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Arc-Evaporated Nanoscale Multilayer Nitride-Based Coatings for Protection Against Wear, Corrosion, and Oxidation

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PACSnumbers:62.20.Qp,62.25.-g,68.65.Ac,81.15.-z,81.40.Pq,81.65.Kn,81.65.Mq

Arc-Evaporated Nanoscale Multilayer Nitride-Based Coatings for Protection Against Wear, Corrosion, and Oxidation

A. D. Pogrebnjak*, O. M. Ivasishin, and V. M. Beresnev

*

Sumy State University, 2 Rymsky-Korsakov Str., UA-40007 Sumy, Ukraine

G. V. Kurdyumov Institute for Metal Physics, N.A.S. of Ukraine, 36 Academician Vernadsky Blvd.,

UA-03680 Kyiv, Ukraine

V. N. Karazin Kharkiv National University, 4 Svobody Sq.,

UA-61000 Kharkiv, Ukraine

Thestudiesof thestructureand properties ofnanoscale multilayercoatings

basedonthenitridesofrefractorymetalsaresummarizedinabriefreview.By

theexampleofTiN/MoN,TiN/ZrN,CrN/MoN,andmorecomplex(multilayer)

(TiZrNbTaHf)N/WN:(TiZrNbTaHf)N/MoN obtained by vacuum-arc

deposi-tionofcathode,thedependencesoftheirhardness,wearresistance, friction,

corrosion,andoxidationonconditionsofthedepositionandlayers’thickness

areinvestigatedandanalysed.Theregularitiesofthestructureandbehaviour

properties of such nanoscale multilayer coatings depending on the size of

nanograins,textures,andstressesarisinginthesecoatingsaredescribed.

Укороткомуоглядіузагальненорезультатидослідженьструктурита

влас-тивостей наномасштабних багатошарових покриттів нітридів

тяжкотоп-кихметалів.НаприкладіTiN/MoN-,TiN/ZrN-,CrN/MoN-табільш

склад-них(багатошарових)(TiZrNbTaHf)N/WN:(TiZrNbTaHf)N/MoN-покриттів, одержанихметодоювакуумно-дугового осадження катоду,досліджено та проаналізовано залежностіїхтвердости, зносостійкости,тертя, корозії й окисненнявідумовосадженнятатовщинишарів.Відмічено закономірнос-тіструктуритавластивостейповедінкиданихнаномасштабних багатоша-ровихпокриттіввідрозмірунанозерен,текстуританапруг,щовиникають уцихпокриттях. Вкраткомобзореобобщенырезультатыисследованийструктурыисвойств наномасштабных многослойных покрытий нитридов тугоплавких

метал-лов.НапримереTiN/MoN-,TiN/ZrN-,CrN/MoN-иболеесложных

(много-Ôîòîêîïèðîâàíèå ðàçðåøåíî òîëüêî

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слойных) (TiZrNbTaHf)N/WN:(TiZrNbTaHf)N/MoN-покрытий, получен-ныхметодомвакуумно-дуговогоосаждениякатода,исследованыи проана-лизированызависимостиихтвёрдости,износостойкости,трения,коррозии и окисления от условий осажденияи толщины слоёв. Отмечены законо-мерности структуры и свойствповедения таких наномасштабных много-слойныхпокрытийотразмерананозёрен,текстурыинапряжений, возни-кающихвэтихпокрытиях.

Keywords:nanoscalenanocompositecoatings,structure,wear,corrosion

re-sistance,hardness. Ключовіслова: нанорозмірнінанокомпозитні покриття, структура, зно-шування,антикорозійнастійкість,твердість. Ключевые слова: наноразмерные нанокомпозитные покрытия, структу-ра,износ,антикоррозийнаястойкость,твёрдость. (Received January 19, 2016) 1.INTRODUCTION

One of themost promising applications ofnanomaterials is the

crea-tionofprotectivecoatingsforproductsandtoolswithdifferent

func-tionalpurposes. Suchmaterialcharacteristicsas hardness,elasticity,

adhesiveandcohesivestrength,durability,thermalandchemical

sta-bilityandothersareparticularlyimportantinthisregard[1–15].

Results of scientific researches showthe tendency ofactive useof

nitridesandboridesoftransitionmetalsandtheircombinationinthe

developmentofprotectivematerials[16–25].Whilenitridesofsingle

elementsarestudiedwellenough,theirmultilayermodificationsneed

studythatismoredetailed[26–35].Therefore,thestudyoffeaturesof

structure,elementalandphasecompositionofmultilayercoatings

de-pendingonthedepositionconditionsisanimportanttaskinsolid-state

physicsandmaterialsscience[36–51].

2.RESULTSANDDISCUSSIONS 2.1.TiN/ZrNCoatings

We clearly see layers with cubic TiN and ZrN phases (of the NaCl

structuretype)withoutpreferredorientationofcrystallitesinthe

sur-facelayers.IncreasingoftheperiodledtoincreasingoftheTiN

lay-ersspecificcontribution.Itisseenfromthechangesintheintensityof

thepeaksofTiNandZrNphases(seeFig.1).Increasingofthe

deposi-tion timeand, as aresult, bilayerthickness as wellas total period of

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phases in the layers. Lattice period decreased with increasing of the

TiN-layers’thicknessfrom0.4241502nm(depositiontimewas20sec,

70 nm)to0.4238870nm (depositiontimewas40sec,250nm).

ChangesforZrNlayerswerenotsolarge:from0.4581055nm

(deposi-tiontime—20sec)to0.4581046nm(depositiontimewas40sec).

Annealingintheoxygenatmosphereledtooxidationofthesurface

layers [19, 20] and to forming of dioxides as the main phases: TiO2

(withtetragonalrutile-typecrystalsystem;basic,upto95vol.%inthe

layers basedontitanium)and anatase (5 vol.%and less).Wecan

de-termineanatase(DBcardnumber5000223)onthediffractionspectra

(Fig. 1,a) usingthemoststrong firstlineontheangle225.36

de-grees. Diffractionspectrum is shownforthe rutile(DB cardnumber

9007531)(seeFig.1,a).OnlyonetypeofdioxideZrO2 (arkelhavinga

cubiccrystalsystem,DBcardnumber5000038)wasformedinthe

zir-coniumnitridelayersafteroxidation.

