laser melt injection of hard ceramic particles into a1 and ti …€¦ · · 2014-05-20into a1...

Laser melt injection of hard ceramic particles into A1 and Ti alloys - processing, microstructure and mechanical behaviour

V. Ocelik, S. Nijman, R. van Ingen, U. Oliveira & J.Th.M. De Hosson Department of Applied Physics, Materials Science Center and Netherlands Institute for Metals Research, University of Groningen, Nijenborgh 4, 9747 AG Groningen, The Netherlands

Abstract

The conditions for a successful Laser Melt Injection (LMI) of Sic and WC particles into the melt pool of Al8Si and Ti6Al4V alloys were studied experimentally and theoretically by FEM calculations. The laser employed is a high power Ncl:YAG .

The formation of a relatively thick aluminium oxide layer on the AI melt surface was confirmed during in-situ melting in an Environmental Scanning Electron Microscope (ESEM) at temperatures above 900 "C. This oxide layer acts as a barrier for particle penetration but it can be dissolved in the melt at temperatures above 1250 "C and consequently it opens a "window" for particle injection. The finite element analysis of the laser melting process of Al-Si alloy leads to the conclusions that the laser scan velocity has only a small influence on the substrate temperature distribution in the vicinity of the laser beam and that the size of extended part of laser melt pool (which is the best place for injection) is extremely small. Therefore, in contrast to Ti-alloys an extension of a melt pool size behind the laser beam, which serves as an effective instrument for successful LMI of ceramic particles, is not a successful processing route in the case of Al alloys.

The relationship between microstructure, tensile and wear properties has been investigated for SiCIAl-Si and WCRi-Al-V metal matrix composite layers. Although the presence of hard and brittle phases formed during the laser

Transactions on Engineering Sciences vol 39, © 2003 WIT Press, www.witpress.com, ISSN 1743-3533

142 Surface Deatment V1

processing decreases the tensile strength in comparison with substrate materials, a massive improvement of wear resistance of both layers is confumed. As in-situ (ESEM) tensile tests show the crack-initiation process in the WC/Ti-Al-V layer strongly depends on the type of WC powder used in the laser melt injection process.

1 Introduction

The laser surface engineering technique is a suitable instrument for the preparation of thick protective coatings on the substrate of metals and alloys. Recently it has been confumed, that laser coatings with properties of functionally graded materials could be prepared. The term "Functionally Graded Materials" @GM) is widely used for a class of materials exhibiting spatially inhomogeneous microstructures and properties [l]. Gradients can be designed at a micro-structural level to tailor specific materials for their functional performance in possible applications. In particular, we have recently produced FGM by laser cladding [2] and by laser melt injection [3] technique, in which a progressive change in both microstructure and related properties is achieved as a function of depth.

During the formation of FGM by laser cladding process, the non- homogeneous temperature field of the melt pool mainly controls the number of strengthening particles since the growth rate and the time available for growth at different depths affects the final size of strengthening particles. On the other hand in the Laser Melt Injection (LMI) process the temperature/viscosity field of the melt pool explains the formation of FGM by injecting of ceramic particles into the elongated part of the pool behind the laser beam [3]. Therefore the presence of elongated part of the melt pool seems to be an imperative condition for the formation of FGM coating by LMI process in the form of a metal matrix composite layer with a gradual change of the volume fraction of ceramic particles from the top to the bottom of the coating layer.

In general, the performance of a coating depends in many applications on the bond strength between coating and substrate material. In the case of FGM metal matrix composite (MMC) coatings prepared by LMI the bonding between strengthening particles and matrix plays even a more important role for mechanical properties. The goal of the presented work is to study the conditions for the formation of MMC coatings on the Al and Ti alloys substrates by laser melt injection technique as well as the description of microstructural features formed during injection. An experimental evaluation of the behaviour of such layers during severe mechanical testing is the final goal of the work presented. The adhesion between the particulates and metal matrix in MMC surface layers has been tested using in-situ straining in an environmental scanning electron microscope (ESEM) in a combination with orientation imaging microscopy (OIM). The adhesion between ceramic particles and matrix was also tested during wear tests, where MMC surface layers were subjected to boundary lubrication wear conditions.


