esd Pa, U

d, G

ratel aspeessin tiffend pstates. Additionally, a Ti/TiAlN multilayer coating and a reference TiAlN monolayer

ction fponenequireman beas TiN,

maximum hardness of about 2832 GPa [9]. In addition, the aluminum chemical processes, such as plasma nitriding. Thus, grinding with geo-

e stresses. The use ofthe heat generation inrains [14,15]. Plasmaouter region of the

Surface & Coatings Technology 260 (2014) 369379

Contents lists available at ScienceDirect

Surface & Coatin

l secomplicated to achieve a good adhesion. The residual stresses inthe coating and in the substrate affect the adhesive and cohesive dam-age processes at the coating/substrate interface, which can either

substrate, and thus to provide a sufcient supporting effect of the basematerial for a subsequently deposited layer [16,17]. It has been demon-strated in many studies that the plasma nitriding process optimizes theDue to the possibility to produce near net-shaped coatings utilizing thePVD technique, forming tools with a very high surface quality can becoated. Besides the mostly low surface roughness of the tools,unfavorable residual stresses in the composite system make it more

al stresses, leading to a reduction of compressivharder grain materials can nevertheless reducethe contact zone since less wear occurs in the gnitriding is a common method to harden thecontent in the coating ensures a high temperature resistance up to800 C [10]. The formation of a thin, dense andwell adhering protectivelayer of aluminum oxide, which acts as a diffusion barrier and thusmin-imizes diffusion-inducedwear, is responsible for this behavior [9,11,12].

metrically undened cutting edges generally causes the induction ofcompressive residual stresses in the boundary layer of the substratema-terial due to supercial plastic deformations. However, friction-inducedtemperature increases can also promote the formation of tensile residu-promote or prevent a failure of the coating

Corresponding author at: Institute of Materials EnginDortmund, Leonhard-Euler-Str. 2, 44227, Dortmund, Germ

E-mail address: tobias.sprute@tu-dortmund.de (T. Spr

http://dx.doi.org/10.1016/j.surfcoat.2014.08.0750257-8972/ 2014 Elsevier B.V. All rights reserved.TiCN, or TiAlN, arewidelyear resistance [18]. Inating properties with a

cleaning the substrate surface (elimination of possible oxide layers),and adjusting the surface topography, structural changes of thesubstrate material are achieved by means of thermal- or plasma-used in the industry due to their high wparticular, TiAlN possesses outstanding cosingle or multilayer PVD coatings cTi-based ceramic coating systems, such1. Introduction

In order to ensure a sufcient proteagainst wear, a functionalization of comlayers are applied. Depending on the rTheir initial and nal residual stress states were measured by means of X-ray diffraction. In addition to theresidual stress analyses, tribo-mechanical tests, such as nano-indentation, ball-on-disc, and scratch tests wereperformed in order to correlate the results with these residual stress states.

2014 Elsevier B.V. All rights reserved.

or industrially used toolst surfaces, PVD protectiveents of the application,used for this purpose.

pre-treatments, selective coating architectures as well as adjustmentsof the deposition process enable to consciously alter the systemproperties and thus increase the life of PVD coated tools.

The substrate pre-treatments are used for the modication of theupper substrate region and include various sequential processingsteps. While mechanical pre-treatments are primarily responsible forwere deposited on each pre-treated substrate.Residual stressesTribo-mechanical behavior in their nal residual stressInuence of substrate pre-treatments on rtribo-mechanical properties of TiAlN-base

Tobias Sprute a,, Wolfgang Tillmann a, Diego Grisalesa Institute of Materials Engineering, Technische Universitt Dortmund, Germanyb Rif e.V. - Institut fr Forschung und Transfer, Joseph-von-Fraunhofer-Str. 20, 44227 Dortmun

a b s t r a c ta r t i c l e i n f o

Available online 30 September 2014

Keywords:PVDSubstrate pre-treatmentTiAlN PVDMultilayer coatings

Residual stresses in the substlifespan of machining as welstresses will lead to a betterconducted in the eld of strstresses in the substrate andIn this investigation, three dgrinding, and SiC grinding a

j ourna l homepage: www.e[13]. Suitable substrate

eering, Technische Universittany.ute).idual stresses andVD coatings

rsula Selvadurai a, Gottfried Fischer b

ermany

and in the PVD coating have a signicant inuence on the coating adhesion andforming tools. Therefore, the understanding and control of the system's residualrformance of the coated components. Moreover, although investigations wereanalysis of PVD coatings, they do not focus on interdependencies of residualhe coating.rent metallographically prepared substrates were used. SiC grinding, diamondlasma nitriding preparations were selected, due to the substantial differences

gs Technology

v ie r .com/ locate /sur fcoatadhesion and signicantly improves the wear resistance as well as thefriction behavior [1825]. Furthermore, it could be proven that the in-crease of the compressive residual stresses in the substrate is betterfor a good adhesion [26]. Selvadurai et al. found out that the residualstresses in the base material can be changed by the subsequent

370 T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379deposition of a PVD layer. The previous high compressive stresses areconverted into tensile residual stresses after coating with a Ti/TiAlNmultilayer system. Here, it is assumed that a shift of the compensatorytensile residual stresses towards the coating/substrate interface oc-curred in order to compensate the high compressive residual stressesin the multilayer coating system [27,28].

