Surface & Coatings Technology 203 (2009) 1833–1843

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Laser engineered surfaces from glass forming alloy powder precursors:Microstructure and wear

D.T.A. Matthews a, V. Ocelík a, D. Branagan b, J.Th.M. de Hosson a,⁎a Department of Applied Physics, Materials Innovation Institute, University of Groningen, Nijenborgh 4, Groningen, 9474 AG, The Netherlandsb The NanoSteel Company, 505 Lindsay Blvd, Idaho Falls, ID 83402, United States

⁎ Corresponding author.E-mail address: j.t.m.de.hosson@rug.nl (J.T.M. de Ho

0257-8972/$ – see front matter © 2009 Elsevier B.V. Aldoi:10.1016/j.surfcoat.2009.01.015

a b s t r a c t

a r t i c l e i n f o

Article history:

Fe-based metallic glass for Received 14 August 2008Accepted in revised form 6 January 2009Available online 19 January 2009

Keywords:LaserMetallic glassWear resistanceNanosteelMicroscopy

ming powders have been deposited on mild steel substrates using high powerlaser cladding. Coatings microstructures have been analysed by scanning- and transmission-electronmicroscopy and at varying substrate dilutions, have been found to comprise a 100 to 500 nm interdendriticaustenitic phase and a dendritic dual-phase of ferrite/martensite. The application of double layer coatings hasshown microstructural refinement. This leads to a needle-like microstructure resulting in a nanoindentationtested hardness increase from ~11 GPa up to almost 15 GPa. The layers have been subjected to both drysliding wear and 3-body microscale abrasive wear testing. The dry sliding results show the layers to exhibitexcellent wear resistance – particularly at high speed (50 cm s−1) with wear rate values of ~1×10−8 mm3/Nmbeing recorded for the double layer coatings. The single layer coatings reveal a micro-wear mechanismconnected with the slip between the ferrite and martensite in the dendritic dual-phase. Microscale abrasivewear testing also reveals that the layers have a good wear resistance, with wear scars exhibitingcharacteristic material removal by micro-chipping. There is no preferential abrasion of any one phase, nor aretrack over-lap areas, cracks or pores found to result in varying wear scar dimensions.

© 2009 Elsevier B.V. All rights reserved.

1. Introduction

Along with a number of functionally poor metals, mild steels arewell known to exhibit low wear resistance and therefore for manyyears much research has been aimed at forming harder steels ordepositing hard surface layers to improve their functionality. Amongrecent advances in the formation of advanced, harder, more wearresistant materials is the concept of amorphous and nanocrystallinemetals. A possibility for the formation of metallic amorphous layers byharnessing the high cooling rates of high power lasers was presentedthe as early as 1987 [1]. Standard metallic glass alloy PdCuSi was cladon a copper substrate using a high power CO2 laser and a 20 µm thickcoating was formed [2]. Later, laser cladding of coatings with typicalbulk metallic glasses compositions were reported for Zr–Al–Ni [3],Fe–Co–Ni–Zr–Si–B [4], Ni–Cr–Mo–Zr–P–B [5] and Zr–Al–Ni–Cu [6]alloys. Laser surface cladding of Fe–B–C, Fe–B–Si and Fe–BC–Si–Al–Con plain carbon steel [7] has been shown to result in coatings with finedispersions of nano/microcrystalline intermetallic and interstitialcompounds and phases in ferritic matrix.

The current authors have shown in recent publications that thehigh power laser may be used to fabricate metallic glasses as thickamorphous, amorphous matrix/nanocrystalline layers on Ti-sub-strates even from multi-element powder mixes [8,9]. Amorphous

sson).

