Effect of microstructure on abrasive wear behavio

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Wear268 (2010) 1309–1319Contents lists available at ScienceDirectWearj o u r n a l h o m e p a g e:w w w.e l s e v i e r.c o m/l o c a t e/w e arEffect of microstructure on abrasive wear behavior of thermally sprayed WC–10Co–4Cr coatingsKanchan Kumari a,∗,K.Anand a,Michelangelo Bellacci b,Massimo Giannozzi ba Materials Research Laboratory,GE India Technology Centre,EPIP-II,Whiteﬁeld,Bangalore560066,Indiab GE Oil&Gas,Nuovo Pignone SPA,Via Felice Matteucci,2-50127Florence,Italya r t i c l e i n f oArticle history:Received12December2008Received in revised form27January2010 Accepted2February2010Available online 10 February 2010Keywords:Thermal sprayingWC–10Co–4CrAbrasionWear mechanismMicrostructure a b s t r a c tThermally sprayed WC–Co coatings have been widely used in the coatings industry for its superior slid-ing,abrasive and erosive wear properties.In applications where corrosion resistance is also required in addition to wear resistance,WC–10Co–4Cr is the preferred coating composition.The coatings produced by different thermal spray processes exhibit a broad range of coating hardness,porosity and microstruc-tural features like grain size and volume fraction of individual phases.In this study,we have evaluated the coating microstructures of various WC–10Co–4Cr coatings produced from different spraying processes, such as high velocity oxy-fuel(HVOF)and pulsed combustion.The objective of our study is to explore the abrasive wear mechanism of WC–10Co–4Cr coatings in great detail,and determine how these mech-anisms are inﬂuenced by the coating microstructure.Dry sand rubber wheel abrasion test rig(based on ASTM G65)is used for evaluating the three-body abrasive wear properties of the coatings,using alumina as the abrasive material.The coating microstructural parameters including WC grain size,volume frac-tion,binder mean free path have been quantitatively measured,and correlation between abrasive wear behavior of coatings and its microstructural parameters is sought.This study shows that binder mean free path of carbides(which is a function of WC grain size and volume fraction)is a very important parameter affecting the abrasion resistance of good quality coatings(containing low porosity).The lower the binder mean free path,the higher is the abrasive wear resistance offered by the coating.This is because the abrasive wear mechanism of these coatings is dominated by preferential removal of the binder phase, followed by pullout of WC grains.© 2010 Elsevier B.V. All rights reserved.1.IntroductionWC–Co coatings have been extensively studied for the last two to three decades because of their superior wear properties,in slid-ing[1–4],abrasion and erosive wear conditions[5–15].They have been widely used for numerous industrial applications like air-craft,oil and gas,mining etc.in solving severe abrasion and erosion problems.In addition to the coated form,they are also used in sin-tered form for structural applications,for making components like cutting tools,dies,plungers,gears,bearings etc.,and their abra-sion behavior have also been widely reported in the literature [10,16–18].Wayne and Sampath[10]have looked at comparison of structure/property relationships in sintered and thermally sprayed WC–Co materials,and have shown that the same set of equations for abrasion and erosion resistance hold good for both sintered and thermally sprayed WC–Co materials,relating them to material’s hardness,fracture toughness and the Co binder content.∗Corresponding author.Tel.:+918040122566;fax:+918028412111.E-mail address:kanchan.kumari@(K.Kumari).Most of the thermally sprayed WC–Co coatings reported in the literature are produced by either air plasma spray process(APS) or various HVOF processes such as JP5000,DJ,Top Gun etc.,and the coating microstructure and wear behavior are described.How-ever,there are very few papers that compare and contrast the microstructure and wear behavior of coatings[5,6,8],which is addressed in this paper.The coating microstructure is a function of the starting feedstock powder,thermal spray gun used for deposit-ing the coating,and the spraying parameters followed[5,13,19]. The choice of the thermal spray gun largely governs(a)tempera-ture exposure of particles,(b)dwell time in the gun and(c)their deposition velocity.The effect of these parameters can vary the porosity levels in the coating,the extent of decarburization,the volume fraction,grain size and distribution of carbides.One of the important attributes of WC–Co coating microstruc-ture is the extent of decomposition of WC grains during spraying, which is a strong function of the particle temperature in theﬂame, and can result in undesirable W2C and W phases,in addition to pri-mary WC phase in the coating.This results in further dissolution of W and C in the binder,which is largely amorphous/nanocrystalline in nature,because of the very high cooling rates experienced by the0043-1648/$–see front matter© 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2010.02.0011310K.Kumari et al./Wear268 (2010) 1309–1319powder particles.This aspect of thermally sprayed WC–Co coating has been extensively studied in the literature[6,8,9,20].Hence,the resultant coating microstructure can have a much lower volume fraction of the wear resistant primary WC phase and a much higher volume fraction of the binder phase compared to the starting pow-der microstructure[1,6,8].In another version of thermal spraying, called cold spraying,where theﬂame temperatures are extremely low(<500◦C),the starting powder characteristics like crystalline Co binder and fully retained WC phase can be preserved in the coating very well[21].WC–Co coatings derive its wear resistant properties from the presence of high volume fraction of hard,wear resistant WC grains in a Co-based metallic binder phase.The presence of the metallic binder provides some toughness in the coating compared to pure ceramic coatings,however,the binder can exhibit some brittleness if high W and C are dissolved in the binder during spraying.We have extensive literature on the abrasive wear behavior of WC–Co coatings[6,7,14,22],while very few of them discuss the abrasive wear behavior of WC–10Co–4Cr coating composition[5,13],which is used in applications that demand some corrosion resistance in addition to wear resistance.In this study,we are interested in evaluating this coating microstructure,produced by off the shelf commercially available thermal spray processes,from reputed ven-dors,for applications in equipments that pump or compress gases. We also studied a few coatings produced by the air plasma process (APS),but they were highly porous,and hence,we chose not to describe them in this paper.Coatings produced by HVOF and APS processes have been studied extensively in the literature.Often the end user does not have control over the coating pro-cess,which is largely controlled by the coating supplier.However, it may be possible to predict the wear performance of coat-ings by evaluating the coating microstructures.The objective of this work is to explore the mechanisms of abrasive wear dam-age of WC–10Co–4Cr coating in great detail and determine how these mechanisms are inﬂuenced by the coating microstructure. In particular,it would be interesting to explore the scale of the microstructural damage relative to low dosage of abrasive par-ticles to understand the wear mechanisms better.The scale of the damage zone relative to the size of the microstructural fea-tures,such as WC grain size,binder mean free path is worth investigating.There have been several investigations on the abrasive wear behavior of WC–Co materials using micro-scale abrasion tests [7,17,23],tests based on dry sand rubber wheel abrasion(ASTM G65)[5,12–14,17,24],tests based on ASTM B611rotating steel wheel with curved vanes[6,18,25],pin/block-on-diamond disk abrasion tests[10,16],in either dry/wet conditions.Micro-scale abrasion tests uses an abrasive slurry,containing veryﬁne abra-sives(alumina,silica,diamond etc.)in the size range of1–10␮m in de-ionized water,to study the abrasion resistance of coatings pressed against a rotating steel ball.These tests representﬁeld con-ditions like in the mining industry where the coating is exposed to ﬁne abrasive contaminants,while the abrasive wear tests based on ASTM G65and ASTM B611look at the abrasive wear behavior of coatings using coarser abrasives like alumina,silica,typically in the size range of100–600␮m,which can be encountered for e.g. in the oil and gas industry.ASTM G65based tests are used in the dry condition,while the tests based on ASTM B611uses coarser abrasives in a slurry medium.In the abrasion tests with diamond disk,the pin/block is coated with the desired WC–Co coating,and the disk is typically covered with a resin bonded30-␮m diamond plate.Bozzi and de Mello[22]have looked at the abrasion wear behavior of WC–12%Co coating as a function of abrasive material hardness(using three different abrasives of silica,alumina and sil-icon carbide),while Stewart et al.[14]have examined the effect of particle size and nature of abrasive(alumina vs.silica)on abra-Table1Details of various coatings used in this study.Coating Coating process Thickness(␮m)Roughness(␮m)W1HVOF(JP5000)2400.157W2HVOF(JP5000)2200.118W3Pulsed combustion2000.151W4Pulsed combustion2700.093sive wear behavior of conventional WC–17Co and nanostructured WC–15Co coatings.The coatings studied here had to meet two speciﬁc requirements for application in dry gas reciprocating compressors:(1)produce low sliding wear of the polymer piston rings and(2)resist abrasion byﬁne hard particles like alumina,silica,iron oxides etc.,that can enter into the compressor,along with the gases.In this work,we report the abrasion resistance of the various WC–10Co–4Cr coat-ings againstﬁne alumina particles in the size range of20–70␮m (average size50␮m),since they have the highest hardness value among the various abrading foreign particles like alumina,silica, iron oxides etc.In our study,we have used modiﬁed dry sand rubber wheel abrasion test rig(based on ASTM G65)to study the three-body abrasive wear behavior of WC–10Co–4Cr coatings using alumina particles.The main reason we have chosen alumina as the abrasive is because alumina can be present in the gases to be compressed as a contaminant,for e.g.in propylene or it can enter during pressure swing adsorption of hydrogen separation/puriﬁcation.Though the actual size of the foreign particles in the gases would be smaller (∼10␮m),we were forced to use50␮m average size alumina par-ticles,since smaller particle sizes(alumina with average particle size of10␮m and25␮m respectively that were also tested)had the tendency to clog the feeder and the nozzle of the machine. Hence,we have used the smallest size of the alumina particles that couldﬂow continuously through the feeder of the dry sand rubber wheel equipment.We wanted to simulate the three-body abrasion condition between the coating,polymer piston ring and the alumina particles in the dry condition,and hence,dry sand rubber wheel test rig was a good option to use.The material rubber is much closer to simulating the polymer material compared to steel wheel or ball,as used in micro-scale abrasion tests based on ASTM B611.In addition,both of these tests use abrasive in a slurry medium,while our requirement was to test the coatings in the dry condition.Also,the preference for the rubber wheel abrasion tests over micro-scale abrasion tests comes from the fact that our tests are more aggressive and produces signiﬁcant wear depth, and hence,can differentiate various coatings much better.2.Experimental procedure2.1.Coating preparationThermally sprayed coatings of composition WC–10Co–4Cr, were obtained on4140steel substrates of different sizes: 3in.×1in.×0.25in.(abrasion test),1in.×1in.×0.25in.(X-ray diffraction,metallography),from reputed external vendors in the industry,using HVOF and pulsed combustion processes.We have neglected the coatings produced by the APS process in this study since they had very high porosity.The coatings were all ground and polished to obtain a surface roughness in the range of0.1–0.2␮m Ra,since the coating application required the coatings to have a smooth surfaceﬁnish.Higher coating roughness can result in much higher wear of the polymeric piston ring materials,and hence,all the coatings were ground and polished.They had a coating thick-ness in the range of200␮m after theﬁnishing operation.Table1 shows a list of the coatings that are investigated in this study.Coat-K.Kumari et al./Wear268 (2010) 1309–13191311 ings W1and W2were deposited by HVOF process using a JP5000gun,while coatings W3and W4were deposited using pulsed com-bustion process,that are not that commonly used in the coatingsindustry.The starting feedstock powder was broadly in the sizerange of−45+15␮m,and was prepared by agglomeration andsintering process.However,the exact details of starting feedstockpowder preparation,WC grain size in the powder,and sprayingparameters followed during the spraying process was proprietaryto the coating vendor.2.2.Coating characterizationSamples for metallography were prepared by cutting the smallertest samples in cross-section,mounting them in epoxy with vac-uum impregnation,followed by emery paper grinding,diamondpolishing andﬁnishing with0.03␮m colloidal silica suspension.The coating cross-sections were examined using Leica opticalmicroscope,equipped with clemex image analysis software,andFEI Quanta400scanning electron microscope.From optical micro-graphs,coating thickness of various coatings were measured(aslisted in Table1),and porosity was determined using image anal-ysis software from500×images.A total of10micrographs wereused to calculate the average porosity for each coating.Coatingparameters like WC grain size,volume fraction and mean free pathwere calculated from high magniﬁcation SEM images of coatingscross-section in the SEI mode(7000×for coatings W1,W3and W4,and20,000–25,000×for coating W2as it contained muchﬁner WCgrains).The high magniﬁcations were necessary to better revealthe WC grain boundaries.SE imaging in preference to BS imag-ing was used for getting the topographical contrast between theWC grains and the metallic binder.Quantitative metallography hasbeen obtained using the linear intercept method,as proposed byLee and Gurland[26],and the equations used are described below:(1)Tungsten carbide grain size:d WC=2V WC2N␣␣+N␣␤(1)(2)Contiguity of tungsten carbide:C=2N␣␣2N␣␣+N␣␤(2)(3)Binder mean free path:=d WC(1−V WC)V WC(1−C)(3)where V WC is the volume fraction of WC phase,(1−V WC)is the volume fraction of the binder phase,N␣␣is the average number of intercepts per unit length of test line with carbide–carbide inter-face,and N␣␤is the average number of intercepts per unit length of test line with carbide–binder interface.SEM micrographs were printed,and a large test grid of18vertical and15horizontal