Layered X-ray analysis showed that after removal of the surface

layer(thicknessofabout5microns)bypolishing,wecouldseedioxides

onlyinthesubsurfacelayerofthecoating.Inmorethicklayers(Fig.1,

b), we did not find oxides whereas nitrides are characterizedby

pre-ferredorientationofcrystalliteswiththe[111]axis,perpendicularto

the plane of growth. Therefore, preferred orientation with the [111]

axiswasformedatthebeginningstageofgrowthforbothTiNandZrN

crystallites. Increasing of the total thickness of the coatings and

re-laxation of the compressive stresses led to disorientation of

crystal-lites,i.e.preferredorientationwasnotobserved.

UsingXRDanalysis data(Fig.1),wecouldassumethatdueto

dis-orientation of crystallites in the surface layers and low compressive

stresses, oxygen fromthe atmosphere penetrates intothe subsurface

layersduringdeposition.Itformedstabledioxidephasesofmetalsdue

Fig.1.DiffractionpatternsoftheTiN/ZrNcoatings(70nm)afterthermal

annealingunderthetemperatureof700Cforonehour:a—fromthesurface

withoutpolishing,b—afterpolishingoftheoxidizedsurfaceonthedepthof5

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toeasierdiffusionbyintercrystalliteways.Inmoredeeplayers

corre-sponding to thebeginning stages of growth, [111] textureof growth

appearedduetocompressivestresses[7,8].Thistextureprevents

dif-fusion of oxygen into such layers due tohigh packing densityof the

plane(111),sothereisnotenoughoxygentoformdioxidephases.

Transmissionelectronmicroscopy(TEM)allowedprovidingdetailed

analysisofchangesinthesurfacelayersafteroxidation.Electron

im-agesofthecross-sectionsoftheTiN/ZrNsamplesfromdifferentseries

arepresentedinFig.2.Goodplanaritywasobservedevenforthe

thin-nestlayersofthecoatingsfromthefirstseries(Figs.2,aandb).High

continuityofthecoatingsandthelackofinhomogeneity,suchas

drop-letfractions,arealsotypicalforinvestigatedsamples.

Comparisonofelectronmicroscopyimagesofthestructuralstateof

the layers in the multilayer TiN/ZrN coatings with total amount of

layers134(Fig.3)showedthatincreasingofthevolumefractionledto

bending oflayers, stratification with separation and loss ofstrength

during oxidation (Figs.3, c and d). We observed dome-like

disconti-nuitiesintheareasofpartialseparationoflayers(Fig.3,d)onthe

sur-face.Comparisonofstructures1and2inFig.3showedthatmain

vol-umechangestookplaceinthetitanium-basedlayers,whosethickness

increased from 80 nm to 110 nm, i.e. by 37.5%. Thickness of

zirco-nium-based layersincreasedfromanaveragevalue120nmbefore

an-nealing to135nmafterannealingduringoxidation(Figs.3,aand3,

c),i.e.by12.5%.Wherein,columnarcharacterofthegrain structure

isclearlyseeninzirconium-basedlayers,thatiswhytheybecomequite

fragile.Highdensitywasobservedintitanium-basedlayers,sowecan

Fig.2.Electronimagesofthecross-sectionsoftheTiN/ZrNsamples(general

viewofthecross-sectionandmagnifiedfragment)withamountoflayers533

(aandb),233(candd)and134(eandf)[15].

a b c

e f

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assumethattheselayersaresubjectedtocompressionandcompaction

because of oxidation, despite the relatively high increasing of the

thickness.Thus,compensatingtensilestraininthelayersoftheplane

of growth should be created in titanium-based layers, which

deter-mines their tendency to brittle fracture. Interphase boundary is the

mainareaofseparation.Itcanbeexplainedbydecreasingofthe

adhe-sivebondsduringformingofphaseswithdifferentcrystallattices

(cu-bicinzirconium-basedlayersandtetragonalintitanium-basedlayers)

duetooxidation.

Thus, unlike single-layer coating or coating with few layers, only

subsurfacelayersaresusceptibletophasechanges(inmultilayer

coat-ingwithamountofnanoscalelayersmorethan100)eveninthecaseof

severeoperationconditionsinactiveoxygenatmosphere,thereby

pre-ventingmainstructuralstateofinnercarrierlayersfromchanges.

Hardness is well known as a universal characteristic, which allow

rapidlyestimatingofmechanicalpropertiesofthecoatings[8].We

de-fined harness using microindentation method, and hardness was

around H42GPa for the first series of samples, H38GPa and

H36GPaforthesecondandthethirdseriesofsamplesaccordingly.

Fig.3.Imageofthecross-sectionandsurfaceofthecoatingofthethirdseries

before(a,b)andafter(c,d)annealing[15].

a b

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Thus, highhardnessistypical forall seriesofsampleswith different

bilayerthickness,sosuchcoatingsareperspectiveforusingas

protec-tiveones.Inthisconnection,itwasalsonecessarytoprovide

tribologi-calteststodeterminebasicmechanicalpropertiesuponcontactofthe

coatingwithcounterbody.

Resultsoftribologicaltestsofmultilayersampleswithdifferent

bi-layerthicknessunderroomtemperaturearepresentedinTable1.Itis

clearly seen from the table, that fabricated multilayer coatings have

highwearcoefficientpairedwithAl2O3 counterbody.

Figure 4shows the imagesof friction tracks forTiN/ZrN samples

withdifferentbilayerthickness.

One could see that friction tracks are characterized by absence of

burrs,cleavagesandradialcracks,indicatingthehighqualityand

ad-hesivestrengthofcoatings.

