asuku

oatducesshe msticvenmeoppmof thatt

cermet/metal laminate coatings can be one alternative approach to further improvement of the mechanicaled cermet coatings.

1. Introduction

s haveprotectsive wexy-fuelniquesby HVthanis muc

decarburization of WC and dissolution of W and C into Co binder Another approach to improve the mechanical properties of cermet

Surface & Coatings Technology 205 (2011) 53585368

Contents lists available at ScienceDirect

Surface & Coatin

l sduring spraying.In order to improve the fracture toughness and overall mechanical

performance of WCCo coatings, mainly two approaches have beentaken. One is to depositWCCo coatings at lower process temperatureand to suppress those detrimental reactions during ying of particles.Jacobs et al. [4,5] deposited WCCo and WCCoCr coatings by highvelocity air-fuel (HVAF) spraying. The coatings did not show the

coatings is addition of metallic phase. Osawa et al. [10] deposited amixture ofWC20Cr3C27Ni agglomerated-sintered powder and Ni orNiCr powder by HVOF and investigated the abrasive wear and impactresistance of the coatings. The deposited coatings contained theuniformly distributed Ni or NiCr particles surrounded by WCCrCNicermet matrix. The coatings showed a signicant improvement inwear and impact properties due to energy absorption by the ductileformation of brittle phases. While the hardndecreased, the WCCoCr coating showedwear resistance. Kim et al. [6] successfullydeposition to fabricate WCCo coatings w

Corresponding author. Tel.: +81 29 859 2469; fax:E-mail address: watanabe.makoto@nims.go.jp (M. W

0257-8972/$ see front matter 2011 Elsevier B.V. Adoi:10.1016/j.surfcoat.2011.05.054perties are formations ofM12C and M6C) due to

between the toughness of WS WCCo and that of a correspondingsintered body.materials. The origins of the poor fracture probrittle phases such as W2C and phase (Thermally sprayed WCCo coatingindustrial components in order toenvironmental damages such as abraobject damages [13]. High velocity oof the most popular deposition techrecent days. While coatings depositedhigher hardness of 10001300 Hvmaterials, their fracture toughnessbeen applied to variousthem from aggressivear, erosion, and foreign(HVOF) spraying is oneof cermet coatings in

OF show similar or evensintered bulk WCCoh inferior to the bulk

phases mentioned above. While the coatings exhibited extremelyhigh hardness of 18202050 Hv, the overall performances of thesecoatings such as toughness and wear properties are still unknown.Watanabe et al. [7] and Chivavibul et al. [8,9] investigated severalmechanical properties of WCCo coatings with various Co contentsfabricated by warm spray (WS) and high velocity oxy-fuel (HVOF)depositions and found the improved toughness for WC-17 and 25% Coby WS. The wear performances of WS coatings such as abrasive wearresistance were also better than those of HVOF coatings for a givenhardness value. However, there still remained a substantial gapess of HVAF WCCo wasimproved hardness andapplied cold spray (CS)ithout the detrimental

metallic particlesimpact behaviorwhich a metalliwiched by thedeposited by eitthe coatings wascermet coatingperformance tha

+81 29 859 2401.atanabe).

ll rights reserved. 2011 Elsevier B.V. All rights reserved.

properties of thermal sprayMultilayered WCCo/Cu coatings by warm

Makoto Watanabe , Masayuki Komatsu, Seiji KurodHigh Temperature Materials Unit, National Institute for Materials Science, 1-2-1 Sengen, T

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

Article history:Received 8 March 2011Accepted in revised form 31 May 2011Available online 15 June 2011

Keywords:Warm sprayWCCoMultilayered coatingsFracture toughnessDamage tolerance

WCCo/copper multilayer cto investigate the effect offracture, and surface hardnremoval of the substrates. Tin theWCCo layers and plafraction of copper showed eand no benecial feature inhigher volume fraction of cbending strength than theamong all samples instead oby the intact WCCo layers

j ourna l homepage: www.espray deposition

ba, Ibaraki, 305-0047, Japan

ings consisting of 8 layers were fabricated by warm spray deposition in ordertile layer inclusion onto their fracture behavior. Bending strength, work ofof freestanding coatings were examined by three point bending tests afterultilayered samples showed non-linear stressstrain curves due to crackingselongation of the copper layers. The multilayered samples with lower volumelower bending strength than the monolithic samples of WCCo and copperchanical performance was found. On the other hand, the samples containinger exhibited more than twice higher work of fracture and moderately betternolithic WCCo coatings, while the surface hardness was almost identicale monolithic copper. The ductility of copper layers and the plastic constraintributed to enhance their mechanical properties. It has been concluded that

gs Technology

ev ie r.com/ locate /sur fcoat. Hadad et al. [11] investigated adhesion strength andof HVOF cermet coatings consisting of three layers, inc intermediate layer (Ni, NiCr, CoCr) was sand-WCCrCo layers. The intermediate layers wereher HVOF or electroplating. The impact resistance oftested by an experimental shooting device and thewith the Ni-electroplated layer exhibited bettern the monolithic coating. The results implied the

Table 1Spray parameter.