lines (with a grid size of the order of WC grain)printed on a trans-parency sheet was placed on top of the micrograph,to count the number of intersections with WC grains.V WC is the ratio of the number of intersections with WC grains to the total number of intersections.Similarly,N␣␣and N␣␤were calculated for each of the micrographs,and an average value of WC grain size,volume fraction and binder mean free path were calculated from a set ofﬁve representative micrographs for each of the coatings.For coatings W1,W3and W4at7000×magniﬁcation,a total area of 42␮m×41␮m was covered in each micrograph,while a much smaller area of6␮m×6␮m for coating W2at25,000×was covered in each micrograph.Though this represents a rather small portion of the total coated area,but various important observationscan Fig.1.Schematic diagram of dry sand rubber wheel abrasion test apparatus(ASTM G65)[24].be drawn using these microstructural parameters,and its effect on abrasive wear behavior of the coatings.We understand that the accuracy of microstructural parameters can be improved by using a higher number of micrographs,however,we limited ourselves to 5micrographs per coating.Microhardness tests on coating cross-sections were performed using a CLARK,CM-400AT machine with a300gf load and a dwell time of15s.Hardness values reported represent the average of 10individual indentations made on a coating cross-section.For phase analysis of various coatings,a Philips X-Pert X-ray diffrac-tor with Cu-K␣radiation was used.2Âscans were run between20◦and90◦using a step size of0.002◦and a dwell time of10s.The surface roughness of the various coatings was determined using a Zeiss-TSK Surfcom1800D surface proﬁlometer,using a diamond tip of radius2.5␮m.Both roughness and contour measurements of coatings could be performed using this equipment.2.3.Abrasive wear testingThe abrasive wear resistance of the various coatings was evalu-ated in three-body abrasive wear condition using a dry sand rubber wheel abrasion test rig based on ASTM G65[24].Fig.1shows the schematic diagram of the dry sand rubber wheel apparatus,as taken from this ASTM standard.While this standard proposes to use 200–300␮m size Ottawa silica sand as the abrasive,we have used alumina particles of50␮m average size,as per our requirement discussed in Section1.The3in.×1in.face of the coated sample was pressed against the rubber lined wheel,and is abraded by the ﬂowing particles of alumina causing three-body abrasion of the coating.Based on several trial experiments with various sizes of alumina abrasive,to get a continuousﬂow rate of particles,alu-mina of average particle size50␮m(with90%of particles falling in the range of20–70␮m as per particle size analysis)was used for all the tests.Fig.2shows the morphology of the alumina abrasive particles,which are highly angular in nature,and can cause severe abrasion of the coatings.Parameters like load and duration of test in the initial experi-ments were varied to get a measurable weight loss of coatings,and a load of210N and test duration of10min wasﬁnalized for all the tests.The rubber wheel was made of Neoprene rubber,and showed a hardness of Durometer A-60.Theﬂow rate with alumina particles was measured to be134g/min,and a test time of10min allowed a dosage of1.34kg of alumina to cause the wear of the individual coatings.This is a very high dosage of abrasive,and is meant to produce differential wear for the different types of coatings stud-ied.This resulted in a total sliding distance of1440m,with the1312K.Kumari et al./Wear268 (2010) 1309–1319Fig.2.SEM micrograph showing the morphology of alumina abrasive. rubber wheel rpm of200,and rubber wheel diameter of228.6mm, in10min test duration.No recycling of the abrasive took place and all tests were performed dry.The test samples were ultrasonically cleaned in acetone before and after conducting the tests,and the abrasive weight loss of the coating was determined.An average of 3–4samples were tested for each coating,and their mean values were reported as the abrasive wear data.The worn surfaces of various coatings were examined by SEM on the top surface,in the middle of the wear scars.Also,surface proﬁlometer was used to study the worn surface roughness,and the maximum depth of coating removal for each of the coatings. In order to understand the wear mechanisms better,a few exper-iments were done later on,for each polished coating,with a very low dosage(∼7g)of alumina abrasive,which were quite useful to reveal the abrasive wear mechanism of these coatings.3.Results and discussion3.1.Microstructural characterizationThe coating thickness for various coatings as measured from the coating cross-section,are reported in Table1.The table also shows the surface roughness values for the various coatings in their as-received conditions,as measured by the surface proﬁlometer.Most of the coatings obtained had a thickness of∼200␮m,and a surface roughness in the desirable range of0.1–0.2␮m.Fig.3shows the optical micrographs of the coatings cross-section.All of these coatings have quite low porosity.Coatings W1 and W2,produced by JP5000process,have a very high volume frac-tion ofﬁne,uniformly dispersed WC grains with a low porosity(1%), while coatings W3and W4,produced by the pulsed combustion process,though have low porosity,but have much lower volume fraction of WC grains which are not uniformly distributed in the coating.It is interesting to note that the coating W4show close to zero porosity.So,we can see that the coating microstructure is very much governed by the type of the thermal spray gun used for depositing the coating.In other words,coatings W1and W2 show a high retention of primary WC grains,while coatings W3and W4show a much lower retention of WC grains,and an increased volume fraction of the binder phase with dissolved W and C in it. Coating W2shows muchﬁner WC grains than observed in coat-ing W1.The porosity content of these coatings is determined using clemex image analysis software on optical images,and is presented in Table2.Legoux et al.[5]have reported the microstructures of var-ious APS and HVOF WC–10Co–4Cr coatings,with porosities in the range of5–16%for the APS coatings,and in the range of0.6–8.5%for various HVOF coatings,by varying the spraying parameters.How-Fig.3.Optical micrographs of WC–10Co–4Cr coatings:(a)W1;(b)W2;(c)W3;(d)W4.K.Kumari et al./Wear 268 (2010) 1309–13191313Table 2Microstructural parameters measured for various coatings.Coating Porosity (%)WC grain size (␮m)WC volume fraction Mean free path (␮m)W1 1.10.630.610.27W2 1.00.330.730.08W3 1.40.750.12 3.67W40.20.960.400.96ever,Liao et al.[6]have reported porosity in the range of 1%for the APS coating,and 0.6%for the HVOF coating.Zhao et al.[13]have reported porosity in the range of 0.2–4%for WC–CoCr coat-ing deposited by HVOF (DJ 2600)process.The porosity values we obtained are in the range of 1%,and hence,are dense,good quality coatings.For studying the coating microstructural features,such as grain size,volume fraction,binder mean free path etc.,SEM micrographs of the coating cross-sections are also obtained at 7000×,except for W2coating,where 20,000–25,000×magniﬁcation was needed to reveal the ﬁner WC grains (Fig.4).From this ﬁgure,we observe that coating W1shows a very wide distribution of WC grain size,which was not evident from the opti-cal micrograph.Coatings produced by HVOF process,W1and W2show high volume fraction of WC grains.It is interesting to see that coating W4,produced by the pulsed combustion process,has a bimodal distribution of WC grains (containing two sets of grain sizes;one in the range of a few microns,and another in the range of <1␮m),which is resolved only by going to high magniﬁcation SEM images.The ﬁner WC grains could not be resolved from the optical micrograph.Hence,the total volume fraction of carbide grains in coating W4is effectively much higher than in coating W3,since coating W3has very low volume fraction of more or less uniform WC grains,in the size range of 1␮m.Microstructural parameters,such as grain size,volume fraction and binder mean free path,for the various coatings have been obtained from several SEM images (Fig.4),by using Eqs.(1)–(3),and the obtained values are tabu-lated in Table 2,which also show their porosity values.We see that coating W2has an average WC grain size of 330nm,which is the smallest among the four coatings examined,and would fall in the category of nanostructured WC–Co–Cr coatings,which has been widely reported in the literature [14,19,27].The measured volume fraction of the WC grains vary a lot among the various coatings:it is lowest for coating W3(12%),followed by coating W4(40%),followed by coating W1(61%),and coating W2shows the highest volume fraction of 73%.The calculated WC grain size and volume fraction result in smallest binder mean free path of 0.08␮m forcoating W2,0.27␮m for coating W1,0.96␮m for coating W4,and a very large binder mean free path of 3.67␮m for coating W3.It is expected that the coating mechanical properties like hardness willFig.4.SEM SEI images showing details of microstructure of WC–10Co–4Cr coatings:(a)W1;(b)W2;(c)W3;(d)W4.1314K.Kumari et al./Wear268 (2010) 1309–1319Fig.5.Microhardness data for various coatings obtained at300g load.be a strong function of microstructural parameter like binder mean free path,since we know that WC phase hardness(∼2400–2500Hv) is much higher than the hardness of the binder phase(<1000Hv) since it is metallic.One of the other coating properties,fracture toughness,can also affect the wear performance of coatings,and it was of our inter-est to measure it.However,our coatings were relatively thinner (∼200␮m thick),and the effort to produce indentation cracking on the coating cross-section,for generating fracture toughness data, was not successful.For example,when we used a5kg load on coat-ing W4(the thickest of all the coatings observed),the indentation diagonal covered more than60%of the coating thickness,without producing any cracks.No cracking could be seen with the use of lower loads.Hence,we were limited by the coating thickness to obtain any fracture toughness data on the coatings.The microhardness of the various coatings are obtained from the coatings cross-section using a standard load of300g.Fig.5shows the average hardness for all the coatings,which is obtained from10 measurement values.The minima and maxima values reported for each of the coating,is indicated by error bars.The coatings obtained by JP5000process,W1and W2,are very much superior in their hardness compared to coatings produced by pulsed combustion process.As per the literature,most of the thermally sprayed WC–Co family of coatings,sprayed by HVOF process have hardness in the range of900–1330Hv[5,6,13].The hardness of most of our coat-ings is in agreement with the reported literature.The large spread in the coating hardness of thermal sprayed coating is a function of the non-homogenous coating microstructure.Among the coatings produced by the pulsed combustion processes,slight improvement in the microhardness of W4coating is seen compared to W3coat-ing,since it has negligible porosity and higher volume fraction of WC grains.However,these coatings are lower in hardness com-pared to coatings produced by JP5000process,and could perform differently in abrasion tests.Fig.6shows the X-ray diffraction patterns for all the coatings studied.The most important peak seen in these patterns is that of WC from the starting powders used for coating deposition.How-ever,the peaks for W2C and W are seen because of decarburization. The extent of decarburization is different in coatings obtained by different processes like HVOF(JP5000gun)and pulsed combus-tion.Coatings obtained by JP5000gun show low decarburization(as indicated by very small intensity W2C peak).However,the coatings obtained by pulsed combustion process(W3and W4)have signif-icant amounts of W2C phase,in addition to the primary WC phase. Coating W3,also produced by pulsed combustion,shows the max-imum extent of decarburization(signiﬁcant presence of elemental W peak,in addition to W2C peaks),and the presence of WC phase is reduced to a minimum.This explains the microstructure observed in our coatings where we saw signiﬁcant reductions in volume frac-tion of hard,wear resistant phase for coatings W3and W4,and a corresponding increase in the volume fraction of the binder phase. The absence of Co peak in these XRD patterns indicate that the Co binder is amorphous or nanocrystalline innature.Fig.6.XRD patterns for various WC–10Co–4Cr coatings examined in this study.There is vast literature available on the XRD patterns of starting WC–Co family of powders,and their thermally sprayed coat-ings[1,6,9,20],which conﬁrms similar observations.It has been reported[14,19,27]that WC–Co coatings from nanostructured powders have higher extent of decarburization.In these papers, the coatings are produced by Top Gun,DJ etc.,but in our study, coatings are produced by JP5000gun,and weﬁnd that coating W2 containing330nm WC grains show slightly lower intensity of W2C phase,compared to the coating W1showing higher WC grain size (630nm).This illustrates that the coating vendor has a very good control on the spraying parameters using the nanostructured pow-ders to control decarburization,and conﬁrms to the possibility of reducing nanocomposite degradation with liquid fuel systems like JP5000gun,as pointed out by Stewart et al.[14].Yandouzi et al.[21] have shown the phase evolution of WC–Co based coatings obtained by cold spraying process,where theﬂame temperature is very low (<500◦C),and hence the resulting coating has crystalline peaks of Co binder,in addition to WC peaks.No W2C peaks are seen here, which suggests that decarburization is a result of the higherﬂame temperatures encountered in the thermal spray guns.3.2.Abrasive wear behavior of coatingsThe weight loss and standard deviation data for the various coatings tested in abrasion is presented in Fig.7.We observe that coatings obtained by JP5000process,W1and W2,have much lower weight loss in abrasion testing,compared to the coatings obtained by pulsed combustion.We obtained a very low standard devia-tion for coatings W2and W4,while coatings W1and W3showed high standard deviation,as indicated by error bars.The large dif-ference between the abrasive wear loss of coatings W3and W4can be explained in terms of their volume fraction and mean free path of carbide grains,as reported in Table2.Since the coatings were all ground and polished to a surface roughness of∼0.1␮m,the corresponding coating wear loss val-ues were small,compared to the scenario if the coatings were tested in the as-deposited condition with roughness of3–5␮m.We observed a wear coefﬁcient of∼2×10−5mg/(N-m)for the coatings obtained by JP5000process when abraded by50␮m alumina par-ticles,while Stewart et al.[14]have observed a wear coefﬁcient of ∼4×10−4mg/(N-m)when abraded by125–150␮m alumina parti-cles in their modiﬁed“dry sand rubber wheel technique”,since the。