In addition to the pre-treatments of the substrate, the choice of thecoating design is further crucial for the performance of the coatingsystem. In contrast to monolayers, multilayer coating systems typicallyhave an increased toughness and thus extend the life span of thecompound system [29,30].Metallic interlayers allow the energy adsorp-tion by plastic deformation [31] and consequently promote relaxationprocesses, which can degrade the high residual stresses [32,33] andsimultaneously prevent delaminations of the coating [34]. At the sametime, the boundaries of the different individual layers inhibit crackgrowths and crack propagations by deecting the cracks in the layertransitions [4,29,35]. For example Castanho and Vieira investigated theinuence of the numbers of interlayers on the mechanical propertiesand residual stresses of alternating Ti/TiAlN systems and found outthat even one interlayer of titanium can halve the residual stresses inthe coating [36]. In addition, it was shown that the adhesion of thePVD coating is improved with an increasing Ti interlayer thickness ora decreasing proportion of ceramic [24], while the wear has increasedas a result of the declining hardness [26,37].

In order to further investigate these correlations, TiAlN monolayerand Ti/TiAlN multilayer coatings are deposited on different, pre-treated steel substrates and examined regarding their tribo-mechanical behavior as well as residual stresses.

2. Material and methods

Three different pre-treatments were performed on the hot worktool steel substrate X37CrMoV5-1, used for this research. The rstpre-treatment consisted of a metallographic preparation of thesubstrate using SiC grinding papers; the second pre-treatment wasa similar metallographic preparation, which, however, uses diamondgrinding discs instead of SiC grinding papers. Finally, a plasma nitrid-ing process was carried out on the rst pre-treatment in order to ob-tain the third pre-treatment. Herein, the different pre-treatmentswill be called SiC, Diamond and Nitrided, respectively. Theplasma nitriding process was carried out by an Arc PVD device(Metaplas, Germany), in which the samples were heated to 560 Cunder a bias voltage of650 V and a controlled atmosphere with aconstant pressure of 200 Pa for 8 h. The composition of the atmo-sphere consisted of 75 vol.% hydrogen and 25 vol.% nitrogen heldconstant during the complete nitriding process.

The dimensions of the samples were kept constant throughout theinvestigation and discs with a diameter of 40 mm and a thickness of4 mm were used.

The deposition process of the PVD coatings was carried out by anindustrial magnetron sputtering device (CemeCon MLsinox800,Germany). The samples rotated in the center of the chamber under a con-stant heating power of 5 kW during the deposition. This heating outputequates a temperature of about 400 C. The inert gases argon(295 sccm) and krypton (200 sccm) were used as plasma gases duringthe coating processes. The gas ow of the reactive gas nitrogen was con-trolled at a constant gas pressure of 580 mPa. Moreover, for all the depo-sition processes, a constant bias voltage of100 V was applied. Threealloyed TiAl and one pure Ti targets were mounted on the PVD deviceas cathodes. Furthermore, at the time of the deposition, a target powerof 9.5 kW and 4.0 kW, respectively, was applied.

Monolayers of TiAlN and multilayers of Ti/TiAlN ((50 nm + 500 nm) 5), with a total thickness of 3000 and 2750 nm respectively, were de-posited on the three different substrates and their thickness was evaluat-ed by using a scanning electron microscope (FE-JSEM 7001 JEOL, Japan).

In order to ensure a goodadhesive bondingof the coating and an excellentwear protection, themultilayer starts with a Ti-layer on the surface of thesubstrate and endswith a hard TiAlN layer.Moreover, the topography andmorphology of the deposited coatingswere evaluated by using a scanningelectronmicroscope aswell. An additional EDXdetectorwas used in orderto analyze the chemical composition of the ceramic layers.

Additionally, mechanical properties such as the coating hardnessand Young's modulus were determined by means of nanoindentationtests (G200 Agilent Technology, USA). A depth controlled penetrationwas performed at room temperature and, in order to avoid the inuenceof the substrate on the properties of the thin lm, the results were eval-uated in a range from10 to 15% of the total coating thickness, whichwasin this case between 100 and 400 nm.

