Surface & Coatings Technology 205 (2010) 2197–2208

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Damage tolerant functionally graded WC–Co/Stainless Steel HVOF coatings

Alfredo Valarezo a,⁎, Giovanni Bolelli b, Wanhuk B. Choi a, Sanjay Sampath a, Valeria Cannillo b,Luca Lusvarghi b, Roberto Rosa b

a Center for Thermal Spray Research, Materials Science and Engineering Department, 130 Heavy Engineering Bldg., Stony Brook University, Stony Brook, NY, 11794-2275, USAb Department of Materials and Environmental Engineering, University of Modena and Reggio Emilia, Via Vignolese 905, 41125 Modena (MO), Italy

⁎ Corresponding author.E-mail address: alfredo.valarezo@gmail.com (A. Vala

0257-8972/$ – see front matter © 2010 Elsevier B.V. Aldoi:10.1016/j.surfcoat.2010.08.148

a b s t r a c t

a r t i c l e i n f o

Article history:Received 23 April 2010Accepted in revised form 31 August 2010Available online 8 September 2010

Keywords:Damage tolerant coatingsFunctionally graded materialsHigh Velocity Oxygen-Fuel (HVOF) sprayingWear resistanceImpact testingResidual stress

In this paper, effective damage tolerance of a functionally graded coating (FGC) deposited by high velocityoxygen fuel (HVOF) spraying is observed. The thick FGC (≈1.2 mm) consists of 6 layers with a stepwisechange in composition from 100 vol.% ductile AISI316 stainless steel (bottom layer) to 100 vol.% hard WC–12Co (top layer) deposited onto an AISI316 stainless steel substrate. Damage tolerance is observed via 1) anincrease in compliance with depth, and 2) an increase in fracture resistance by containment, arrest anddeflection of cracks. A smooth gradation in the composition and hardness through the coating thickness isfound by scanning electron microscopy and depth-sensing microindentation, respectively. The in-situcurvature measurement technique reveals that during the deposition of the FGC, compressive stresses exist inthe lower, metallic layers owing to peening effect of successive impact, and these gradually evolve to hightensile, in the top layers. Tensile stresses appear to be due to quenching alone; thermal stresses are lowbecause of the gradation. All of this is beneficial for the deposition of a thick coating.The FGC structure shows the ability to reduce cracking with increased compliance in the top layer duringstatic and dynamic normal contact loading, while retaining excellent sliding wear resistance (ball-on-disktests). Results are discussed in comparison to the behavior and properties of coatings of similar individualcompositions and thicknesses, as well as a thick monolithic WC–12Co sprayed coating. Further improvementsin the processing are proposed to enhance the adhesion strength and avoid coating delamination under highload contact-fatigue conditions.

rezo).

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© 2010 Elsevier B.V. All rights reserved.

1. Introduction

Functionally graded materials (FGMs) are a special category ofcomposite materials, the composition and microstructure of whichare not spatially homogeneous but vary according to a predeterminedprofile [1], designed in order to improve certain characteristics(usually mechanical), over monolithic materials [2–5]. FGM manu-facturing is non-trivial, but is well-suited for thermal spray (TS)deposition methods, largely due to the ability of TS to process dis-similar materials concurrently, in sequential layers of differentcomposition [6]. Layer-by-layer compositional gradients can indeedbe easily manufactured by simultaneously feeding different powdersin the hot flame of the spraying torch [7–9], and gradually varying theratio of the two or more powders [10–16]. This gradation can beobtained by the simultaneous control of the feed rate in variouspowder feeders, by strategic injection of the various powders in theplume, or by previously blending the powder materials in differentproportions, and periodically replacing the feedstock (as it is pre-sented here).

The application of the FGM concept to TS has been extensivelystudied for thermal barrier coating (TBC) applications [14–16], in anattempt to suppress the generation of cracks along the bond coat(MCrAlY) / top coat (Y-PSZ) interface. As a result of those investiga-tions, very complex architectures have been proposed to address thenumerous synergistic factors leading to such cracking (including themismatch in thermo-mechanical properties, but also the oxidation ofthe metallic bond coat material) [14]. Few studies have, up to now,addressed the subject of thermally-sprayed FGMs for wear andcorrosion protection [17–22]. As traditional single-layer coatings forwear protection generally exhibit much higher hardness and elasticmodulus than the substrate (the typical combination being a cermet,like WC–Co or WC–CoCr, deposited onto steel or light alloy sub-strates), the stress distribution produced under localised contactconditions concentrates inside the coating and along the coating/substrate interface, where cracking and delamination become likely[23–25]. Anti-corrosion coatings should be as thick as possible, inorder to minimise or eliminate the interconnected porosity [26,27];however, during the deposition of thick single-layer coatings, exces-sively large residual stress is often built up [28,29] producing verticalmacrocracks or premature delamination. These issues have led to there-examination of thermal sprayed FGMs as potential solutions. Inaddition, advances in process science, such as the in situmeasurement

2198 A. Valarezo et al. / Surface & Coatings Technology 205 (2010) 2197–2208

of stress states, have enabled new insights into materials andprocesses providing new opportunities of layered and gradedmaterials design.