Brake Drums制动鼓

Corresponding author: Yu Zhao E-mail:zhaoyu_0223@ Journal of Bionic Engineering Suppl. (2008) 20–27Frictional Wear and Thermal Fatigue Behaviours of BiomimeticCoupling Materials for Brake DrumsYu Zhao1, Lu-quan Ren1, Xin Tong2, Hong Zhou2, Li Chen11. Key Laboratory of Terrain-Machine Bionics Engineering (Ministry of Education, China),Jilin University, Changchun 130022, P. R. China2.The Key Lab of Automobile Material (Ministry of Education, China),Jilin University, Changchun130022, P. R. ChinaAbstractIn order to enhance frictional wear resistance and thermal fatigue resistance of brake drums, two kinds of biomimetic coupling materials are prepared by laser surface melting and laser coating technology respectively. These materials are all compounded by base metal and biomimetic coupling units, and have different coupling characteristics, such as the variation of the unit shape, different microstructures and hardness, different chemical compositions of units. The frictional wear resistance and thermal fatigue resistance of biomimetic coupling materials and gray cast iron used for brake drum are compared. The results indicate that frictional wear resistance and thermal fatigue resistance of biomimetic coupling sample is better than that of untreated sample, and among the biomimetic samples, laser coating treated sample has superior resistance to wear and thermal fatigue comparing with laser melting treated sample.Keywords: frictional wear, thermal fatigue, biomimetic coupling, brake drum, gray cast ironCopyright © 2008, Jilin University. Published by Elsevier Limited and Science Press. All rights reserved.1 IntroductionBrake drum is one of the most important compo-nents in brake system of trucks. According to the prin-ciples of braking, brake drum receives powerful pres-sure and static or kinetic friction force from brake disk. Especially the sliding velocity and specific pressure increase with the speed of automobiles, which result in heavier wear and higher temperature. It is reported that the surface temperature of brake drum can reach as high as 900 ÛC, and the alternating heating and cooling caused by frequent braking also lead to initiation and propagation of thermal cracks[1í4]. Fig. 1 shows the wear and thermal fatigue characteristics of a disabled brake drum. The above damages to brake drum do a lot of harm to the braking performance and even the safety of automobiles.Compared with frictional materials, the material for brake drum has a single category and slow development. There are mainly two kinds of braking materials, cast iron and aluminium base composite. Gray cast iron has long been used because of its low cost and desirable properties such as low melting point, good fluidity and cast-ability, excellent machinability etc. Concretely, HT150 and HT200 are more commonly used by in-boards, while outboards prefer gray cast iron grade 250and high-carbon gray cast iron. However, these materialsFig. 1 Thermal fatigue cracks and wear patterns of a failbrake drum.Zhao et al.: Frictional Wear and Thermal Fatigue Behaviours of Biomimetic Coupling Materials for Brake Drums21can not satisfy applications due to their low hardness and intensity. In order to refine and increase pearlite micro-structure, some alloy elements are added to cast iron. For example, the high carbon low alloyed cast irons wherein carbon content is more than 7.36 % and silicon content is less than 1.5 %, containing molybdenum, cuprum or nickel are widely applied in Britain and America; while in Germany the added alloy elements are cuprum, chromium and stannum[5,6]. Although the intensity and hardness of cast iron can be enhanced by this method, yet its disadvantages are weak resistance to thermal fatigue and high waste of alloy elements. For lightening weight and improving thermal conductivity of braking material, the ceramic reinforced aluminum base com-posite has become a new researching point, and the usually used hard grains are SiC, TiC, Si3N4 and Al2O3. But the high cost and difficult selection of brake disk limit their applications[7í11]. Therefore, it is urgent to develop new materials for brake drums.Bionics is a science that imitates the principles of bio-systems to build new technological systems of use-ful components. Over the past two decades, bionics has had a profound influence on engineering material re-search, because the unique structures, compositions and correspondingly excellent properties of biological and natural materials gave researchers many clues to im-prove the properties of engineering materials or increase the reliability of structural components. In fact, biology can not only adapt to the environment, but also satisfy its survival with the lowest energy cost. Therefore, the new components or materials, which are designed and manufactured by biomimetic methods, also have ad-vantages such as high efficiency and low consumption of energy. Research on biomimetic non-smooth surface found that the bulldozer blades with such surfaces can reduce sliding resistance against soil[12]. In recent years, the technology of biomimetic non-smooth surface has been applied to brake disks for trains, brake drums for trucks and a lot of hot-work dies, aiming to solve the problem of frictional wear, abrasive wear or thermal fatigue[13,14]. The efforts are all about changing the morphologies of biomimetic non-smooth surface such as shape, size and convex-concave of non-smooth units. With further investigation, we discover that some com-plex problems in engineering systems are difficult to solve completely by our invented biomimetic method which only has single changing factor.In this research, two biomimetic coupling materials are prepared by laser surface melting and laser coating technology respectively. These materials are all com-pounded by base metal and biomimetic coupling units, with different coupling characteristics, such as varied units’ shape, dissimilar formed structure between base metal and units, different chemical compositions of units and so on. The frictional wear and thermal fatigue re-sistance of the coupling materials and gray cast iron are compared.2 Experiments2.1 Experimental materialsA gray cast iron codenamed HT200, which is widely used for brake drums, is applied as the base metal material whose microstructure consists of flake graphite and pearlite shown in Fig. 2. According to different molding technology, biomimetic coupling units are di-vided into Laser Surface Melting Coupling Unit (LMU) and Laser Coating Coupling Unit (LCU) respectively. Among them, the LCU is made through laser coating of Fe-based self-fluxing alloy (Fe30A). The chemical compositions of the materials are given in Table 1 and Table 2.Fig. 2 Microstructure of HT200 gray cast iron.Table 1 Chemical compositions of HT200 gray cast iron (wt%) Compositions C Si Mn P S Cu Cr Fe Content 3.250 1.5700.920 0.060 0.059 0.5000.270Bal.Journal of Bionic Engineering (2008) Suppl.22Table 2 Chemical compositions of Fe30A alloying power (wt%)Compositions C B Si Cr Ni Mo FeContent 0.5–0.6 1.0–2.0 2.0–3.512–14 30–37 4.0–6.0Bal.2.2 Sample preparationA solid state Nd-YAG laser of 1.06 μm wavelengthand maximum power of 300 W is employed for lasersurface melting and laser coating. The base metal ismechanically polished, using progressively finer gradesof silicon carbide impregnated emery paper. By con-trolling the laser tracks, LMU is formed to latticed shapewith spacing of 3 mm. For processing LCU, the basemetal has some notches of 1 mm width and 1 mm depth,which are used for filling Fe-based self-fluxing alloy.Laser tracks are along the longitude direction of notch,and a shielding gas of argon is used to minimize con-tamination of the treated surface. Other laser processingparameters are listed in Table 3.Table 3 Laser parameters for processing biomimetic unitsPulse duration FrequencyElectriccurrentLaser spotdiameterScanningspeed7.0 ms 10 Hz 120 A 0.8 mm 1 mm·sí1Wear samples of 14 mm length, 10 mm width and 10 mm thickness and thermal fatigue samples of 40 mm length, 20 mm width and 6 mm thickness are cut by an electric spark machine, and a 3 mm diameter round hole is drilled at one side of each thermal fatigue sample, so that they can be fixed onto the plate of the thermal fa-tigue experimental machine. According to different materials, samples are divided into three kinds: Un-treated Sample (US), Laser Surface Melting Sample (LMS), and Laser Surface Coating Sample (LCS).2.3 Experimental methodTransverse sections of biomimetic coupling units are cut parallel to the length direction, and standard methods of metallography are followed. The micro-structure is characterized by JSM-5500LV scanning electronic microscope. A XTL-2400 stereomicroscope is used to study the initiation and propagation of thermal cracks. Phases and compounds formed in the coating zone are identified by D/max-RC X-ray Diffraction (XRD) with Cu KĮ radiation operated at a voltage of 40 KV, a current of 40 mA, and a scanning rate of 40Û·miní1. The alloying element’s distribution of mi-cro-zones is analyzed by an EDAX-Falcon Energy Diffraction Spectrometer (EDS) which is equipped on a JSM-5500LV scanning electronic microscope. A Vick-ers Microhardness Tester (model 5104, manufactured by Buehler Co. Ltd., USA) is used for the microhard-ness measurement.Dry sliding tests are performed using an MM-200 block-on-ring wear tester (made by the Xuanhua Testing Machine Factory). All the friction and wear tests are carried out at room temperature. Experimental samples slide on a rotating GCr15 steel ring. The GCr15 steel has an average hardness of 61–63 HRC and the sizes of 16 mm inside diameter, 50 mm outside diameter and 10 mm thickness. A normal load of 8 kg is used for wear tests. The rotational speed of the ring is 400 r·miní1 and the sliding time is 30 min. Before each test, the surface of the steel ring is polished to a roughness of about 0.1 mm, while samples were ultrasonically cleaned in an-hydrous alcohol and dried before and after wear tests. The mass loss is measured by a sensitive electronic balance with an accuracy of 0.0001 g. The difference in mass of three test blocks before and after the experiment gives the average mass loss. To assist the analysis of wear mechanisms, the worn surfaces of block specimens are examined by LEXT-OLS 3000 OLYMPUS laser confocal scanning microscope.Thermal fatigue tests are carried out by a self- restrain thermal fatigue testing machine. The samples are heated by a high temperature electric resistance furnace, and cooled by running water. This machine can record thermal cycle times automatically. The furnace temperature is controlled by thyristor control of type (KSY-12-16), and the sample temperature is monitored by a thermocouple attached to the sample at the centre of its length. A complete thermal cycle includes heating to 600±5 ÛC in 70 seconds, and then cooling to 25±5 ÛC in 3 seconds. The samples are free from any externally applied load, and are taken out to observe the cracks every 200 cycles, then the thermal fatigue resistance of samples is evaluated.Zhao et al.: Frictional Wear and Thermal Fatigue Behaviours of Biomimetic Coupling Materials for Brake Drums233 Results and Discussion3.1 Structure characteristics of biomimetic couplingunitsFig. 