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Ti6Al4V functionally graded materials produced by laser melt injection
Acta Materialia 50 (2002) 2035–2051 www.actamat-journals.com
SiCp/Ti6Al4V functionally graded materials produced by laser melt injection Y.T. Pei, V. Ocelik, J.Th.M. De Hosson ∗ Department of Applied Physics, Materials Science Center and the Netherlands Institute for Metals Research, University of Groningen, Nijenborgh 4, 9747 AG Groningen, The Netherlands Received 2 January 2002; received in revised form 25 January 2002; accepted 25 January 2002
Abstract With a well-controlled laser melt injection (LMI) process, for the first time the feasibility is demonstrated to produce SiC particles (SiCp) reinforced Ti6Al4V functionally graded materials (FGMs). SiCp are injected just behind the laser beam into the extended part of the laser melt pool that is formed at relatively high beam scanning velocities. The process allows for the minimization of the decomposition reaction between SiCp and Ti6Al4V melt, and also leads to FGMs of SiCp/Ti6Al4V instead of a homogeneous composite layer on Ti6Al4V substrates. An injection model is designed based on the temperature/viscosity field of the laser pool for a deeper understanding of the mechanism of formation of the FGMs with LMI. The model is based on finite element calculations of the temperature field in the melt pool, physical considerations of the LMI process and it is supported by experimental observations. Three types of reaction layers are observed around SiCp, namely a thin monocrystalline TiC layer, a cellular polycrystalline TiC layer and a thick mixed layer of TiC with Ti5Si3. Among them, only the monocrystalline TiC layer exhibits particular orientation relationships (ORs) to the SiCp lattice, i.e. (111)TiC(0001)SiC and具110典TiC具11¯ 00典SiC or (111)TiC(101¯ 2)SiC and 具11¯ 0典TiC具12¯ 10典SiC. These two kinds of TiC reaction layers act as a barrier against the interfacial reaction and its swift formation during rapid cooling hinders the dissolution of SiCp in the Ti-melt. 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Functionally graded materials; Laser treatment; Microstructure; Metals; Carbides
1. Introduction The recent resurgence of interest in materials with spatial gradients in composition and structure has been driven by the fact that these materials exhibit characteristics not attainable by conven∗ Corresponding author. Tel.: +31-50-363-4898; fax: +3150-363-4881. E-mail address: [email protected] (J.Th.M. De Hosson).
tional materials. Recent theoretical and experimental work has established that controlled gradients in mechanical properties offer attractive challenges for the design of surfaces with resistance to contact deformation and damage. In particular the damage and failure resistance of surfaces to normal and sliding contact or impact can be changed substantially through such gradients. For a recent review see Ref. [1]. SiC reinforced metals are appropriate candidates for enhanced tribological applications, e.g. the automotive industry, but
1359-6454/02/$22.00 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 2 ) 0 0 0 4 9 - 6
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the mechanical performance strongly depends on the degree of dissolution of SiC in the matrix and the type of reaction layer [2]. Abrasive wear and sliding wear tests show a good relative wear resistance of SiC particles (SiCp) embedded in Ti-6Al4V produced by laser treatments [2,3]. The wear rate has been found equal to 10⫺8 mm3/Nm. Injection of ceramic particles into a laser melt pool provides an advanced method to tailor a composite layer onto a surface of a metallic substrate [2–5]. The most important factors that dominate the success of the injection process and consequently the properties of the composite layer are the following: (a) particles dissolution and reaction with the matrix melt at high temperatures; (b) distribution and volume fraction of the injected particles; and (c) thermal stresses built up in the composite layer during cooling of the melt pool. Regarding the first aspect, heating of ceramic particles under direct irradiation of a laser beam promotes particle dissolution/reaction and therefore it is undesirable. Further, it has been shown that a thick reaction layer around the particles can be formed and the reaction products in the matrix result in a brittle behavior of composite layer that affects the transfer of load from the matrix to the particular reinforcements [2–5]. In a special case, e.g. laser melt injection (LMI) of SiCp into Al-alloys [4], the particles could be injected only near the center of the Al melt pool where the oxide skin of the Al melt was partially broken under the irradiation of the beam. In such a situation, heating up of the particles by the laser beam could not be avoided. Regarding the second factor, a homogeneous particle distribution along the injection track is a primary requirement for LMI, which can be obtained with a stable particle stream from an injection nozzle and a constant powder-feeding rate. Due to the fast solidification process of the laser pool, clustering of the particles in the melt does not appear under optimum injection conditions [6]. In order to reduce the residual stresses involved a gradual increase in volume fraction of the injected particles from the bottom toward the surface of the composite layer is highly desirable, which, to the best of our knowledge, has never been obtained by a single-step injection process. In our previous work in the field of LMI [2,4]
the dissolution of the SiCp and the amount of the reaction products in the SiCp/Ti6Al4V composite were clearly reduced by positioning the powder stream partially out of the beam. This idea is extended further in this paper not only allowing for a minimization of the reaction between SiCp and Ti melt, but also to obtain a functionally graded material (FGM) of SiCp/Ti6Al4V. The work is divided into two parts. The focus is on the laser injection process of the FGMs and the interface reaction between SiCp and Ti6Al4V melt. For a better physical understanding of the mechanism of formation of the FGMs a model is designed which is based on the temperature/viscosity field of the laser pool. The model is based on finite element calculations of the temperature field in the melt pool, physical considerations of the LMI process and it is supported by experimental observations. In addition, the microstructure and crystallography of the reaction layer around SiCp are scrutinized with orientation imaging microscopy (OIM) and scanning electron/scanning Auger microscopy (SEM/SAM). In a subsequent paper, the defect structure in the reaction layer and the interface bonding are emphasized.