Surface Treatment V1 143

2 Experimental procedures

A Rofin Sinar CW020-type 2kW Nd:YAG laser was used in this work to produce FGM coatings tracks. The focal length of the lens is 120 mm, but the lens is operated at a defocusing distance of 6 mm with a 2.5 mm spot in diameter on the surface of the substrate. A numerically controlled X-y table executes the movement of the specimen under the laser beam. The powder feeder apparatus is a Metco 9MP instrument.

Commercially available A18Si cast alloy and Ti6A14V alloy were used as substrates in the LMI technique to prepare single tracks of SiCIAl-Si, WCIAl-Si and WCITi-AI-V MMC layers. Sic particles with the mean size of 60 pm, globular multigrain WC particles with the mean size of 80 pm and single grain WC particles with the size of 10-30 pm were injected. A more detailed description of LMI process, in particular the process parameters as well as a detailed description of microstructures and particle distributions formed inside SiCIAl-Si and multigrain WCITi-AI-V layers can be found elsewhere [4-61.

Optical microscopy (Olympus Vanox-AHMT), scanning electron microscopy (Philips XL30 FEG with EDX), orientation imaging microscopy (OIM attachment to Philips XL30s FEG) and XRD diffraction (Philips PW- 1830) were used to study the particle distribution, microstructure and microstructural relationships in WCIAI-Si and WCITi-AI-V layers prepared by injection of multigrain and single crystal WC particles into AI-Si and Ti-6A1-4V alloy substrates.

Flat specimens with a thickness of 0.8 mm and a width of 4 mm were cut from TiAlV substrate with WCITiAlV laser track along the longitudinal direction using laser cutting. The form of V-notched tensile specimen with approximately a working length of 15 mm was chosen where the notches touched the laser track in the middle of the sample. In-situ tensile test samples were loaded by a special stage (Kammrath & Weiss) inside a FEG-E-SEM Philips XL30 (field emission gun - environmental - scanning electron microscope), which allows the observation of the specimen surface simultaneously with the sample loading. The same FEG-E-SEM Philips XL30 microscope has been used in the so-called "wet" mode to observe the surface of aluminium during heating up to 1300 "C.

Samples for wear tests were cut from 10 mm thick substrates (5 mm in case of Ti-AI-V substrate) by spark erosion with dimensions of 10 X 10 X 1 mm3 (10 X 5 X 1 mm3 for Ti-A1-V substrate). Tests were performed at a Plint tribometer TE 67 using pin-on-disk set-up. Stationary specimen was pressed under a known normal force against the surface of rotating disk from 100Cr6 hardened (60 HRc) steel with an initial surface roughness of = 0.1 pm. Specimen and ring were completely submerged in a oil bath (BP Transcal N) at room temperature. The friction force and the change of relative position between pin and ring were measured continuously during 20 hours of wear test. To exclude hydrodynamic lift effects and to perform wear tests under boundary lubrication conditions, wear experiments were executed at a low constant sliding speed (0.01 mls) and with a contact pressure of 20 and 100 MPa. Wear surfaces



were inspected after the test by SEM and confocal optical microscope (pSurf Nanofocus Messtechnik).

3 Results

3.1 Particle injection process

The LMI process is schematically depicted in Fig. 1. Defocused laser beam with the spot size of 2.5 mm in diameter on the substrate is used. Ceramic particles are injected into the laser pool just behind the laser beam. This permits the particles to penetrate in the melt to certain depths and also to avoid undesirable heating up of particles during contact with the laser beam. The key point for successfd and stable injection conditions is therefore to extend the laser melt pool behind the beam, which will provide the necessary space for injection without touching particles to the beam.

a- 10 mmls 1200 W

d g l000

1 .O 1.5 Distance [cm ]

Fig. 1: a) Sketch of the laser melt injection process. b) Temperature profile on the surface of AI-Si block along the laser track axes at the moment when the fiont of laser beam reaches a middle of the block. A dashed line depicts actual position of laser beam. Different scanning speeds and laser power are used.