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layers fabricated by high power laser have been shown to exhibitattractive properties such as low friction [10], high hardness and goodwear resistance [8,11]. However, multi-powder processing is not soconvenient for industrial application and with the development ofglass forming alloys, comes the desire to process them for commercialapplication. The NanoSteel company has worked on harnessing theproperties of nanocrystalline materials through the development ofiron-based Glass Forming Alloys (GFA) [12] in the form of gas atomisedpowders with the goal of depositing (for example) plasma transferredarc welds [13] or thermal spray coatings [14]. These powders areinteresting in the scheme of research concerned with laser cladding ofhard, wear resistant, amorphous or nanocrystalline coatings. NanoS-teel powders pose the possibility of being able to form thick, nano-crystalline layers on iron/steel substrates. Of course, the possibility forproducing fully amorphous layers with these powders is diminisheddue to dilution from a substrate, whichwill shift the composition froma fully glass forming one to a “diluted” variant. In this case, it isexpected that amorphous matrix or nanocrytalline layers can bedeposited in order to harness the attractive properties alreadyoutlined [7–11] in the surface of a functionally poor material (A63mild steel). Furthermore, the deposition of thick coatings, provide abetter opportunity by which to fabricate functionally graded layersand to provide a gradual decrease in hardness from the surface to thesubstrate.

To this end, the upcoming research focuses on the microstructuralevolution of several novel GFA powder coatings deposited by high power

Table 1Test parameters for the micro-scale abrasion test [10]

Test parameter Value

Ball diameter (mm) 25Relative sliding speed (ms−1) 0.05Normal load on sample (N) 0.25Slurry Aqueous suspension of silicon carbideAbrasive size (µm) 4Abrasive concentration (%) 20

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laser. The research also addresses, relative to the mild steel substrate, thecoatings' hardness, sliding and abrasivewear resistance and examines theassociated surface wear morphologies.

2. Experimental procedures

A high power (2 kW) continuous pulse Rofin-Sinar Nd:YAG laserwasused to perform the laser cladding process, whichwas conducted over arange of processing parameters as will be specified where appropriate.The chosenpowder isNanoSteelNSSHS7574gasatomisedpowder in the+53 to −180 µm size range. The powder has a given composition inatomic percent of FebalanceCrb25Mob15 Wb10Cb3Mnb5Sib2Bb5 (meltingpoint=1388 K). In the results upcoming, the substrate is A36 steel (a USstandard structural mild steel, melting point=1643 K). The overlappingrate for all multi-pass coatings is 30%.

The resultant layers were cut in longitudinal and perpendicularcross-sections, mounted and polished to a mirror-finish before lightmicroscopy (Olympus Vanox-AHMT) and scanning electron-micro-scopy with energy dispersive spectroscopy (EDS) capability (PhilipsXL30 Series FEG SEM) was conducted to assess the coating micro-structural features. In conjunctionwith the latter, orientation imagingmicroscopy (OIM) permits information to be gathered as regards themicrostructural texture. Additional microstructural analysis wasperformed using (high resolution) transmission electron microscopy((HR)TEM) (FEG Jeol 2010) with in-situ heating and EELS capabilityand X-ray diffraction (XRD) (Phillips PW1710).

Mechanical analysis surrounded micro- and nano-indentationtesting, conducted on a CSM Revetester with a Vickers indenterhead and an MTS Nanoindenter XP (with CSM/LFM control and aBerkovich head) respectively as well as tribo-testing. Two forms ofwear testing were investigated, namely dry sliding wear and 3-bodymicroscale abrasive wear testing.

Fig. 1. NanoSteel SHS7574 powder for laser cladding. (A) SEM overview of the particles (seesections of the particles (seen here as white) embedded in plastic (seen as gray and black)revealing significant inherent porosity (seen as dark areas within the white regions).

The sliding wear tribo-testing is conducted on a CSM HT (HighTemperature) Tribometer. The technique is commonly referred to aspin-on-disk. The test piece can either take the role of pin or disk. In allcases for these investigations, the sample is a pin of dimensions(length×breadth×height) 2mm×4mm×3mmor 4mm×4mm×3mm.The counter-piece is a hardened (63Rc) 100Cr6 steel disk in all cases. Thepin ismounted in a stiff lever, designed as a frictionless force transducerand since the sample is always in contact with a sliding face, resultantfrictional forces acting between the pin and the disk, measured by verysmall deflections of the lever. Wear coefficients may be calculated fromthe resultant material volume loss during sliding and normalised forsliding distance and load to give wear rate values with dimensions andmagnitude (for awear resistant material) ~10−6 mm3/Nm. The effect onthese rates of variances in contact stress, wear test speed andcounterface roughness were investigated during wear testing. The testspeeds were 10 cm/s, 20 cm/s or 50 cm/s, will be noted whereappropriate. Loads of 5, 10 and 15 N are used and the resultant contactpressures for those loads are given as appropriate. Confocal opticalmicroscopy (µSurf Nanofocus Messtechnik) and SEM was additionallyimplemented in the characterization of the worn surfaces.