Parameter Cu WCCo

Fuel (dm3/min) 0.35 0.38Oxygen (dm3/min) 713 779Nitrogen (dm3/min) 1000 500Powder feed rate (g/min) 63, 32, 16 85Spray distance (mm) 180Barrel length (mm) 203.2

Table 2Sample list.

SampleID

Number of layer Thickness of layer(m)

Amount of Cu(vol. %)

WCCo Cu WCCo Cu

A 1 0 600.0 0.0B 4 4 98.9 23.7 19.4C 4 4 93.7 43.6 31.8D 4 4 74.0 121.0 62.0E 0 1 1970.0 100.0

5359M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 53585368possibility of further enhancement of fracture properties in cermetcoatings by introducing multilayer structures. Valarezo et al. [12]investigated HVOFWCCo/stainless steel functionally graded coatings(FGC) consisting of six different layers with a stepwise change incomposition from 100% stainless steel to 100% WCCo. The FGCsexhibited high fracture resistance for vertical crackings due to energyabsorption of the distributed metal phases and due to relaxation ofresidual stress in the coatings. In all the three cases, the essential pointis adding the ductile phases in the coatings.

Based on the previous studies, the design strategy to maximize thefracture properties of WCCo coatings can be summarized as follows.The rst is to suppress the detrimental reaction during spray and thesecond is to add the metal reinforcements such as particles, bers, orlayers. Fiber reinforced composites [13] or laminates [14,15] are wellknown to have greater fracture toughness than their monolithiccounterparts. Although it is of great interest to develop berreinforced coatings as studied in [16,17], multilayer coatings hasbeen chosen in the present study because of the difculty ofreinforcing coatings by continuous bers. However even for multi-layer system, since it requires rather complex operations to depositmany layers, only few works have been attempted in the past.Ravichandran et al. [18] investigated thermal conductivity ofmultilayer coatings of alumina and yttira-stabilized zirconia, whichwere deposited by plasma spraying. The maximum number of layerswas 16. Although the effectiveness of multilayer structure onreduction of thermal conductivity was small, it assured the potentialof thermal spray technology to make a multilayer structure. In thepresent paper, WCCo/copper multilayer coatings containing 8 layershave been developed by warm spray deposition aiming at damagetolerance of cermet coatings with keeping high hardness on theirsurfaces. WS copper coatings showed high strength with moderateductility in previous study [19,20] and hence copper was chosen asmetal reinforcements in the current work. In order to reveal the basicand fundamental aspects of fracture characteristics of multilayercoatings deposited by warm spraying, the strength and fracturebehavior of the freestanding coatingswere investigated by three pointbending tests.

2. Experimental procedure

Tungsten carbidecobalt (WCCo) and copper (Cu) layers weredeposited one after the other onto a pure aluminum plate with adimension of 501002 mm3 by warm spray (WS) deposition.Powder ofWC12wt.%Co (Fujimi Inc., Aichi, Japan) with a particle sizeof 520 mand a carbide size of 0.2 m, and copper powder (Cu-ATW,Fukudametal foil and powder co. ltd., Japan)with a particle size under45 m were used. The WCCo powder was manufactured by spraydrying and light sintering, and the copper powder was fabricated bywater atomization. The aluminum substrate was chosen because it iseasy to dissolve in alkali solution in order to obtain a freestandingcoating.

AWS system [21] has been developed bymodifying a conventionalHVOF equipment (JP5000, Praxair Technology Inc., USA) by adding amixing chamber between a combustion chamber and powder feedports. It is capable to control temperature of combustion ame byinjecting nitrogen gas at the mixing chamber so that this process cankeep temperature of sprayed particles under their melting point withmoderately heated and thermally softened states. The detail of theprocess can be found elsewhere [22]. The spraying conditions of WSprocess are listed in Table 1. In these conditions, both powders ofWCCo and copper were sprayed as solid states and bonded by ultra-highspeed impact. Surfaces of the substrates were blasted by alumina gritand degreased by ultrasonic cleaning in acetone before spraying. Fivecombinations of multilayer coatings were fabricated as listed inTable 2. Samples were labeled as A, B, C, D, E in terms of Cu volume