3 shows the cross-section morphology of biomimetic coupling unit. Observation indicates that the unit contains two areas. One is the melted zone at the outer layer which is bright in Fig. 3, because cast iron is melted completely in this zone, and subsequently re-crystallized under a fast cooling rate. Due to thermal conduction, the materials around the melted zone takes solid phase transformation, thus this area is called transformed zone. Fig. 4 shows the SEM micrographs taken from the transverse section of LMU. It can be seen that the melted zone (see Fig. 4a) is composed of Fe3C, M and Ar, and the Fe3C is reticular. The microstructures of the transformed zone, as shown in Fig. 4b, consist of M, a little Ar and some un-dissolved G. The crystal grains in the whole unit zone are very fine and compact, and these results agree well with the findings of Shen and Li[15], which can be explained by the rapid cooling rate during solidification caused by laser processing method. Fig. 5 shows the SEM micrographs taken from the transverse section of LCU. It can be seen that the melted zone (see Fig. 5a) is composed of Ar, martensite, M7C3, M23C6, and Fe2B. Therefore, the unit is strengthened by solution strength and dispersion strength. Because the heating temperature of the transformed zone is between eutectoid reaction temperature and eutectic transforma-tion temperature, the microstructures in this zone of LMU and LCU are similar, and they all have martensite, Ar and un-solved graphite. The X-ray diffraction results also identify the above microstructure.Fig. 3 Cross-section morphology of biomimetic coupling unit.Fig. 4 Microstructure of the LMU.Fig. 5 Microstructure of the LCU.Journal of Bionic Engineering (2008) Suppl.243.2 Micro-hardness of biomimetic coupling unitsThe curves of hardness distribution in the LMU and LCU are shown in Fig. 6. The reported micro-hardness is measured along the direction of laser irradiation (from the top of unit to the bottom), which are obtained at dif-ferent locations on the polished and etched cross-sections of the biomimetic unit for each sample. The average micro-hardness of the substrate is 280HV, while that of biomimetic unit is much higher. It can be seen that the micro-hardness of LMU varies from 561 HV to 603 HV, and that of LCU varies from 567 HV to 632 HV. The micro-hardness of LCU is higher than that of LMU, which can be attributed to the solution of carbon, the formation of martensite, the refinement of structure, the large proportion of cementite, especially the solution strength and the precipitation of alloying carbide. Mi-cro-hardness is of particular importance in assessing wear resistance coatings, because it is considered as a good indication of coating integrity and wear resistance.Fig. 6 Micro-hardness distributing of LMU and LCU.3.3 Frictional wear resistance3.3.1 Wear mass lossThe wear resistance samples are divided into 3 groups. One group is untreated, the others are processed to biomimetic coupling samples by laser melting and laser coating technology respectively. Three samples of each group are chosen for wear tests. The average mass losses of different samples are shown in Fig. 7. It can be seen that the wear losses of the biomimetic coupling samples is 48.6% less than that of the untreated sample in the same sliding way against the steel blocks. Among biomimetic coupling groups, the sample treated by laser coating with Fe30A has less wear mass loss.Fig. 7 Wear mass loss of test samples.3.3.2 Wear patternFig. 8 shows the worn surface of US, LMS and LCS after sliding wear tests against the wearing ring GCr15. As seen in Fig. 8a, some materials are lifted and sheared off in the wear process; the wear grooves with irregular shapes are in various depth. Besides, macroscopic ob-servation reveals that some red-brown and black wear debris, whose chemical nature is Fe2O3 and Fe3O4[16], clearly appear on the worn surfaces and some materials adhere to the worn surfaces and other materials adhere to the steel counterparts. This phenomenon indicates that there is severe plastic deformation on the surface of US, some materials are pulled out and transferred to the counterpart surface resulting in adhesive wear. Deep grooves on the wear surface are generated as a result of cutting action, which is produced by wear debris. Some of the debris is iron oxides, which primarily originated from the contacting surface of GCr15 ring and are oxi-dized during the dry sliding wear process. Figs. 8b and 8c compare the morphologies of the worn surface of the biomimetic samples under the same sliding conditions. Compared with the US, the extent of the abrasive wear is lower and there is less red-brown and black wear debris on the steel counterpart observed by macrographs. The dominating wear mechanism is adhesive wear, whereas the abrasive and oxidation signs are not obvious. Moreover, the LCS displays better wear resistance than the LMS, because of its greater hardness.Zhao et al.: Frictional Wear and Thermal Fatigue Behaviours of Biomimetic Coupling Materials for Brake Drums25(a) US (b) LMS (c) LCSFig. 8 Worn surface of different samples after a sliding wear test.3.4 Thermal fatigue resistance3.4.1 The number and length of thermal cracksDuring thermal fatigue tests, the number of thermal cracks on each sample surface is counted after 2000 thermal cycles, and in the counting, if any crack is longer than 0.5 mm it is defined as one crack. Fig. 9 shows the number of thermal cracks, indicating that the number of cracks on US is always the most, then the LMS, and the LCS. Based on the above results, it can be deduced that biomimetic coupling units on sample’s surface have superior resistance to the initiation of thermal crack, the biomimetic coupling sample, especially formed by laser coating with Fe30A has higher resistance. Fig. 9 also shows the measured length of crack of different samples. It can be seen that after the same thermal cycles (2000), the length of crack on US surface is always the longest, then the LMS, and the LCS. I n fact, the biomimetic coupling samples, especially formed by laser coating with Fe30A, have the higher resistance to crack propa-gation. Based on the above number and length of cracks, it can be deduced that the thermal fatigue resistance of biomimetic coupling sample is better than that of un-treated sample, and among biomimetic samples, laser coating treated sample has superior resistance to thermal fatigue compared with laser melting treated sample.3.4.2 Crack initiation resistance of biomimetic coupling unitI t is discovered that macroscopical place where thermal crack initiated is mainly from the fixed hole and the edge of sample. After 50 thermal cycles, microcos-mic place where crack initiated is also observed, and from Fig. 10 it can be seen that the place is at the sharp point of graphite. Due to nearly no intensity, the graphite is just like a hole distributing in the matrix, thus a crack initiated there first. According to the microstructure in melted zone, where graphite is not found, the source of crack initiation disappears. Therefore, processing units on sample surface, especially at the edge of sample can decrease the number of thermal cracks obviously.Fig. 10 Microcosmic place of crack initiation.Fig. 9 Crack number and length of different test samples.Journal of Bionic Engineering (2008) Suppl.263.4.3 Crack propagation resistance of non-smooth unit Many micro-cracks extend along the graphite as their initiation source, and these micro-cracks connect each other to form a main crack. Observations of the propa-gation routes of main crack indicates that main crack propagates along the graphite and the direction of their shortest distance, showing that propagation of main crack is realized by bridge connections with mi-cro-cracks in front of it (see Fig. 11). In fact, biomimetic coupling unit not only causes such an advantageous inhibition to crack initiation, but also resists crack propagation. It can be explained as that, in the unit area there is no graphite phase, hence there is no bridge connection between main crack and micro-crack which grows relying on graphite. That is to say, the route of crack propagation is cut off by the unit.Fig. 11 Bridge connections between main crack andmicro-cracks.3.4.4 Effects of alloying elementsThough microstructure of biomimetic unit obtained by laser melting also has characteristics of no graphite phase and refined grain sizes, which can improve ther-mal fatigue resistance of samples, it is still worse than that treated by laser coating of Fe30A. It is well known that crack will lanate biomimetic unit, then gets further propagation when fatigue damages are bigger than the rupture strength of the unit. I n addition, the oxidation occurs badly during thermal tests, and repeated building and falling of oxide film at crack point promotes the destruction of the unit. So the higher strength and oxi-dation resistance of the unit, the better thermal fatigue resistance that biomimetic sample has. Due to laser coating, non-smooth unit is rich in Cr, Ni, Mo, austenite is strengthened by the solid solution of them, and eutec-tic carbide M7C3 has higher strength and hardness comparing with cementite Fe3C. On the other hand, Cr is an element which can improve heat-durability of cast iron, because it forms an oxide film rich in Cr with oxygen, like Cr2O3 or FeCr2O4 which can restrain oxi-dation again[17]. Based on the above reasons, thermal fatigue resistance of biomimetic coupling sample can be improved further by laser coating of Fe30A powder.4 ConclusionBiomimetic coupling unit contains melted zone and transformed zone. The melted zone of laser melting treated unit is composed of Fe3C, M and Ar; while the melted zone of laser coating treated unit consists of Ar, martensite, (Cr,Fe)7C3, (Cr,Fe)23C6, Fe2B and some austenite is strengthened by the solid solution of Cr, Ni, Mo elements. The heating temperature of transformed zone is between eutectoid reaction temperature and eutectic transformation temperature, and only solid-state phase change occurs in this zone.The average micro-hardness of substrate is 280 HV, while that of biomimetic unit is much higher. The mi-cro-hardness of laser melting treated unit varies from 561 HV to 603 HV, and that of laser coating treated unit varies from 567 HV to 632 HV.Frictional wear and thermal fatigue resistance of biomimetic coupling sample is better than that of un-treated sample, and among the biomimetic samples, laser coating treated sample has superior resistance to wear and thermal fatigue comparing with laser melting treated sample.AcknowledgementThis work is financially supported by the Project 985-Automotive Engineering of Jilin University, the National Natural Science Fund of China (Grant no. 50635030) and the Science and Development Founda-tion of Jilin (Grant no. 20060196).References[1]Naji M, Al-Nimr M. Dynamic thermal behaviour of brakesystem. International Communications in Heat and Mass Transfer, 2001, 28, 835í845.[2]Mackin T J, Noe S C, Ball K J, et al. Thermal cracking indisc brakes. Engineering Failure Analysis, 2002, 9, 63í76.[3]Lejeail Y, Kasahara N. Thermal fatigue evaluation of cyl-Zhao et al.: Frictional Wear and Thermal Fatigue Behaviours of Biomimetic Coupling Materials for Brake Drums27inders and plates subjected to fluid temperature fluctuations.International Journal of Fatigue, 2005, 27, 768í772.[4]Dai W S, Ma M, Chen J H. The thermal fatigue behaviourand cracking characteristics of hot-rolling material. Materi-als Science and Engineering A, 2007, 448, 25í32.[5]Yamabe J, Takagi M, Matsui T, Kimura T, Sasaki M. De-velopment of disc brake rotors for trucks with high thermal fatigue strength. JSAE Review,2002, 23, 105í112.[6]Yamazaki T, Shibuya T, Jin C J, Kikuta T, Nakatani N. Lin-ing of hydraulic cylinder made of cast iron with copper alloy.Journal of Material s Processing Technol ogy, 2006, 172, 30í34.[7]Noguchi M, Fukizawa K. Alternate materials reduce weightin automobiles. Advanced Material s and Processes, 1993, 143, 20í26.[8]Shaw M L, Tsang P H S, Rhee S K. Study of the friction andwear behavior of aluminum composites sliding against polymer composites.I n: Ludema K C, Bayer RG (eds), Proceedings of International Conference on Wear of Mate-rials,ASME, New York, 1991, 167í175.[9]Hosking F M, Folgar Potillo F, Wunderlin R, Mehrebian R.Composites of aluminum alloys: fabrication and wear be-havior. Journal of Materials Science, 1982, 17, 477í498. [10]Long T T. Mechanical properties and wear resistance of6601 alloy reinforced with a hybrid of Al2O3 fibers and SiC whiskers. Transactions of Japan Institute Metals, 1988, 29, 920í927.[11]Alpas A T, Zhang J. Effect of SiC particulate reinforcementon the dry sliding wear of aluminum-silicon alloys (A356).Wear, 1992, 155, 83í104.[12]Ren L Q, Han Z W, Li J Q, Tong J. Effects of non-smoothcharacteristics on bionic bulldozer blades in resistance re-duction against soil. Journal of Terramechanics, 2003, 39, 221í230.[13]Chen B C. Morphology of Vehicle Running Gear and Theoryof Reducing Adhesion and Scouring Soil, Mechanical I n-dustry Publisher, Beijing, 2001. (in Chinese)[14]Ren L Q, Xu D S, Qiu X M, Zhao Y G. Research on wearresistance composite with bionic unsmooth surface. Trans-actions of the Chinese Society of Agricultural Engineering, 2001, 17, 7í9. (in Chinese)[15]Shen L, Li C L. Effect of laser melting processing on themicrostructure and wear resistance of gray cast iron. Wear, 1991, 150, 195í206.[16]Habig K H. Wear Behavior and Hardness of Material s,Mechanical Industry Publisher, Beijing, 1987. (in Chinese) [17]Chen J J, Yu Z S, Xu G K. High Cr Alloyed Cast Iron andApplication, Metallurgy Industry Press, Beijing, 1999.。