Allfabricatedcoatingshavegoodwearresistanceaveragevaluesof

the reduced wear were (1.3–1.5)105mm3N1mm1. Wear of the

counterbodywasrathersmall(1.9–2.2)106mm3N1mm1.Chipping,

cracking and peeling of coatings were not observed during friction

tests. We found good adhesion of the coatings to substrates. There

were noplasticdeformation duringtests;theobservedwear israther

typicalforsoftmetals[5,6].

TABLE1.TribologicalpropertiesofthemultilayerTiN/ZrNcoatings.

Series FrictionCoefficient

Wearfactor,

mm3N1mm1

No.

Startingmoment Duringtests Counterbody

(103) Samples (105) 1 0.59 1.0 1.9 1.3 2 0.62 1.2 2.0 1.5 3 0.62 1.1 2.2 1.4

Fig.4.FrictiontracksofthemultilayerTiN/ZrNcoatings:sampleofthefirst

(a),second(b),andthird(c)series[15].

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Figure 5 shows values of friction coefficients for multilayer

TiN/ZrNcoatings.Samplesfromthesecondserieshadthehighest

val-uesofthefriction coefficient(curve2inFig. 5).Therewasagreat

differencebetweenvaluesoffrictioncoefficientforthefirstandthird

series onthedistancefrom0to300meters.Frictioncoefficientfor

thesamplesfromthethirdseriessharplyincreasedtothevalue1.2and

stayedonthislevelalongthelengthoftestingarea.Valuesofforthe

samples of the first series monotonically increase along the friction

distanceandareequaltothevaluesforthethirdseries.Friction

coef-ficient of the samples from the second series sharply achieved the

value1.2andthenslowlymonotonicallyincreasedtothevalue1.3.

Friction coefficient  significantly depended on bilayer thickness

andtotalthicknessofthecoatings.Thelowestfrictioncoefficientwas

observed for the samples from the first series with 20 nm bilayer

thicknessand40nm.Thehighestvaluesoffrictioncoefficientwere

observedforthesamplesofthesecondserieswith70nm.

2.2. TiN/MoN Coatings

Multilayer nanostructured TiN/MoN coatings were fabricated using

vacuum arc evaporation of two cathodes at atmosphere of molecular

nitrogen.Forthisprocedure,weusedunit‘Bulat-6’,whichallows

ob-taining coatings both for scientific purposes and for industry.

Sub-strate holder rotatable at predetermined speed allows alternating

depositionoftitanium andmolybdenumnitrideslayersfromtwo

dia-metrically arranged evaporators [16]. Thus, it is possible to obtain

coatings with differentelemental and structural–phasecompositions

byadjustingthecurrentandvoltageofthesubstrateandcathodes,

ni-Fig.5.FrictioncoefficientsoftheTiN/ZrNcoatings:samplesofthefirst(1),

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trogenpressure inthechamberand otherparameters. Thedeposition

parametersarelocatedinTable2.

Becauseofdepositionofsamples,themultilayercoatingswith

vari-ousthicknessesoftheindividuallayersfrom8nmto100nmwere

pre-pared.Thetotalcoatingthicknesswasreached8.4microns.

Ifweconsiderthepresenceofonlytwoelements,MoandTi,weget

the data shown in Fig. 6. The alternation of TiN and MoN layers is

tracedquitewellinthisgraph.Unfortunately,surfaceroughnessand

relief layers of coating make it difficult to separate the layers more

clearlybySIMSanalysis.

Used parametersof depthprofile analysisforsampleNo.4leadto

sputteringrateof2.0nm/minforSiO2.Duetotheroughness,wedid

notmeasurethecraterdepthwithstylusprofilemeter.

InordertoestimatethesputteringrateofTiN/MoN,wecarriedout

SRIM simulations (see Table 3). Itis most likely that the sputtering

TABLE2.ParametersofmultilayerTiN/MoNdeposition.

Sample Period,nm tlayer,

s Idep, A Ibias, A Ubias, V f, kHz PN, Pa Expected Measured No.1 4 8 2 95–100 0.8 40 7 0.5 No.2 4 8 230 No.3 20 25 10 0.9 40 No.4 20 25 230 No.5 40 50 20 0.9 40 No.6 80 100 40 1.0 40

Fig.6. SIMS depth profile analysis of normalized secondary ion currents:

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rateofTiNandMoNwaslowerthanforSiO2 (lowerthan2.0nm/min).

HavingcomparedresultsofSRIMsimulationwithdataobtainedfrom

RBS spectra and SEM images ofthe cross-sectionof multilayer

sam-ples, weassumed that theaverage sputteringrate wasabout 1.1–1.3

nm/min(or0.02nm/s).

Thedepthprofileanalysiswascarriedoutforfivehours.Thecrater

of2.52.5mmwassputteredduringthis time.Takingthesputtering

rate of 0.02 nm/s as a basis, it is possible to determine that crater

depth was360–400nm. Figure7shows amodernizedpreviousgraph

withthecalculatedthicknessesofthefirstfewlayersfromthesurface

ofTiN/MoNcoating.

XRD patterns of TiN/MoN multilayer coatings with double layer

thicknessof25,50,100nmareshowninFig.3,a.Mainpeaksare

locatedaround236.5and242.5.Amoredetailedstudyofthe

spectral lines gave it possible to detect the asymmetric shape of the

peaks.

Peaksfoundat242.5canbedividedintotwocomponents,which

correspond to(200)f.c.c. TiNand(200)cubic -Mo2N planes(seeFig.

TABLE3.ParametersofSIMSanalysisandSRIMsimulations.

500nA,1.72keVArprimaryionbeamat45incidenceangle

Material SiO2 TiN MoN

S*(atoms/ion) 3.46 3.27 4.37

Ionrange*(nm) 4.4 3.3 2.7

Sputteringrate(nm/min) 2.0 — —

*

ValuesobtainedfromSRIMsimulation.