fraction, vm=0.0, 19.4, 31.8, 62.0, and 100.0%. The Cu volume fractionwas controlled by varying the powder feed rate of Cu powder(Table 1) and thus changing their thickness ratio. Samples A and E aremonolithic WCCo and copper coatings respectively. In all casesexcept for sample A, the rst layer on the substrate was Cu and theouter most surface was WCCo. The total number of layers was xedat 8 containing 4 WCCo and 4 copper layers in all laminate samplesBD. Since the number of layers was xed, the total thicknesses ofcoatings were varied among samples as shown in Table 2. Afterdeposition of coatings, the samples were cut into rectangular barswith a dimension of 505 mm2 (lengthwidth). All the bars exceptfor sample E were immersed into NaOH aqueous solution (NaOH:15 g, distilled water: 500 ml) at a temperature of 60 C, and thealuminum substrates were dissolved in order to obtain freestandingcoatings. For sample E, the substrate was mechanically removedbecause it was very easy to obtain Cu coatings with more than 2 mmthick and thus the machining was handy for this sample.

Three point bending tests were conducted on those freestandingcoatings. The outer span L was 40 mm and the crosshead speed of0.1 mm/min was applied. Three specimens were tested for each case.For the cases of themultilayer coatings, the copper side surface, whichwas the rst layer during deposition, was placed in the compressionside and thus the WCCo surface, which was the eighth layer, wasplaced on the tension side. The load P and crosshead displacement were recorded. Fracture behaviors were monitored during thebending tests by a CCD camera. The strain t and the apparentbending stress t on the outer most tension side for arbitrary werecalculated as [23]

t =6tL2

1

t =32

Lwt2

P 2

wherew and t is the width and thickness of the specimen. Note that tis only valid for elastic deformation and cannot represent correctstress states after yielding, and thus this value should be considered asthe apparent parameter. In addition, apparent work of fracture wasdened and calculated as the area under the load-displacement curvedivided by the twice the cross section area of the specimen [24,25],

= max0

Pd

2wt3

where max stands for the maximum displacement of the crosshead.The microstructure and crack patterns were examined with ascanning electron microscope (SEM, JSM-6500F, JEOL Ltd., Japan)and the phases in the monolithic coatings (samples A and E) weredetermined by X-ray diffraction analysis (XRD, RINT2500, RigakuCorp., Japan) on the surface of as-sprayed samples. Microhardnesstests were also carried out on the surfacewith a 300 g load and a dwelltime of 15 s. At least ten measurements were made for each sample.Oxygen contents in the copper powder and coatings were analyzed bythe inert gas fusion method (LECO TC-600, Leco corp., US).

3. Results

3.1. Microstructural characterization of coatings

Fig. 1 shows the cross-sectional images of the monolithic coatingsof sample A (Fig. 1a,b) and sample E (Fig. 1c,d). The WS WCCocoating has very dense microstructure with tightly packed blockycarbides. In Fig. 1b, some carbide are larger than 1 m possibly due tothe crystal growth during light sintering after spray drying in thepowder manufacturing process. The images of the copper coatingswere taken after removal of the substrate. Although the microstruc-ture of the Cu coating appears very clean and dense, some interfacesamong the deposited particles contain voids (Fig. 1d). This implieslower mechanical properties of the WS copper coating than bulkcopper metal. XRD analysis results are plotted in Fig. 2. Sample Aconsists of onlyWC and Co phases indicating low process temperature

samples while the thickness of WCCo layers was almost constant(Table 1). The vertical cracks were observed in someWCCo layers inall three samples (Fig. 2a,c,e). The cracks seem to have occurred at thesecond WCCo layer from the interface between the coating and

1 m

10 m

d

a

b

Fig. 2. Cross-sectional image of laminated coatings deposited by warm spraying: (a)(b)sample B, (c)(d) sample C, (e)(f) sample D.

5360 M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 53585368in WS and suppression of the formation of detrimental phases such asW2C as already reported [79]. On the other hand, sample E showednot only the peaks of copper but also of a copper oxide CuO. Itindicates that the particle temperatures were high enough to causeoxidation during ight. The oxygen content in sample Ewas 0.24 mass% while the value in the feedstock powder was 0.096 mass%.