In this paper, the approach of damage tolerance, which has beenextensively used in the design of brittle components that withstandload without failure despite the presence of defects (e.g. cracks), isapplied to TS coatings for wear resistance. Given that TS coating layersare built by individual molten particles, the presence of defects (e.g.intersplat interfaces and pores) and, potentially, of cracksmakes themsingular systems that can prematurely fail at the weakest regions viamacro- or micro-cracking under applications of contact loads in wearsituations. Thus, damage tolerance applied to TS coatings, in thisresearch, considers the combination of process and microstructuraldesign concepts by grading structures to achieve enhanced perfor-mance. Here, i) crack inhibition is proposed to be achieved in hardmetal-based cermet coatings (e.g. usually embrittled WC–Co, CrC–NiCr due to decarburization) by exploiting the process conditions inorder to induce compressive residual stress by peening in the bottomlayers, and ii) crack arrest and deflection is achieved by addition of atough metallic phase.

2. Experimental procedures

2.1. Coating deposition

The WC–12 wt.%Co material (hereafter WC–Co) used in theseexperiments, with particle size −62+42 μm, was an experimentalbatch of fine carbide size produced by Osram Sylvania [30] withdesignation SX-432. The stainless steel 316 powder (hereafter SS316)was produced by ANVAL, particle size −46+24 μm.

The WC–Co/SS316 functionally graded coating (FGC) was depos-ited onto grit-blasted AISI316 stainless steel plates of228.6×25.4×1.6 mm size, using a Sulzer-Metco DJ 2700 torch withpropylene as fuel. The WC–Co and SS316 powders were mixed invarious proportions to produce a FGC. Starting from a 100 vol.% steellayer, the amount ofWC–Cowas increased by 20 vol.% until a top layerconsisting of pure WC–Co was deposited, thus obtaining a 6-layersystem of about 0.2 mm per layer. The parameters employed for thedeposition of each layer are listed in Table 1. For comparativepurposes, monolithic single-layer coatings corresponding to eachlayer of the FGC (and having identical thickness) were also depositedonto the stainless steel plates, using the same deposition parameterslisted in Table 1. They will be hereafter referred to as “single-layercoatings”. Additionally, a thick (≈0.6 mm) pure WC–Co layer wasproduced, in order to compare the thick FGC to a homogeneouscoating of similar thickness. This layer could not be grown thickerthan ≈0.6 mm without debonding and severely distorting thesubstrate.

2.2. Microstructural and micromechanical characterization

X-ray diffractometry (X'Pert PRO, Panalytical, Almelo, TheNetherlands) was performed on the top layer of the FGC and on thesingle-layer coatings, using Cu–Kα radiation in the 30°b2θb90°range.

Table 1Parameters employed in the HVOF-deposition of the various layers of the FGC.

Nominal powder blendcomposition (vol.%)

Layer thickness(mm)

Numberof passes

Fuel(l/min)

Oxygen(l/min)

100% SS316L 0.188 7 67 25180%SS316L-20%WC–Co 0.184 8 78 25160%SS316L-40%WC–Co 0.265 14 78 26640%SS316L-60%WC–Co 0.164 9 78 26620%SS316L-80%WC–Co 0.189 13 78 280100% WC–Co 0.313 20 78 280

Cross-sectional samples of all of the coatings were prepared bycutting, hot-mounting in phenolic resin, grinding with SiC papers (upto 2500 mesh) and polishing with diamond slurry (up to 0.5 μm). Thepolished sections were observed by scanning electron microscopy(SEM, Quanta-200 and XL30, FEI, Eindhoven, The Netherlands) andthe composition of the various layers was assessed by image analysis(NIH ImageJ version 1.37): at least 5 micrographs at 400× magnifi-cation were employed for each layer.

The residual stress build-up during the deposition of the FGC andof each single-layer coating was monitored by the in situ coatingproperty (ICP) sensor [31–33]. ICP uses a laser displacement sensor tomeasure continuously the stress-induced curvature of a thin platesubjected to layer-by-layer deposition: quenching, peening, andthermal stress arise during spraying. A simultaneous measurementof temperature is carried out via multiple thermocouples in contact tothe back side of the sample.

The microhardness of each layer in the FGC was measured bydepth-sensing Vickers microindentation on the polished cross-section(15 indentations for each layer, 3 N load, 2.4 N/min loading andunloading rate, 10 s loading time), according to the Oliver–Pharrprocedure [34]. The results were compared to the hardness of thecorresponding single-layer coatings, in order to verify whether theinclusion in the FGC structure could produce significantmodifications.

The coating elastic modulus for the SS316 and WC–Co single-layers was obtained by instrumented microindentation (Micro-Materials Limited, Wrexham Technology Park, Wrexham, UK, WC–Co Berkovich indentor at 5 [N] load). The modulus for the othercompositions was obtained by rule of mixtures.

2.3. Sliding wear testing

The dry sliding wear resistance of the FGCwas assessed by ball-on-disk testing, using a pin-on-disk tribometer (CSM Instruments,Peseux, Switzerland) with sintered WC–6%Co spherical pins (diam-eter: 3 mm) as counterparts. Two different experimental configura-tions were employed: 10 N normal load, 0.20 m/s relative slidingspeed, 5000 m overall sliding distance, 5 mm wear track radius(hereafter referred to as “test-1”); 10 N normal load, 0.40 m/s relativesliding speed, 10000 m overall sliding distance, 8 mm wear trackradius (hereafter referred to as “test-2”). The friction coefficient wasmonitored during the test by a load cell attached to the pin-holdingarm, the wear rate of the sample was assessed by measuring thevolume of the wear scar using an optical confocal profilometer(Conscan Profilometer, CSM Instruments), and the wear scars wereobserved by SEM. The test was also performed on the 100 vol.% WC–Co single-layer coating, to study whether the underlying layers in theFGC can modify its intrinsic sliding wear behaviour. Also, the 80 vol.%WC–Co single-layer coating was tested to ascertain whether theaddition of a low amount of steel to the hard but brittle WC–Comaterial would improve or impair its wear performance.