With the purpose of determining the tribological properties of thecoatings, ball on disc tests at room temperature were conducted usinga tribometer (high temperature tribometer CSM, Switzerland) equippedwith aWC/Co ball as counterpart to the rotating samples with the aim ofanalyzing the wear coefcient and with 51CrV4 pins as a counterpart toobtain the friction coefcients of the different systems. 51CrV4 is a steelused for quenching and tempering according to DIN EN 10083, and it isusually used in the automotive and mechanical engineering industryon components formed of metal such as gear parts, pinions, and shafts.

During the experiment, which consists of 8000 rotations, no externallubricant was used and both the normal force and linear velocity of therotating discs were kept constant at 5 N and 40 cm/s, respectively. Thewear coefcientwas evaluated by analyzing thewear trackswith an op-tical 3D surface analyzer (Innite Focus Alicona, Austria) that consists ofa confocalmicroscope connected to an image analyzer software, and thewear mechanisms evidenced by SEM.

The scratch tester Revetest (CSM, Switzerland) was utilized to ex-amine the adhesion between the PVD coatings and the steel substratesat room temperature. For this, scratch tracks were generated with atotal length of 10 mm. The force was steadily and linearly increasedfrom 0 to 100 N and the results were analyzed using SEM in combina-tion with an EDX-detector (Oxford Instruments, UK).

The residual stress evaluation was executed by means of a diffrac-tometer (Bruker Advance D8), using the sin2method [38]. Fe-K radi-ation was used instead of the usually used Cu-K radiation, in order toavoid the uorescence radiation from the substrate [39].

Before proceeding with the residual stress measurements, phaseanalyses of themetallic substrate, TiAlNmonolayer, and Ti/TiAlNmulti-layer were performed to determine the present phases and to establishthe 2 angles related to the Bragg law, which thereafter are used for thedetermination of the residual stresses. Consequently, the Fe 220 peakfound at an angle of 2 equaling 145was selected tomeasure the resid-ual stresses in the substrate. This specic reection was chosen insteadof other Fe reections, due to the high value of the angle 2which ben-ets the sensitivity of the sin2method, where small changes in the lat-tice spacing, d, result in a corresponding shift of the diffraction angle 2.

For the analysis of the residual stresses in the coating systems, theTiAlN 220 peak was selected and Fe-K radiation was used as well.The reectionswere scanned and afterwards tted by the Pearson func-tion to determine the positions of the peak. The X-ray elastic constantsof Fe were calculated by elastic single crystal constants. For the TiAlNcoatings, the XEC were calculated after Voigt by utilizing the experi-mentally determined macroscopic elastic moduli of these coatings.

To analyze the residual stresses in the substrates after the depositionof monolayer and multilayer coating systems with X-ray diffraction,several conditions need to be considered. In homogeneous materials,the following relationship between the beam intensities at themomentwhen the X-ray enters and leaves the specimen (hereinafter denoted asI0 and If, respectively) holds true:

I f z I0 exp k z : 1

Here, stands for the linear attenuation coefcient of the material

and z is dened as the depth. The variable k denotes a geometry factor

bstr

371T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379which can be calculated from the incidence () and exidence () angleor from the goniometer angles , and the Bragg angle 2. In thinmultilayer coatings, combining two materials, the penetration depth isinuenced by the attenuation coefcient of the selected materials andthe thicknesses of the individual layers. Hence, special formulas weredeveloped by Fischer et al. to calculate the penetration depth in thecompound systems, depending on the goniometer angles [40]. Thefollowing formulas were used to calculate the penetration depth inthe substrate, coated with a monolayer (Eq. (2)) and with a multilayer(Eq. (3)), respectively:

1c k

ac1

ta 2

1c k

n ac1

ta

bc1

tb

: 3

Eqs. (2) and (3) stand for the phases a, b and c, respectively.

Fig. 1. Residual stress depth proles in the suHere, n is the total number of layers of phases a and b and t standsfor the thickness of the single sub-layers. In this particular case, acorresponds to TiAlN, b to Ti metallic interlayers, and c to thesubstrate material.

Afterwards, using Eqs. (2) and (3), the penetration depths of theX-rays, for the reection Fe 220, in the different systems used for this

Fig. 2. Coating thickness from TiAlN monolresearch at the different angles, were calculated. Here, attenuationcoefcients of Fe-K radiation in Fe was 552.6 cm1 [41], in Ti was1731.9 cm1 [41] and in ceramic layer of TiAlN, determined proportion-ally from TiN and AlN values, was 800.084 cm1 [42].