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8). The peak at 2  36.5 is attributed to (111)-oriented TiN and

-Mo2Ngrains.ThevolumefractionsofTiNand-Mo2Nphaseswere

ex-tracted from the XRD line fitting procedure of the (200) and (111)

peaksusingthe‘New_profile’software.

AlloftheidentifiedpeaksaremarkedinFig.9.Theseresultsshow

thatTiN/MoNcoatingsconsistofhighlytextured(200)cubiclayers.

X-rayanalysisshowstheformationofonlyonephasewiththef.c.c.

lattice(structuraltypeNaCl)incoatingwith8nmwhensubstrate

voltageis40V.Theformationoftwo-phasesystemofTiNwith

NaCl-Fig.9.Theseparationofdiffractionspectraintocomponentspeaksfromthe

two phases: curve 1—TiN (200), curve 2—γ-Mo2N (200).Left—sample #5,

right—sample#6[16].

Fig.8.Thediffractionpatterns (XRD),obtainedforcoatings withdifferent

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typef.c.c.latticeandhigh-temperature-Mo2N isobservedwhen

sub-stratevoltageisincreasingto230V.Thevolumetricratioof

TiN/-Mo2Nphasesis90/10accordingly.

ThepresenceofonlyonephaseatUb 40Vcanbeexplainedbythe

alleged epitaxialgrowthofthin layers, whichgrowthperiodis

deter-minedbystrongerbondsinTiNlayer.Increasingofsubstratepotential

to230Vleadstotheappearanceoftwo-phasebecauseof

intensifica-tion of ion bombardment, which contributes tothe grain refinement

andthebeginningofinterfacesformation.Formationofseparate

lay-ersMo2Nwithacubiclatticeandtheappearanceofinterphase

bounda-ries leadtogrowthof stressintheTiN phaseandincrease lattice

pe-riodinunstressedsection.Thestructureofthesecoatingsiscolumnar.

Figure10 showsTEM darkfieldimage (sampleNo.5),which

dem-onstratescolumnargrowthinmultilayer nitride.Itstartsfrom

inter-facebetweentextured(111)steelsubstrateandTiN/MoNlayers.

Coat-inghasa100nmthin interlayer.Accordingtothedata,itconsistsof

Ti,Mo,CandtracesofN.

Measurements ofhardness and elasticitymodulusof mostsamples

wereconductedtodeterminetheirmechanicalpropertiesandabilityto

durability. Amoredetailedstudyofhardnessandelasticityaregiven

forsamplesNo.5andNo.6inFig.11.Theyshowtypicalregularities

ofH(L)andE(L).Inmeasurements,theindenterhasreachedadepth

ofcoveragealmost3m.Penetrationoftheindenterwasalmostlinear

onthewholestageofloadapplication.

Results ofhardnessandelasticitymodulusmeasurementsfor

coat-ingswithdifferentlayerthicknessesareshowninTable4.Asseen,the

hardness ofthe samplestends todecrease withgrowth of period in

the coating. The maximum value of hardness, H47GPa, was

achieved at the minimum of the resulting coatings period 8nm.

Modulusofelasticitychangedwithchangingthethicknessofthe

dou-Fig. 10. Cross-section TEM bright field image for sample #5 with λ  50 nm [16].

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blelayer.ThemaximumvalueofE470GPawasobtainedatthesame

period8nm.

However,knowingthevalues ofhardnessor elasticityofthe

mate-rial is not sufficient to predict its protective capacity. The graph is

showninFig.12,whereareaisdivided intosections:1—section with

H/E0.1, which has no good plasticity of the material, 2—section

with goodplasticityofthematerial.Ascanbeseen,thesampleNo.2

with8nmisflaggedexactlyonthelineofplasticity.

The general trend of graph testifies to expediency of the studied

multilayer TiN/MoNcoatings at 8nm andatotherthicknesses of

layers.

Thesedataindicate agoodchance ofproducingTiN/MoNcoatings

withhighplasticityand,hence,thewearresistance.

Fig.11.Physical and mechanical properties of coatings: dependence of the

penetrationdepthontheload(a),thehardnessontheload(b),andtheYoung

modulusontheload(c)[16].

TABLE4.Hardnessandelasticitymodulusmeasurements.

λ,nm H,GPa Е,GPa H/E

8 47 470 0.1

25 31.8 456 0.07

50 26.5 418 0.063

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2.3.MoN/CrNCoatings

As seenfrom theside cutarea,a multilayercoating(Fig. 13) differs

with a sufficiently high planarity of layers and the absence of drop

phaseintheinteriorareasofthecoating.

The results of elemental analysis show that for small thicknesses,

when the layers are the thinnest and the mostsignificant portion of

timeduringprecipitationhavegoesforhigh-speedrotationofthe

sur-face andinteractionwith residualgases inthe workingchamber;the

depletion of thelayers ofthecoating with lightnitrogen atoms(Fig.

14)occurs.

Itshould benotedthatatthicknesses greaterthan50nm, the

con-tentofelementsinthecoatingcomestovaluesclosetoconstant,andat

a pressure of 3103 Torr makes a proportion close to 1 between the

metal atoms, and about 33% of nitrogen, which corresponds to the

stoichiometry of the phases Me2N (whereMe are metal atoms:Mo or

Cr).Atalesserpressureof7104Torrand2.4104Torr,thenitrogen

contentsdropssharplyto17.09and6.33at.%,respectively.

Analysisofdiffractionspectraofthecoatingsshowsthatinthecase

ofasmallvalueofthenegativebiaspotentialappliedtothesubstrate

during thedeposition(20V)inthespectra(Fig. 15)forall thelayer

thicknessesintherangeof5–200nm,phaseswithcubiclatticesbased

Fig.13.Theimageofmultilayercoating[17].

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onf.c.c.onewithaweaktexturewiththeaxis[311],typicalforgiven

regimesinmonolayerstatefor-Mo2Nphase[21].

Increasing thebias potential to150V and 300V leads to

forma-tionoftexturewith[100]axisinthelayersandtoincreaseinits

inten-sitywithincreasingthethicknessofthelayer.