The microstructures of samples B, C, and D are indicated in Fig. 3ab, cd, and ef respectively. They consist of four WCCo (lighter gray)and four copper (darker gray) layers. Since the feed rate of copperpowder has been varied in order to control the volume fraction ofcopper, the average copper thicknesses are different among the three

Cu

WC-Co

a

100 m

500 m

cFig. 1. Cross-sectional images of monolithic coasubstrate, and propagated toward the coating surface. Such crackingwas not observed in the monolithic WCCo sample. There are twopossible mechanisms behind, causing this phenomenon. One is thelarger thermal expansion of the Cu layers than WCCo. In order tospray multilayer structures, it was necessary to change the powderfeeding lines for each layer and correspondingly the spray conditions,which it took some time so in the present spraying system. As a result,the coating-substrate couple was cooled down and reheated againbefore the deposition of a next layer. Since the thermal expansion ofCu [26] is much larger than WCCo [27,28], these thermal cyclingmight cause such cracking. The other possible explanation is the

btings of (a)(b) WCCo and (c)(d) copper.

particles is subjected to substantial impact energy and it can be

copper is much larger than that of WCCo, thus higher shrinkage canoccur in the copper layer during cooling stage after spraying. Thisshrinkage can induce compressive stress in the intact WCCo layers.However, the real residual stress states in the layered coatings are

100 m 100 m

20 m 20 m

e

f

warm spraying: (a) (b) sample B, (c)(d) sample C, (e)(f) sample D.

5361M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 53585368cracked especially for a brittle material. Since such pre-existing crackswill reduce the mechanical properties of coatings, it is necessary tounderstand the mechanism and to avoid crackings in the future. Inmost of the cases, cracks were not recognized in the copper layersindicating higher fracture toughness of the copper layers thanWCCoas expected.

3.2. Mechanical properties of coatingsimpact of solid Cu particles. Although sprayed particles are heated andsoftened in WS process, the surface impacted by the high velocity

a

100 m

20 m

c

b d

Fig. 3. Cross-sectional image of laminated coatings deposited byExamples of typical stressstrain curves obtained from the mono-lithic coatings of WCCo (sample A) and copper (sample E) by Eqs. (1)and (2), are shown in Fig. 4. In the case of WCCo coating, the stressincreased linearly with the increase of strain, and fracture occurred in abrittle manner at t=433 MPa and t=0.19%. On the other hand, theCu coating exhibited signicant plastic elongation after elastic defor-mation similar to bulk Cu. The 0.2% proof stress was 301 MPa. Afterreaching an apparent ultimate strength of 317 MPa at t=0.87%, ductilefailure initiated and the stress gradually decreased for furtherdeformation. Mechanical properties of the coatings were summarizedin Table 3. The reference data of bulk WCCo and copper materials[28,29] are also presented in the table. The bending modulus andstrength of WS WCCo coating are about 40% and 20% of the bulkcounterpart with an average carbide size of 1 m. On the other hand,mechanical properties of WS copper coatings are comparative to thoseof rolled copper [29]. However, it should be noted that the mechanicalproperties ofmaterials strongly dependon theirmicrostructures suchascarbide size, grain size, oxide content, degree of hardening, and so on.Thus the reference values in the table should be interpreted as a roughmeasure of bulk properties. The Vickers hardness on the surface ofsamples is also shown in Table 3. The laminate coatings have evenhigher values than the monolithic WC12Co coating. At the moment, itis not clear whether this difference was originated from the scatteringdue to the spray based process or the intrinsic properties due tolaminate structures. One possible explanation, for the latter case, mightbe the introduction of higher compressive stress inWCCo layers due tothe existence of copper ones. Since the thermal expansion coefcient ofunknown. This requires further investigation bymeasuring the residualstress development through the deposition process. But it is expectedthat the laminates have at least comparable wear properties to themonolithic coatings until the top WCCo layer is worn out.a) Sample A

b) Sample E

Fig. 4. Stressstrain curves of monolithic coatings of WCCo (sample A) and copper(sample E) in three point bending tests.

Typical stressstrain curves of laminated coatings with differentcopper layer thicknesses, samples B, C, and D, are given in Fig. 5. Insample B (Fig. 5a), the load did not increase linearly with strainincrease but showed gradual reduction of the gradient resulted in thenon-linear increase of stress. This behavior was possibly caused bycracking in the WCCo layer and following plastic deformation of Culayer on the tensile side. After several small dropping and recoveringas shown in the inset of Fig. 5a, the maximum stress of t=183 MPaat t=0.17% was reached. Beyond the load maximum, the stressstrain curves exhibited large load drops. However it is worth notingthat load did not drop catastrophically like sample A (Fig. 4a) butdecreased in a step like manner with failure proceeding. Samples Cand D generally showed similar trend in stressstrain behaviors(Fig. 5b,c). Both of the ultimate strength and the strain to failureincreased as the increase of volume fraction of copper. Interestingly,the ultimate strength of sample D, which contained 62% volumefraction of copper, was higher than sample E (monolithic copper) butthe strain to failure was only 20% of the sample E. This is perhapscaused by restriction of plastic deformation in the copper layers due tostrong bonding with adjacent WCCo layers.