2. Experimental details The experimental setup of LMI has been documented in detail in Ref. [2]. A Rofin Sinar 2 kW Nd:YAG laser is employed for the injection experiments. A numerically controlled X–Y table executed the specimen movement. The powder feeder used is a Metco 9MP-type commercial instrument. The processing parameters varied in the range of 600–1000 W laser power, 500–1200 mm/min beam scanning velocity and 1–3 g/m powder feeding rate. Selection of SiC/Ti6Al4V is based on the fact that Ti6Al4V has a very low thermal conductivity, which leads to the possibility of elongating the laser pool along the beam scanning direction. As a consequence it allows for the injection of SiCp outside the beam and also to obtain FGMs. Ti6Al4V substrates of dimensions 50 × 40 × 5mm3 are cut with electrical discharge
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erosion from a rectangular bar. The surface of the substrates did not experience any further post finishing after cutting and it was just cleaned by acetone prior to LMI. Monocrystalline 6H-SiC powder with a size distribution ranging between 50 and 90 µm is used. The particles are injected into the melt pool at an angle of 35° with respect to the surface normal. The injection velocity of the SiCp particles is measured by recording their flight out of the powder nozzle with a Kodak 4540 Ektapro highspeed camera at speed of 27 × 103frames / s and turns out to lie in the range between 2 and 4 m/s. No pre-heating of the substrate is necessary for the injection, which is different from the situation of SiCp injection into Al [4]. Transverse and longitudinal sections of the injected composite tracks are cut for microstructural studies and quantitative metallographic analysis. The polished specimens are etched with Keller’s reagent for 10–20 s at room temperature. A Philips XL-30s FEG scanning electron microscope (SEM) equipped with an electron backscattered pattern-orientation imaging microscope (EBSPOIM) and an energy dispersive spectrometer (EDS) is employed for the combined microstructure and crystallography analysis. For the operational principle of OIM as well as the definition of the sample reference frame reference is made to [7]. The composition of the reaction layers is analyzed by a dedicated small-spot and ultrahigh vacuum (UHV) scanning electron/scanning Auger microscope (SEM/SAM) derived from a JEOL JAMP 7800F.
3. Results 3.1. Particle injection process The LMI process is schematically depicted in Fig. 1. An over-defocused beam is used and the focal spot of the beam lies 9 mm above the substrate surface, which represents a spot of about ⭋3mm. SiCp are injected into the laser pool just behind the beam in such a way that the powder stream is positioned close to the beam, but without touching it. This permits the particles to penetrate in the melt to certain depths and the method also
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Fig. 1. Sketch of the laser particle injection process. The white arrows schematically indicate the maximum injection depth of SiCp at different positions from the starting point A to the end B of the injection process.
avoids reaction of the particles with the melt at higher temperatures. A special powder nozzle developed for laser powder cladding and welding [8] satisfies all the demands of a stable injection process. The nozzle possesses a co-axial flow of helium shielding gas not only to protect the laser pool from the atmosphere but also to focus simultaneously the powder stream. A stable flow and straight wire-like shape of the powder stream from the nozzle are essential for controlling the injection process and obtaining a homogeneous distribution of the injected particles along the track. The key point for the realization of the injection conditions is to extend the laser pool backwards behind the beam as far as possible, which will provide the necessary space for particles to inject without touching to the beam. This is achieved with the aid of relatively high scanning velocities of the beam in combination with appropriate laser powers. For understanding the dependence of the size of the pool on the beam scanning velocity, finite element analyses (FEA) are executed with a three-dimensional model attaining an estimate of the shape and the temperature profile of the laser melt pool. The simulation is performed for a YAG laser beam with a top-hat energy distribution and a Ti6Al4V plate with the actual dimensions of the substrates used in the experiments. The model takes into account the latent heat but it does not account for convection and particles in-take. The results are summarized in Fig. 2 and in Table 1. It appears that both the length (L) and the extended part length (Le) of the pool behind the beam increase with increasing beam velocity at constant