TBM滚刀刀圈材料合金成分对组织和耐磨性能的影响

第51卷第10期2020年10月中南大学学报(自然科学版)Journal of Central South University (Science and Technology)V ol.51No.10Oct.2020TBM 滚刀刀圈材料合金成分对组织和耐磨性能的影响贾连辉1，尚勇1，龙伟民2，夏毅敏3，薛广记1(1.中铁工程装备集团有限公司，河南郑州，450016；2.郑州机械研究所有限公司，河南郑州，450001；3.中南大学机电工程学院，湖南长沙，410083)摘要：为提高TBM 滚刀刀圈在高磨蚀地层的使用寿命，研究刀圈合金成分和组织中耐磨相种类及状态对耐磨性的影响规律，调配新型TBM 刀圈材料进行微观组织观察，并通过磨粒磨损试验、缩尺滚刀破岩试验及掘进破岩试验进行磨损性能对比。

研究结果表明：C ，Cr 和Mo 等合金元素的增多有助于共晶碳化物耐磨相的形成，对材料的耐磨性有一定的提升作用，而对材料的抗冲击性却是不利的；耐磨相的尺寸和形貌对材料耐磨性能及硬度的影响各不相同，大块的共晶碳化物对耐磨性的贡献比粒状的共晶碳化物的高，大块的共晶碳化物对硬度的贡献比粒状的共晶碳化物的略小。

在掘进高磨蚀、完整性较好的弱冲击岩层时，相比于常规的高韧性C-1材料，C-2新材料的刀圈寿命可提升30%以上，C-3新材料的刀圈寿命可提升70%以上。

3种材料的综合性能能够满足不同岩层的掘进要求，可为刀圈地质适应性选型提供参考。

关键词：滚刀刀圈；合金成分；共晶碳化物；耐磨性能中图分类号：U455.31文献标志码：A开放科学(资源服务)标识码(OSID)文章编号：1672-7207（2020）10-2730-09Effect of alloy composition of materials for TBM cutter rings onmicrostructure and wear resistanceJIA Lianhui 1,SHANG Yong 1,LONG Weimin 2,XIA Yimin 3,XUE Guangji 1(1.China Railway Engineering Equipment Group Co.Ltd.,Zhengzhou 450016,China;2.Zhengzhou Research Institute of Machinery Mechanical Engineering Co.Ltd.,Zhengzhou 450001,China;3.School of Mechanical and Electrical Engineering,Central South University,Changsha 410083,China)Abstract:In order to improve the service life of disc cutter rings in TBM in high erosion lithology,the effects of alloy composition ingredient and type of wear-resistant phase in microstructure on wear resistance were studied.New TBM cutter ring material was blended for microstructure observation.The wear performance was compared through abrasive wear test,scale-reduction rock-breaking test and tunneling rock-breaking test.The results show that the increases of C,Cr,Mo and other alloying elements enhance the formation of wear-resistant phases of eutectic carbides,and have a certain effect on improving the wear resistance of the material,but reduce the impactDOI:10.11817/j.issn.1672-7207.2020.10.005收稿日期：2020−01−25；修回日期：2020−04−29基金项目(Foundation item)：国家重点研发计划项目(2017YFB1302601)；河南省重大科技专项(181200210100)(Project(2017YFB1302601)supported by the National Research and Development Program of China;Project(181200210100)supported by the Significant Science and Technology Program of Henan Province)通信作者：贾连辉，教授级高级工程师，从事隧道掘进机设计制造研究；E-mail:**********************第10期贾连辉，等：TBM滚刀刀圈材料合金成分对组织和耐磨性能的影响resistance of the material.The effect of the size and morphology of the wear-resistant phase on the wear resistance and hardness of the material is also rge eutectic carbides contribute more to wear resistance than the granular eutectic carbides,and large eutectic carbides contribute slightly less to the hardness than the granular eutectic carbides.In driving high abrasion and weak integrity rock formations,compared with the conventional high toughness C-1material,the life of the cutter ring of the new C-2material can be increased by more than30%, and C-3new material ring life can be increased by more than70%.The comprehensive performance of the three materials can meet the requirements of different rock formations,which provides reference for the geological adaptability selection of cutter ring.Key words:disc cutter ring;alloy composition;eutectic carbide;wear-resisting performance岩石隧道掘进机(tunnel boring machine,TBM)掘进过程中伴随刀盘转动，刀盘上安装的盘形滚刀刀圈直接与岩体相互作用，滚刀贯入岩体滚动挤压，岩石表面产生局部变形并出现微观裂纹，随着挤压力增大，微观裂纹发展成为主裂纹并贯通，形成岩片破碎[1−2]。

激光熔覆马氏体

第27卷第2期粉末冶金材料科学与工程2022年4月V ol.27 No.2 Materials Science and Engineering of Powder Metallurgy Apr. 2022DOI:10.19976/ki.43-1448/TF.2021090激光熔覆马氏体/铁素体涂层的组织与抗磨耐蚀性能张磊1, 2，陈小明1, 2，霍嘉翔1，张凯1, 2，曹文菁1, 2，程新闯3(1. 水利部产品质量标准研究所浙江省水利水电装备表面工程技术研究重点实验室，杭州 310012；2. 水利部杭州机械设计研究所水利机械及其再制造技术浙江省工程实验室，杭州 310012；3. 绍兴市曹娥江大闸管理局，绍兴 312000)摘要：为提高液压活塞杆的耐腐蚀和抗磨损性能，在45号钢表面采用激光熔覆技术在不同激光功率下制备具有马氏体/铁素体组织的Fe基合金熔覆层。

利用X射线衍射仪、扫描电镜、X射线能谱仪等手段表征涂层的物相组成、微观形貌和元素分布，采用维氏硬度计和干滑动摩擦试验机对涂层的显微硬度和抗磨损性能进行测试，并通过电化学工作站研究熔覆层的耐腐蚀性能。

结果表明：Fe基合金熔覆层的主要物相为α-Fe、Ni-Cr-Fe、γ-(Fe,C)和Fe9.7Mo0.3等，主要组织为马氏体、铁素体和少量残余奥氏体。

熔覆层的枝晶态组织均匀致密，无裂纹和孔隙缺陷，涂层与基体呈冶金结合。

涂层的硬度与耐磨性能随激光功率增大而提高，当功率为2.4 kW时，涂层的平均显微硬度(HV)为647.64，耐磨性能为45号钢的9.37倍，磨损机制为磨粒磨损。

随激光功率提高，Fe基合金熔覆层的耐腐蚀性能先升高后降低，当激光功率为2.0 kW时涂层具有最佳耐腐蚀性能，显著高于活塞杆常用碳钢、不锈钢以及电镀硬铬等材料，可在相关领域替代电镀铬。

关键词：激光熔覆；Fe基合金；组织；磨损；腐蚀；活塞杆中图分类号：TG174.44文献标志码：A 文章编号：1673-0224(2022)02-196-09All Rights Reserved.Microstructure and wear-corrosion resistance performance oflaser cladding martensite/ferrite coatingZHANG Lei1, 2, CHEN Xiaoming1, 2, HUO Jiaxiang1, ZHANG Kai1, 2, CAO Wenjing1, 2, CHENG Xinchuang3(1. Key Laboratory of Surface Engineering of Equipment for Hydraulic Engineering of Zhejiang Province, Standard &Quality Control Research Institute, Ministry of Water Resources, Hangzhou 310012, China;2. Water Machinery and Remanufacturing Technology Engineering Laboratory of Zhejiang Province, HangzhouMechanical Research Institute, Ministry of Water Resources, Hangzhou 310012, China;3. Shaoxing Municipal Cao’e River Floodgate Construction Administration Committee, Shaoxing 312000, China)Abstract: To improve the corrosion resistance and wear resistance of piston rod, Fe-based coatings with martensite andferrite structure were prepared on 45# steel by laser cladding. The phase compositions, microstructure and elementsdistribution of the coatings were characterized by X-ray diffractometer, scanning electron microscope and X-ray energydispersive spectrometer. The microhardness and wear resistance of the coatings were tested by Vickers hardness testerand dry sliding friction wear tester. Furthermore, the corrosion resistance of laser cladding Fe-based coatings was studiedby electrochemical workstation. The results show that the phase of laser cladding Fe-based alloy coating is mainlycomposed of α-Fe, Ni-Cr-Fe, γ-(Fe,C), Fe9.7Mo0.3. The main microstructure is martensite, ferrite and a small amount ofresidual austenite. The dendritic structure of coating is uniform, compact, without cracks or pores. The coating and thesubstrate are bonded metallurgically. The hardness and wear resistance of the coatings increase with increasing基金项目：浙江省“一带一路”国际科技合作项目(2019C04019)；浙江省公益性技术应用研究计划资助项目(GC22E017317，LGC19E090001，2018C37029)收稿日期：2021−11−02；修订日期：2021−12−23通信作者：张磊，工程师，硕士。

