1
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
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.IncreasingoftheperiodledtoincreasingoftheTiN
lay-ersspecificcontribution.Itisseenfromthechangesintheintensityof
thepeaksofTiNandZrNphases(seeFig.1).Increasingofthe
deposi-tion timeand, as aresult, bilayerthickness as wellas total period of
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 firstlineontheangle225.36
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
annealingunderthetemperatureof700Cforonehour:a—fromthesurface
withoutpolishing,b—afterpolishingoftheoxidizedsurfaceonthedepthof5
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
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 H42GPa for the first series of samples, H38GPa and
H36GPaforthesecondandthethirdseriesofsamplesaccordingly.
Fig.3.Imageofthecross-sectionandsurfaceofthecoatingofthethirdseries
before(a,b)andafter(c,d)annealing[15].
a b
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)105mm3N1mm1. Wear of the
counterbodywasrathersmall(1.9–2.2)106mm3N1mm1.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,
mm3N1mm1
No.
Startingmoment Duringtests Counterbody
(103) Samples (105) 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].
Figure 5 shows values of friction coefficients for multilayer
TiN/ZrNcoatings.Samplesfromthesecondserieshadthehighest
val-uesofthefriction coefficient(curve2inFig. 5).Therewasagreat
differencebetweenvaluesoffrictioncoefficientforthefirstandthird
series onthedistancefrom0to300meters.Frictioncoefficientfor
thesamplesfromthethirdseriessharplyincreasedtothevalue1.2and
stayedonthislevelalongthelengthoftestingarea.Valuesofforthe
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
thicknessand40nm.Thehighestvaluesoffrictioncoefficientwere
observedforthesamplesofthesecondserieswith70nm.
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),
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:
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(or0.02nm/s).
Thedepthprofileanalysiswascarriedoutforfivehours.Thecrater
of2.52.5mmwassputteredduringthis 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
thicknessof25,50,100nmareshowninFig.3,a.Mainpeaksare
locatedaround236.5and242.5.Amoredetailedstudyofthe
spectral lines gave it possible to detect the asymmetric shape of the
peaks.
Peaksfoundat242.5canbedividedintotwocomponents,which
correspond to(200)f.c.c. TiNand(200)cubic -Mo2N planes(seeFig.
TABLE3.ParametersofSIMSanalysisandSRIMsimulations.
500nA,1.72keVArprimaryionbeamat45incidenceangle
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.
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)incoatingwith8nmwhensubstrate
voltageis40V.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
typef.c.c.latticeandhigh-temperature-Mo2N isobservedwhen
sub-stratevoltageisincreasingto230V.Thevolumetricratioof
TiN/-Mo2Nphasesis90/10accordingly.
ThepresenceofonlyonephaseatUb 40Vcanbeexplainedbythe
alleged epitaxialgrowthofthin layers, whichgrowthperiodis
deter-minedbystrongerbondsinTiNlayer.Increasingofsubstratepotential
to230Vleadstotheappearanceoftwo-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
ofcoveragealmost3m.Penetrationoftheindenterwasalmostlinear
onthewholestageofloadapplication.
Results ofhardnessandelasticitymodulusmeasurementsfor
coat-ingswithdifferentlayerthicknessesareshowninTable4.Asseen,the
hardness ofthe samplestends todecrease withgrowth of period in
the coating. The maximum value of hardness, H47GPa, 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].
blelayer.ThemaximumvalueofE470GPawasobtainedatthesame
period8nm.
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/E0.1, which has no good plasticity of the material, 2—section
with goodplasticityofthematerial.Ascanbeseen,thesampleNo.2
with8nmisflaggedexactlyonthelineofplasticity.
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
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 3103 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).Atalesserpressureof7104Torrand2.4104Torr,thenitrogen
contentsdropssharplyto17.09and6.33at.%,respectively.
Analysisofdiffractionspectraofthecoatingsshowsthatinthecase
ofasmallvalueofthenegativebiaspotentialappliedtothesubstrate
during thedeposition(20V)inthespectra(Fig. 15)forall thelayer
thicknessesintherangeof5–200nm,phaseswithcubiclatticesbased
Fig.13.Theimageofmultilayercoating[17].
onf.c.c.onewithaweaktexturewiththeaxis[311],typicalforgiven
regimesinmonolayerstatefor-Mo2Nphase[21].
Increasing thebias potential to150V 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%)ath100nmto1.4%atlargethicknesses(seeFig.16,a).