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Fig. 2. Temperature field and dimensions of the laser pool obtained by FEA modeling at a constant power density (I0) and different beam velocity of: (a) 0.6 m/min; and (b) 0.9 m/min. Table 1 Lengths and surface temperature of laser melted Ti6Al4V pool obtained from the FEA modeling vb (m/min) 0.6 Le (mm) L (mm) Le /D L /D Temperature at X=0 (°C)
1.02 2.95 0.34 0.98 2418
0.9
1.2
2.45 4.53 0.82 1.51 3007
3.83 5.93 1.28 1.98 3717
laser energy density (I0), which is defined as the product of the laser power and the irradiation time on a local area divided by the area of the beam, according to: I0 ⫽
4P πDvb
(1)
where P is the laser power, D the spot size of the beam and vb the scanning velocity of the beam. For instance, the ratios of L/D and Le/D change from 0.98 to 1.98 and 0.34 to 1.28, respectively, by increasing the laser beam velocity from 0.6 to 1.2 m/min. Although the FEA result may differ slightly from the real temperature field due to the simplicities made in the model, it gives a clear indication that a powder stream of approximate the
same diameter as the beam spot can be injected in the extended part of the pool at high beam velocities without making contact with the beam. This set-up does not disturb the formation of a stable laser pool and favors a well-controlled injection process. Therefore, all the problems associated with the flight of SiCp through the beam, such as heating up the particles, shadowing the pool and increasing the temperature of the melt, will not arise anymore. It should be realized that the extension of the laser pool strongly relies on the low thermal conductivity 6.7 W/m K of the Ti6Al4V alloy with respect to other alloys or metals. 3.2. Particle distribution and FGM formation Fig. 3a shows a polished surface of a SiCp/Ti6Al4V composite track produced with LMI, which is prepared for tensile tests in such a way that the convex surface (part) of the track has been just polished away to make it flat and coplanar with the substrate surface. The picture gives an overview on the overall quality of the composite material in which SiCp are embedded over the whole molten area of the substrate. The particles are distributed homogeneously along the track. One of the striking findings is the fact that the particle distribution at different depths can be controlled by the laser beam scanning velocity and the injection position. As shown in Fig. 3b, SiCp are
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Fig. 3. (a) Plane view of a SiCp/Ti6Al4V FGM track produced with laser particle injection; (b) cross-section of a track injected at 0.5 m/min beam scanning velocity; (c) and (d) cross-section and longitudinal section, respectively, of an FGM track produced at 0.8 m/min beam scanning velocity.
distributed much more evenly over the whole cross-section of the composite track produced at 500 mm/min. In contrast, the track injected at 800 mm/min has more SiCp on the top area and fewer particles at the bottom as shown in Fig. 3c, that is to say, a FGM is formed. A typical longitudinal cross-section of the FGMs displayed in Fig. 3d clearly demonstrates the gradual distribution of SiCp along the depth direction of the composite
and the smooth transition between the composite layer and the substrate. To characterize the particle distribution as a function of depth, quantitative metallographic analysis is carried out with the aid of computer software that works on the principle of image analysis. A digitized picture of a longitudinal cross-section of the FGM tracks, e.g. Fig. 3d, is first divided into a series of horizontal strips with
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a width of 100 µm, which is just slightly larger than the maximum particle size. The software measures the pixels occupied by SiCp and subsequently the ratio of the particle-occupied pixels to the total pixels of a strip area is calculated. This ratio is then assigned as the equivalent volume fraction of SiCp to the depth, where the middle of the strip is located on the longitudinal cross-section. In this way, a profile of the particle distribution along the depth direction of a composite track can be drawn as displayed in Fig. 4. It is clear that there is a transition from a homogeneous composite material, produced at 500 mm/min beam velocity, to the graded composite, i.e. FGM obtained at beam velocities over 800 mm/min. The higher the beam scanning velocity the steeper is the slope of the profile. Because the powder-feeding rate (g/m) is fixed for different beam scanning velocities, however, the total amount of the injected particles in those tracks maintains nearly the same, which is equal to the corresponding area of volume fraction versus depth for the profiles in Fig. 4. The formation of the FGMs is most likely controlled by the laser beam scanning velocity, if the powder stream is strictly focused onto the extended part of the laser pool and the injection velocity of the particles is appropriate.
Fig. 4. Profile of SiCp volume fraction as a function of the depth in the FGM tracks produced at different velocities of beam scanning.