Usually, this can be achieved by relatively high scanning velocities of the beam v b in a combination with appropriate laser power. To understand the dependence of the size and shape of the melt pool on the laser beam scanning velocity, the finite element analysis (FEA) was performed in 3D. Calculation was realized for YAG laser beam with a top-hat energy distribution and AI-Si plate substrate with actual dimensions used in the injection experiments. The model takes into account the latent heat, but it does not account for melt convection and particles presence. Figure l b summarizes the results of FEA calculations in the form of the surface temperature profile inside a beam and its surrounding for different scanning speeds and laser powers. It can be concluded that for standard values of used power the size of melt pool behind the laser beam is quite small (0.5-1.0 mm) and it does not increase when scanning speed is redoubled. Only the increase of laser power to the value of 2400 W (which is too large from a technological point of view) leads to a formation of reasonable extension, as Fig. lb also demonstrates. This result is substantially distinct from the results



obtained for Ti6A14V where the same increase in scanning speed induced an increase of size of extended melt pool by a factor 4 reaching values of 3-4 mm [3]. It should be realized that the extension of the laser melt pool strongly relies on the value of the thermal conductivity (6.7 w ~ - ' K - ' for Ti-AI-V alloy and 121 wrn-'K1 for AI-Si alloy).

The second obstacle for a successful laser melt injection in aluminium and its alloys may be the presence of a strong oxide layer on the top of the melt. Kaptay [7] studied the behaviour of this oxide skin, as a fimction of temperature and concluded that the oxide skin is present on the melt till about 850 'C. Above this temperature chemical reactions between A1 and the oxide at the interface occur, leading to the formation of gaseous sub-oxides. Consequently, the oxide gradually disappears and the A1 melt at the surface takes over with increasing temperature by the growth of A1 islands. At about 1100 'C the oxide is totally removed, according of analysis of indirect observations [7]. Environmental scanning electron microscopy provides the possibility to realize direct experimental observation. We focussed on the surface of a 1 mm thick pure A1 (0.999 at%) disk ,3 mm in width, during the heating inside E-SEM chamber at a pressure of 1.9 102 Pa of Ar. The image is detected with a positively biased Gaseous Secondary Electron detector (GSE mode). Results of these observations are summarized in Fig. 2.

Fig. 2: FEG-SEM from in-situ melting of Al. a) Formation of aluminium oxide skin b) Al-melt bubbles penetration c) Oxide skin dissolution

The temperature in the chamber increases in steps of 50 "C keeping the temperature fixed during 10 min so that the reactions on the substrate can equilibrate. At the highest temperatures the temperature is further increased manually to reach the maximum value, at which it is still possible to obtain



images despite the presence of strong thermal electrons current. The latter increases exponentially with temperature exponentially, according Richardson- Dusham formula [8]. Figure 2a demonstrates the growth of an oxide skin layer that is observed above 900 "C and covers the whole surface at 1050 "C. At this temperature the bubbles of A1 melt start to penetrate to the surface at some intersections of oxide skin facets as Fig. 2b shows. The last three micrographs in Fig. 2c confirm the dissolution of oxide skin layer with time at the highest reachable temperature of 1280 "C. FEG electron beam placed on the dagger mark stimulate the dissolution at this temperature.

The difficulties with the extension of the melt over a small size could be solved by preheating the substrate, as it was demonstrated in the study of S ic particles injection in A1 and its alloys [4]. On the other hand the presence of a strong oxide skin layer seems to be a more serious obstacle for the successful injection with a wider process parameter window.

b) longitudinal cross-section in the centre of laser track.

Figure 3a clearly demonstrates, that at the sides of the laser track, where the temperature of the melt did not exceed 1200 "C substantially, the oxide skin covers the melt sufficiently. As a result 30 pm sized WC particles cannot penetrate. Longitudinal cross-section on Fig. 3b shows instabilities formed on the surface during the laser melt injection, that result in irregularities in particles density along the laser track.