Influence of micro-blasting on the microstructure and residual stresses of CVD κ-Al2O3 coatings

In ﬂuence of micro-blasting on the microstructure and residual stresses of CVD κ-Al 2O 3coatingsC.Barbatti a ,b ,J.Garcia c ,d ,⁎,R.Pitonak e ,H.Pinto a ,A.Kostka a ,A.Di Prinzio f ,M.H.Staia f ,A.R.Pyzalla caMax-Planck-Institut für Eisenforschung GmbH,40237Düsseldorf,GermanybCorus Research,Development and Technology,P.O.Box 10000,1970IJmuiden,The Netherlands cHelmholtz Zentrum Berlin fuer Materialien und Energie GmbH,14109Berlin,Germany dRuhr Universität Bochum,Materials Technology,44780Bochum,Germany eR&D,Boehlerit GmbH &Co.KG,8605Kapfenberg,Austria fSchool of Metallurgy and Materials Science,Universidad Central de Venezuela,P.O.Box 47885,Los Chaguaramos,1041Caracas,Venezuelaa b s t r a c ta r t i c l e i n f o Article history:Received 10February 2009Accepted in revised form 3June 2009Available online 17June 2009Keywords:CVD κ-Al 2O 3Cemented carbides Micro-blasting Residual stressesSynchrotron X-ray diffractionCemented carbides coated with CVD multilayers are commonly used in metal cutting turning and milling operations.For many applications,a micro-blasting ﬁnishing procedure based on an impact treatment of the surfaces is carried out in order to smooth the coated surface and reduce sharp cutting edges.In this work,micro-blasting with corundum in aqueous solution at pressures between 0.05and 0.3MPa was applied to CVD TiN/Ti(C,N)/κ-Al 2O 3multilayer coatings deposited onto cemented carbides in order to investigate its in ﬂuence on the micro-topography,microstructure and residual stresses.The results showed that the micro-blasting reduces the surface roughness and affects the coating thickness.TEM investigations revealed no signi ﬁcant changes on the microstructure of the κ-Al 2O 3top layer.Synchrotron X-ray investigations showed that the residual stress state of the as-deposited κ-Al 2O 3top layer is not affected by the micro-blasting treatment under the conditions investigated.©2009Elsevier B.V.All rights reserved.1.IntroductionCoated cemented carbides are commonly used in the metal cutting industry for machining of iron-base alloys.In order to increase the tool life of cutting inserts,cemented carbides are coated with wear resistant coatings by chemical vapour deposition (CVD)or physical vapour deposition (PVD).Al 2O 3is considered an ideal coating layer in high-speed metal cutting due to its high chemical stability,high hardness and favourable thermal properties at the high temperatures reached during metal cutting [1–3].Due to the mixed ionic and covalent bonding,the material is highly insulating [4,5],thus being usually applied as a top layer to act as a thermal and diffusion barrier.The CVD technique is commonly employed since it is the only deposition method that allows the economical large-scale production of high-quality alumina coatings [6].In addition to the stable α-Al 2O 3,alumina can be obtained in two other metastable modi ﬁcations by CVD,namely κ-Al 2O 3and γ-Al 2O 3[7,8].A typical multilayer combination is that of Ti(C,N)to provide wear resistance,followed by an Al 2O 3top layer,which enhances the thermal resistance of the tool.Due to adhesion problems between the Ti(C,N)and the Al 2O 3coating,the use of an intermediate layer made of Ti(C,N,O)is proposed[9,10].These intermediate layers of Ti-based compounds favour the nucleation of κ-Al 2O 3[11,12].More recently,a novel method for producing an adequate bonding of the Al 2O 3top layer onto the Ti(C,N)by the formation of a nano-transition needle-like phase arrangement has proved to be advantageous [13].This method also leads to the formation of κ-Al 2O 3.Micro-blasting is a ﬁnishing procedure based on an impact treatment of the surfaces.Generally,this procedure is applied to coatings to clean,smooth the surface and reduce sharp cutting edges [14–16].This process involves the use of high pressure and abrasive powders.Besides,there is still the possibility of achieving an attractive side effect,namely,the introduction of bene ﬁcial compressive residual stresses.Previous works showed that micro-blasting of ground and polished hard metal substrates can increase the cutting performance of coated tools mainly due to ﬁlm adhesion enhancement [17–20].Moreover,it has been recently shown that micro-blasting treatment of PVD coatings deposited on hard metal substrates improves its residual stress with regard to fatigue strength and cutting performance of the coated tool [21].In [22,23]it is reported that compressive residual stresses were induced on PVD/CVD multilayers by micro-blasting with pressures between 0.2and 0.6MPa.Holzschuh reports in [24]a method to introduce compressive stresses in CVD multilayers by removing the TiN top layer of a Ti(C,N)/Al 2O 3/TiN multilayer coating by micro-blasting,however,without any mention of the blastingSurface &Coatings Technology 203(2009)3708–3717⁎Corresponding author.Helmholtz Zentrum Berlin fuer Materialien und Energie GmbH,14109Berlin,Germany.Tel.:+492343228229.E-mail address:jose.garcia@helmholtz-berlin.de (J.Garcia).0257-8972/$–see front matter ©2009Elsevier B.V.All rights reserved.doi:10.1016/j.surfcoat.2009.06.021Contents lists available at ScienceDirectSurface &Coatings Technologyj o u r n a l h om e p a g e :w w w.e l s ev i e r.c o m /l o c a t e /s u r fc o a tconditions employed.In[22–24]improvements of tool life of micro-blasted Al2O3coated cutting inserts in metal cutting are presented. The reason for the increment of the tool life is reported to be related to the introduction of compressive residual stresses on the Al2O3top layer.Although not mentioned in their publications,it is to be understood from the results presented that the authors worked with α-Al2O3coatings.Nevertheless,in spite of the many publications,no systematic study on the inﬂuence of micro-blasting pressure on Al2O3coating is found in the literature.In this work,we applied micro-blasting to CVD multilayer TiN/Ti (C,N)/κ-Al2O3coatings.The micro-blasting pressure range varies between0.05and0.3MPa.We aim at investigating systematically the inﬂuence of the blasting pressure on the micro-topography and the possibility of inducing compressive residual stresses on theκ-Al2O3 top layer.The microstructure and residual stress state of theκ-Al2O3top layer were investigated for the as-deposited and micro-blasted conditions by scanning electron microscopy,transmission electron microscopy,synchrotron as well as laboratory source X-ray diffraction and roughness measurements.2.Experimental2.1.Substrate and coatingsThe coatings were deposited onto cemented carbide substrates (90wt.%WC–10wt.%Co)with a size of14mm×14mm×5mm, produced by state-of-the-art powder metallurgy sintering techniques.A TiN/Ti(C,N)/κ-Al2O3multilayer was applied on the hard metal substrate by CVD in a standard hot-wall CVD reactor working at 1015°C and15kPa.The TiN and Ti(C,N)layers were deposited from the gas mixture TiCl4–CH4–N2–H2.The Al2O3layer was deposited on top of the Ti(C,N)layer from the gas mixture AlCl3–ClH2–CO2.Due to the deposition conditions employed,κ-Al2O3was produced as a top coating layer.2.2.Micro-blastingAfter deposition,the coatings were subjected to a micro-blasting treatment performed using an injector-type system in a commercial industrial blasting furnace.Standard industrial parameters were employed.The grit blasting material used was320-mesh(average diameter within~50µm)corundum shots in aqueous solution.The micro-blasting treatment was carried out for30s to ensure full coverage of the surfaces.The pressure was varied within the range from0.05MPa to0.3MPa in steps of0.05MPa.2.3.MicroscopyScanning electron microscopy(SEM)was performed on non-polished surfaces and fractured cross-sections.Specimens for transmission electron microscopy(TEM)analyses were prepared by using a gallium source JEOL JEM-9320focused ion beam(FIB)system operating at30kV.A JEOL2200F TEM/STEM operating at200kV and equipped with an EDX system was used for studying the microstructure of coatings in the as-deposited and micro-blasted state.2.4.RoughnessThe surface roughness(R a and R z)of as-deposited and micro-blasted coatings was determined from atomic force microscopy pictures.Four areas were measured in order to determine an average value.The measurements were carried out in a Nano-scratch tester equipped with an atomic force microscope(CSM Instruments).2.5.Phase and residual stress analysesIn order to assess the phase constitution within the as-deposited and micro-blasted coatings,X-ray diffraction(XRD)diffractograms were obtained using synchrotron radiation at the experimental station G3at DORIS III at HASYLAB at DESY,Hamburg.The radiation energy was 6.9keV(λCo=1.7904A˚´).The beam size was set as 4mm×1mm.The residual stresses within theκ-Al2O3coatings were determined on a laboratory four-circle diffractometer operating with Co–K a radiation and parallel-beam geometry using the sin2ψmethod[25]. Owing to the weak macroscopic texture of theκ-Al2O3coating, unrestricted lattice spacing measurements could be performed for24 sample directions deﬁned by the inclination angleψbetween0°and 89.8°at the azimuthsφ=0°and180°.Considering the stress state in the coatings to be biaxial and of rotational symmetry,the fundamental relationship of X-ray stress analysis can be solved to provide the depth-dependence of the in-plane residual stressesσ//[26]:σ==τeffðÞ=eψτeffðÞ12s21ÀÁð1Þwhereɛψis the measured lattice strain,σ//is the in-plane residual stress,τeff is the information depth,and1/2s2and s1are the phase-speciﬁc diffraction elastic constants(DEC).The concept of effective information depthτeff is usually applied in diffraction studies of thinﬁlms to deﬁne the depth from which the diffraction information comes,since the penetration depthτof the radiation into a certain material may be much larger than theﬁlm thickness.According to its deﬁnitionτeff can be tuned within theﬁlm thickness by changing theψ-angle[27]:τeff=R DzÁe−zτd zRe−zτd z=τ−De−Dτ1−e−Dτð2Þwhereτis the average penetration depth for an inﬁnitely thick sample,which is a function of theψ-tilt,and D is the thickness of theκ-Al2O3layer.In this approach the maximum achievable information depth approaches the limit of half layer thickness(D/2) as the ratio D/τin Eq.