3.3. Microstructure of FGM and reaction layers One of the main effects by shifting the powder stream fully out of the beam is the minimization of the dissolution of SiCp in the melt. Nevertheless, as SiCp slightly dissolve during penetration in the melt the matrix composition changes depending on the temperature of the melt. The matrix consists mainly of α-Ti dendritic grains that exhibit a lenticular martensite structure and interdendritic eutectics. In addition, some faceted Ti5Si3 particles in a micrometers range as well as a few TiC freedendrites are observed in the top area of the composite track. The latter are rarely found near the bottom of the track. Indeed, the total amount of these reaction products is considerably reduced in comparison with the results described previously [2]. A reaction layer surrounds always the injected SiCp. Three types of reaction layers can be distinguished: (a) a thin monocrystalline TiC layer with a particular OR to SiCp; (b) a cellular polycrystalline TiC layer; and (c) a mixed layer of TiC and Ti5Si3 particles. The monocrystalline TiC layer shown in Fig. 5a appears relatively smooth, few hundred nanometers thick and it may grow on one or a few flat facets of a SiCp. A SiCp is not fully surrounded by this type of TiC layer due to the orientations of different facets of the particle. The SiCp that exhibit a monocrystalline TiC reaction layer are mostly located in the top region of the injection track. The cellular TiC layer depicted in Fig.5b forms the majority the reaction layers (also see Ref. [2]), and can be found on most of the SiCp that are distributed everywhere over the whole cross-section. The third albeit minor type of reaction layer is more complicated and its microstructure is revealed in Fig. 6. It is composed of more or less spherical TiC particles and faceted Ti5Si3 particles together with a broken TiC ‘skin’ at the outside. In particular, this complex structure does not form a dense layer of TiC and Ti5Si3 particles in the Ti matrix. The increase in size of the TiC particles from the inner interface towards the outer TiC skin clearly indicates a continuous nucleation and growth feature of the layer, as seen in Fig. 6a. Such a mixed reaction layer is often found on the SiCp near the surface of the composite tracks and
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Fig. 5. SEM micrographs showing: (a) monocrystalline TiC reaction layer formed at a flat facet of SiCp; and (b) cellular polycrystalline TiC layer surrounding SiCp.
may partially cover a SiCp. Moreover, it frequently shows a transition from the monocrystalline TiC layers, as shown in Fig. 6b and c. In another words, the mixed layer originates at the sites of the monocrystalline TiC layer previously formed. It is interesting to note that the mixed layer stimulates the SiCp to dissolve continuously, whereas the monocrystalline TiC layer prevents further dissolution of the SiCp as soon as it forms on the facets. As revealed in Fig. 6c and d, the protrusions on the SiCp facets indicated by arrows, are still covered by a monocrystalline TiC layer and the concave parts adjacent on both sides of the protrusions are exposed to the mixed layers. The step height of 1 µm between the protrusion and the adjacent concaves points to the stimulated dissol-
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ution of the SiCp. Obviously, the mixed reaction layer is much thicker than the other two types of reaction layers due to the different mechanisms of formation. Because TiC is not a phase that has a stoichiometric composition, the carbon content of the TiC reaction layers is analyzed with SEM/SAM in order to reveal its formation. Before taking Auger electron spectra, the surface of the specimens is cleaned in situ by Ar+-ion sputtering under UHV. The intensities of Si, C and Ti elemental peaks in the Auger electron spectra are analyzed. Both the area mapping and line profile of carbon intensity presented in Fig. 7 clearly show a rapid decrease of the carbon content in the monocrystalline TiC layer from the SiCp side to the matrix side. In contrast, the polycrystalline TiC layer exhibits less change in carbon content over the thickness, as seen in Fig. 8. This implies different conditions of carbon exhaustion/supplement during the growth of the two layers, which are most likely associated with a different growth of the layers. Temperature dependence of the compositional change during solidification is not expected to play a role in this case, because the monocrystalline TiC layer is even thinner than the polycrystalline TiC layer. Besides the morphology of the reaction layers, its thickness is another important factor that may affect the interfacial bond and load transfer from the matrix to the SiCp. The TiC reaction layers change in thickness depending on the beam velocity and the depth to which the SiCp penetrate in the melt as shown in Fig. 9. At the low beam velocity namely 500 mm/min, the layer thickness is relatively large and decreases as SiCp distribute from the top to the bottom of the track. In contrast, the layer thickness is in general much thinner at higher beam velocities but it increases as the particle locations change from the top towards the bottom of the tracks. Such a dependence of the layers thickness indicates a different reaction circumstance of the particles in the melt, as discussed in the next section. 3.4. OIM study of reaction layers In order to obtain crystallographic information of the reaction layers OIM has been employed. A
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Fig. 6. SEM micrographs showing: (a) microstructure of the mixed reaction layer, and the evolution of the mixed reaction layer formed at the damaged place of the monocrystalline TiC layer from a single site (b); multiple sites (c) to the formation of a continuous layer along the interface (d).
typical orientation map of the cellular TiC reaction layer is presented in Fig. 10a. The reaction layer on the left facet of the SiCp consists of more than 20 TiC grains that are displayed in different colors in Fig. 10b, which are assigned on the basis of the local details of the lattice orientation as indicated by the inserted unit triangle of the inverse pole figure (IPF), Fig. 10c. The TiC grains are connected to each other with high angle boundaries. In the pole figure shown in Fig. 10c, 具001典 poles of all the TiC grains are plotted on the projection plane of the sample coordinates TD and RD, with corresponding colors of the grains. Together with the [0001] pole of the SiCp and Fig. 10c it was clear that there is no specific OR between the polycrystalline TiC reaction layer and the SiCp.