The maximum penetration depth of the uncoated substrates wasequal to 8.64 m. After the deposition of the monolayer, this depth isreduced to a value of 4.3 m from the interface monolayer/substrate, avalue obtained by using Eq. (2). In the end, theX-ray penetration depthsfor the multilayer/substrate systems were also computed by applyingEq. (3). Here, the X-rays reach a maximum depth of 4.24 m, which isa lower value when compared to the previously obtained range of themonolayer/substrate compound, due to the high attenuation coefcientof Ti metallic interlayers. These X-ray penetration depths can be seen inFig. 1. Moreover, the residual stress depth proles were obtained byelectrolytic polishing of the surface of the substrate (LectroPol-5 Struers,Denmark) and measuring the electro-polished depths with an optical3D surface analyzer (Innite Focus Alicona, Austria).

3. Results and discussion

ate for the different substrate pretreatments.3.1. Substrate pre-treatments and their inuence on residual stresses

The aim of this work is to study the inuence of different substratepre-treatments on the tribo-mechanical coating properties as well ason the residual stresses in composite systems. This includes the

ayer and Ti/TiAlN multilayer systems.

bstr

372 T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379induction of different residual stress states in the substrate as well. Inthis case, it has to be taken into account that conventional pre-

Fig. 3. a. SEM images of the TiAlNmonolayer systemdeposited on the different pretreated susubstrates.treatments are used, which provide reproducible results aswell as com-parable surface qualities and roughnesses. Therefore, two differentgrinding processes and a combination of grinding and nitridingwere se-lected. The different abrasive materials cause varying degrees of me-chanical machining or deformations in the surface of the substrate onthe one hand, while generating an uneven heat input on the otherhand. This results in two different residual stresses in the substrate.Further pre-treatment includes an additional nitriding process, whichis responsible for the hardening of the substrate and for the extremelyhigh compressive stresses in the upper substrate region. The differentresidual stress states are shown in Fig. 1 as residual stress depth proles.Compared to the diamond grinded substrates, the SiC grinded

Fig. 4. Arithmetic average roughness Ra of the substrates before coatinsubstrates exhibit signicantly lower compressive residual stresses. Inaddition, the inuence zone of the machining process is not as deep as

ates. b. SEM images of the Ti/TiAlNmultilayer systemdeposited on the different pretreatedin the diamond prepared substrates. This can be mainly explained bythe choice of the abrasive material. In general, the mechanical machin-ing of the fringe results in the formationof compressive residual stressesdue to supercial plastic deformations. However, friction-induced tem-perature increases can lead to the reduction of compressive stresses.Softer grain materials, such as SiC abrasive grains, wear out faster(blunting effect) and thus induce more heat into the contact zonewhile harder grain materials (diamond grains) lead to a deeper me-chanical deformation and consequently induce compressive residualstresses [14]. During the nitriding process, the crystal lattice is extreme-ly distorted and clamped in the diffusion zone. Finely dispersed nitridesand dissolved nitrogen atoms in the iron lattice cause compressive

g deposition and after deposition of both mono- and multilayer.

373T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379residual stresses in the substrate material [13,17,43,44]. The area ofinuence (diffusion zone), which is also visible in Fig. 1, is much deepercompared to the grinded substrates. Compensating tensile residualstresses can only be observed beginning at a depth of approximately100 m.

3.2. Coating thickness and structure

Measured thicknesses of all the different systems as well as thedeviations of the desired coating thicknesses are shown in Fig. 2. It isclearly visible that the compound systems, except from the monolayerdeposited on the SiC grinded substrate, achieve the planned lm thick-nesses. Nevertheless, the outlying thickness value obtained for themonolayer system deposited onto SiC is denitely in the range of toler-ance accepted, using an industrial magnetron sputtering device.

In order to analyze the rst layer/substrate bonding as well as thecoating structure, fracture patterns were evaluated with a scanningelectron microscope. This procedure further allows the detection of

Fig. 5. Topography SEM pictures of the mono- and mult

Table 1Chemical composition of the TiAlN ceramic layer by mean of an EDX analysis.

Element at.-% with N at.-% without N

Ti 19.98 0.43 43.66 0.21Al 25.18 0.35 56.34 0.21N 54.84 0.78 possible defects in the coating. The facture patterns of the monolayerand of the multilayer are shown in Fig. 3a and b. In both gures, theleft columns display the coatings in the secondary electron mode inorder to examine structural and morphological differences, whereasthe right columns show fractured surfaces in the backscattered electronmode to showdifferentmaterial contrasts. All TiAlNmonolayer systemsin Fig. 3 show the same structure due to the same parameter during thedeposition processes. The coatings are very dense and free of defects.Furthermore, there are no spallations or delaminations recognizable atthe interfaces, which probably indicate a remarkable adhesion of thecoatings. Themorphologies exhibit a predominantly glass-like structurewith a very slight orientation in the growth direction. However, a clearcolumnar structure cannot be observed.