(2)decreases towards zero.The lattice strains were measured for the{135}diffraction line of κ-Al2O3.This reﬂection was chosen based on a compromise between the ease for separating this line from neighboring ones,its intensity and the highest possible2θ-angle.The DEC were estimated using the mechanical values of Young's modulus E=361.7GPa and Poisson's ratioν=0.24forκ-Al2O3at room temperature available in[5].The integral breadth of the selected reﬂection was also determined as a function of the micro-blasting conditions.This quantity is related to the density of crystalline defects(dislocations,stacking faults)and small size of crystallites(i.e.domains over which diffraction is coherent)[28].The integral breadth for each condition represents an average on the results obtained at differentψ-angles byﬁtting the {135}diffraction lines using pseudo-Voigt functions.3.Results3.1.Microstructure of the as-deposited coatingsThe as-deposited layer system is shown in Fig.1.The CVD TiN/ TiCN/κ-Al2O3layers show good adhesion to the cemented carbides, hence no pores,brittle phase formation nor spallation between the substrate and the coating are observed.The layer thickness is homogeneous across the coating,the thickness of the TiN is of about 0.2µm,for the intermediate Ti(C,N)layers is about3µm and the thickness of theκ-Al2O3is about4µm±0.05.3709C.Barbatti et al./Surface&Coatings Technology203(2009)3708–3717The alumina layer is composed of columnar grains that grow directly from the nucleation surface (here,the intermediate Ti(C,N)layer).The length of the columnar grains is mostly equivalent to the thickness of the alumina layer.A change in the morphology towards the surface can be observed,with the formation of hexagonal-shaped grains (Fig.1b and c).Some cracks (mainly transgranular)can also be observed.They are typical cracks as a consequence of the different thermal expansion coef ﬁcients between the substrate and the coating and cannot really be avoided in commercial CVD cutting inserts.Several isolated large grains with average diameter larger than 5µmformed as a single crystal or a polycrystalline structure in a random distribution throughout the surface (Fig.2).Their number increases towards the edge of the inserts.Crack initiation seems to be associated with these outgrowths.The average surface roughness R a measured was 0.110±0.008(Table 1).The penetration depth of the X-rays in alumina is about 18µm.Hence,we have full penetration of the radiation into the coating system.The X-ray diffraction (XRD)patterns for the as-deposited sample are shown in Fig.3.The contribution of the substrate is also registered,indicated by the peaks associated with WC re ﬂections.The XRD measurements also reveal that the intermediate layer is composed of Ti (C 0.3N 0.7).The intensity ratio between peaks shows no marked texture in the layers.Although the observation of the κ-→α-Al 2O 3phase transformation already during the deposition process is generally reported in literature [29],there is no evidence for the occurrence of such transformation,since the XRD diffractograms only revealed re ﬂections associated with κ-Al 2O 3.This indicates that the super ﬁcial cracks observed in the as-deposited coatings arise upon cooling due to the differences in thermal coef ﬁcient between the coating and the substrate,and not originating from volume contraction that occurs in phase transformation.Fig.4shows a STEM micrograph of the as-deposited coating.The darkest areas at the bottom correspond to the cemented carbide substrate (WC grains are seen in black).The TiN layer with a thickness of 0.5µm lies adjacent to the substrate.Onto the TiN layer,the Ti(C,N)layer shows a changing grain morphology (Fig.4),making it appear as two distinct structures.The ﬁrst sub-layer consists of columnar grains with a width of about 0.1µm.The following sub-layer has signi ﬁcantly smaller grains with a needle-like morphology and appears randomly oriented.The uppermost layer consists of κ-Al 2O 3(in lighter contrast)and shows grains grown along the CVD deposition direction perpendicular to the substrate.The columnar microstructure istypicalFig.2.SEM image (BSE mode)of κ-Al 2O 3surface topography (top view)after coating.Isolated large Al 2O 3grains (with average size of 5µm)form in a random distribution throughout the surface.Crack initiation seems to be associated with these large grains (see arrows around the large Al 2O 3grain).Table 1Values of R a ,R z and thickness for the κ-Al 2O 3top layer as a function of micro-blasting pressure.Micro-blasting pressure [MPa]R a roughness [µm]R z roughness [µm]Thickness [µm]00.110±0.0080.690±0.04 4.20±0.050.050.096±0.0050.540±0.04 4.20±0.050.10.079±0.0060.480±0.04 4.00±0.050.150.038±0.0040.190±0.03 3.40±0.050.20.045±0.0030.250±0.03 3.00±0.050.250.074±0.0040.440±0.03 2.50±0.050.30.084±0.0040.540±0.042.50±0.05Fig.1.SEM images of (a)fracture cross-section of the coated cemented carbide,(b)top view of the surface topography of as-deposited CVD κ-Al 2O 3coating and (c)higher magni ﬁcation (top view)of the surface topography showing hexagonal-shaped grains.3710 C.Barbatti et al./Surface &Coatings Technology 203(2009)3708–3717for coatings deposited at low gas pressures and low temperatures compared to the melting point temperature [30].Besides the STEM mode,conventional Bragg diffraction contrast TEM was applied to investigate the evidence of defects within the κ-Al 2O 3top layer.A TEM micrograph in bright ﬁeld (BF)mode is shown in Fig.5a.No dislocations but a high concentration of planar defects (twin-related faults)were observed (Fig.5a).All obtained electron diffraction patterns(e.g.Fig.5b)only reveal the presence of the κ-Al 2O 3phase.No evidence of phase transformation was obtained.The two possible variations of alumina (i.e.the metastable κ-Al 2O 3and the stable polymorph α-Al 2O 3)have strikingly different symmetry,which could have been easily distinguished if the latter were present.EDX analysis reveals almost no chemical intermixing,which suggests good separation of the layers.3.2.Microstructure of micro-blasted coatingsMicrographs of surface morphology and fracture cross-section of the samples micro-blasted at pressures in the range 0.05–0.15MPa are presented in Fig.6.Lower micro-blasting pressures,i.e.0.05and 0.1MPa cause no signi ﬁcant changes in grain morphology,although a decrease in the average roughness is measured.At blasting pressures as low as 0.05MPa few isolated areas already appear deformed.As from 0.15MPa,signi ﬁcant loss of coating material begins to occur and the number of cracks increases.Crack propagation typically occurs in a transgranular mode,but seems to be con ﬁned to the near-surface zone of the upper coating.Micro-blasting at 0.2MPa causes a ﬂattening of the grains and ﬂaking of the coating as a consequence of the strong abrasive effect due to micro-blasting (Fig.7).Higher micro-blasting pressures (0.25and 0.3MPa)lead to strong changes in the surfacemicro-topographyFig.3.X-ray diffractogram of the as-deposited κ-Al 2O 3/Ti(C,N)coatings taken at ψ=0°with synchrotron radiation energy of 6.9keV (λCo =1.7904A˚´).The arrow indicates the position of the κ-Al 2O 3(135)re ﬂection used for the residual stressevaluation.Fig. 4.(a)STEM micrograph showing the cross-section microstructure of the as-deposited coating.(b)Higher magni ﬁcation of the top κ-Al 2O 3layer.Fig.5.TEM micrograph in bright ﬁeld (BF)mode showing planar defects within the as-deposited κ-Al 2O 3layer.(b)Selected area diffraction pattern (SADP)corresponding to κ-Al 2O 3.3711C.Barbatti et al./Surface &Coatings Technology 203(2009)3708–3717and an increase in the number and length of the cracks.The microstructure of the Ti(C,N)intermediate layer does not seem affected by the impact of the alumina shots.STEM micrographs of the cross-section microstructure of coatings micro-blasted at 0.1MPa and 0.3MPa are shown in Fig.8.By comparing the microstructure of the micro-blasted samples at different conditions,the only difference corresponds to the signi ﬁcant reduced thickness of the alumina layer micro-blasted at 0.3MPa.The in ﬂuence of micro-blasting pressure on the thickness of the κ-Al 2O 3layer is shown in Table 1.On the other hand,the results clearly show that the micro-blasting does not affect the adherence of the layer.Moreover,the Ti(C,N)is not affected by the mechanical treatment,the reduction in layer thickness being con ﬁned to the top alumina layer.The TEM micrograph in BF mode (Fig.8c)shows the typical planar defects also observed within the alumina layer of the micro-blasted coatings (here we only show the example of the sample micro-blasted at 0.3MPa).The concentration of the twin-related faults observed in samples in all conditions seems to be identical,which directly suggests that the micro-blasting process has no in ﬂuence on the defect density observed in the top κ-Al 2O 3layers.Besides,not a single dislocation was observed in the analyzed micro-blasted top κ-Al 2O 3layers.Micro-blasting leads to a decrease in the roughness of the alumina coating,showing a minimum at blasting pressures between 0.15and 0.2MPa (Table 1).Increasing the blasting pressure above 0.2MPa increases the roughness due to deterioration of the coating layer.ThisFig.6.SEM micrographs showing the changes on the κ-Al 2O 3thickness (fracture cross-sections)and surface topography (top view images)as a function of micro-blasting pressure:(a)and (b)at 0.05MPa;(c)and (d)at 0.1MPa;(e)and (f)at 0.15MPa.3712 C.Barbatti et al./Surface &Coatings Technology 203(2009)3708–3717is related to the extensive loss of the top layer material that occurs with increasing pressures.The removal of top coating material causes an increase in the line intensity of the WC re ﬂections,since more contribution from the substrate is being registered (Fig.9).No changes in the intensity ratio between peaks with increasing micro-blasting pressures are observed.3.3.Residual stress analysesThe sin 2ψcurves of the κ-Al 2O 3layers in the as-deposited and micro-blasted (at 0.1and 0.3MPa)conditions are presented in Fig.10.The obtained d –sin 2ψdistributions are shown to be linear with positive slopes up to sin 2ψ~0.92in all conditions,indicating the presence of average residual stresses of tensile nature within the κ-Al 2O 3layer,which still predominate in the micro-blasted state.On top of the coatings (at large ψangles,thus,at small penetration depths),however,the d ψhkl–sin 2ψdistributions tend to curve downwards,suggesting the presence of near-surface gradients of the in-plane residual stresses in the as-deposited as well as in the micro-blasted conditions.ψ-splittingin the d ψhkl–sin 2ψplots does not occur,indicating the absence of shear stress components in all cases investigated.The in-plane residual stresses averaged over the entire κ-Al 2O 3layer are 355±41MPa,349±45MPa,and 392±39MPa in the as-deposited,0.1and 0.3MPa conditions,respectively (Table 2).Based on the sin 2ψdistributions (Fig.10),the in-depth distribution of the residual stresses in the κ-Al 2O 3coatings was determined using Eq.(1).The stress-depth pro ﬁles for the as-deposited and micro-blasted coatings are similar as presented in Fig.11.In the as-depositedstate,Fig.7.SEM micrographs showing the changes on the κ-Al 2O 3thickness (fracture cross-sections)and surface topography (top view images)as a function of micro-blasting pressure:(a)and (b)at 0.2MPa;(c)and (d)at 0.25MPa;(e)and (f)at 0.3MPa.3713C.Barbatti et al./Surface &Coatings Technology 203(2009)3708–3717the residual stress in the ‘bulk ’of the layer was found to be σ//(τ≈1.5µm)=699MPa,whereas the obtained value at the outer surface decreases down to σ//(τ≈0)=−32MPa.This indicates thatthe stresses become gradually less tensile towards the free surface,eventually turning into low compression at the surface.A similar behaviour is observed for the specimen micro-blasted at 0.1MPa:the tensile residual stress within the coating is slightly lower σ//(τ≈1.5µm)=603MPa,whereas at the surface low compressive residual stresses σ//(τ≈0)=−42MPa are also obtained.For the specimen micro-blasted at 0.3MPa,the stress values are σ//(τ≈1.5µm)=546MPa and σ//(τ≈0)=153MPa within the coating and close to the surface,respectively (Table 2).The analysis of the average integral breadth of the {135}diffraction line of κ-Al 2O 3suggests that the density of lattice imperfections does not signi ﬁcantly increase with increasing micro-blasting pressures.This information corroborates the TEM investigations,which reveal low dislocation densities,even with increasing micro-blasting pressures.4.Discussion4.1.Effect of micro-blasting on microstructure and roughnessOur results indicate that the micro-blasting procedure only affects the κ-Al 2O 3top coating layer.SEM and TEM images show that there are no disturbances in the integrity of the intermediate Ti(C,N)layer.Strong deformation of the coatings appears as a ﬂattening effect on the near-surface zone,mainly at pressures higher than 0.1MPa.ThisFig.8.STEM micrographs showing the cross-section microstructure of coatings micro-blasted at (a)0.1MPa and (b)0.3MPa.