OnasubstructurallevelatthelowestUb 20Vwiththeincreaseof

thickness of the layer,the growth of theaverage crystallite size and

nonmonotonic microstrain behaviour are observed: from high values

(1.5%)withalayerthicknessoflessthan20nm,throughaminimum

(1.1%)ath100nmto1.4%atlargethicknesses(seeFig.16,a).

WithanincreaseofUb upto150Vintheabsolutemagnitude,a

de-creaseinmicrostraininthelayersbothbyabsolutevalue(0.8–1.05%),

Fig.14.Dependenceofthecontentofatomsofnitrogen(1),molybdenum(2),

and chromium(3)onthethicknessof thelayersofthemultilayercomposite

materialMoN/CrN[17].

Fig.15.AreasofdiffractionspectraofthecoatingsobtainedatPN 310

3Torr

andUb 20Vatthicknessofthelayersof6nm(1),13nm(2),25nm(3),50

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andbyamplitudetakesplace.

The average crystallite size was varied nonmonotonically,

increas-ing proportionally tothe thicknessof the layertoa thickness of100

nm,andthenwasreducedby40%withafurtherincreaseofthelayer

thickness.

For highest Ub 300V used inpaper,the valueof microstrainin

thelayersdidnotexceed0.4%(seeFig.17,b),andthecrystallitesize

wasthesmallestoftheconsidered fortherespective thicknesses.The

observed decrease in microdeformation indicates on recombination

processesstimulatedbyhigherdensityofradiationdefectsalongwith

theincrease ofthemeanenergyofthefilm-forming particlesbecause

ofincreaseofU.Thedecreaseoftheaveragesizeofcrystallitescanbe

linked with the intense action of the defects, which increases the

growthofcentresofformation.

Fig.17.Changeintheaveragevaluesoftheamplitudeoftheacoustic

emis-sion(spectrum1,rightscale)andthecoefficientoffriction(spectrum2,left

scale) for the coatingsproduced at PN  310

3 Torr and

Ub  300 Vat the

thicknessofthelayers13nm(a)and400nm(b)[17].

Fig.16.Dependenceofsubstructuralcharacteristics(microdeformationε(1)

andthesizeofcrystallitesL(2))onthethicknessofthelayersinthecoatings

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The smallersize of thecrystallites andhence alargeraverage

spe-cificvolumeofthebordersdefinesahigherrelaxationcapacityforthe

randomlyformeddislocationdefects,whichdefinemicrostrain.

Thepresenceofasmallmicrodeformationandgrainsizeof

crystal-lites may be factors of increasing the adhesivestrength of the

mate-rial. Thus, for adhesive strength tests, the coatings obtained at

Ub 300Vwithdifferentlayerthicknesseswerechosen.

The conducted studies have shownthat for theentire rangeof the

usedthicknessesofthelayers,theuniformwearofthecoatingoverthe

entirerangeofappliedloadstakesplaceinthecompositecoatings;this

is manifestedin the homogeneity of acoustic emission spectrum (see

Fig.17,spectrum1).Alongwiththis,thecoefficientoffrictionforall

thethicknessesofthecoatingsissufficientlycloseandisintherange

of0.18–0.24(seeFig.17,spectrum2).

Alongwith this,thenatureofwear withthedecreaseof the

thick-ness ofthe layers becomes moreuniformthat is especiallyevident at

theareasofthefirstcriticalloadLC1 (seeFig.18).Thisindicatesa

de-crease in brittleness (ductility increase) of the layers along with

de-creasing h. Suchchanges may be relatedto adecrease ofthe average

crystallitesizeandmicrostraininthiscase(seeFig.18,b).Inthiscase,

thecriticalloadvalueisdeterminedbyadecisiveroleofcoating

hard-ness value, which with theincrease of thickness ofthe layers varied

from23to35GPa.

Alongwiththis,themostductileformofwear(seeFig.19)is

char-acteristic forthe coatings obtained atlow pressure and reducing the

nitrogen content, which determines strong covalent interatomic

bonds.Thehardnessofthesecoatingsis4–7GPa.

SummarydataonthecriticalloadabrasionofthecoatingsLC5 shown

inTable5forthecoatingsproducedatPN 310

3TorrandU

b 300V

show that the adhesion strength is great for all layer thicknesses of

Fig.18. Wear tracks at critical loads, LC1, for the coatings produced at

PN310

3TorrandU

b 300Vatlayerthicknesses:400nm(a),200nm(b),

13nm(c)[17].

(17)

such coatings, however, the increase in the thickness and hardening

allowsreachingthemaximumvalueLC5 185–187Natthethicknesses

ofthelayersof200–400nm.

Itshouldbenotedthatfortheinitialstageofwear,thecriticalload

forthecoatingsproducedatPN 310

3Torrisintherangeof13–18N,

whereas fortheplastic coatings obtained atlowpressure ofnitrogen

atmosphere,thevalueLC1 ismuchlower,andhasvaluesof6and4N,

respectively,forPN 710

4TorrandP

N 210

4Torr.

2.4.MetalNitride(GroupVI)Coatings

The studies of morphology of growth of the multilayer coatings

showed their sufficiently large homogeneity and planarity in all the

usedmodesforbothtypesof(TiZrNbTaHfMo)Nand(TiZrNbTaHf

W)Nsystems.Thedropinhomogeneity wasrevealedonthesurface,

which did not lead to a significant change in planarity and average

thickness(lessthan0.5%,seethedataonthicknessinFig.20).

The analysisofelement composition (EDXmethod)hasshownthat

theincreaseofnegativebiaspotential(Ub)ledtodepletionofthe

coat-Fig.19. Wear tracks at critical loads LC1 for the coatings produced at

PN 710

4Torr(a)and

PN210

4Torr(b)[19].

TABLE5. Critical load LC5 for the coatings obtained at PN 310

3Torr,

Ub 300Vanddifferentthicknessoflayers.