4: a) Optical micrograph of perpendicular cross-section of single laser track of WClTi6A14V FGM coating prepared by laser melt injection. b) FEG-SEM of longitudinal cross-section of WCITi6A14V FGM laser track (back scattered mode after 'inversion" and "threshold' image manipulations).

Next, a hctionally graded WCITi6A14V MMC coating was prepared by laser melt injection of single grain WC particles in Ti6A14V alloy substrate. Figure 4a shows a polished cross section of this layer observed in optical microscope. A successful combination of laser melt injection parameters was: the laser power density of 79 w/mm2, the scanning speed of 10 mm/s and the powder feeding rate of 50 mgls. Similarly, as in the case of injection multigrain



WC powder into the same substrate [6] a wider range of processing parameters window exists, which allows to obtain tracks of different characteristics and a production of wider coatings by an overlap of adjacent laser tracks. The width of a single laser track is about 1.9 mm and the maximum injection depth reaches a value of 0.7 mm. To describe the distribution of WC particles inside the laser track the longitudinal cross-sections were also made along the laser track. An example is shown in Fig. 4b. The method of quantitative metalographic analysis was performed, which is described elsewhere [3]. The profile of volume fraction of WC particles on both perpendicular and longitudinal cross-sections shows a similar behaviour. It is almost constant (35-40%) from the surface until a depth of about 200 pm and then a linear and parabolic decrease follows in the perpendicular and longitudinal cross-sections, respectively.

3.2 Microstructure of WCITi6A14V FGM coating

The microstructural analysis reveals the formation of new phases, similarly as during the injection of fused multi-grained WC particles into the same substrate [9]. The W2C (Trig. ~ 3 m 2 ) layer surrounds each WC (Hex. ~ 6 m 2 ) particle and it is followed by the T i c (Cub. Fm3m) layer. Individual T i c dendrites surrounded by Ti-alloy matrix are observed in the vicinity of injected WC particles. The thickness of W2C and T i c reaction layer depends on the injection depth. Near the FGM surface they are both about 0.1-0.2 pm thick and their thickness increases almost linearly with the injection depth to reach the maximum values at the bottom of laser track; 1 pm for W2C and 1.8 pm for T i c layer, respectively. These microstructural features are shown in Fig. 5.

Fig. 5: a) 1:EG-SEM micrograph of the microstructure formed around individual WC particle injected into Ti6A14V matrix. W2C and Tic reaction layer surround WC particle. Individual Tic dendtrites are present in vicinity of injected particle. Orientation relationships detected by OIM are depicted. b) FEG-SEM micrograph of the crack initiation in WCJTi6A14V FGM coating during in-situ tensile test.

Figure 5a also demonstrates the crystal orientation relationships detected by OIM. Quite often the orientation relationship in which the close-packed planes and directions are parallel between W2C and Tic crystals is observed:


148 Surface neatment V1

The single crystal WC particles have their facets oriented in accordance with the crystallographic planes and therefore another orientation relationship is observed quite often, between WC particles and W2C layer:

As Fig. 5a demonstrates, both orientation relationships are fulfilled on the left and the right side of WC particle, but only orientation relationship (1) is observed on its upper and lower side.

3.3 Mechanical performance of SiCIAl-Si and WUTi6A14V layers

As in-situ observations in ESEM during tensile tests of WCITi6A14V FGM indicate, the new phases formed at the vicinity of WC particles play a very important role in the crack initiation and the crack propagation processes. The crack always initiates through cleavage of T ic at a macroscopic tensile stress of about 60% of final fracture stress (of, - 650 MPa). Sometimes an individual dendrite close to WC particle fractures first and subsequently a new crack appears in Tic layer, which surrounds the particle. Sometimes cleavage fracture appears directly in the Tic layer. Figure 5b shows the initiation of brittle cracks on both sides of the strengthening particle inside Tic layer. Both cracks propagate further into W2C phase without any change of direction. As it was c o n f i e d in subsequent OIM observations, further crack propagation depends mainly on the presence of orientation relationship between W2C and WC phases (2). The crack usually propagates without any change of direction into WC particle and it cleaves completely, when an orientation relation is present. In the case, when the orientation relation is not fulfilled, the crack is very often deflected at the W2C/WC interface and the cleavage changes into the brittle fracture of this interface, as Fig. 6 clearly demonstrates.