(c)TEM micrograph (BF)showing planar defects within the alumina layer of the coating micro-blasted at 0.3MPa.Fig.9.X-ray diffractograms of micro-blasted coatings as a function of micro-blasting pressure showing no change in the intensity ratio between peaks with increasing micro-blasting pressures.3714 C.Barbatti et al./Surface &Coatings Technology 203(2009)3708–3717ﬂattening is produced on the one hand by the overlapping of the several localized deformations caused by the impact of the alumina shots on the surface.On the other hand,the impacts also remove material from the alumina coating,which also results in a smoothing of the protruding irregularities on the surface (Table 1).The lower roughness in the contact zone of the coating decreases the mechanical contact between asperities of the tribological pair tool-workpiece.This leads to a lower friction coef ﬁcient and thus,a contribution for extended tool life of the inserts.4.2.Origin of stresses and defects4.2.1.Stress development upon CVD depositionCVD coatings are known to evolve relevant thermal mismatch residual stresses upon cooling from the high deposition temperatures [31].The strain and stress state in the layers change due to thermal strain introduced by differences in the coef ﬁcients of thermal expansion between the individual layers and the substrate.The maximum thermal stresses induced by cooling (without taking stress relieving processes into account)can be estimated by using the Dietzel equation [32]:σc =E c αc −αs ðÞΔT f g =1−v c ðÞ+1−v s ðÞ=E s ½ d c =d s f gð3Þwhere αis the linear coef ﬁcient of thermal expansion,νis the Poisson ratio,E is the Young's modulus,d is the thickness,c stands for coating and s for substrate.The thermal expansion coef ﬁcient(TEC)Fig.10.The sin 2ψdistributions of the κ-Al 2O 3(135)re ﬂection for φ=0°of the coatings (a)as-deposited,(b)micro-blasted at 0.1MPa and (c)micro-blasted at 0.3MPa.Table 2Values of internal stresses in the κ-Al 2O 3top layer as a function of micro-blasting pressure.Micro-blasting pressure [MPa]Average in-plane residual stresses [MPa]Stress value near the outer surface of κ-Al 2O 3[MPa]Stress value in the “bulk ”of κ-Al 2O 3[MPa]0355±41σ//(τ≈0)=−32σ//(τ≈1.5µm)=6990.1349±45σ//(τ≈0)=−42σ//(τ≈1.5µm)=6030.3392±39σ//(τ≈0)=153σ//(τ≈1.5µm)=546Fig.11.Depth pro ﬁle of the in-plane residual stresses obtained from the evaluation of the sin 2ψdistributions shown in Fig.11for the coatings (a)as-deposited,(b)micro-blasted at 0.1MPa and (c)micro-blasted at 0.3MPa.3715C.Barbatti et al./Surface &Coatings Technology 203(2009)3708–3717ofκ-Al2O3and Ti(C,N)usually lies within8.0–9.0×10−6K−1[33], whereas in the case of WC–Co cemented carbides it is in the range within 4.0–7.0×10−6K−1depending on the composition[34]. Hence,tensile stresses will necessarily develop in theκ-Al2O3coating sinceακ-Al2O3Nαs.Our results clearly show that average tensile stresses develop within theκ-Al2O3layer,indicating that thermal strains represent the major source of residual stresses in the present case.The low compressive stress values towards the surface could be attributed to the relaxation of the tensile residual stresses existent in the‘bulk’of the layer,since the interaction between the substrate and the coatings tends to decrease when approaching the free surface. Besides,the crack propagation,which seems to be conﬁned to the zone closest to the surface,might also play a role in the relaxation mechanism.4.2.2.Stresses and defects induced by micro-blastingThe residual stresses are generally classiﬁed into three types: macro-stresses,micro-stresses,and root mean square(r.m.s.)stresses [25].Macro-stresses often develop during sample manufacturing, extending homogeneously over macroscopic distances,i.e.several grains.In contrast,micro-stresses arise over microscopic volumes, typically grain sizes,due to the mismatches in thermal and mechanical properties of the individual phases or grain orientations. Both cause diffraction line shifts since they are related to average lattice strains of grains with speciﬁc orientations.As discussed in the Section4.2.1,they are of thermal origin in the present case,and are not affected by the applied micro-blasting treatment.The third type of stresses(r.m.s.stresses),which is inhomogeneous within subgrains or crystallites due to the presence of lattice imperfections and vary on the atomic scale,does not induce line shifts but only diffraction line broadening,since any disturbance in the regularity of the crystal lattice is reﬂected in lattice parameter ﬂuctuations.Typical causes for diffraction line broadening are therefore lattice distortions that arise from crystal defects(e.g.dislocations and planar faults),changes in the diffraction coherence length(domain size)and domain-size distribution[28].Based on evidences provided by both TEM and XRD investigations, we conclude that theκ-Al2O3top layer does not undergo plastic deformation during micro-blasting.In particular,TEM imaging indicates practically the absence of dislocations,even at a micro-blasting pressure as high as0.3MPa,which promotes extensive loss of theκ-Al2O3layer and strong changes in layer micro-topography.Planar defects in the form of twin-related faults are the only type of crystalline defects observed and already predominate in the as-deposited state,the density of which is not affected by the micro-blasting process.The crystal structure of the metastableκ-Al2O3is well established. It has an orthorhombic crystal structure(space group Pna21)with lattice parameters a=4.844Å,b=8.330Å,c=8.955Åformed from a pseudo-closed-packed stacking ABAC of oxygen atoms with Al in octahedral and tetrahedral environments in a3:1ratio[35,36]. Supposedly,due to the rather complex crystalline structure,distortion of the lattice occurs differently according to the crystallographic direction leading to twinning as the predominant deformation mechanism inκ-Al2O3phase in contrast toα-Al2O3,in which dislocation is the type of recurrent defects observed.5.ConclusionsMicro-blasting at pressures ranging from0.05up to0.3MPa was applied to CVDκ-Al2O3/Ti(C,N)/TiN multilayer coatings in order to investigate the potentiality of this procedure to control the surface micro-topography and modify the residual stress state of the coating. Our investigations reveal that:1.Theκ-Al2O3growth occurs in columnar direction,perpendicular tothe substrate.Noκ→αphase transformation of the Al2O3isinduced by micro-blasting according to both synchrotron X-ray diffraction and TEM studies.2.The effects of the micro-blasting treatment on the TiN/Ti(C,N)/κ-Al2O3multilayer under the conditions investigated are con-ﬁned to theκ-Al2O3top layer,as revealed by the TEM and STEM investigations.3.Increasing the micro-blasting pressure leads to the formation andpropagation of cracks.SEM and TEM investigations show that crack propagation is conﬁned to the zone closest to the surface.4.Increasing micro-blasting pressure leads to a reduction in roughnessfrom R a:0.110±0.008/R z:0.690±0.04in the as-deposited coatings up to R a:0.038±0.004/R z:0.190±0.03in coatings micro-blasted at0.15MPa,which is mainly associated with the loss of coating material.Blasting pressures above0.2MPa lead to coating degradation.5.TEM investigations showed high density of planar defects in theκ-Al2O3in the form of twin-related faults(Figs.5and8).TEM and XRD investigations reveal that theκ-Al2O3top layer does not undergo plastic deformation during micro-blasting.6.The residual stresses are of tensile nature and thermally induced andare not affected by the applied micro-blasting treatment.The average stress values over theκ-Al2O3layer are355±41MPa,349±45MPa, and392±39MPa in the as-deposited,0.1MPa and0.3MPa conditions.These tensile stresses ease towards the free surface.It can be in general concluded that,under the conditions investigated,the more remarkable effect of the micro-blasting treat-ment on theκ-Al2O3top layer was a reduction of the surface roughness for blasting pressures up to0.15MPa and a strong coating degradation for blasting pressures above0.2MPa.Neither the morphology nor the residual stress state of theκ-Al2O3top layer was affected by the micro-blasting.This is a main difference in the response to the treatment compared with the reported values forα-Al2O3[22,24,26],pointing out that the generation of compressive internal stresses by micro-blasting depends strongly on the crystalline structure of the Al2O3phase.This observation may be supported by previous investigations indicating a higher degree of plasticity ofα-Al2O3[37]compared toκ-Al2O3.On the other hand,the different physical characteristics of the metastable CVD κ-Al2O3compared to the CVDα-Al2O3,such as a smaller grain size,a lower pore density and an epitaxial growth,may also have an inﬂuence on the internal stress conditions after micro-blasting treatment[35]. AcknowledgementsThe authors(J.Garcia and H.Pinto)thank theﬁnancial support of the DFG project444Bra-113/25/0-1to carry out part of this work.Anna di Prinzio thanks ADEMAT Network,Alfa-project Nr.II-0240-B1-AT-RT-CT forﬁnancial support.Dr.A.Rothkirch(Hasylab),Dr.C.Juricic and Mr.P. Brito(MPIE)are kindly acknowledged for their support during the synchrotron X-ray measurements.Thanks to Mr.U.Föckeler(Ruhr Universität Bochum)for atom force microscopy measurements. References[1]F.Bunshah,Handbook of Hard Coatings,Noyes Publications/William AndrewPublishing,LLC,Norwich,New York,USA,2001.[2]P.K.Mehrothra,D.T.Quinto,High Temp.High Press.18(1980)199.[3]B.M.Kramer,N.P.Suh,J.Eng.Ind.102(1980)303.[4]R.H.French,J.Am.Ceram.Soc.73(1990)477.[5]B.Holm,R.Ahuja,Y.Yourdshhyan,B.Johansson,B.I.Lundqvist,Phys.Rev.,B59(1999)12777.[6]S.Ruppi,Int.J.Refract.Met.Hard Mater.23(2005)306.[7]S.Ruppi,rsson,Thin Solid Films388(2001)50.[8]S.Ruppi,rsson,A.Flink,Thin Solid Films516(2008)5959.[9]H.Holzschuh,Patent US6,436,519B2,20.08.2002.[10]H.Halvarsson,S.Vuorinen,Surf.Coat.Technol.56(1993)165.[11]M.Halvarsson,H.Nordén,S.Vuorinen,Surf.Coat.Technol.68–69(1994)266.[12]M.Halvarsson,J.E.Trancik,S.Ruppi,Int.J.Refract.Met.Hard Mater.24(2006)32.[13]R.Pitonak,J.Garcia,R.Weissenbacher,K.Udier,Patent EP1948842,30.07.2008.[14]D.M.Kennedy,J.Vahey,D.Hanney,Mater.Des.26(2005)203.[15]J.Qu,A.J.Shih,R.O.Scattergood,J.Luo,J.Mater.Process.Technol.166(2005)440.3716 C.Barbatti et al./Surface&Coatings Technology203(2009)3708–3717。

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