The microscale abrasive wear test procedure entails a series ofprogressive tests, to form a sequence of wear craters, or scars. Initially,in accordance with the NPL report compiled by Gee et al. [15–17], thesample is ultrasonically cleaned in acetone for 15 min, followed by anacetone rinse. Once the sample, typically 25 mm×50 mm×1 mm insize, is secured in place, and the motor speed correctly configured, theslurry drip feed is started. A hardened steel (100Cr6) ball is rotatedabout an axis perpendicular to, and pressed against, a coated samplein the presence of a pool of slurry of small abrasive particles. Theabrasive slurry, which is a suspension of 1200 grade silicon carbideparticles in distilled water, is maintained and replenished at thecontact region by the slow drop feed.

After a pre-determined number of revolutions, awear scar or crater isformed in the test sample which has the shape of the counter body,(spherical). The wear scar volume, V, can be completely characterised bythe radius of the sphere R and the wear scar diameter, b for bbbR:

V =πb4

64R

This procedure is repeated at intervals of 200, 400, 600 and 900revolutions. In between these intervals, the test piece is removed,ultrasonically cleaned, and rinsed as before. Once this sequence has

n as gray) adhered to a conductive carbon disc (seen as black) and (B) polished cross-and subsequently ground to reveal random cross-sections of a number of the particles,

Table 2Laser processing parameters for selected NanoSteel coatings

Coatingname

Power(W)

Scanspeed (m/s)

Feedingrate (g/s)

Defocus(mm)

Carrier gas(Ar) (l/min)

Shielding gas(Ar) (l/min)

STA 600 0.01 0.0667 +6 2 15STD 1000 0.015 0.1 +6 3 15MTA 1000 0.02 0.1 +6 3 15MTD 800 0.02 0.1 +6 3 15MTE 800 0.02 0.1 +6 3 15

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been completed, the wear scars are measured using an SEM, and theresults recorded. A summary of the test conditions is shown in Table 1.

3. Results and discussions

3.1. Microstructural characterisation

Initial productionwas somewhat limited since, although the shapeand size of the powders is ‘optimal’ for laser processing (spheroids ofØ~50–180 µm) (Fig. 1A), initial cladding experiments revealed highporosity. Fig. 1B shows that this porosity is induced from the powderitself. This does not complicate the thermal spray process too much,for which the powders were originally developed, since, firstly, a finersize cut is utilized typically (+15 to −53 µm) with negligible porosityfound, and secondly, splat formation opens the powder particles andallows the gas to escape.

Glass forming alloy compositions are close to eutectic, whichimplies a (relative) low melting point. The interaction time betweenthe powder and the laser beam is not sufficient to fully melt theimpinging particles, and therefore many of the particles onlyexperience melting of the outer shell.

If the particles are however deposited into a molten pool which isat a higher temperature than the melting temperature of the particlesthemselves, an opportunity is present for the particles to become fully

Fig. 2. (A) BSE micrograph of the microstructure of coating ‘STA’ and (B) SE micrograph of tinterfacial bonding for coatings STA and STD, respectively.

liquid and the gasses in the particles may escape from, not only theparticles themselves, but also the solidifying melt pool.

In order to combat this problem, a number of variations on theprocessing were investigated including varying the scan speed, laserdensity, volume of carrier gas and substrate and/or powder pre-heating. It was found in these investigations that the best method ofprolonging the meltpool life and thus avoiding coating porosity,coincided with an extremely high percentage dilution.

High coating dilutions obviously affect the coating chemistry andgrain size and in order to analyse the microstructural evolution ofthese coatings, herewe address the deposition of single track coatings.Processing parameters and names for all coatings under investigationare given in Table 2.