Although it is uncertain whether the left facet of the SiCp orients edge-on or parallel to the basal plan of its lattice, which is inclined at an angle of about 10° to the edge-on direction, it is found that most of the TiC grains grow in the [001] direction nearly perpendicular to the facet, within a maximum inclined angle of 15° assuming the facet is edge-on. This is evidenced in Fig. 10c where the [001] poles of the TiC grains are clustered mainly into two groups, close to the edge-on position on the projection plane marked by the arrow. These two groups of the [001] poles are actually located within a deviating cone of the maximum inclined angle around the facet normal in the stereo-projection. This is indicated by the fact that one group of the poles is orientated above the projection plane
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Fig. 7. SEM/SAM microscopy of the monocrystalline TiC reaction layer: (a) micrograph of the observation area; (b) carbon content mapping; and (c) intensity profiles of carbon and titanium elements along the scanning line indicated in (a).
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Fig. 8. SEM/SAM microscopy of the polycrystalline TiC reaction layer: (a) micrograph of the observation area; (b) carbon content mapping; and (c) intensity profiles of carbon and titanium elements along the scanning line indicated in (a).
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case, the OR is recognized as (111)TiC(101¯ 2)SiC and 具11¯ 0典TiC具12¯ 10典SiC. The ORs recognized by OIM are considered to be the same as those observed with HRTEM [9,10], although it is impossible for OIM to distinguish (111)TiC(0001)SiC and 具11¯ 0典TiC具112¯ 0典SiC from (111)TiC(0006)SiC and [12¯ 1]TiC[101¯ 0]SiC. 4. Discussions 4.1. Formation mechanism of FGM
Fig. 9. Average thickness of the TiC reaction layers as a function of the depth to which SiCp are injected in the FGM tracks, produced at different beam-scanning velocities.
and the other below the plane. Those TiC grains having their 001 pole tilted far away from the facet normal, e.g. grain No. 2, 7 and 25, are smaller in size and are probably suppressed by the neighboring grains during growth. Therefore, the preferred orientation of the TiC grains is a growth phenomenon related to the steepest thermal gradient, as the thermal conductivity of TiC is maximum in the [001] direction. Associated with this growth phenomenon, geometrically necessary tilts of the 001 growing direction of the TiC grains on the slightly curved facet can still be distinguished in Fig. 10b, where the TiC grain Nos. 1–14 nucleated on the top half of the facet slightly tilt towards topleft. In another words, their 001 poles are located above the projection line of the facet normal if the poles are on the left to ND-RD or below the line if the poles are on the right to ND-RD. The TiC grains No. 15–28 on the bottom half of the facet just tilt their 001 pole in the opposite direction. Another OIM map shown in Fig. 11 reveals not only the specific OR between the monocrystalline TiC layer and SiCp lattices, but also twinning in the layer. The monocrystalline TiC layer forms on a facet almost parallel to the basal plane of SiC lattice. All the TiC twins have the same OR with the lattice of SiCp, i.e. (111)TiC(0001)SiC and 具11¯ 0典TiC具112¯ 0典SiC. Other examples of the monocrystalline TiC layer are observed on (101¯ 2) facets of SiCp and show denser twining inside. In this
The FGMs of SiCp/Ti6Al4V are formed in situ during a one-step process of LMI. To understand the formation mechanism of the FGMs with LMI, a simple model will be designed. The particle injection process can be divided into two stages: penetration through the melt surface and movement of the particles inside the laser pool. To simplify the problem the particles are assumed to be of spherical shape, having a radius R and a vertical component of velocity v0 approaching the melt surface. If v0 is larger than the minimum velocity vmin that is necessary to overcome the surface barrier, the particles will penetrate into the melt with a reduced velocity vr [4] vr ⫽ 冑v02⫺vmin2.
(2)
The forces acting on a moving particle in the melt are Stokes’ force F(1): F(1) ⫽ ⫺6πRhv(t)
(3)
and a buoyant force F(2) according to Archimedes’ principle 4 F(2) ⫽ πR3(rSiC⫺rTi)g, 3
(4)
where h is the viscosity of the melt, rTi ⫽ 4150kg / m3 and rSiC ⫽ 3217kg / m3 are the densities of the melt and SiCp respectively, and g the acceleration due to gravitation. According to Newton’s second law of motion, one can write dv 4 3 πR rSiC ⫽ F(1) ⫹ F(2) 3 dt with
(5)
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Fig. 10. Combined microstructure and crystallographic analysis of a cellular polycrystalline TiC reaction layer: (a) SEM micrograph; (b) [001] IPF map showing TiC grains connected by high angle grain boundaries inside the layer; (c) 具001典 pole figure of the TiC grains. The poles with corresponding color and numbers to the grains are clustered into two groups c1 and c2 that are orientated close to the facet normal of the SiCp, represented by the arrow.