Thethicknessoflayersh,nm LC5,N

13 130.9 25 152.1 50 156.7 100 157.3 200 185.7 400 187.6 a b

(18)

ing by light atoms. The main reason of this is selective spraying of

light atoms during the spraying from thesurface of growth [12]. To

thelargestextentitaffectsnitrogenatoms,thecontentofwhichinthe

coatingwiththeincreaseofUb,inthestudiedragedecreasesmorethan

1.5times(seeFig.21).

Itisworthnoting,that atthesametimesaturationofthecoatings

(TiZrNbTaHf)N/WN with nitrogen is larger in comparison with

(TiZrNbTaHf)N/MoN by modulus. In addition, the spray character

(secondary selective spraying) is, apparently, a basis of change of

metalcomponentsoflayers.

In (TiZrNbTaHf)N/WN coatings, in thelayers with WN, with the

increase of Ub from 90 to280 V,the content ofheavy W(with

re-specttothecontentofmetalelementsinTiZrNbTaHflayer)increases

from33%to53%.

Fig.21.Dependenceofnitrogenatoms’contentinthecoatingonthevalueof

(Ub): 1—(TiZrNbTaHf)N/WN (PN  410 3 Тоrr), 2—(TiZrNbTaHf)N/MoN (PN 410 3Тоrr),3—(TiZrNbTaHf)N/MoN (PN 1.510 3Тоrr)[18].

Fig.20. SEM-pictures of the side surface of the ‘coating-substrate’ of the

multilayercoatings:а—(TiZrNbTaHf)N/WN (PN 410

3 Torr, Ub 90V), b—(TiZrNbTaHf)N/MoN (PN 410 3Torr, Ub 50V), c—(TiZrNbTaHf)N/ MoN(PN 1.510 3Torr,U b 50V)[18].

(19)

Forthecoatings(TiZrNbTaHf)N/MoN,theratiobetweentheatomic

contents ofMo and metal of thesecond layer(TiZrNbTaHf) with the

increaseofUb ispracticallynotobserved,remainingatalevelof41–42

at.% byitsmagnitude (Fig.22,a, c)shows typicalenergy-dispersive

spectraofthecoatingsandcalculatedatomiccomposition).

Annealingpracticallydoesnotchangetheratioofmetalcomponents

andleadstoasubstantialchangeincontentofnitrogenatomsof

nitro-genandimpurityoxygenatomsinthecoating.

If,atalowbiaspotential,thecontentbynitrogenatomsisdecreased

bytheabsolutevaluebyavalueofabout2%(Fig.22,a,b),thaninthe

case oflarge Ub  200V,thedecrease is moresignificantand is 5%

(Fig. 22, c, d). This can be linked with the additional formation of

pathsoflightdiffusionduringtheformationofsolidsolutionofHEA

atomsMo(W)aboundaryareabecauseofradiationstimulatedmixing.

The effect of bias potential and the pressure of working nitrogen

atmosphere also greatly influenced phasecomposition and structural

stateofthecoatings.

Fig.22.Energy-dispersionspectraandelementcontentscalculatedbyitinthe

coatings(TiZrNbTaHf)N/MoN(PN 410

3Тоrr)obtainedat

(a)Ub 100V

(beforetheannealing),(b)100V(aftertheannealing),(c)200V(beforethe

(20)

Areas of X-raydiffraction spectra ofthe coatings, obtained under

differenttechnologicalconditionsareshowninFig.23.

Forcomparison,thespectraofthecoatingbeforetheannealingand

afterthehightemperaturevacuumannealing areshowninonefigure

forcomparison.

The analysis of theobtained diffraction spectra shows that for all

depositionmodestheformationofphaseswithcubic(f.c.c.)lattice

oc-cursinbothlayersofmultilayercoatings.

In thelayers ofhigh-entropyalloy, itisa disorderedsolidsolution

(TiZrNbTaHf)NwiththecrystallatticeofNaCltype[22],inthelayers

oftheMo–Nsystem,itis-Mo2N,butinthelayersoftheW–Nsystem,

itis -W2N (PDF 25–1257). Thesimilarityof structuralstates inthe

layers basedonhighentropy alloyandnitridesofVIGrouptransition

metals(closerelationofpreferredorientationofcrystalliteslayers)

in-dicatestherelationshipbetweenstructureoflayersandtheirgrowth.

From theobtained spectra, itcan also beseen that inthe coatings

receivedduringthedepositionatlowUS,thepost-condensation

anneal-ingdoesnotleadtoasignificantchangeofthetypeofdiffraction

spec-Fig.23.AreasofX-raydiffractionspectraof(TiZrNbTaHf)N/MoNcoatings:

а—PN  1.510 3Тоrr, Ub 50V; b—PN  410 3 Тоrr,Ub 50 V; c—PN  4103 Тоrr, Ub  200V; and d—(TiZrNbTaHf)N/WN, PN  410 3 Тоrr,

(21)

tra(let’scompare1and2ontheFig.23,a,b,c,d).Duringtheincrease

of pressure ofthe working nitrogen atmosphere, theincrease of

tex-turing rate occurs (relative increase of intensity of reflexes). Thus,

under relatively low values of Us, in the (TiZrNbTaHf)N/MoN

coat-ings,suchatexturehasanaxis[311](Fig.23,a,b).

ApplyingalargenegativepotentialUb 200Vleadstoanincrease

inthedegreeof‘chaotization’ofthestructure(thestructureinherent

tosmallUs isnotmanifestedforhighUs),aswellastoincreaseof

dis-persionofcrystallineformationsinthelayersofthecoatings,whichis

manifestedthemostforthelayers-Mo2Nforwhichwiththeincrease

of Us theaveragesize of crystallitesdecreasesfrom54 nm to37nm.

For thecoatings (TiZrNbTaHf)N/WN,the formationoftexturewith

theaxis[111](Fig.23,d)occursatrelativelylowUb 90V,whichis

typicalforthepreferredminimizationofdeformationintheprocessof

growth,thedegreeofperfectionofwhichincreases(spectrum2inFig.