If we view the coatings by SEM, we see that despite the higherscanning speed for coatingSTD (Fig. 2B), thedendrite size is larger than forSTA (Fig. 2A). This is a result of the higher laser power used for coating STDwhich induces a higher substrate dilution. One of the most attractiveaspects of laser cladding over other coating deposition techniques is theexcellent coating adhesion resultant from metallurgical bonding. This isseen microscopically in Fig. 2C,D. The STA coatings exhibit excellentinterfacial bonding which can be seen in Fig. 2C between a well-defined50 µm band of large dendrites comparative to the ‘bulk’ coating and thesubstrate. The high dilution in coating STD, Fig. 2D, is evident through thelow contrast in BSE mode. Nonetheless, the interfacial bonding is againexcellent and in the design of functionally graded coatings, this is animportant observation.

Themicrostructure appears to bemade up of a dendritic phase and aninterdendritic phase in both coatings. Rather strikingly, while a fewmacro-pores exist for STA, nomacro-pores are seen for STD. Both coatings,however, exhibit “nano”-pores, which are only present in the interden-dritic phase due to shrinkage during solidification. The reason for thereduced amount of nano-pores in the STD coating, compared to the STAcoating is explained because of the higher dilution in the STD coating.

The microstructures actually appear rather simple, but closerinspection reveals that this is not the case. For TEM investigation, wefocus on two single laser tracks with differing microstructures (STA andSTD). TEM shows us that indeed the interdentric phase is ‘small’ – less

he microstructure of coating ‘STD’. (C) and (D) are BSE micrographs revealing excellent

Fig. 3. TEM analysis of NanoSteel single track coatings. In (A) and (C), coating STD is shown and coating STA is shown in (B), (D1,D2) are EDS analyses for the dendritic and interdendriticregions respectively shown in (B).

Fig. 4.OIManalysisof coatingSTDrevealingSEM, ImageQuality (IQ), austenite,martensiteand ferriteaswell as theorientationanalysisof coatingwithanaccompanying inversepolefigurekey (IPF).

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Fig. 5. BSE micrographs revealing of longitudinal cross-sections of coatings MTA, MTD and MTE, with reference to Table 2. The coating has poor contrast between the substrate andcoating due to high dilution.

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than one or two hundred nanometres in thickness in both cases (Fig. 3A,B,C). From the same figures, it is also clear that the dendrites themselves aremultiphase. Phase identificationbyOIM, analysis coupledwithEDS (Fig. 4)of a microstructure characteristic of higher dilution (coating STD) revealsthat the interdendritic phase is made up of austenitic iron. No austenite ispresent in the dendritic phase, but instead a mixture of ferritic iron and

Fig. 6. (A) BSE image of coating MTA revealing two adjacent microstructu

untemperedmartensitic iron exists. Thismartensitic content supports theobservation of the ~100 nm plate-like volumes in TEM analysis of thesame dendrites.

While many single track layers have been produced, in an industrialapplication, the contact area of an engineered part can be much largerthan the width of a single laser track would permit. It is also possible to

res. (B) is an 2θ XRD scan from MTA coating and the original powder.

Fig. 7. BSE image (left up) overlain with EDS maps for the elements Si (green), Cr (yellow), Fe (light blue), Mn (magenta) and Mo (royal blue). (For interpretation of the references tocolour in this figure legend, the reader is referred to the web version of this article.)

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deposit overlapping tracks and multi-layer coatings. Actually, althoughthe coatings presented so far of high dilution appear “crude”, the resultsshow high dilution to be an extremely effective method for eradicating

Fig. 8. OIM micrographs for 2 adjacent areas: left micrographs are image quality maps; ce(bottom) alpha (red) and gamma (green) iron as well as martensite (yellow). The right imreferences to colour in this figure legend, the reader is referred to the web version of this a

(powder induced) porosity. It is proposed that subsequent laserprocessing can be applied for several benefits. Firstly, to produce thickerfunctional layers andsecondly tonegate the substratedilutioneffectof the

ntre images are phase maps featuring (top) alpha (red) and gamma (green) iron andages relay textural information according to the key insert. (For interpretation of therticle.)

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first clad track on the layer composition. It is also well known that thedilution of the second and all subsequent laser tracks is substantiallyreduced when coating is formed by laser track overlapping [18]. Singletrack processing is important in order to ascertain suitable initialprocessing parameters. However, multi-pass processing is necessary inorder to bring theuppermost surface layer closer to the original GFA.Witha focus on the results upcoming, a fewof the coatings are exhibited here –MTA,MTDandMTE inTable 2 and Fig. 5. The coating sections shownhereexhibit excellent interfacial bondingandrevealminimalporosity.Actually,porosity in multilayer coatings was found to be produced by trappedparticles in the overlap areas, and rarely in the meltpool itself.