Table 3Mechanical properties. *The values of bulk materials (ID: F and G) were taken from Refs.[2

ID Sample type Bending modulus (GPa) Apparent U

A WC12Co monolithic coating 2349 40145B Laminate coating 13356 17626C Laminate coating 17612 30173D Laminate coating 16624 42839E Cu monolithic coating 657 31761F* WC12Co bulk material [28] 570 1800G* Rolled copper [29] 3900.06

5362 M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 53585368a) Sample B

b) Sample Cc) Sample D

Fig. 5. Stressstrain curves of laminated coatings in three point bending test.3.3. Fracture surface of monolithic coatings

SEM micrographs of the fracture surfaces of samples A and E arepresented in Fig. 6ac and df respectively. The fracture surface ofsample A is very smooth and at indicating brittle failure (Fig. 6a).There exist many small pores, which were difcult to recognize inFig. 3b due to plastic deformation of binder phases during mechanicalpolishing. These pores could be an origin of lower strength of sampleA than a sintered bulk WC12Co material. Although some cleavagefacets of carbides can be observed as indicated by the arrow in Fig. 6c,fracture mainly occurred at the interfaces of WC/Co or WC/WC. In thefracture surface of sample E, the deformedmorphologies of individualcopper particles due to high velocity impact can be clearly recognizedin Fig. 3d. There are two fracture modes, one is brittle fracture atparticle boundaries as shown in Fig. 6e and the other is ductilefracture with elongated dimple formation (Fig. 6f). These local plasticdeformations contribute the plastic elongation of sample E during thebending test in Fig. 4b.

3.4. Fracture paths in laminate coatings

The typical cross-sectional images of crack paths at the centerregions in samples B, C, D are presented in Fig. 7ac, df, and girespectively. Please note that these images are not always taken fromthe same test pieces in Fig. 5. In all pictures, the tension side is placedat the bottom and thus cracking directions are from bottom to top. Inaddition, in all cases, the order of spray was from the top to bottom,and thus the bottom layer, which was sprayed at last, is WCCo, andthe top layer, which was sprayed at rst, is copper in all cases. Insample B, two cracks were formed in the WCCo layer in tension sideand it seems that those cracks merged into one (Fig. 7a), propagatedin a straight line (Fig. 7b), and then deected at the interface betweenWCCo and copper layers (Fig. 7c). The small load drops and non-linear stress increase observed in the stressstrain curve (Fig. 5),corresponds to the formation of cracks in the outer WCCo layer, andthe load reduction in a step like manner after maximum point can bethe results of the deections at the interface between layers. Since thethickness of the copper was very thin and was not strong enough toprevent crack propagation, the maximum load was the lowest(Fig. 5a). In the case of sample C, the main crack did not go straightbut curved (Fig. 7d). The plastic elongation of a copper layer (Fig. 7e)

8] and [29].

ltimate strength (MPa) Strain to failure (%) Hardness on surface (Hv)

0.170.02 14091250.340.07 1542620.310.06 15651280.550.04 1588831.870.74 9240.003 13002.630.06 and bridging of the crack propagation (Fig. 7f) can be observed. Fromthese pictures, it is understood that theWCCo layer at the crack frontfractured before the copper layer was broken and that the main crackpropagated and curved so as to connect to the crack formed inWCColayer in front due to the stress concentration around the cracks. Thisfracture behavior is more pronounced in sample D, which containedthe largest volume of copper (Fig. 7gi). The crack branching can beobserved in the second layer from the bottom. This probably occurreddue to the crack generations in the third WCCo layer prior to thefracture of the second copper layer. Large deection in the fourthcopper layer was also due to cracking in WCCo layer in advanceimplying the large fracture resistance in this sample (Fig. 7h and i). Itcan be clearly said that the highest fracture strength and strain of

)(c)

5363M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 53585368sample D among all samples (Figs. 4, 5) is due to higher fractureresistance from crack bridging, deection, and branching in copperlayers. Fig. 8a exhibits the microstructure in the Cu layer in front ofmain crack in sample D. The main crack propagated from the bottomand the vertical crack formed in the WCCo layer before the maincrack reached into it, and themain crackwas deected into Cu layer tothe direction parallel to the interface of two layers. Large shear

a

10 m

20 m

b

d e

Fig. 6. Fracture surface of: (a)(bdeformation is manifest from the distorted pore shapes at particleboundaries in the ligament (Fig. 8b). In addition, during crackpropagation in the Cu layer, it appears that the particle boundarieswere preferentially fractured and well bonded interfaces stillremained bonded and bridging the crack opening in the crack tip(Fig. 8a). The delamination can be also observed at the interfacebetween the Cu and WCCo layers in front of the deected crack tip.