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冉
冊
rTi b ⫽ g 1⫺ rSiC and c⫽
9 h . 2 R2rSiC
From Eq. (7) the injection depth of SiCp depends on the viscosity (h) of the melt, the injection velocity (v0) of SiCp, the minimum velocity (vmin) of SiCp necessary for penetration through the melt surface and the time t. The value of vmin can be estimated according to [2,4]: vmin ⫽
冪2s Rr 3
lv
Fig. 11. OIM of a monocrystalline TiC layer showing the twin boundaries inside the layer.
v(t) ⫽
dz(t) , dt
(6)
where z(t) is the penetration depth of the particle in the melt pool at time t. Integration of Eqs. (5) and (6) and substitution of the initial conditions v ⫽ vp and z(t ⫽ 0) ⫽ 0 yields: z(t) ⫽ where
1 [bct⫺(b⫺vrc)(1⫺e⫺ct)] c2
(7)
.(slv ⫹ slp⫺spv),
(8)
SiC
where s is the interface tension between the phases indicated by the subscripts l=liquid, p=solid SiCp and v=vapor. For slv ⫽ 1650mJ / m2, spv ⫽ 1920mJ / m2, slp ⫽ 1000mJ / m2 [4,11–12] and R ⫽ 40µm, vmin becomes 1.94 m/s, which lies in the same range of the experimental value of vi, that is 1.5–2 m/s. In the following calculations, vmin is assumed to be constant in the melt. Reliable data are scarcely available in literature on the temperature effects on the surface tensions. According to Eq. (7) the time dependence of the injection depth of a moving particle in the melt is shown in Fig. 12 for various values of the viscosity. The particle reaches the maximum depth independent of this viscosity range within a time lapse of 5 ms, which is much shorter than the cooling time of the melt (80–150 ms at the used beam velocities). Afterwards, the particle may be pushed back by the buoyant force over a few micrometers, which is, however, so small that it can be neglected with respect to the maximum injection depth in the range of millimeter. Therefore, for simplicity the time for SiCp to penetrate in the melt is fixed at 5 ms. The most sensitive parameters that dominate the final injection depth of SiCp are likely the viscosity of the melt and the injection velocity of SiCp. To correlate the viscosity of the melt to the temperature field of the laser pool, as demonstrated in Fig. 2, it is necessary to find a relationship between the viscosity of the melt and the temperature. Although up to now little work is published on the
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and hm into Eq. (9) C is found to be 0.168 cP. Therefore, the correlation between h0 and T is fixed and depicted in Fig. 13a. On the other hand, from the rheological point of view the actual injection conditions are rather different over the whole extended part of the pool. In another words, particles are injected at the starting point A (see Fig. 1) into an infinitely dilute Timelt at a temperature around 3000 °C, and at the end B injected into a dense suspension containing
Fig. 12.
Injection depth of SiCp as a function of time.
thermophysical properties of Ti6Al4V, the following equation [13] relates the viscosity h0 of a metal to the temperature T: h ⫽ Cexp(E / RT),
(9)
where C is a constant, R the molar gas constant and E the activation energy of the viscosity. It has been observed in about 20 different liquid metals [14] that the activation energy (in kcal/mol) exhibits a power-like behavior as a function of the melting point Tm, according to: E ⫽ 0.431T1.348 m
(10)
Further, at Tm the viscosity of a liquid metal reads (in cP) [15]: hm ⫽ 5.7 × 10⫺6(AwTm)1/2 / V2/3 m ,
(11)
where Aw is the atomic weight and Vm is the atomic volume at Tm. Combining Eqs. (9)–(11) provides a relationship between viscosity and temperature. For Ti6Al4V melt, E and hm are found to be 11.585 kcal/mol and 3.41 cP, respectively, by substituting the melting temperature 1935 K, atomic weight 46.75 g/mol and atomic volume 11.265 10⫺6 m3/mol in Eqs. (10) and (11). The calculated value of hm is consistent with the measured viscosity of Ti–C alloys [16]. Next, by substituting E
Fig. 13. (a) Calculated viscosity (h0) of liquid Ti6Al4V versus the temperature profile of the laser pool according to the FEM results. The origin of the X-coordinate coincides with the starting point of injection process. heff ⫽ h0exp(2.5f/ 1⫺1.5f) is also plotted assuming SiCp concentration increases linearly from 0 at X ⫽ 0mm to 35% at X ⫽ ⫺2.3mm; (b) maximum injection depth (d⬘max) of SiCp at different injection positions determined by heff. dmax is an approximate estimation of the maximum injection depth that takes into account the viscosity change of the suspension in depth direction.