It should be noted that, under these test conditions, the contactstress distribution mainly involves near-surface regions of thecoatings. An approximate computation of the Hertzian stressdistribution that occurs in the contact of a 3 mm-diameter sphereon a flat surface having an elastic modulus of about 300 GPa and

Air(l/min)

Carrier gas(N2) (l/min)

Feed rate(g/min)

Average particletemperature (°C)

Average particlevelocity(m/s)

280 12 35 1926 702306 12 39 1917 735350 12 39 1943 680350 12 40 1764 667350 12 38 1672 613350 12 41 1636 562

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Poisson's ratio of 0.25 (typical properties of a HVOF-sprayed WC-Cocoating [35,36]) indicates that the maximum sub-surface shear stress(in both test conditions) is about 850 MPa and it is located about20 μm below the surface, i.e. the dominant contact-induced stressdistribution occurs within the top most layer of the FGC.

2.4. High-load indentation and ball drop cyclic impact testing

To observe the effect of the deep compliant layers on themechanical behaviour of the top harder layers, high load indentationand cyclic impact tests were performed on the FGC as to probe itsload-carrying capability and damage tolerance. The results werecompared to those obtained on the 100 vol.% WC–Co single-layer andon the 0.6 mm-thick WC–Co coating as a reference.

The indentation tests were carried out using a Mitutoyo AVK-C2hardness tester with a load of 50 kg applied for 15 s, monotonicallyand cyclically (1, 10, and 100 times) in the same location. A WC–Cospherical indenter of 1/8″ (3.125 mm) diameter was used in thesetests. The bonded-specimen technique was used to prepare thesamples. A detailed description of the sample preparation is describedelsewhere, thus a brief description is mentioned here [37]. In thistechnique, two 7 mm×2.5 mm specimens were cut, and the cross-sections were polished to 0.05 μm with Al2O3 solution. The two sliceswere then put together with a mechanical press and the top surfacewas polished. The assembly was indented from top at the interface.Afterwards, the faces were separated, and optical micrographs weretaken from the cross-section.

The ball drop cyclic impact test was carried out by dropping anX200Cr13 steel ball (diameter: 39 mm), attached to an overall weightof 12 N, at a frequency of 45 impacts/min on the sample's surface. Thedrop height was 95 mm, and the overall number of impacts was set at1000. The samples were inspected by optical and scanning electronmicroscopy after the test. To observe the cross-section of the coatingsafter the cyclic impact test, the sample was mounted in resin,dissectedwith ametallographic cuttingmachine and then ground andpolished with diamond slurries, as described above.

3. Results and discussion

3.1. Coating microstructures

Fig. 1 (a–g) shows XRD patterns of the specimens. Significantoxidation and phase transformations occurred during HVOF proces-

Fig. 1. XRD patterns of the single layers and of the FGC coating. Legend: a=100%stainless steel single layer; b=20% WC–Co single layer; c=40% WC–Co single layer;d=60%WC–Co single layer; e=80%WC–Co single layer; f=100%WC–Co single layer;g=FGC top layer. □=WC; ○=W2C; ■=Co3W3C; ●=Co6W6C; △=W; ▲=Fe3O4;◊=γ-Fe (austenite); ♦=δ-Fe (ferrite-δ).

sing of the stainless steel powder. The X-ray diffraction pattern of theAISI316 single-layer coating (Fig. 1, pattern a) shows presence ofoxides (mainly Fe,Cr -spinel oxide) and δ-ferrite phase, in addition tothe peaks related to f.c.c. γ-Fe structure (austenite). The remarkablesensitivity of austenitic stainless steel powders to phase transforma-tions during spraying has been reported by other authors [38,39]. Thechemistry of the WC–Co powder was also modified during spraying,through the WC decarburization reaction with the formation of theW2C and W phases, as well as the dissolution of WC in the metalmatrix resulting in the formation of mixed carbide phases, such asCo3W3C and Co6W6C [40]. This is reflected by the XRD patterns of allthe layers that contain WC–Co (Fig. 1, patterns b-g).

Both the oxidation of stainless steel and the decomposition ofWC–Co can be observed by SEM micrographs of both the single-layercoatings and the FGC. In the stainless steel layers, oxides mainlyappear as dark stringers along the boundaries of lamellae or aredispersed within the splats (Fig. 2A). Image analysis reveal that theoxide content is about 10%, both in the single-layer coating and in thebottom layer of the FGC (Fig. 3). No significant microstructuraldifferences are observed between these two layers as they weresprayed under identical deposition conditions. They also containnumerous large unmelted particles, which indicates that the smallerparticles were oxidized and decomposed predominantly. In the WC–Co layer, the WC decomposition is revealed by the presence ofbrighter lamellae in the backscattered electron micrographs (Fig. 2F):the brighter contrast is produced by the overall increase in theaverage atomic weight due to the dissolution of tungsten in the Cometal matrix and to the loss of C. Once again, the WC–Co single-layercoating and the top layer of the FGC (Fig. 3) exhibited no appreciabledifferences in the amount and nature of the decomposed regions.