In Fig. 9, the SEM images of the fracture surfaces of laminates sampleB (Fig. 9ac) and sample D (Fig. 9df) are shown. In sample B, thefracture surface is smooth indicating the occurrence of brittle failure.The thickness of the copper layer is about 20 min Fig. 9b and theplasticdeformation cannot be observed clearly (Fig. 9c). Since the Cu layers aresubstantially thin comparedwithWCCo, their plastic deformationwasrestricted by the neighboringWCCo layers, which resulted in the verylimited elongation of those layers. In sampleD, the fracture surface has astepdue to crack deection (Fig. 9d) and substantial plastic deformationcan be found in the Cu layers (Fig. 9e,f). This elongation is the reason ofhigh strain to failure of sample D observed in the bending test as shownin Fig. 5. Fig. 10 is the sequential images of fracture propagation from(a) to (f) during the threepoint bending test of sampleD. Before loading,pre-cracks can be observed on the right side of the fth and seventhWCCo layers (dotted circle in Fig. 10a). As the load increased, a verticalcrack was initiated in the outer three layers of WCCo/Cu/ WCCo(Fig. 10b). Right after the crack initiation, delamination occurred at theinterface of between the fourth and fth layers of Cu and WCCo(Fig. 10b,c), and the fourth layer bridged and prevented the main crackfrom propagating straightly (Fig. 10c). The delamination propagatedalong the interface and connected with the pre-existing crack on theright side in the fthWCCo layer (Fig. 10c). At the center region, a newvertical crack also formed from the delamination (Fig. 10d). Finally, thefourth Cu layer was fractured and then sixth Cu layer was broken(Fig. 10e). The sample lost most of load sustainability (Fig. 10f). It isconcluded that the laminate coatings with higher volume fraction of Cuexhibited better fracture strain because the WS Cu layers has betterresistance for crack propagation than WCCo due to higher elongationto failure.

1 m 250 nm

5 m

c

f

10 m

sample A, (d)(e)(f) sample E.4. Discussions

4.1. Fracture strength

Measured values of the apparent bending strength, which isdened as the peak value in the stressstrain curve, are comparedagainst Cu vol.% in Fig. 11a. Black symbols stand for the monolithiccoatings of WCCo and Cu, respectively and white circles for thelaminates. The error bars represent the standard deviation of theobtained values. The laminate with the lowest metal fraction (19.4%)exhibited the lowest strength, and as the volume fraction of metalincreased, the bending strength increased. The highest value wasobtained for the highest amount of Cu in laminates. Also it isimportant to note that the highest value among the laminates isalmost comparable to the monolithic WCCo and higher than themonolithic Cu coating. When a simple mixing law is considered, alaminate with higher amount of WCCo, thus lower amount of Cu isexpected to have higher strength. This discrepancy may be caused bytwo factors. One is the initial cracks in WCCo layer in the laminates.As can be observed in Fig. 3, most of the as-deposited coatings containvertical cracks in the WCCo layers (Fig. 3a,c,e). It is apparent thatstress concentration occurs during loading at the tip of such initialcracks and the strength of entire body is reduced signicantly. That iswhy the fracture strength of sample B (vm=19.4%) is substantiallylower than the monolithic sample A. The other factor is the constraintof plastic deformation in Cu layer by the intact WCCo layers. Theselaminates showed much lower strain to failure (Table 3) indicatingthe suppression of plastic deformation in the laminate samples. InFig. 11, the linear least square t for the laminates is also shown andthe bending strength at vm=100% can reach to 658 MPa. This value

5364 M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 53585368a

100 m

b

d ecan be considered as the effective constrained strength e of thecopper layer. By dividing it by the strength of the monolithic Cucoating m, e/m=2.1 is obtained. Hwu et al. [30,31] carried out thew-notched tensile tests of the constrained metal mono-layer in themetal/ceramic laminates fabricated by diffusion bonding processes.They reported the ratio of e/m=1.81, 1.86, and 1.72 for Al/Al2O3,Cu/Al2O3, Ni/,Al2O3, where e and m corresponds to the effectiveyield strength of the constrained metal layer and the metal foil yieldstrength (unconstrained), respectively. The constrained factorobtained in the present study has similar values to them. Thus itcould be concluded that the high bending strength obtained in sampleD (vm=62.0%) is attributed to the plastic constraint of the Cu layer bythe intact WCCo layers.