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about 35 vol.% SiCp (measured value, see Fig. 4) at a temperature close to the solidification temperature of the melt. Therefore, our considerations should focus not only on the temperature dependence of the viscosity, but also on the interactions with the injected particles inside the melt. That is to say, rather the effective viscosity (heff) of a SiCp/Ti suspension is relevant for the injection processing mechanism. Since Einstein proposed a viscosity prediction theory for the well-dispersed dilute suspension with uniform rigid spheres, numerous equations have been developed for suspensions at higher concentrations. Among others, Mooney’s model is recommended because of its applicability. Especially, the crowding effects considered in his model fit in generally with the situation of particles injection. The following relationship between heff and particle concentration f is proposed [17]: heff ⫽ h0exp
冉 冊 2.5f 1⫺kf
(12)
where h0 is the viscosity of the melt at different temperatures, and k is a constant, the crowding factor, estimated to be in the range of 1.35 ⬍ k ⬍ 1.91. Here, k is set to 1.5 according to Mooney’s suggestion. Assuming that SiCp are homogeneously distributed over the cross-section of the powder stream, the SiCp concentration in the melt increases linearly from 0 at the starting point to 35 vol.% at the end of injection. According to Eq. (12), heff is plotted in Fig. 13a as a function of relative position (X-coordinate) over the elongated part of the laser pool. By replacing h with heff, one is now able to calculated with Eq. (7) the maximum injection depth (d⬘max) of SiCp, the results of which are illustrated in Fig. 13b. It should be realized that in this way the calculated value of the maximum injection depth does not take into account any influence of the viscosity change due to the suspension in the vertical direction, i.e. as a function of depth. To include these influences, the actual depth (dmax) can be approximately estimated by using the middle-value theorem of a definite integral, as shown in Fig. 13b. The calculated dmax is somewhat larger than the measured maximum injection depth on the cross-
sections. It may be attributed to the fact that the injected SiCp particles actually also cool the melt. The last point of consideration is the effect of the injection velocity of SiCp on the injection depth. Fig. 14 presents the injection depth as a function of the injection velocity of SiCp. The steepness of the curve reveals the fact that the particles may be distributed over the whole distance of penetration, i.e. from the molten surface to the maximum injection depth due to the wide interval (2–4 m/s) of v0. Provided the particles are evenly distributed along the injection depth, the total amount of the injected SiCp at different depths of a composite layer is simply equal to the sum of the particles injected at different positions from the starting point to the end of injection process. In another words, the total amount of SiCp at different depths is proportional to the lengths of the gray bars shown in Fig. 13b. Therefore, it is understandable that the SiCp concentration at the top of a FGM layer is much higher than the value at the bottom, and an FGM is thus formed with a gradual profile of SiCp volume fraction along the depth as shown in Fig. 4. 4.2. Interfacial reaction and formation of reaction layers Titanium matrix composites are commonly fabricated by solid-state diffusion-bonding processes,
Fig. 14. Calculated injection depth of SiCp versus injection velocity.
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because severe reactions exist between liquid titanium and reinforcing fibers/particles. Scant work [18,19] on liquid infiltration process suggests that methods of minimizing the reaction at the interface during fabrication must be developed in order to ensure the success of the molten metal route. These may include shortening the processing time, providing protective coatings on the reinforcements, or incorporating alloying elements that retard the reaction. The LMI process as shown in this work is one of the potential solutions for minimizing the reaction, with which no any other process can compete in shortening the processing/reaction time. The processing time is typically shorter than 150 ms in this case. The observed interfacial microstructures suggest the following elementary steps of the interaction between the injected SiCp and the molten Ti-alloy referring to the phase diagrams [20]: SiC→Si∗ ⫹ C∗(T ⬎ 1650°C)
(13)
Ti ⫹ C →TiC(1650°C ⬍ T ⬍ 3067°C)
(14)
∗
∗
and 5Ti∗ ⫹ 3Si∗→Ti5Si3(1650°C ⬍ T
(15)
⬍ 2130°C) where ∗ refers to the element in the molten Ti6Al4V material. First, when SiCp comes into contact with the liquid Ti-alloy, SiCp starts to dissolve in the liquid metal into carbon and silicon atoms according to reaction (13). As the concentrations of C and Si atoms in the adjacent melt increase and the melt rapidly cools down, TiC will form first by reaction (14), because its melting point is higher than that of Ti5Si3. The surface of SiCp serves as a preferential heterogeneous nucleation site of TiC reaction layer. Reaction (15) may occur only when Si content in the melt is high enough. The essential factor that governs the interfacial reaction is the formation of a TiC reaction layer surrounding the particles, which separates SiCp from a direct contact with the liquid metal. Dissolution of SiCp in Ti-melt at temperatures higher than 1650 °C is a fast process, indicated by the exothermal character according to simple thermal dynamic analysis. What should be kept in mind, however,