The coating design as well as the coating structure of the multilayersystems can be seen in Fig. 3b. Especially the images with different ma-terial contrasts clearly visualize the alternating metallic and ceramiclayers. The thin Ti interlayers have constant thicknesses within thecomplete coating system. In the structural images, left column, themetallic interlayers are only shown as breaks of the TiAlN layers,which are probably responsible for the inhibition of crack growth andpropagations [24,45,46]. In addition, all multilayer systems also showno imperfections in the layer or spallations and delaminations at theinterfaces. Although all coatings have very dense structures, supposeddifferences in terms of orientations are identied. Themultilayer depos-ited onto SiC grinded and nitrided substrates shows a randomlyoriented structure like the monolayer coatings, while the multilayer

ilayer deposited on different pretreated substrates.

Fig. 6. X-ray phase diffractogram of the Ti/TiAlN multilayer system, Fe-K radiation.

374 T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379deposited on the diamond grinded substrate presents a slightly colum-nar structure. However, the selected process parameters were constantfor all systems.

In addition, the comparable structures and morphologies of alldeposited coatings can also be attributed to the similar nucleation onthe surface of the substrate. It can be assumed that the surface rough-nesses of the differently treated substrates prior to the deposition areresponsible for this behavior [4749]. The values are in the range ofRa = 101 nm and Ra = 149 nm (Fig. 4). Due to the slight rougheningduring the nitriding process by sputtering the substrate material, thenitrided substrate has the highest roughness with Ra = 149 nm [43].The SiC and diamond grinded substrates possess lower roughnessvalues of Ra = 116 nm and Ra = 101 nm, respectively. The trend ob-served for the different roughness values is constant both before andafter the deposition of the coatings. Thus, it can be noted that the rough-ness of the substrates before the deposition affects the roughness of thedifferent coatings. However, in general, the roughness values increaseafter the coating deposition of mono- and multilayer systems due tothe sum of the substrate roughness and the intrinsic roughness of thecoating itself [47].

The initial roughness of the surface of the substrates is not onlyreected in the roughness values after the coating application, but alsoin the topography images of the systems (see Fig. 5). Here, the sametrend as in the roughness values can be observed. The coated surfaceof the nitrided base material has a coarser texture than that of thegrinded and coated samples.

The chemical composition of TiAlNwas determined by an EDX anal-ysis. The chemical composition of themonolayer and the ceramic layersof the multilayer system are shown in Table 1. The alternating Ti layers

used in the multilayers consist of pure Ti. Due to the fact that light

Fig. 7. Residual stress evolution for the different substrates beforelements are difcult to determine quantitatively by EDX, the composi-tion is examined with and without N. According to the literature,sputtered coatings with an Al/Ti ratio of about 3/2, generally exhibitfcc NaCl-type crystallographic structure with excellent tribo-mechanical properties [9,50]. Nonetheless, different investigations asthe ones performed by Makino [51] and Greczynski [52] should beconsidered, since both authors have attributed the obtaining of cubicor hexagonal phases to the use of different deposition techniques,instead of different Al contents in the TiAlN structure.

3.3. Development of residual stresses

Residual stress depth proles for each substrate pre-treatment weredeveloped in order to determine the residual stress state at differentdepths from the surface of the substrate and to evaluate their gradient,Fig. 1. Hereby, the substrate was electrolytically polished one step afterthe other to a depth of approximately 120 m, and residual stress eval-uations were conducted for every step. In the detailed graph, shown inFig. 1, it can be observed that the residual stress states for the three dif-ferent substrate pre-treatments present a very low gradient in depthsbetween 0 and 10 m. The penetration depth of the X-rays during thedifferent residual stressmeasurements in the substrate of these samplesreaches up to 10 m, being located within the above described region.The residual stress measurements of the three different pre-treatedsubstrates before and after deposition, in this case, do not depend onthe deposited coating and thus are comparable.

Moreover, in order to measure the residual stresses in the coatings,phase analyses were performed on themono- andmultilayers to denethe most suitable reection. Fig. 6 shows a phase diffractogram of a Ti/

TiAlN multilayer system, deposited onto the substrate which was

e and after deposition of TiAlN monolayer, Fe 220 reection.

Fig. 8. Residual stress evolution for the different substrates before and after deposition of Ti/TiAlN multilayer, Fe 220 reection.

375T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379metallographically prepared with SiC grinding discs. The reectionTiAlN 220 at the 2 angle equal to 82, showing a sufciently high inten-sity and no overlapping, was chosen for the stress evaluation.

Figs. 7 and 8 show the evolution of the residual stresses in thesubstrate for each substrate pre-treatment. The residual stresses weremeasured prior to the metallographic preparation (here namedunprepared), after each preparation (identied with the name of eachpre-treatment), and after the deposition of the mono- or multilayercoatings.