23,d)duringtheannealing.Atthesametime,reductionoftheperiod

of a lattice is observed in the annealed coatings in both layers: in

(TiZrNbTaHf)Nfrom0.443nmto0.439nm,andin-W2Nlayersfrom

0.425nmto0.421nm.

In thecoatings (TiZrNbTaHf)N/MoN,annealing leads toa

signifi-cant changeofthelattice parameteralmostexclusively inthenitride

layersofhighentropyalloy.ThelayersoftheMo–Nsystemare

char-acterized by the slight change of the grating period remaining at a

rangeof0.418–0.419nminthecoatingsobtainedatlowPN 1.510

3

Torrandatthelevelof0.425–0.424nmatPN 410

3Torr.An

excep-tionisthecoatingobtainedwithalargeUs 200V,forwhichevenat

PN 410

3 Torr grating period does not exceed 0.420 nm. One more

characteristic feature of this typeof thecoatings is formationof

ni-tridephasesofhigh-entropyalloywithasmallerperiod.Therefore,if

themainnitride phaseinthelayershasaperiodof0.44–0.47 nm,in

thecaseofhighvaluesofappliedUb 200V,theformationofanew

finecrystallinephaseoccurs.Itismanifestedonthediffraction

spec-trabeforetheannealingintheasymmetryofreflexes(seespectre1in

Fig. 23,c),whichincaseofannealedcoatingsisfoundintheformof

independentreflectionsfromlatticeplanesofthephasewithaperiod

of0.434–0.435nm(seespectre2inFig.23,c).

Presumably, this effect may be associated with the formation of

mixed solid solution phase based on the high entropy nitride of the

type(MoTiZrNbTaHf)N(includingMoatomsfromthesecondlayeras

acompoundelement)onadisorderedinterphaseborder.Formationof

suchalayermayleadtoadecreaseoffunctionalpropertiesofthe

coat-ingand,inparticular,mechanicalproperties.

The universal mechanical characteristics taking in account their

sufficienteaseofdefinition andgoodreproducibility are

(22)

The coatingsofthegreatestfirmnessaccording tothe

microinden-tationdataarethecoatingsdepositedatarelativelylowpotentialbias.

For the coatings (TiZrNbTaHf)N/WN, thehardness reaches 44 GPa,

andfor(TiZrNbTaHf)N/MoN,41GPa.

WhenincreasingUb,thehardnessslightlyfallsdownto39GPa,

ap-parently duetotheradiation-stimulatedmixing.Thepressure

reduc-tionalsoleadstoadecreaseinhardness.

In the case of the coatings obtained at low Ub, post-condensation

annealingleadstoincreaseofhardnessofsuchcoatingsbecauseof

or-dering at high temperatures in high entropy nitride layers. To the

greatest extent,itaffects thecoatings(TiZrNbTaHf)N/WNobtained

atUb 90V,wherethehardnessincreasesfrom44GPato59GPa.In

the case of coatings (TiZrNbTaHf)N/MoN, the largest increase in

hardnessisobservedatUb 50V:from40.5GPabeforetheannealing

to48.5GPaaftertheannealing.

ForthecoatingsobtainedatlargeUb [280,200]V,theannealing

isfollowedbyaslightnotonlyleadstoincreasedhardness,butalso

ac-companiedbyaslightdropofhardnessfrom39–40GPabeforethe

an-nealingto38–37GPaaftertheannealing.

Resultsofscratchtestsalsoshowthatthegreatestpressurepriortothe

LC1 4.950(a)

LC2 32.00(b)

LC3 47.40(c)

LC4 65.00(d)

LC5 117.9(e)

Fig.24.Viewof weartracksandtheresultingcriticalloadsforthecoatings

(TiZrNbTaHf)N/MoN(PN 410

3Тоrr,Ub 150V)[18].

b

a c

(23)

failureisinherenttothecoatingsobtainedatlowvaluesofUb and,forthe

coatings(TiZrNbTaHf)N/WN,reachesthevaluesLC5 117.9N(Fig.24),

andforthecoatings(TiZrNbTaHf)N/MoN,LC5 124.9N(Fig.25).

Theannealingat700Cofthecoatingsof(TiZrNbTaHf)N/MoN

sys-tem obtained at PN  310

3 Torr, U

b  150 V, for which the initial

hardness was relatively high (35 GPa), which then increased to41.5

GPa aftertheannealingledtoenhancedwearresistanceforallvalues

LC (see Figs. 24, 25), and the wear is inherent to abrasion, which is

manifestedby theabsenceoflarge amplitudepeaks (whichare

inher-enttobrittlefailure)onthecurve ofdependenceofacousticemission

onpressure(seeFig.26,a).

Slightlylargerbyitsmagnitudehardness(48.5GPa)ofthe

compos-itecoating(TiZrNbTaHf)N/MoNobtainedatPN 410

3TorrandU b

50VaftertheannealingincreasesthecriticalvaluesofLC1,LC2 and

LC3 (seeFigs.24,25),but,atthesametime,thevaluesofLC4 andLC5

are decreased, i.e. critical stresses responsible for the formation of

multiplecracksandwearofmaterialofthecoatingtakesplace.

The effect of formation of multiple cracks and brittle fracture of

high hardness coating is observed even better for a steel-based

sub-strate with a higher plasticity for the system (TiZrNbTaHf)N/WN,

obtained atPN 410

3Torr, U

b  90V,the hardnessofwhich is

in-creasedasaresultofannealingupto59GPa.Inthiscase,theincrease

LC1 6.060(a)

LC2 31.30(b)

LC3 45.50(c)

LC4 81.60(d)

LC5 124.9(e)

Fig.25.Viewof weartracksandtheresultingcriticalloadsforthecoatings

(TiZrNbTaHf)N/MoN(TiZrNbTaHf)N/WN(PN 410

3Тоrr,Ub 90V)[18].

a b c

(24)

of value of the critical load LC1 (see Figs. 24, 25) takesplace only, a

slightrelativedecreaseofLC2 andLC3 andastrongdecreaseofLC4 and

LC5.Itissignificant thatinthisarea,strong peakstypical tothe

for-mationofmacroareaswithbrittlefailureappearonthedependenceof

theacousticemissionsignalontheload(Fig.26,b).