The index of carbide retention (I) in a WC-based coating is definedin [41] as:

I =IWC

IWC + IW2C+ IW

ð1Þ

where: IWC=intensity of the WC peak at 2θ=35.6°; IW2C=intensityof the W2C peak at 2θ=39.6°; IW=intensity of the W peak at2θ=40.6°. This ratio is found to be equal to 0.86 in both cases. Thisvalue is consistent with the typical carbide retention in standardmicrostructured WC–Co-based cermet coatings by HVOF [41],although values as high as 0.91 could be achieved in some studies[42]. The WC decomposition is favoured by the very fine carbide sizeof the feedstock (Fig. 2F, see inset), which increases the specificsurface area of the carbide phase and favours its reaction with themetal matrix and/or the oxygen [30].

TheWC–Co/stainless steel composite layers (Fig. 2B–E) allow for asmooth gradation in composition of the FGC through the thickness(Fig. 3). The FGC exhibits a reasonably dense microstructure, with noapparent defects within the various layers or between them (Fig. 3).Significant differences between the nominal and actual compositions(Table 2) were, however, observed for the WC–Co-rich layers, wherethe actual steel content was lower than the nominal one. A possiblereason for this discrepancy can be inferred by observing that theunmelted steel particles, which were seen in the steel-rich layers, arescarce in the WC–Co rich ones (Figs. 2D,E and 3): as these particleshave lower cohesion strength to the surrounding material, they couldhave been removed during the impact of the WC–Co particles, whichexert an erosive action on them. That is to say, the WC–Co particlesthemselves could eject some of the unmelted, weakly bonded steelparticles. The very high local strain field produced during the impact,indeed, inserts enough energy in the underlying layer to cause severeplastic deformation, compaction, residual stress and even debonding[43,44].

Fig. 2. Cross-sectional SEM micrographs (backscattered electrons) of the single-layer coatings: 100% steel (A), 20% WC–Co (B), 40% WC–Co (C), 60% WC–Co (D), 80% WC–Co(E), 100% WC–Co (F). The inset in panel F is a high-magnification detail of the fine carbide structure of the WC–Co coating.

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3.2. Residual stresses

In TS coatings, particles forming the coating inherently developtensile residual stresses (quenching) due to the solidification andcontraction during cooling. These stresses are high but can be relievedduring splat formation via inelastic mechanisms such as yielding,microcracking, or creep. In high impact velocity processes like HVOF,the stress state of the deposited splats changes as a result of plasticdeformation and peening effect from the incoming particles. Thiseffect imposes compressive stress in the already deposited layers.Therefore, the stress state during deposition of this type of coatingsis a combination of quenching and peening stresses, here named“evolving stress” [45]. Metal/alloy materials are more prone todevelop overall compressive stress during HVOF processing comparedto ceramics or cermet materials. For this reason, SS316 was chosen for

the inner layers to reduce tensile stresses that contribute to crackopening and delamination as the coating grows thicker [46]. As theparent substrate material has the same chemical composition as thefeedstock, good wettability is expected in the coating/substrate inter-face, which may improve coating adhesion. Using the same parentmetal material also minimizes the thermal mismatch residual stressesbetween coating and substrate when cooling the system from thespraying temperature to room temperature.

In Table 3, the curvature gradient per deposited coating thicknessfrom ICP sensor is reported for each single-layer coating. As the WC–Co volume fraction increases, the curvature gradient also increases asa result of larger quenching stresses — a typical behaviour of WC-based cermet coatings. SS316 develops overall low tensile stress dueto the higher effect of peening. Evolving stresses are calculated in thesix individual experiments of curvature evolution of the single-layers

Fig. 3. Cross-section of the functionally graded coating (FGC): image obtained byjoining 5 adjacent micrographs (backscattered electrons).

Table 3Parameters used in residual stress analysis: curvature evolution per depositedthickness for each single layer as obtained via ICP sensor, evolving stress calculatedfrom Stoney Formula, and coefficients of thermal expansion (CTE) and elastic modulusby rule of mixtures for thermal stress calculation.

Nominal powderblend composition(vol.%)

Curvature change perdeposited thickness(m-1 /mm)

Evolvingstress(MPa)

CTE(×10−6 C°−1)

Elasticmodulus(GPa)

100% SS316L 0.214 37.7 16.0 9780%SS316L-20%WC–Co

0.376 66.3 13.6 114

60%SS316L-40%WC–Co

0.778 137.1 11.5 129

40%SS316L-60%WC–Co

1.189 210.1 8.5 152

20%SS316L-80%WC–Co

1.967 349.1 6.7 164

100% WC–Co 2.015 351.3 5.2 175

-1.5 -1.0 -0.5 0.0 0.5 1.0-600

-400

-200

0

200

400

600Evolving StressThermal StressResidual Stress

CoatingSubstrateA

Str

ess,

σ [M

Pa]

S

tres

s, σ

[MP

a]

-200

0

200

400

600

Evolving StressThermal StressResidual Stress

Thickness, t [mm]

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coatings, and not in the FGC as the consideration of thin coating on asubstrate is not valid in the FGC. Stoney formula was used to computeevolving stress from curvature gradient per deposited coatingthickness, under the following assumptions: 1) each deposited passof a single-layer coating (each single-layer coating is built by severalpasses to reach the desired thickness, see Table 1) is much thinnerthan the substrate; and 2) the stiffness of the substrate does notchange much by the progressively deposited passes.