Before the substrate pre-treatments, the anisotropic character of theresidual stresses in the substrate is evident due to the differencebetween the values of the stress tensor 11 and 22. Thus, 11 presentscompressive residual stresses due to the transversal contraction derivedfrom thepreferential cutting direction of the samples,while22 exhibitsan absence of stresses.Moreover, after the pre-treatments conducted onthe different metallic substrates, no big differences are found between11 and 22, corroborating the isotropic behavior of the resultingresidual stresses after each preparation. Subsequently, three differentresidual stress states were identied: SiCwith low compressive residualstresses, diamondwith slightly higher compressive residual stresses, andlastly, nitrided samples presenting the highest value of compressive re-sidual stresses. These results were correlated to each other and areshown in ascending order, from left to right in Figs. 7 and 8.

In contrast to the uncoated substrates, the residual stresses in all thesubstrates after depositing the different coating systems are reduced,this behavior was also found by Tnshoff et al. [53]. Nitrided substratesamples show the greatest reduction of residual stresses after theFig. 9. Residual stresses in tdeposition of the coatings whereas diamond and SiC grinded substratesshow the smallest reduction which is comparable with the results ob-tained by Selvadurai et al. [28]. Moreover, a slightly greater reductioncan be evidenced in the substrates depositedwith TiAlNmonolayer sys-tems than for those depositedwith Ti/TiAlNmultilayer systems. This re-duction cannot be only attributed to a single factor, as differentparameters affect the induction of residual stresses in the compoundsystems, coating/substrate. One of these reasons is the temperature ofapproximately 400 C, at which the deposition process is carried out ap-proximately, andwhich can cause a relaxation of residual stresses in thesubstrate [13]. Other reasons which are supposed to affect residualstresses are for instance the different thicknesses between the mono-and multilayers [54], and the lattice mist between the substrate andthe growing coating [13].

Because of the previously described randomly orientated structurein the deposited coatings, a strong linear regression is obtained betweensin2 the lattice spacing dwhich is the base for the residual stress calcu-lations in the coatings [38]. In Fig. 9, the inuence of the substrate prep-arations on the residual stresses of the coatings is evident. Hence, thesame coatingmaterial, which is supposed to have the same intrinsic re-sidual stresses, presents higher residual stresses in those samples inwhich the substrate shows higher residual stresses, caused by the pre-treatment of the substrate. At the same time, clarifying that both TiAlNand Ti/TiAlN deposited lms follow the trend described above, highercompressive stresses were found in the multilayer systems depositedon substrates with the same pre-treatment. This behavior was also evi-denced by Selvadurai et al. [28]. The presence of higher compressivehe coatings, TiAlN 220.

Fig. 10.Mechanical properties of the TiAlN monolayer and Ti/TiAlN multilayer deposited on different pretreated substrates.

376 T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379residual stresses in the multilayer systems can also be attributed to thepresence of thermal stresses, which can be found in PVD coatings due totemperature changes during the deposition, and differences on thethermal expansion coefcients of the coating system constituents [13,55]. For instance, the thermal expansion coefcient for X37CrMoV5-1steel is equal to 13 106 C1 [56], for Ti interlayers it is 8.5 106 C1 [57], and 7.37.5 106 C1 for TiAlN ceramic layers[5860]. Compressive residual stresses are developed in the materialwith the lower thermal expansion coefcient [13]. This rule applies formetallic Ti interlayers and ceramic TiAlN layers,where titaniumexhibitsa higher thermal expansion coefcient, increasing the compressive re-sidual stresses in TiAlN layers of the multilayer coatings. Additionally,as the multilayer systems present more dissimilar material interfaces(substrate/Ti and Ti/TiAlN), the dislocation density is increased inthese areas as well, causing an increase of the compressive residualstresses at the proximity to the interfaces [13].

Finally, analyzing Figs. 7, 8, and 9 as a single entity, the following hy-pothesis can be assumed: The compressive residual stresses in the outerregions have to be compensated by tensile residual stresses in themate-rial core (substrate or substrate/coating system). After the deposition ofa PVD coating, a displacement of the tensile residual stress region to aposition closer to the interface substrate/coating occurs, causing thetranslation of the stresses between the coating and the substrate.