Fig. 6: FEG-SEM micrograph of the crack deflection on W2C/WC interface without orientation relationship between W2C layer and WC particle



In the next step of failure process the same is repeated around a particle, which lies on the macroscopic path of the main crack. Finally, when all T ic dendrites on the path between two cracked particles are already broken and when the local stress concentration exceeds the strength of Ti-alloy matrix a macrocrack is formed through ductile fracture of Ti matrix. This macrocrack forms a new stress concentration at its tip and thus the cracking of the near embedded particles appears and the whole process is repeated and accelerated. Macroscopically the failure mode is brittle fracture.

Both SiCISi-A1 and WCITiAlV layers were tested in wear performance. Figure 7 shows the sliding distance dependence of friction coefficient and wear observed during pin on disk wear tests on SiCIAl-Si MMC layer, WCITi-AI-V MMC layer with multigrain WC injected particles and WCITi-AI-V MMC FGM coating with single grain WC injected particles. For a comparison the same tests were performed on not treated AI-Si and Ti alloys surfaces and on the same surfaces just remelted by laser beam, using the same laser processing parameters as during the laser melt injection. Some of data curves are smoothed using the Savitzky-Golay linear smoothing procedure [l01 to enhance the clarity. Some of them are plotted as recorded to demonstrate experimental data spread.

0 l 0 0 200 300 400 500 600 700 Sliding distance [ m ] Sliding distance [ m ]

Fig. 7: Friction coefficient and hear versus sliding distance for SiCIAI-Si and WClTi6Al4V coating prepared by laser melt injection in comparison with non-treated surfaces: l-basic alloy surface; %-basic alloy surface remelted by laser beam; 3-MMC layer tested at 20 MPa; 4-MMC layer tested at 100 MPa; 5-WClTi6A14V FGM layer with single grain WC particles tested at 100 MPa; e -experimental data, s -smoothed data.

Because of very intensive wear, the tests for untreated and laser beam remelted Ti6A14V substrate were aborted after 3 and 4 hours, respectively. It is possible to conclude, that for non-treated (marked as 1) and laser beam remelted (marked as 2) surfaces, a transition period is present during which the wear rate is not constant andtor the friction coefficient is substantially changed. This behaviour is not observed in MMC samples (marked as 3, 4 and S), where both the friction



coefficient and the wear rate seem to be not substantially changed during the wear process. Table 1 summarizes the test conditions and the results for all wear measurements and shows the specific wear rate coefficient ks [10-~ mm3/~m] estimated from the slope of wear curves in Fig. 7 inside the parts with "constant" wear rate. The value of specific wear rate coefficient km can be determined also from the mass loss of the pin measured by a weighting it before and after the test, using known densities of the substrate alloy and particles and assuming the known volume fraction of ceramic particles in the composite layer. Although the k values determined by these two methods seem to be slightly different for the same samples, they both clearly demonstrate a substantial increase in wear resistance of the composite layer in comparison with AI-Si or Ti-AI-V substrate. The benefits in reduced wear rate in MMC can be described by a "relative wear resistance" (RWR) defined as the wear rate of the unreinforced matrix divided by that of the coating under the same wear conditions [ll]. RWR improvement reaches values of 28-32 for SiCIAlSi and 500-1500 for WCITi-AI-V metal matrix composite, respectively. The values of friction coefficient in Table 2 show, that according to its aim, the boundary lubrication condition prevailed. Wear loss of disks in all experiments were under the detectable value, e.g. kdsk < 0.01 X 10'~ mm3/~m.