An analytically-computed stress profile is reported in Fig. 4 forcomparison of a thick FGC and a thick pure WC–Co coating. Threeprofiles are obtained for each case, as follows: 1) the evolving ordeposition stress profile (quenching+peening) is obtained accordingto Tsui-Clyne analytical model [47], 2) the thermal stress profile, due

Table 2Actual composition of the composite layers, as determined by image analysis. Identicalresults were obtained on the single-layer coatings and on the FGC layers.

Nominal vol.% of WC–Co Actual vol.% of WC–Co

20 22.1±2.840 41.4±3.360 69.9±7.180 86.3±2.8

to thermal expansion mismatch during cooling from deposition tem-perature to room temperature, is obtained using a multiple-layeranalytical model for layers of significant thickness described in [48],and 3) the residual stress profile is obtained by the sum of the pre-vious two profiles.

The evolving stress profile is obtained under the assumption of aprogressive addition of layers at certain steady temperature (depo-sition temperature ~270 °C, as measured by ICP), each one causing a

-1.5 -1.0 -0.5 0.0 0.5-600

-400

CoatingSubstrateB

Thickness, t [mm]

Fig. 4. Through thickness residual stress profiles of A) FGC, and B) thick WC–Co sample.Three profiles are presented: evolving stress (stress developed during the depositionsession), thermal stresses due to thermal contraction mismatch, and final residualstress state.

Fig. 5. Vickers microhardness of the single-layers and of the FGC layers. The rule ofmixtures prediction for the composite layers of the FGC is shown by the solid line.

2202 A. Valarezo et al. / Surface & Coatings Technology 205 (2010) 2197–2208

curvature change in the coating-substrate system produced by themisfit strain. The curvature was measured by the ICP sensor andreported in Table 1. By feeding in the model the curvature changeproduced by each single-layer, the balancing forces and therefore, thestresses can be calculated for the layer being deposited and thematerial beneath. The progressive addition of layers in the FGC, eachone having different stiffness and deposition stress, causes the stressprofile to change each time a layer is deposited.

The thermal stresses are calculated as thermal mismatch duringcooling from the spraying temperature to room temperature(temperature gradient ~250 °C). Due to the significant thickness ofthe layers to the substrate, the elastic thermal stresses are affected bythe stiffness of each single-layer (the thermal strain is not driven bythe substrate only, as it is the case for thin films). The analytical modelused in this elastic multilayered system is described in [48]. The valuesof thermal expansion (CTE) used in the calculation are reported inTable 3. The coating modulus used in the stress calculations for eachsingle-layer is also reported in Table 3, as obtained by instrumentedmicroindentation.

During deposition, the discontinuity of the stress profile in thecoating/substrate interface is much smaller for FGC than for the pureWC–Co, which reduces the stored energy in the interface that mightresult in interface fracture during the spraying process itself [46], andin this case allows the build up of a thicker coating. Upon cooling,compressive thermal stresses benefit the WC–Co-rich layers becauseof the larger thermal strain mismatch with the substrate. Bothcoatings therefore show significantly high tensile stresses in the outerlayers and progressively turn in neutral and compressive stressesinside the coatings due to the balance of forces of the incoming layersto the already present layers. Therefore, a much steeper profile changeis observed in the pureWC–Co coating due to the highermagnitude ofevolving stress of this composition.

The FGC stress profile can benefit the system for crack arresting,provided that if a crack initiates in the top layer (WC–Co), as the crackprogresses deeper in the coating, it would find a region of lowertensile stresses and even compressive stresses much sooner in theFGC than in the WC–Co assuming similar thicknesses. It should benoted that as a crack initiates and grows in the coating, the adjacentregions relieve stress, and the underlying layers will rearrange thestress profile. In this regard, a crack growing from the outer surface inthe pure WC–Co coating is more prone to grow in presence of tensilestresses than in the FGC.

The analysis considers the system in the elastic regime, althoughsome yielding may have occurred in the substrate (SS316 YieldStrength: 290 MPa) if superposition of peening stresses from grit-blasting are considered.

Table 4Wear rates and friction coefficients of the coatings after ball-on-disk tests.

80% WC–Co 100% WC–Co FGC

Test-1 Wear rate[10−8 mm3/(Nm)]

3.57±0.14 0.879±0.148 0.889±0.253

Friction coefficient 0.38±0.02 0.40±0.02 0.37±0.01Test-2 Wear rate

[10−8 mm3/(Nm)]1.50±1.20 0.736±0.109 0.579±0.120

Friction coefficient 0.38±0.02 0.37±0.01 0.40±0.02

3.3. Microhardness

The microhardness values of the various layers of the FGC are notmuch different from those of the corresponding single-layer coatings(Fig. 5), indicating that the embedment in a thick graded coating doesnot affect the mechanical characteristics of the various layers, despitethe different residual stress state and peening effect exerted by thesubsequent layers. The peening effect of a layer of the new compo-sition would only modify the top-most portion of the depositedcomposition and may not have a significant effect on the overallhardness, whereas the residual stress is not significant at the stresslevels that the indentation produces.

The hardness increases quite linearly from the full SS316 to thetop-most WC–Co layers, in spite of the deviations from the nominalcomposition in the WC–Co-rich layers. The presently-adopted HVOF-spraying technique was therefore capable of forming a FGC, close tothe intended structure, although some improvements to the deposi-tion process control are still required.