The same statement has been made by Denkena and Breidenstein[61], Tnshoff [53] and Bunshah [62], who have conrmed the relief ofcompressive stresses in the substrate in order to balance the residualstress state in the compound coating/substrate.Fig. 11. H/E ratio of the TiAlN monolayer and Ti/TiAlN mu3.4. Mechanical properties

Hardness and Young's modulus values are presented in Fig. 10. Asdescribed in Material and methods section, the nanoindentation exper-iments were performed with a Berkovich tip; and the data were ana-lyzed at a depth of 1015% of the overall coating thickness.Consequently, the hardness and Young's modulus values are constantfor all TiAlN monolayers and thus for all Ti/TiAlN coatings depositedonto the three different pre-treated substrates. It has been previouslyreported by different researches that TiAlN monolayers, deposited bymeans of magnetron sputtering, present a hardness value around30 GPa, which is comparable to the values obtained for this paper [63,64]. According to Hrling et al., Ti1 xAlxN monolayers with x = 0.66,have as a consequence of the grain size (HallPetch relation),deposition-induced point defects, and precipitation hardening fromcompositional inhomogeneities [65] a hardness equaling 33.1 GPa.

However, higher values for both hardness and Young's modulus arefound in the monolayer coatings when compared with the multilayersystems for the same substrate pre-treatment. This behavior can be ex-plained by the high metallic content of Ti in the multilayer bi-layer sys-tem, in this case 10% of the adjacent TiAlN layer. Moreover, the totalamount of hard ceramic material is also reduced in the multilayer sys-tem in relation to the TiAlNmonolayer. In further investigations, the in-uence of the Ti interlayer thickness on the hardness has to beevaluated.

Additionally, the relation H/E is shown in Fig. 11. This relation repre-sents the plasticity index or elastic strain to failure and it is a suitableltilayer deposited on different pretreated substrates.

Fig. 12. Critical load of the TiAlN monolayer and Ti/TiAlN multilayer deposited on different pretreated substrates.

377T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379parameter to predict thewear resistance and to explain the deformationproperties ofmaterials by considering the elastic rebound [66]. Thus, forthe present research, high H/E values for the Ti/TiAlN multilayer areobserved on all pre-treated substrates, which guarantee a higher tough-ness of the system and might be an indication for a better tribologicalprotection.

Critical load is an evidence of a better adhesion of the coatings to thesubstrate [66]. Fig. 12 shows the critical loads for the investigated sys-tems. Here, Lc3 represents the load valuewhere the coating is complete-ly removed from the center of the scratch. SiC and diamond preparedmetallic substrates presented the lowest adhesion for both mono- andmultilayer coating systems. During the scratch test, coatings depositedon substrates with a low hardness result in a plastic deformation ofthe substrate, promoting the crack formation and fragmentation of thecoating [63]. In contrast, a higher critical load value for those coatingsdeposited onto substrates with higher hardness, nitride substrates, isevidenced (Fig. 7). Values of Lc3 equal to 48.6 and 56.1 N for TiAlN andTi/TiAlN, respectively, were obtained. According to the abovementionedresults, an improvement of the adhesion of the coatings is observed dueto the presence of a metallic interlayer in the Ti/TiAlN multilayers [26,32]. The great inuence of the substrate conditions on coating scratchbehavior can be observed as well. Nonetheless, the inuence of theroughness of the substrate on the critical load of the coatings beforedeposition should be considered. For substrates with high roughness(in this case a nitrided substrate), the contact area between thesubstrate and the coating at the interface substrate/coating is largerthan for substrates with a low roughness. This implies an enhancedFig. 13.Wear coefcient of the TiAlN monolayer and Ti/TiAlNmechanical and physical bonding between the coating and the sub-strate, subsequently increasing the critical load value [67]. However,the various hardness results of the substrate materials exhibit greaterdifferences than the roughness values. Thus, a major effect on thecritical load value is rather attributed to the hardness than to the rough-ness differences, even though these twoparameters are closely linked inthis case.

3.5. Tribological behavior

Fig. 13 shows thewear coefcient obtained for the different coating/substrate systems evaluated within this research. Hereby, the abrasivewear mechanism was identied as a consequence of the high hardnessof the tribological counterpart used, WC/Co. Moreover, both the TiAlNmonolayer and Ti/TiAlN multilayer have shown a reduction of thewear coefcient with increasing hardness and compressive residualstresses in the substrate. In addition, the wear resistance correlateswith critical load and thus adhesive bonding of the coating on the sub-strate material. Therefore, the highest wear values were obtained atthe coatings deposited on SiC prepared substrates and the lowest valuesonto nitride substrates.

Furthermore, Ti/TiAlN coatings have shown a poor tribologicalbehavior during wear resistance tests, which is evident in the higherwear coefcient than this obtained for TiAlN monolayers, deposited onthe same type of substrate. This trend might result from the Ti metallicinterlayer thickness which, despite stopping the crack propagation, fur-ther reduces the system stability, as Ti interlayers can be plasticallymultilayer deposited on different pretreated substrates.

osi

378 T. Sprute et al. / Surface & Coatings Technology 260 (2014) 369379deformed [26,37]. Furthermore, it has to bementioned that the Ti/TiAlNmultilayer is thinner (2750 nm) than the monolayer (3000 nm) andpossesses a lower amount of total ceramicmaterial, a fact that undoubt-edly inuences the total wear coefcient.