Fig. 6A shows an area along the longitudinal direction (relative tothe cladding direction) within coating MTA, which has two differingmicrostructures adjacent to one another due to cooling ratedifferences from the melt. XRD analysis (Fig. 6B) allows us to gain ahandle on the phases which may be present; however the resolutionin standard “lab” XRD is not sufficient to pinpoint such a small region.If this area is viewed using OIM and EDS, however, we can addressboth the chemical compositions and the phases present in thecoatings and confront the microstructural evolution.

EDS analysis (Fig. 7) shows that the principal alloying elements ofthe NanoSteel powder under investigation form the interdendriticphase, while the dendrites have a higher iron content. Since the(nominal) composition of A36 steel is 0.29C max, 0.80–1.2Mn, 0.04P,0.05 S, 0.15–0.3Si bal Fe, it is insufficient to say whether the Mn and Sireally show a greater affinity for Fe, or whether their 'stronger'appearance in the dendrites of these elements is due to the substratedilution. Certainly, when viewed in conjunction with the TEM/EDSresults in Fig. 3, it would appear that the latter is the case, and thusarea 1 can be considered as an areawhere the mixing of the NanoSteelpowder and the Fe substrate does not take place.

With reference to the fact that the dendrites appear to be multi-phase in both STA, MTA andMTD, OIM analysis was performed onMTA.The additionalOIManalysis shows that in area1 (Fig. 8), thedendritesdo

Fig. 9. Secondary electron micrographs revealing microstructural overviews for 3 NanoStee‘D’=MTD and (C) BSE-mode SEM for coating ‘E’=MTE with hardness profiles for all three sh

not form martensite. This is indicative of a reduced cooling rate in thisfraction of the coating. The “macroscopic” microstructural evolution inthe coatings consists of uniform ferrite/martensite dendrites and aninterdendritic austenitic phase, made-up of the principle alloyingelements of the NanoSteel powder. The texture of the interdendriticphase is quite remarkable, with the austenitic phase appearing here inlongitudinal cross-section close to (101) (Fig. 8, right).

When the microstructure is refined, by double layer coatings, thedistribution of the elements of the NanoSteel powder form a homo-geneous, almost needle-like microstructure, which will be showntogether with the specific microstructures for MTA and MTD in thecoming results (Fig. 9) and the effects of these variousmicrostructures onmechanical properties and performance will also be exhibited.

For comparison, extensive XRD analysis on the plasma sprayedcoatings [19] after devitrification has shown three phase structurewith ferrite, austenite and M23(BC)6 phases. In laser fusing, since theglass forming region is missed due to cooling rate and effects ofdilution, another solid state transformation can occur leading tomartensite formation.

3.2. Dry sliding wear

The non-amorphous coatings produced using GFA powders have alsobeen investigated with respect to their tribological performance. A largenumber of coatings were produced at various parameters as discussedalready. Three of those differing coatings/microstructures with sampleidentifications, MTA,MTD, andMTE are shown in Fig. 9A,B,C respectively.They were chosen for sliding wear investigations, similar to those [8] forfully amorphous Ti-layers; although here, all counterpart 100Cr6 diskswere polished to rab100 nm. In line with the analysis of the fullyamorphous coatings, micro-hardness profiles are given for the 3 coatingsin Fig. 9D. It can be seen that there is a reduction in dendrite spacing size,according MTANMTDNMTE, but, perhaps surprisingly, there is not such asignificant difference in hardness between coatings MTA and MTD (both

l coatings in: (A) mixed-mode SEM for coating ‘A’=MTA (B) BSE-mode SEM for coatingown in (D).

Fig.10. A)Smoothedslidingwear and friction curves for coatingsMTA,MTDandMTEat1.7MPa contact stress and10 cms−1 sliding speed. B) reveals smoothedhigh-speed (50 cms−1)wearand friction curves for coating MTE. WR refers to the specific wear coefficient for a given coating.