Failure strain is also plotted in Fig. 11b as a function of Cu vol.%. Inall laminates, the strain to failure is larger than the monolithic WCCosample. As the volume fraction of copper increases, the value tends toincrease. Although the sample C has higher volume fraction andthicker copper layers than sample B, it has almost same value withsample B. As can be seen in Fig. 5, samples B and C sometimesexhibited the step like reduction of load after themaximumpeak. Thisimplies that a main crack was bridged and deected so as to connectthe pre-existing cracks in WCCo layers when it propagated, whichresulted in larger strain to failure. But it is not always the case and

200 m

200 m

g h

Fig. 7. Cross section after three point bend test of (a) (b)50 m 50 m

c

fsometimes load quickly dropped like brittle failure. In addition, themetal layer thickness is not perfectly uniform in one sample, and thusthe thinner metal regions might be preferentially fractured at thepropagation of amain crack due to limited plastic deformation leadingto lower strain to failure. Perhaps these scatterings in fracturebehaviors would be the reason for similar values in strain to failurebetween samples B and C.

4.2. Work of fracture

Thework of fracture is plotted in Fig. 12 for the volume fraction ofCu layer. When the volume fraction of Cu is small, is lower than themonolithic WCCo. As the volume fraction increases, monotonicallyincreases, and sample D, which contains thickest Cu layers, exhibited2.53 times higher than the monolithic WCCo. Thus signicantimprovement of toughness can be achieved by laminate structures.This higher is attributed to the ductile elongation of Cu layers andbridging at the main crack as observed in Figs. 7 and 10. However, itappears that the size of bridging zone is limited in the present studydue to relatively low ductility of WS Cu layers as can be seen in Figs. 7and 10. Hence small scale birding conditions is applicable and theincremental toughening of the laminates for a metal layer can berelated to the work of stretch and fracture of the bridging layer [32

50 m

100 m

50 m

100 m

i

(c) sample B, (d)(e)(f) sample C, (g)(h)(i) sample D.

a20 m

5 m

b

Fig. 8. Crack bridging and deection in Cu layer (a) and higher magnication image of circle region (b).

5365M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 5358536834]. The total fracture energy of the laminates Glam can be estimatedas [32,33]

Glam = Gcer 1vm + vmu

0 u du 4a

100 m

200 m

WC-Co

WC-Co

WC-Co

Cu

Cu

Cu

WC-Co

Cu

Cu

WC-Co

b

d e

Fig. 9. Fracture surfaces of (a)(b)(c)where (u) is the nominal stress required to stretch a bridging metallayer by u, and u* is its failure stretch. Gcer and vm stand for the fractureenergy of WCCo layer and the volume fraction of metal layer,respectively. The integration term corresponds to the energy requiredto fracture a metal layer Gmet for the given thickness. The degree of10 m 5 m

4 m 50 m

c

f

sample B, (d)(e)(f) sample D.

plastic constraint and the resulted Gmet can be largely varieddepending on its layer thickness, and thus the Gmet is not constantfor the samples with different layer thickness. By following Mataga[35] and Marakaki et al. [36], Eq. (4) can be rewritten as,

Glam = Gcer 1vm + vmhm2Yw 5

w = uhm

0

Y

duhm

6

where hm is the thickness of metal layer and Y is metal yield strength.w is the work of fracture parameter which can be interpreted as the

a

400 m

400 m

400 m

400 m 400 m

400 m

b c

d e f

Fig. 10. Crack propagation behavior from side views during three point bending test of WCCo/Cu layered coatings (sample D). Dotted circle in (a) shows the pre-crack.

a

5366 M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 53585368bStra

in to

failu

re (%

)

Fig. 11. (a) Variation of bending strength as a function of Cu volume fraction and(b) variation of failure strain as a function of Cu volume fraction.normalized-fracture energy of a constraint metal layer [35]. Bysubstituting hm=(vm/1vm)hc into Eq. (5), Glam can be expressedas [36],

Glam = Gcer 1vm +w2hcY

v2m1vm

!: 7

where hc is the thickness of WCCo layer. By substituting Gcer=400 J/m2, hc=90 m (Table 3), and Y=300MPa (Fig. 4) into Eq. (7), thevariation of Glamwas predicted as a function of vm for variousw (Fig. 13).Gcer=400 J/m2 was chosen to be a half of the value for the monolithicWCCo coating (817 J/m2) because of the existence of crackings in theWCCo layers in the laminates. Since Eq. (7) diverges at vm=1, thepredicted value at highermetal fraction such as vm=0.9 are not realistic,but the trend of the predicted Glam appears to have a good agreementwith the experimental data especially for w=2. From Eqs. (4) and (5),the fracture energy of the metal layer Gmet can be written as,

