would then be employed for converting the surface

ardness into cross section hardness values, which would be sub-

equently employed, together with the indentation depth data, inq. (1) for determining the parameters k1 and k2 and therefore,he cross section hardness profile of the nitrided material. In thisay, there would be no need for conducting hardness measure-ents directly on the cross section of the specimen, but only on

Fig. 9. Change in hardness along the cross section of the nitrided D2 tool steel as afunction of the surface hardness determined with different indentation loads.

its surface, provided that the relationship between both measure-ments, presented in Fig. 9, remains unaltered due to the nitridingtreatment.

Regarding the behavior of the coated systems under indentationloads, Fig. 3 shows clearly that Eq. (2), coupled with Eq. (3) for thedescription of the behavior of the uncoated nitrided substrate, isquite satisfactory for the analysis of the experimental data. Fig. 3illustrates that the duplex system with a TiAlN film (DTiAlN) dis-plays the best behavior, as far as the capability of the system forenduring indentation loads. The curve corresponding to this sys-tem is shifted towards the right of the other curves, indicating thatthe system can withstand a given indentation load to higher rela-tive indentation depths, in comparison with the DTiN system andthe uncoated nitrided substrate. The rate of decrease of hardnesswith ˇ for both systems is observed to be similar and therefore thecoated materials tend to achieve progressively the core substratehardness at large ˇ values, which points out the satisfactory loadsupport provided by the nitrided substrate.

However, the analysis of static indentation curves cannot beextrapolated directly for the prediction of the behavior of the coatedsystems under wear conditions. For this purpose, load-carryingcapacity tests under increasing load or scratch tests could providemore information and better predictions. As shown in Fig. 8, atroom temperature the best wear behavior is displayed by the DTiNsystem, which has less capability of enduring indentation loadsin comparison with the DTiAlN system. At 300 ◦C the wear resultsare compatible with the outcome of the indentation tests but at600 ◦C the behavior is inverted entirely, which points out that thewear mechanisms are dominated by other phenomena, such as par-tial delamination of the coatings, presence of third-body particlesand oxidation, different to those that determine the behavior of thematerials under indentation loads. Figs. 4–6 clearly show the occur-rence of such phenomena, leading to distinct changes in the frictioncoefficient with running distance.

At 25 ◦C the wear tracks of both the DTiN and DTiAlN systems wereobserved to have a similar depth and the difference in the wear vol-

ume arises mainly from the wear width, which is less for the firstsystem. At 300 ◦C, the presence of a smaller a track for the DTiAlNsystem determines its better behavior in comparison with the DTiN,which displayed a narrower but deeper wear track. However, at600 ◦C, the wear results seem to indicate that the behavior is deter-

ar 267

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M.H. Staia et al. / We

ined mainly by the fracture, delamination and oxidation of theoatings. Such phenomena apparently lead to the formation of par-icles that contribute to the wear of the underlying substrate, whose

echanical strength decreases as temperature increases, giving riseo the development of deeper wear tracks for the DTiN and DTiAlNystems in comparison with the uncoated nitrided substrate, whichhows the best behavior.

. Conclusions

The wear experiments that have been conducted at tempera-ures of 25, 300 and 600 ◦C indicate that coating the nitrided D2ool steel substrate with TiN and TiAlN films deposited by PVDives rise to an improvement in the wear behavior up to a testemperature of 300 ◦C, in comparison with the uncoated substrate.t a temperature of 600 ◦C, besides oxidation of the coatings, theechanical strength of the substrate decreases leading to fracture

nd delamination of the films and the formation of third-body par-icles that contribute to its significant wear. At this test temperature,he uncoated steel substrate exhibited the highest sliding wearesistance, which gives rise to a significant reduction of the frictionoefficient. It has been shown that the nitriding process appliedo the steel substrate prior to PVD coating provides a satisfactoryoad support which contributes to the improvement of the coatedystems capability to withstand indentation loads at room tempera-ure. In this regard, the coated system with a TiAlN coating displayed

better behavior than that shown by the system with a TiN coating.n experimental procedure has been proposed in order to predict

he hardness profile of the nitrided tool steel, along the cross sec-ion of the material, just from hardness measurements taken on theurface of the sample, employing different indentation loads.

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(2009) 1452–1461 1461

Acknowledgements

The authors wish to acknowledge the financial support receivedfrom Venezuelan National Foundation for Scientific and Technolog-ical Research (FONACIT) through the projects UCV F-2001000600,PI No. 39270 and S1-2001000759 and to CDCH from UniversidadCentral de Venezuela.

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