Fig. 11. Mixed mode (50%SE/50%BSE) micrograph revealing micro-wear mechanisms inthe dendritic phase of the 'D' coating tested at contact stress=1.7 MPa, 10 cm s−1 after1800 m. The black lines identify the differing orientations of the micro-wearmechanism. The sliding direction is from left to right as the figure is viewed on the page.

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around800Hv0.2). Thedouble layer coating,MTEhas a significantlyhigherhardness (around 1000 Hv0.2). As stated previously, this second layer wasdeposited with the same processing parameters as MTD, but due to asmaller amount of steel dilution (i.e. all dilution comes from the first-deposited layer), the composition is closer to that of ‘pure NanoSteel’, andhence forms the nano-crystalline (~400 nm) coating shown in Fig. 9C.

According to the Archard wear equation, the coatings MTA andMTD should exhibit the same wear rates. Fig. 10A, shows that this isnot the case, since coating MTA exhibits a higher wear rate. Thefriction coefficient of 'MTA' is higher than that of ‘MTD’ also. Thewear rates of these coatings are, however, quite remarkable. As a

Table 3Hardness (H), elastic modulus (E) and H/E ratios for 3 nanosteel coatings (MTA, MTD,MTE)

H (GPa) E (GPa) H/E

Coating MTA 11.0 250 0.044Coating MTD 11.5 248 0.046Coating MTE 14.2 274 0.052

Fig. 12.Wear observations for high-speed sliding wear of coating MTE: A) reveals debriscollection in pre-existing pores and also behind pre-existing cracks, B) shows the thickoxide layer formed during high-speed sliding.

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comparative measure, we see that the sliding wear resistance of theamorphous layers is around 2×10−6 mm3/Nm, while under similarloads and speeds, these coatings exhibit wear rates around 1 order ofmagnitude smaller.

Fig. 13. BSE micrographs revealing 3-body wear characteristics for (A and C) Coating MTD aentrance, (B) reveals the same phenomenon at the exit, while (B and D) exhibit dense, randdeformed contact area.

Even more remarkable still is the wear behaviour of the doublelayer coating when subjected to high sliding speeds (50 cm/s). As canbe seen in Fig. 10B, this coating was able to withstand 10 km ofunlubricated sliding wear testing without seizure, at a wear rate ofonly 3×10−8 mm3/Nm, which is comparable to Sprayed TungstenCarbide/Cobalt, Sprayed Chrome Oxide, PVD TiN and just shy oflubricated hardened steel [20]. This reduced wear rate is mirrored by areduced, and incredibly stable, friction coefficient of 0.7.

If the wear surfaces are examined by SEM (Fig. 11), one of the mostinteresting features can be seen for coating ‘D’. This coating reveals amicro (or even nano) scale wear mechanismwithin the dendrites. Thisphenomenon appears to be independent of wear direction (note theblack indicators of this in Fig. 11). Since the orientation of the shear isrelated to the dendrite and not the wear direction, this wearmechanism is closely related to slip between martensitic plates (andferrite). This is supported by the TEM and OIM microscopy presentedin Figs. 3 and 8 respectively which reveals fine plate-like volumes inthe dendritic phase, which, when mapped by OIM as martensiteindicates that, the cooling rate is high enough during the laserprocessing to quench γ-phase iron (austenite) as hard martensite.

No discernable preferential wear of the constituent phases wasfound, and therefore equal proportions of the phases were always incontact with the counterface. This plastic micro-wear mechanismobservation may indicate, however, that the shear modulus/Young'smodulus of the dendritic martensitic phase is lower than that of theaustenitic phase. Unfortunately, as the interdendritic phase is only ofthe order of a few hundred nanometers, it is not possible to conductmeaningful nanoindentation experiments in order to discern thedifference in the moduli of the two phases. Average nanoindentationvalues taken from the top 150 µm of the coatings to a penetrationdepth of 500 nm are shown in Table 3. H/E values are included in thistable since they have been cited as possible indicators of good wearresistance (high H/E values should tend to better wear resistance) [21].