WithanincreaseofUb upto150Vintheabsolutemagnitude,a
de-creaseinmicrostraininthelayersbothbyabsolutevalue(0.8–1.05%),
Fig.14.Dependenceofthecontentofatomsofnitrogen(1),molybdenum(2),
and chromium(3)onthethicknessof thelayersofthemultilayercomposite
materialMoN/CrN[17].
Fig.15.AreasofdiffractionspectraofthecoatingsobtainedatPN 310
3Torr
andUb 20Vatthicknessofthelayersof6nm(1),13nm(2),25nm(3),50
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 310
3 Torr and
Ub 300 Vat the
thicknessofthelayers13nm(a)and400nm(b)[17].
Fig.16.Dependenceofsubstructuralcharacteristics(microdeformationε(1)
andthesizeofcrystallitesL(2))onthethicknessofthelayersinthecoatings
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 310
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
PN310
3TorrandU
b 300Vatlayerthicknesses:400nm(a),200nm(b),
13nm(c)[17].
such coatings, however, the increase in the thickness and hardening
allowsreachingthemaximumvalueLC5 185–187Natthethicknesses
ofthelayersof200–400nm.
Itshouldbenotedthatfortheinitialstageofwear,thecriticalload
forthecoatingsproducedatPN 310
3Torrisintherangeof13–18N,
whereas fortheplastic coatings obtained atlowpressure ofnitrogen
atmosphere,thevalueLC1 ismuchlower,andhasvaluesof6and4N,
respectively,forPN 710
4TorrandP
N 210
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(TiZrNbTaHfMo)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 710
4Torr(a)and
PN210
4Torr(b)[19].
TABLE5. Critical load LC5 for the coatings obtained at PN 310
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
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 to280 V,the content ofheavy W(with
re-specttothecontentofmetalelementsinTiZrNbTaHflayer)increases
from33%to53%.
Fig.21.Dependenceofnitrogenatoms’contentinthecoatingonthevalueof
(Ub): 1—(TiZrNbTaHf)N/WN (PN 410 3 Тоrr), 2—(TiZrNbTaHf)N/MoN (PN 410 3Тоrr),3—(TiZrNbTaHf)N/MoN (PN 1.510 3Тоrr)[18].
Fig.20. SEM-pictures of the side surface of the ‘coating-substrate’ of the
multilayercoatings:а—(TiZrNbTaHf)N/WN (PN 410
3 Torr, Ub 90V), b—(TiZrNbTaHf)N/MoN (PN 410 3Torr, Ub 50V), c—(TiZrNbTaHf)N/ MoN(PN 1.510 3Torr,U b 50V)[18].
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 410
3Тоrr)obtainedat
(a)Ub 100V
(beforetheannealing),(b)100V(aftertheannealing),(c)200V(beforethe
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 1.510 3Тоrr, Ub 50V; b—PN 410 3 Тоrr,Ub 50 V; c—PN 4103 Тоrr, Ub 200V; and d—(TiZrNbTaHf)N/WN, PN 410 3 Тоrr,
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 1.510
3
Torrandatthelevelof0.425–0.424nmatPN 410
3Torr.An
excep-tionisthecoatingobtainedwithalargeUs 200V,forwhichevenat
PN 410
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
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 410
3Тоrr,Ub 150V)[18].
b
a c
failureisinherenttothecoatingsobtainedatlowvaluesofUb and,forthe
coatings(TiZrNbTaHf)N/WN,reachesthevaluesLC5 117.9N(Fig.24),
andforthecoatings(TiZrNbTaHf)N/MoN,LC5 124.9N(Fig.25).
Theannealingat700Cofthecoatingsof(TiZrNbTaHf)N/MoN
sys-tem obtained at PN 310
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 410
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 410
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 410
3Тоrr,Ub 90V)[18].
a b c
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 410
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 410
3Тоrr,
Us 150V)and
b—(TiZrNbTaHf)N/WN(PN 410
3Тоrr,
senceofbarbs,chips,andradialcracks,whichindicatesthehigh
qual-ityofthecoatinganditsadhesivestrength.
Theaveragewidthofthefrictiontrackinacaseofthecounterbody
madeofAl2O3 hasavalueof654.88m(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)105 mm3N1mm1. Wear of the counterbody of Al
2O3 is also
sufficiently small—2.25106 mm3N1mm1 (Table 7), in contrast
withasteelcounterbodyAc100Cr6,forwhichweardiffersbyoneorder
andhasavalueof2.59104mm3N1mm1.
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-hourannealingat700C.
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
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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