Gmet = u

0 u du = hm

2Yw 8

For w=0.2, Gmet can be predicted as 711, 1308, and 3630 J/m2 forsamples B, C, and D indicating that thicker metal layers result intougher laminate coatings. In addition, it is manifest that larger u* isvery effective to increase fracture energy and thus it would be highlyeffective to improve the ductility of Cu layers by changing sprayconditions or by applying post heat treatments.Fig. 12. Variation of work of fracture as a function of Cu volume fraction.

w=0.3

w=0.05

w=0.2

w=0.1

Experiment

Fig. 13. Variation of predicted work of fracture Glam as a function of Cu volume fraction

5367M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 535853684.3. Property map

The variations of work of fracture were re-plotted for the fracturestrength of samples BE in Fig. 14. As elaborated so far, sample D(vm=62%) has far better work of fracture and moderately improvedstrength over sample E (monolithic WCCo). Thus it clearly demon-strates the advantages of laminateWCCo/metal coatings. On the otherhand, the laminates which contain the lower amounts of Cu (samples Band C) did not exhibit any advantages over the monolithic WCCo dueto the existenceof pre-cracks in theWCCo layer in as-sprayed state andthe limited plastic deformation of Cu layer. Even for no degradation inthe WCCo layers, it has been reported that the crack extensionresistance of laminates is enhanced as the metal layer thicknessincreases and that there is no benet for thinner metal layers whenthe volume fraction of metals is xed [14,37,38]. However, for coatingsfor wear resistance, it is obvious that thicker metal layers lead toreduction of total wear resistance of the coating because of poor wearresistance of soft metals. Thus there should be an optimized volumefraction ofmetal layers and thickness ratio ofmetal and cermet layers interms of overall performance of the coating.

In the present study, the interfaces between Cu andWCCo have astrong bonding and the delamination between layers was onlyobserved near the neutral axis. There are two most effectivetoughening mechanisms in laminates. One is absorption of thefracture energy with highly ductile layers. In the present study, this

and w.mechanism improved the mechanical properties of the coatings. Theother is the introduction of debonding and/or sliding at the interfaces

Fig. 14. Distribution of work of fracture and fracture strength.among layers. Clegg [39,40] developed ceramic multilayers with veryweak interfaces and demonstrated their extremely high toughness.Folsom et al. [41] provided a theoretical explanation of themechanisms and the experimental demonstrations [41,42]. It is ofgreat challenge for thermal spray process to introduce such extremelyweak interfaces because moderate bondings are required to pile upsprayed particles. But if such weak interface were introducedsuccessfully, it should be possible to form very tough coatings againstthe large impact loadings. On the other hand, such weak interfacesmay cause the spallation of layers and thus the interfacial strengthwould need to be optimized for practical applications.

5. Conclusions

WCCo/copper multilayer coatings containing 8 layers have beendeveloped by warm spray deposition. While the coatings with lowervolume fraction of copper showed no benecial feature in mechanicalperformance, the coating with higher volume fraction of copper (62%)exhibited more than two times higher work of fracture andmoderately better bending strength than the monolithic WCCocoatings. The mechanisms of the improved mechanical propertieswere attributed to the ductility of WS copper layers and the plasticconstraint by the intact WCCo layers. The total fracture energy of thelaminates was predicted by the theoretical model for the variation ofwork of fracture parameter w. The experimental data showed goodagreement with the prediction for w=0.2. The analysis alsosuggested the effectiveness of thicker metal layers to improve thetoughness of the laminate coatings. Poor work of fracture observed inthe laminates with lower volume fraction of copper, was possibly dueto the existence of initial crackings in WCCo layers which wasintroduced during spraying. Although the mechanisms to form suchinitial crackings remain to be understood for further improvements, ithas been concluded that cermet/metal laminate coatings can be onealternative approach to enhance the mechanical properties of thermalsprayed cermet coatings, especially for the spray processes whichhave a capability to deposit ductile coatings such as warm spraying.

Acknowledgements

The authors greatly acknowledge Fujimi Incorporated (Japan) forproviding WCCo feedstock powder. This work was supported as apart of Fail-Safe Hybrid Materials project (20062010) in NationalInstitute for Materials Science Japan.

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5368 M. Watanabe et al. / Surface & Coatings Technology 205 (2011) 53585368

Multilayered WCCo/Cu coatings by warm spray deposition1. Introduction2. Experimental procedure3. Results3.1. Microstructural characterization of coatings3.2. Mechanical properties of coatings3.3. Fracture surface of monolithic coatings3.4. Fracture paths in laminate coatings

4. Discussions4.1. Fracture strength4.2. Work of fracture4.3. Property map

5. ConclusionsAcknowledgementsReferences