Thus,theachievementofultrahighhardnessin acaseofrelatively

ductilesubstratemaynotleadtoanincreaseinadhesivestrengthdue

to the brittle fracture of the coating during the wear process in the

boundaryareastotheductilebasematerial.

To determine tribological characteristics, the testing ‘ball–disc’

schemewasused,forwhichtheballswithdiameterof6.0mmmadeof

sinteredcertifiedmaterials,Al2O3 andsteelAc100Cr6,wereused.

Visually, friction tracks (see Fig. 27) are characterized by the

ab-Fig.27.Imagesoffrictiontracksduringthetests(bythe‘ball–disc’scheme

withacounterbody(ball)madeofAl2O3 andsteelAc100Cr6)ofthemultilayer

coating(TiZrNbTaHf)N/MoN(PN 410

3Torr,Ub 50V).

a—withdetailed

frictiontracksbythetypeofthecounterbody,b—withthedetermined

aver-agesizesoffrictiontracksfordifferentcounterbodies[18].

Fig.26.Changesinaveragevaluesofthecoefficientoffriction(spectrum1,

leftscale)andintheamplitudeofacousticemission(spectrum2,rightscale)

forthecoatings:а—(TiZrNbTaHf)N/MoN(PN 410

3Тоrr,

Us 150V)and

b—(TiZrNbTaHf)N/WN(PN 410

3Тоrr,

(25)

senceofbarbs,chips,andradialcracks,whichindicatesthehigh

qual-ityofthecoatinganditsadhesivestrength.

Theaveragewidthofthefrictiontrackinacaseofthecounterbody

madeofAl2O3 hasavalueof654.88m(seeFig.27,b),andinthecase

of steelcounterbody,thetrack hasdifferent thicknessand is

charac-terizedbyanon-uniformwearpattern.

The reason for this heterogeneous nature is sticking of relatively

soft and ductile metal of the counterbody to the coating, which

in-creasestheactualimpactarea,andfurtherfrictionoccursinapairof

worn-outmetalandthemetalofcounterbody.

Loweringofthefixed frictioncoefficientandincrease ofwearofa

steelballisalsolinkedwiththis(Table6).

DuringfrictionwithacounterbodymadeofAl2O3,theuniform

abra-sive wear of the friction pair with the removal of wear products and

theiraccumulationattheedgesofthegroovewasobserved(Fig.27,a).

In this case, the amount of transferred material depends on the

strengthofadhesivebond,whichdependsontheelectronicstructureof

thecounterbodybasedonAl2O3 andmultilayercoating,anddetermines

theability toform solid solutions,intermetalliccompounds witheach

other,andoxidesstableathightemperatures.Thisisrelatedtothehigh

valuesofthecoefficientoffrictionduringthetestswithAl2O3

counter-body,whichincludehighentropynitride(TiZrHfVNbTa)N[14].

At the sametime,the coatings possess good wear resistance: wear

value for both types of counterbodies is within the limits (0.39–

2.12)105 mm3N1mm1. Wear of the counterbody of Al

2O3 is also

sufficiently small—2.25106 mm3N1mm1 (Table 7), in contrast

withasteelcounterbodyAc100Cr6,forwhichweardiffersbyoneorder

andhasavalueof2.59104mm3N1mm1.

Thus,auniversalresistanceofmultilayercoatingsbasedonnitrides

of high entropyalloystovarious typesof counterbodies,different in

hardness andtoughness opens good perspectivestousesuchcoatings

as protective coatingsatcomplexexposuresinconditions ofabrasive

wear[1,15].

TABLE6.The criticalloadLC for compositemultilayercoatings,beforeand

afterone-hourannealingat700C.

Coatingtype Ub,

V Annealing, C LC,N 1 2 3 4 5 (TiZrNbTaHf)N/MoN 150 – 4.52 31.8 48.2 65.8 73.1 700 5.07 40.2 51.4 56.0 82.2 (TiZrNbTaHf)N/MoN 50 – 6.06 31.3 45.5 81.6 124.9 700 7.33 42.8 58.4 63.8 81.7 (TiZrNbTaHf)N/WN 90 – 4.95 32.0 47.4 65.0 117.9 700 8.31 31.3 39.7 41.3 62.7

(26)

3.CONCLUSIONS

Inthisarticle,theinfluenceofthesizeofbilayersdeposition

parame-ters, thestressand microstructureofthe changeofphysicaland

me-chanical properties of the nanostructure (nanocomposite) coatings

produced by vacuum arc deposition from the cathode is written and

shown.Presentedresultsareanintermediatestageofourworkinthe

field of development and study of nanoscale multilayer coatings

formedfrom refractorymetals.Thepresent resultssuggest the

pros-pect of obtaining and studying such nanoscale nitride multilayer

nanostructuredcoatingsinthenearfuture.

This work was done under the aegis of the State budget programs

0113U000137c‘Physicalprinciplesofplasmatechnologiesforcomplex

treatment ofmulticomponent materials and coatings’, 0115 U000682

‘Thedevelopmentofmaterialsciencebasisofastructuralengineering,

vacuumplasmasuperhardcoatingsinordertoreachtherequired

func-tional properties’ and in collaboration with the Institute ‘P’-Prime,

University Poitiers France, National Institute for Materials Science,

TsukubaIbasaki,Japan,andNBMC,AdamMickiewiczUniversity,

Po-land. Theauthors are gratefulto co-authors ofthe originalwork (the

basisofwhich composesthis review),namely,O.V.Bondar,B.A.

Pos-tolnyi, V.A. Stolbovoy, V.Yu. Novikov, S.V. Lytovchenko, O.V.

Sobol’,U.S.Nyemchenko,I.N.Toryanik,D.A.Kolesnikovandothers.

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