Table 1: The specific wear rates k, and km [106 mm3/~m] and the average friction coefficient fa measured during the pin on disk wear test at a different contact stress. Sample 5 represents FGM coating with single grain WC particles injected.

Sample Contact stress WPaI ks km f a

1 AI-Si alloy 20 6 13 0.14 2 Al-Si alloy remelted 20 0.05 0.02 0.10 3 SiCIAl-Si layer 20 0.21 0.4 0.09 4 SiCIAl-Si layer 100 0.24 0.04 0.13 1 Ti6A14V alloy 20 269 189 0.21 2 Ti6Al4V remelted 20 240 40 0.18 3 W C i T i 6 W layer 20 0.50 0.13 0.11 4 WC/Ti6Al4V layer 100 0.05 0.08 0.12 5 WCITi6Al4V FGM 100 0.038 <0.01 0.12

SEM and EDS observations examined the morphology and qualitative composition of worn surfaces of all samples. As expected, the hard aluminium oxide layer formed during laser remelting causes high wear improvement. The worn surface of SiCIAlSi and WCRi6A14V layers tested at 20 MPa contact stress shows an eroded wear surface morphology between particles, with a few of them towering from the worn surface. Figure 8 demonstrates a couple of such particles on both worn surfaces. The black arrow indicates the highest place of


Surface Treatment V1 15 1

the uphill area with the ceramic particle, which was in a direct contact with the disk just before wear test was stopped. For the WCITiAIV layer the total size of such areas forms only a small fraction of the whole worn surface and therefore, the actual contact stress was much higher than the value of 20 MPa calculated from the applied normal force and sample dimensions. This is probably the reason why macroscopic cracks and brittle fracture features are present in the surrounding area of the particle. It happens when local contact stress exceeds 300 MPa, which is approximately the strength of T i c layer [l21 in the particle surrounding. The BSE mode helps us to highlight the presence of iron-based debris collected as a micro-cutting products near Sic particles on SiCIAlSi worn surface (Fig. 8a) and to distinguish between WC particle and its surrounding on WCITiAlV worn surface (Fig. 8b).

Fig. 8: SEM micrographs (BSE mode) of worn surfaces of (a) SiCIAlSi and (b) WCITi6Al4V coating tested at contact stress of 20 MPa. The white arrows indicate the wear direction. Black arrows show the top of strengthening particle areas, which were in a contact with disk just before the end of wear test.

Figure 8 also demonstrates that WC particles are worn mainly by intergranular brittle fracture whereas the wear of Sic particles proceeds by a combination of two mechanisms (the slow abrasive wear on the contact surface combined with the fast cleavage out of outstanding SIC particles). When the same WClTi6A14V layer was tested at a contact stress of 100 MPa, the worn surface shows a relatively flat morphology near the front edge. In the second part of the worn surface a rougher area with deeply eroded zones between embedded WC particles was present again. The total number of WC particles counted on the worn surface is still quite high.

A completely different picture can be observed on the worn surface of WClTi6A14V FGM tested at contact stress of 100 MPa (see Fig. 9b). The density of WC particles on the worn surface corresponds to the density observed in microstructural observations, which means that all embedded WC particles join together in the wear process and stay on the surface without any loss. The insert in Fig. 9b clearly demonstrates that the main wear mechanism is now controlled by a local damage of WC particles through an accumulation of plastic deformation in a system of slip bands. The orientation of these slip systems towards the wear direction differs from one single grain WC particle to another



and therefore the damage "features" on particle surfaces are also different. The insert in Fig. 9b shows the small triangles formed at the intersection of two slip systems as an initial damage place. Traces of localized plastic deformation are also visible near the edge of S ic particle on SiCIAl-Si worn surface, as Fig. 9a clearly demonstrates.

WClTi6Al4V FGM coating tested at contact stress of 100 MPa. The white arrow indicates the wear direction. Insert in (b) is the high resolution SEM micrograph fsom the surface of WC particle.