The experimental values of the hardness of the composite layers ofthe FGC can be compared to those computed according to the rule ofmixtures (2):

HC = VSSHSS + VWCHWC ð2Þ

where

HC computed hardness of the composite layer;HSS, HWC experimentally-measured hardness of the 100% stainless

steel layer and 100% WC–Co layer, respectively (from Fig. 5);VSS, VWC actual volume fractions of stainless steel andWC–Co cermet

(respectively) in the composite layer (from Table 2).

The results are shown as a solid line in Fig. 5. It is interesting tonote that, within the experimental error, the experimentally-measured values agree with the numerical prediction of Eq. (2).This indicates that the mechanisms tested in the hardness measure-ments, such as, among others, strain hardening by peening, intersplatcohesive strength, thermal stresses at the splat level, all average theireffects in a weighted manner as the composition changes. This alsocorroborates to the use of the rule of mixtures for the computation ofthe elastic modulus of the composite layers (Table 3), as mentioned inSection 3.2.

3.4. Ball-on-disk tests

The FGC and the 100 vol.% WC–Co single-layer coating exhibitalmost identical wear rates under both test conditions, whereas the80 vol.% WC–Co single-layer coating suffers more severe wear(Table 4). No remarkable difference exists between the averagefriction coefficients recorded in all tests (Table 4).

For both FGC and the 100 vol.% WC–Co single-layer coatings, twowell accepted mechanisms govern the wear behaviour of suchcoatings [30,49–52], and are briefly described here. Brittle fractureoccurs (i.e. cracks form on or near the coating's surface) and

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propagates until portions of material are removed (Fig. 6A, B). On theother hand, abrasive removal of the metal matrix between the carbidegrains takes place (Fig. 6C). This latter mechanism typically dominatesthe wear behaviour of bulk sintered cermets [49] and of thosethermally-sprayed coatings where the metal matrix content is higher(e.g. WC–17%Co or WC–10%Co–4%Cr) and/or the dissolution of thecarbide phase in the metal matrix is very limited [50,51]. The onset ofcrack propagation, by contrast, is triggered when the matrix isembrittled by some degree of carbide dissolution [30,50,52]: theoccurrence of this phenomenon was shown in Section 3.1. Fine-grained cermets, like the present one, are favourable for limitingmatrix abrasion as the mean free path between carbide grains islowered [30]. Thus, the penetration and inclusion of hard secondary

Fig. 6. SEM micrographs (secondary electrons) of the ball-on-disk wear scars on the FGC (detail) and the 80% WC–Co single layer (D: overview, E: detail).

asperities inside the metal matrix is reduced. They are, however, notequally desirable for preventing the onset of brittle fracture: fine-grained particles are more prone to dissolve in the matrix duringspraying because of the larger overall carbide/matrix interface area[30,52].

These results show that the inclusion in the FGC architecture doesnot affect the intrinsic tribological characteristics of the WC–Co layeras the depth of the stress field is relatively shallow and is not reachingthe more compliant sub-layers.

The 80 vol.% WC–Co single layer coating was tested in order toverify if the addition of a low amount of steel to the brittle WC–Comaterial could improve its wear performance by increasing toughnessand ductility; however, the sliding wear behaviour of the coating is

A), the 100% WC–Co single layer (B: lower-magnification view, C: high-magnification

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clearly impaired by the addition of SS316 (Table 4). Because of thereduction in the overall hardness, indeed, ductile grooving occurs(Fig. 6D); moreover, the SS316 splats are preferentially worn andmayalso be oxidised (Fig. 6E). Consequently, it is confirmed that the choiceof pure WC–Co as the top layer of the FGC is suitable for maximizingthe sliding wear resistance.

3.5. Indentation and ball drop cyclic impact tests

Dense WC–Co coatings by TS are intrinsically brittle due to theirhigh carbide content and the embrittlement of the matrix by carbidedissolution (as observed in Sections 3.1 and 3.4). The combination ofthis phenomenon with process-induced tensile residual stresses (asobserved in Section 3.2) causes crack initiation and a deleteriouseffect on coating cohesive strength. In the present study, the brittleand hard WC–Co layer was engineered for enhancing the damagetolerance via incorporation of tough and compliant sublayers. The aimis to achieve crack containment through designing intended crackdeflection paths with intrinsic splat interfaces, reducing crack-tipstress concentration with the compressive residual stress, andarresting any penetrant cracks with a secondary ductile metal phase[53]. These damage tolerance concepts in the present FGC architectureare therefore not expected to manifest themselves as a “crack preven-tion” methodology, but rather through the achievement of defecttolerant coatings to control failure mechanisms. To verify the fulfil-ment of the damage tolerance requirements, indentation-based con-centrated load tests, namely spherical indentation and ball drop cyclicimpact were carried out on the FGC and compared to that of the HVOFWC–Co coatings.

Under a spherical load, WC–Co coatings show the typical fracturebehaviour of a hard and brittle ceramic coating onto a soft substrate(as described in [24]). The thick single layer (~0.6 mm coating),Fig. 7A, shows near-contact ring cracks (or cone-cracks, label c),whereas the thin single-layer, Fig. 7B, shows radial cracks at thesubstrate-coating interface (label r) and secondary circumferentialouter-contact ring cracks (or peripheral ring cracks, label p) at the topsurface. In both conditions, the applied stress field exceeded thecritical load for crack initiation. In the thin coating, the plasticdeformation of the substrate during loading induces a flexural stresscondition that promotes radial cracking near the interface. Moreover,during unloading, the elastic recovery of the coating adjacent to theplastically deformed zone in the substrate promotes horizontal crackgrowth by cohesive failure (Fig. 7B, label h). Increasing the load and/orrepeating contact cycles induce larger horizontal cracks during theunloading process.