The reduced wear rate and friction coefficient for coating MTE athigh sliding speed is extremely encouraging. Despite the coating being

nd (B and D) Coating MTE after 900 revolutions: (A) reveals micro-chipping at the scarom-orientation micro-chipping at the centre of both wear scars, resulting in a heavily

Fig. 14. A) (left) BSE micrograph of half a wear scar produced at 900 revolutions oncoating MTE; (right) SE micrograph of half a wear scar produced at 900 revolutions oncoating MTA. B) Confocal microscopy profile through a wear scar after 600 revolutionsfor coating MTE.

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cracked after processing, this was not found to be detrimental to thewear behaviour (Fig. 12A), since no large macroscopic materialspallation was found to occur. The debris which formed was sub-micron in size and heavily oxidised and, as can be seen in Fig. 12B,forms a layer on the coating surface during sliding in places, andtherefore may be a contributory factor in the reduced frictioncoefficient. In further accordance with the concepts for good wearresistance outline previously, it is no surprise to fine coating MTEhaving the best wear resistance, as it also offers the highest H/E ratio.

3.3. Microscale abrasive wear

Given the exceptional sliding wear performance of the NanoSteelcoatings, microscale abrasive wear tests were performed on a few of thecoatings, with a brief overview of the results given here. Abrasive wearmay take 2 forms – either 2-body, which is aggressive and characterisedby surface grooving, or 3-bodywear,which is ‘milder’ and is characterisedby a wear scar featuring many small chips, or dimples. A third optionexists, where these two abrasive wear forms co-exist. In microscaleabrasion testing, it is possible to vary the concentration of the abrasiveslurry so that one of the wear mechanisms dominates. In this shortinvestigation, we consider only 3 body abrasive wear for the single anddouble layered coatings studied under sliding wear in Section 3.2.According to theprocedureoutlined inSection2, thewear rate forall threecoatings (to 1 significantfigure) is 1×10−12m3/Nm. These values comparefavourably with the work of Navas et al. [22] who performed similarexperiments on laser clad and flame sprayed-melted NiCrBSi coatings.Wear scar morphologies are shown for coating MTD in Fig. 13 (A and C)and for coating MTE in Fig. 13 (B and D). No directionality is seen in the

wear scars for the micro-chipping, which is a characteristic feature of 3-body abrasive wear.

Further significant findings confirm the observation in Section 3.2that although the coatings may contain cracks and/or porosity afterprocessing, this is not detrimental to themechanical performance of thecoatings (Fig. 14A left and right) for the loading conditions permissiblewith the test equipmentavailable. This imageallowsadirect comparisonbetween thewear scardiameter in thedouble layer,MTEcoatingand thesingle layer MTA coating. The larger wear scar of MTA signifies a largerwear volume after equal test revolutions. Significant material pile-upoccurs around the circumferences of the wear scars, as seen underindentation, further confirming that these coating offer minimal work-hardening and significant ductility. Fig.14B displays the profile of awearscar including pile-up around a wear scar for coating MTE after 600revolutions.

4. Conclusions

Laser cladding of pore-free layers from porous glass forming alloypowder pre-cursors has been achieved, with the parameter offeringthe largest influence being high dilution (N65%). The resultant layersare therefore reduced in glass forming ability, but nonetheless, thick(up to 1 mm) uniform coatings with excellent interfacial bonding maybe produced. Coatings of more than one layer thickness realise a finermicrostructure and are yet still able to be processed both pore- andcrack-free.

These coatings evolve with an interdendritic phase of austeniticsteel and a dendritic phase of ferrite and martensite. Higher dilutionleads to higher ferrite/martensite content since the austenitic phaserelates to the NanoSteel powder. In both cases the interdendritic phase‘thickness’ is found to be 100–200 nm.

Radical improvements are attainable in both sliding (adhesive) andabrasive (3-body) wear compared to the substrate material. So muchso in fact that for double-layer coatings, unlubricated sliding wearrates approach those of lubricated hardened steel. The mechanism forthis is found to be the growth and accumulation of thick oxide layerscomprising nano-scale oxidised wear debris.

Acknowledgements

Financial support from the foundation for Fundamental Researchon Matter (FOM-Utrecht) is gratefully acknowledged. Part of thisresearch was carried out under the project number 02EMM17 in theframework of the research program of the Materials innovationinstitute M2i (Adm. Office in Delft, www.m2i.nl), the former Nether-lands Institute for Metals Research.

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