Moreover, higher residual stress values were found in Ti/TiAlN mul-tilayer systems when compared to TiAlN monolayers, Fig. 9, promotingpartial spallations, which are reected in higher wear coefcients ofthese coating systems.

Friction coefcients of the coatings were evaluated with two differ-ent material counterparts, WC/Co and 51CrV4 tool steel (Fig. 14). Theexperiments carried out using 51CrV4 pins have presented a relativelyconstant value for all the substrate pretreatments and no trend derivedfrom the substrate pre-treatment was recognizable since the externallayer of the coatings always consisted of TiAlN. Nevertheless, frictionvalues obtained for the Ti/TiAlNmultilayer systemswere slightly higherthan those for TiAlN monolayers. Due to the low hardness of the pincounterpart, the wear mechanism for this analysis was identied asthe adhesion of the pin material on the coating, also called micro-coldwelding. Consequently, friction values obtained using 51CrV4 werehighly inuenced by themetallicmaterial, derived from the pins, depos-ited onto the coating systems.

Moreover, the experiments performed using WC/Co balls showedthat substrates possessing higher compressive residual stresses alsopresented lower friction values. Such a behavior, comparable to theresults of the wear test, can probably be attributed to the big differenceof the hardness between the non-nitrided substrates and the coatings[26]. Unquestionably, high elastic and plastic deformations of the softsubstrates limit a good tribological performance of the hard TiAlN and

Fig. 14. Friction coefcient of the TiAlN monolayer and Ti/TiAlN multilayer depTi/TiAlN coatings and assist the plowing effect on the coating systems[63]. Finally, in the present research, no strong correlation betweenthe roughness and tribological behavior of the coatings can be seen.

4. Conclusions

Different substrate pre-treatments were performed and their resid-ual stresses before and after the deposition of TiAlN monolayers andTi/TiAlN multilayers analyzed. This research was carried out in orderto correlate the initial and nal residual stress states with the tribo-mechanical properties of the deposited coatings. The main conclusionsof this investigation are:

Different residual stress states and gradients are founddependingonthe different metallic substrate pretreatments.

Nitrided substrates present the highest compressive residual stress-es, and their hardness results are benecial for the tribo-mechanicalbehavior of the coatings.

Low residual stresses in the substrate lead to high elastic and plasticdeformations of the substrate and, consequently, have a negativeeffect on the tribo-mechanical properties of the coatings. Residual stresses in the coatings are highly inuenced by the residu-al stresses in the substrate before the deposition process. Therefore,coatings deposited onto nitrided substrates present the highestcompressive residual stresses. Moreover, the highest reduction ofthe residual stress values of the substrate after deposition arefound in nitrided substrates, ranging fromhigh compressive residualstresses to a value close to zero.

Ti/TiAlN ((50 nm + 500 nm) 5) multilayer systems, despite allexpectations, present higher residual stresses than TiAlNmonolayersystems and therefore, an inferior tribo-mechanical response thanthe monolayer.

The toughness of themultilayer system is higher than the toughnessof the TiAlN monolayer according to the obtained H/E ratio.

The adhesion of the coatings is improved by increasing hardness ofthe substrate which is also related to an increase of compressive re-sidual stresses. Ti metallic interlayers of the Ti/TiAlN multilayersystem inhibit crack propagations and enhance the critical load.Moreover, the small disparity between the residual stresses of thesubstrate before deposition and the residual stresses of the coatingresults in a benecial adhesion.

With the aim of continuing the residual stress investigations and inorder to nd answers to different behaviors which cannot be explainedat this point, further investigations are planned:

Fatigue tests of different substrate pre-treatments of coating systemsshould be conducted.

ted on different pretreated substrates using WC/Co and 51CrV4 counterparts. Different coating architectures have to be investigated, among themdifferent are Ti thicknesses in multilayers, graded systems, nano-composites and nanolayers, only to mention a few.

Nanoindentation with spherical indents and an analysis of the loaddisplacement curves should be performed in order to obtain moreinformation about the elasto-plastic properties of the compoundsystems. Moreover, conclusions about the toughness and crackpropagation in the coating can be obtained by analyzing crosssections through the spherical indents.

Residual stresses after the deposition and their behavior under aninduced tensile load should be studied, investigating mini tensilespecimens by means of X-ray analysis.

Acknowledgment

The authors gratefully acknowledge the DFG (German ScienceFoundation) for the nancial support for this work within the projectsTi 343/34-1 and Fi 686/8-1.

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