When indenting the FGC, the occurrence of cone-cracks wasalmost completely suppressed (some shallow cone cracks may beobserved from the top view but they can hardly be observed in thecross section) as a result of an increase in compliance with depthpenetration, Fig. 7C, and, therefore, a redistribution of stress fields[18]. Limited damage was observed in the cross-section after oneindentation. As the number of indentations increased, more horizon-tal cracks and cracks growing in tortuous paths were observed,Fig. 7D. Cracks nucleated along inter-splat interfaces on the subsurfaceof the indented region. Hence, flexure and spring-back effects areenabled in the stiff WC–Co layers in response to local yielding in thesublayers enriched with metal, all this leading to radial cracks as wellas horizontal cracks (cohesion failure). In Fig. 7E, the deflection of acrack growing from theWC–Co phase along an inter-splat interface ofWC–Co and SS316 is illustrated. The inter-splat interface is the regionof the weakest fracture toughness and thus, cracks blunt along theinterfaces. In Fig. 7F, a crack branching is shown betweenWC–Co anda thin SS316 splat. The dominant crack path extends even through theSS316 to the adjacent WC–Co splats regions as well as through theinter-splat interfaces. When a crack approaches a thin SS316 splat,this one imposes a shield effect to crack propagation due to plasticity

that may blunt the crack; however, the effective driving force topropagate the crack may be amplified by the presence of the stiff WC–Co underneath the SS316 splat, and depending upon the thicknessesof the layers and the cohesive strength between the splats, the crackmay be allowed to advance through the ductile phase [54].

The cyclic impact tests on the thin and thick WC–Co coatingsresemble the results from the spherical indentation at a much higherload, strain rate, and deeper stress field. In the thin layer, radial cracksstarting from the substrate-coating interface and propagating towardthe coating's surface (Fig. 8A) exist in the middle of the impactedregion. Circumferential outer-contact ring cracks (Fig. 8C) propagatefrom the top surface at an angle of about 45° (Fig. 8B) through thethickness. This cracking mode is similar to the one described byprevious reports of impact tests in WC–Co [55]. Additionally, somesubsurface cracks arise in the middle of the impact area (Fig. 8D).These subsurface cracks propagate parallel to the substrate interface,following the lamellar structure; in particular, they propagate alongthose bright regions where the metal matrix is more significantlyembrittled by larger dissolution of carbides, similarly described in[52].

The thick WC–Co coating was subjected to a larger fraction of theimpact energy as the substrate underwent negligible plastic defor-mation. Consequently, significantly larger damaged site was observedwith the cyclic impact loads (Fig. 9A,B). In addition to circumferentialcone and ring cracks, other cracks, propagating over long distancesoutside the impact region were distinguished on the top surface(Fig. 9A, arrows), which compromises the integrity of the coating overa wider area. Moreover, the cross-sectional micrograph below theimpact crater reveals severe disruption (Fig. 9B) in the coating struc-ture as the damage zone progresses in both the through-thickness andin-plane directions.

Consistently with the indentation test, the graded structure ofthe FGC yielded significant compliance and ductility during the balldrop impact test as well, as compared to the WC–Co top layer. Theimpression left in the coating was therefore larger than on the thickpure WC–Co coating (compare Fig. 10A to Fig. 9A), but crackingwas much less extensive. No cone-cracks were observed; peripheralring cracks (Fig. 10C) were reduced and the middle radial cracks(Fig. 10B) were almost completely suppressed. As mentionedearlier, the SS316 splats in the underlying layers are capable ofarresting cracks, and of absorbing part of the impact energy viaplastic deformation. In Fig. 10C, the cross-section of the 80 vol.%WC–Co layer clearly shows how these cracks are arrested andblunted at the presence of SS316 phase. Contributions from residualstresses are also present as it is at around this depth were a shift tocompressive residual stress occurs.

The overall through-thickness impact damage of the FGC iscomparatively much lower (Fig. 10E) than the pure WC–Co layer(Fig. 9B). It is noteworthy, however, that a large crack arises in the FGCalong the interface between the pure SS316L layer and substrate orjust above it. Its formation is related to the occurrence of plasticdeformation in the SS316L layer and in the substrate. Presumably, thesubstrate exhibited larger plastic deformation as compared to thecomposite layers due to lower hardness. When elastic recoveryoccurs, the system composed by the pure WC–Co top layer and by thecomposite layers behaves as a thick elastic beam with large flexuralstiffness, and exerts large out-of-plane tensile stress on the plastically-deformed SS316 layer and substrate. The out-of-plane tensile stressinduces failure in the weakest locations at the SS316L layer(weakened by the inclusions of oxides and unmelted particles,Section 3.1) and the substrate interface.