4 Conclusions

Laser melt injection is a suitable technique for the production of protective layer on top of light metals in the form of MMC coatings. If the melt pool can be extended behind the laser beam MMC layers can be created with a gradually change of volume fraction of ceramic particles. The formation of the oxide skin on the aluminium melt is a serious obstacle for the successful injection into aluminium and its alloys.

During the injection of single grain WC particles into Ti6A14V substrate new phases are formed on the particlelmatrix interface, i.e. T i c and W2C in the form of a reaction layer. The thickness of this layer increases with injection depth. The presence of orientation relationships between the WC particle and the new phases formed around it plays an important role in the crack initiation and propagation processes.

The excellent bonding between strengthening particles and Ti6A14V matrix was confirmed using in-situ tensile test. The sliding wear test at boundary lubrication conditions confirmed bonding for WCITi-AI-V coatings and for SiCIAI-Si coating. This leads to the improvement of wear resistance (28-32x) of SiCIAlSi coating and to the extreme improvement of wear resistance (500- 1500x) of WCITiAlV coating in comparison with not treated alloys surface.

Acknowledgement

This project was financed by the Netherlands Institute for Metals Research.



References

Shiota, I. & Miyamoto, Y. (eds.) Functionally Graded materials, Elsevier, Amsterdam (1997). Pei, Y.T. & De Hosson, J.Th.M. Functionally graded materials prepared by laser cladding, Acta Mater. 48, pp. 2617-2624,2000. De Hosson, J.Th.M. & Pei, Y.T. Funcionally graded materials with laser cladding, in: Sugace treatment V , Computer Methods and Experimental Measurements for Surface Treatmen Effects, ed. C.A. Brebbia, Wessex Institute of Technology, UK, WIT Press, Southampton, Boston 2001,221- 232. Pei, Y.T., Ocelik, V. & De Hosson, J.Th.M. SiCP/Ti6A14V functionally graded materials prepared by laser melt injection, Acta Mater. 50, pp. 2035-205 1,2002. Ocelik, V., Pei, Y.T. & De Hosson, J.Th.M. Interfacial adhesion of laser clad functionally graded materials, ASM Surface Engineering Congress, October 7-10, 2002, Columbus, Ohio. Vreeling, J.A., Ocelik, V., van Agterveld, D.T.L., Pei, Y.T. & De Hosson, J.Th.M. Laser melt injection in aluminium alloys: On the role of the oxide skin, Acta Mater. 48, pp. 4225-4233,2000. Ocelik, V., Vreeling, J.A. & De Hosson, J.Th.M. EBSP study of reaction zone in SiCIAl metal matrix composite prepared by laser melt injection, J.Mater.Sci., 36, pp. 4845-4849,2001. Vreeling, J.A., Ocelik, V. & De Hosson, J.Th.M. Microstructure characterization of laser melt injected WC particles in Ti-6A1-4V, in: Surface Treatment V; Computer methods and Experimental Measurements for Surface Treatment effects, ed. C.A. Brebbia, WITPress 2001, Southampton, 253-262. Kaptay, G. On surface properties of molten aluminium alloys of oxidized surface, Materials Science Forum, 77, pp. 3 15-330, 1991. Ashcroft N.W. & Mermin N.D, Solid state physics, Holt, Rinehart and Winston, New York 1976, 826 p. Vreeling, J.A., Ocelik, V. & De Hosson, J.Th.M. Ti-6A1-4V strengthened by laser melt injection of WC, particles, Acta Mater. in press.

[l01 Savitzky, A. & Golay, M.J.E. Smoothing and differentiation of data by simplified least squares procedures. Anal. Chem. 36, pp. 1627-39, 1964.

[ l l ] Hutchings, I.M., Wilson, S. & Alpas, A.T., Wear of aluminium based composites, in: Comprehensive Composite Materials, ed. In chief A.Kelly and C. Zweben, Amsterdam 2000, 3.19

[l21 Engineered Materials Handbook, Vo1.4: Ceramics and Glasses, ASM International, 199 1.


laser melt injection of hard ceramic particles into a1 and ti …€¦ · · 2014-05-20into a1...

Documents