It was confirmed that the delamination occurred by contact fatigueof the underlying area. It was verified that after a few hundreds ofcycles, there was no-apparent presence of delamination cracks in thecross-section. This cycling and concentrated loading type of the ball-drop impact test does not represent the scenario of a typical

Fig. 7. SEMmicrographs of spherical indentation results on the cross-sections of bonded-specimen samples. A: WC–Co thick single layer (c-cone cracks); B: WC–Co thin single layer(p-peripheral cone cracks, r-radial crack, h-horizontal crack); C: FGC after a single indentation, ten indentations, and a hundred indentations in the same area; D: FGC, detail of ahundred indentation area E: FGC, detail of crack branching through an intersplat interface in the 80 vol.% WC–Co layer F: FGC, detail of crack blunting when the crack approaches aSS316 region in the 80 vol.% WC–Co layer. Dashed lines illustrates crack paths.

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application of cermet coatings; whereby, the advantage of using agraded structure over a monolithic thick coating still prevails.

The FGC structure has been proven effective in absorbing theimpact energy and limiting the contact damage to both the WC–Cotop layer and to the entire system. The gradation also improves theresponse of the top WC–Co layer of the FGC over the single WC–Colayer as the plastic deformation of the underlying layers (as the case is,in Fig. 7B) is limited, which eliminates flexural stresses and thegrowth of large radial and horizontal cracks.

Apparently, coating adhesive failure would be more prone to takea place than cohesive failure, and therefore, processing improvementsare still needed. Further analysis is needed to compare adhesionstrength between the FGC and pure thick WC–Co layer.

4. Conclusions

A WC–Co/stainless steel functionally graded coating (FGC), con-sisting of six different layers with a stepwise change in composition

Fig. 8. SEMmicrographs of the cyclic impact testing results on the 100%WC-Co single layer. A: cross-sectional view of the middle area; B: cross-sectional view of the peripheral area;C: top surface view; D: cross-sectional view of the middle area, detail of the near-surface region. The circles in panels A and B highlight radial cracks.

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from 100 vol.% AISI 316 L stainless steel to 100% WC–Co, wasdeposited by HVOF-spraying onto a stainless steel substrate. Actualcomposition of the WC–Co/stainless steel composite layers did notperfectly match with the nominal one due to processing effects. Asmooth gradation in hardness (matching quite well with a rule-of-mixture prediction) was obtained.

A notable result is, compared to a thick pureWC–Co coating, the FGCarchitecture significantly reduced the stress buildup and stressdiscontinuity arising at the coating/substrate interface during deposi-tion. The tensile quenching stress, which is very large in pure WC–Co,is indeed mitigated by compressive peening stresses in the stain-less steel-rich layers close to that interface. The risk of delaminationduring deposition of thick coatings is therefore minimized.

Fig. 9. Backscattered electron SEM micrographs of the top surface (A) and cross-

The overall stress distribution in the FGC features a smoothgradient from a tensile stress in the cermet top layer to slightly com-pressive stresses in the stainless steel-rich layers. Owing to such stressdistribution and to its special architecture, the FGC therefore actedquite effectively as a damage tolerant structure under severe contactstress conditions, i.e. the effectiveness of the FGC structure is par-ticularly manifested when loads are large enough to probe the entirethrough-thickness behaviour of the coating. Both static (high-loadindentation) and dynamic (cyclic ball-drop impact) tests indeedshowed that the FGC structure has superior ability to control andmitigate the propagation of cracks and defects, in comparison to purecermet layers of comparable thickness. Cracks can initiate in thecermet top layer, where large tensile stress exist, but the underlying

section (B) of the≈0.6 mm-thick WC–Co coating after cyclic impact testing.

Fig. 10. SEM micrographs of the surface (A) and of the cross-section of the FGC after cyclic impact testing: middle area (B) and peripheral area (C) of the top layer; 80 vol.% WC–Colayer (D) and overview (E).

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layers can arrest or deflect them by three main mechanisms, whichwere theoretically expected and experimentally verified: deflection ofcracks along tortuous paths through the interfaces between steel andWC–Co lamellae, arresting of cracks entering the ductile steel lamel-lae, reduction of crack-tip stress concentration by the compressiveresidual stress in the underlying layers.

The advantages given by the FGC structure are less obvious whencontact stresses are localized in the vicinity of the top surface only: noperceivable difference indeed arose between the behaviour of pureWC–Co and of the FGC during ball-on-disk tests, where most of thecontact pressure distribution involved the near surface region only. Inthe near-surface region, indeed, both the pureWC–Co coating and theFGC feature large tensile residual stresses (due to the large quenchingstresses developed by cermet lamellae during the deposition process)and the contact stress distribution does not extend far enough belowthe surface of the coating for the underlying layers to play an effectiverole on the mechanical response.

Some limitations of the present FGC architecture, however,emerged. One of the main problems seems to be the poor char-acteristics of HVOF-sprayed stainless steel. It suffers from extensiveoxidation and phase alteration during spraying, which compromise itsmechanical strength and probably impair its adhesion to thesubstrate, which made the FGC to come off under the ball drop im-pact test. The use of a fuel-rich flame (oxygen reducing) is proposed asa possible solution.

Acknowledgements

This work was partly supported by PRRIITT (Regione Emilia-Romagna, Italy), net-lab “INTERMECH”, and by MIUR, Italian Ministryfor University and Research (Funds: “Programmi per l'incentivazionedel processo di internazionalizzazione del sistema universitario”).

The research at Stony Brook is sponsored in part by the NationalScience Foundation GOALI-FRG program CMMI0605704 and throughthe consortium for thermal spray technology.

The authors are grateful to Prof. Andrew Gouldstone for hisvaluable comments andMs. Alessia Candeli for the precious assistancewith the experimental activities.

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