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are the facts that the LMI process is a non-equilibrium process and that the reaction happens at a wide temperature interval from approximately 3000–1650 °C. Due to rapid cooling of the melt behind the laser beam, the dissolution of SiCp makes the adjacent melt very quickly constitutionally supercooled by the carbon content. Therefore, TiC phases will precipitate and nucleate heterogeneously on SiCp. Depending on the character of the interfaces/facets of SiCp, the TiC nuclei may exhibit a definite OR to the SiC lattice on the basal plane or on the (101¯ 2) plane, or have no specific ORs on randomly oriented and curved interfaces. Iwamoto et al. [9,10] showed that both the basal planes and the (101¯ 2) planes of 6H-SiC possess a polar character. The exposed carbon-atom plane is suitable for a carbon bridge between TiC and SiC lattices that determines the ORs. They observed in situ during reactive wetting of SiC in a Ti-containing AgCu-alloy that SiC tends to dissolve along the basal plane and to produce an atomically flat plane. The nucleation sites of TiC nuclei are steps or defects on the basal plane. Such nucleated TiC islands are actually surrounded by the most favorable fivefold coordinated sites and will quickly spread over the whole facets of SiCp [7]. The subsequent stacking of TiC atomic layers mainly exhausts carbon atoms from the adjacent supercooled melts and, therefore, ends by the depletion of carbon. This is evidenced in Fig. 7 where a gradual profile of carbon intensity clearly indicates the quick decrease of carbon content in the monocrystalline TiC layer. Solid-state diffusion of carbon atoms through the formed monocrystalline TiC layer is negligible with respect to the rapid growth of the layer. In contrast, the polycrystalline TiC layers grow from randomly oriented nuclei. The grain boundaries between TiC grains, which are the last solidifying part of the layer, provide channels for fast transfer of carbon atoms from SiCp during the later period of growth. Therefore, it yields a longer period of growth that corresponds to a thicker layer and the depletion of carbon in the polycrystalline layer is also less severe, as confirmed in Fig. 8. From Fig. 6 and the preceding discussion, it is clear that the TiC reaction layers actually serve as a barrier against the interfacial reaction and eventu-
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ally stops the dissolution of SiCp. Consider a monocrystalline TiC layer formed at a temperature of 2500 °C, for instance, in the top region of a FGM track. A drop-down of temperature exposes the monocrystalline TiC layer under tensile stresses because of the different thermal expansion coefficients between TiC and SiC phases, 7.7 × 10⫺6 / K and 4.4 × 10⫺6 / K [21], respectively. A simple estimate predicts a tensile stress of 590 MPa in the layer with a temperature drop to 2100 °C, which is much higher than the tensile strength of a TiC layer. As a result, direct fracture or mechanical twinning may be induced in the monocrystalline TiC layers depending on the real stress levels, and both of them have been observed as can be seen in Figs. 6 and 11. Fracture of the monocrystalline TiC layer brings SiCp partially into a direct contact again with Ti-melt and an interfacial reaction starts again. However, the reaction at temperatures below 2130 °C (the melting point of Ti5Si3 phase in Ti-melt) favors the simultaneous precipitation of both TiC phases and Ti5Si3 phases, which produces a mixed reaction layer as is shown in Fig. 6. The broken TiC layer may still play a critical role in the formation of Ti5Si3 phases, which prevents silicon atoms to diffuse freely away into the melt. Higher concentration of silicon in the melt is needed to form Ti5Si3 phases, in comparison with the necessary carbon concentration for the formation of TiC phases at the low temperatures. Probably due to different growth behavior, i.e. faceted and non-faceted characters between the two phases, the mixed layer is not a dense layer and cannot halt the reaction anymore. The reaction will last till the surrounding melt resolidifies. Formation of monocrystalline TiC layers at a lower temperature, e.g. injecting SiCp at lower temperatures or into a deeper depth, will reduce their fracture probability and thus restrains the formation of the mixed reaction layer. Another way to avoid fracture of the monocrystalline TiC layer is to choose other carbide ceramic particles that have a similar thermal expansion coefficient as TiC, e.g. WC as reinforcement. On the other hand, polycrystalline TiC layers may relax via plastic deformation and fracture in polycrystalline TiC layers is hardly observed. Based on the model for the formation of FGMs
and the interfacial reaction mechanism, it is possible to interpret the thickness dependence of the TiC reaction layers on the beam velocity and the injection depth of SiCp as seen in Fig. 9. All the SiCp at the bottom of an FGM track are injected at the beginning of the injection process through a hotter melt, thus, longer reaction/growth time yields thicker TiC layers on these SiCp. In contrast, most of SiCp in the top region of an FGM track are injected at a later stage of injection and therefore the shorter reaction time and lower temperatures make the TiC ration layers thinner. In addition, the appearance of both the mono- and polycrystalline TiC layers in this region leads to a large spread in layer thickness. To conclude, the linear increase in the thickness of the TiC layers with the injection depth of SiCp in the FGM tracks produced at high beam velocities is determined by the particle injection sequence that determines reaction temperatures and growth time. In the case of a homogeneous composite track produced at the beam velocity of 8.3 m/min, a direct interaction with the beam makes SiCp at the top region overheated and heavily dissolved in the melt. The decrease in thickness of the TiC layers with the injection depth of SiCp is mainly attributed to the large temperature gradient between the top and the bottom of the laser pool.
5. Conclusions The main conclusions of this work are the following: 1. For the first time particle-reinforced FGMs are produced with a well-controlled LMI process. 2. A clear physical picture of the particle injection process is obtained on the base of FEA modeling and theoretical calculations. Temperature/ viscosity field of the laser melt pool explains the formation of the FGMs by injecting SiCp into the elongated part of the laser pool behind the beam. The effective viscosity of the SiCp/Ti6Al4V suspension dominates the maximum injection depth of SiCp. 3. Three types of reaction layers are observed around SiCp: a thin monocrystalline TiC layer,
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a cellular polycrystalline TiC layer and a thick mixed layer of TiC with Ti5Si3 particular phases. Both kind of TiC reaction layers act as a barrier against the interfacial reaction and its swift formation during rapid cooling stops the dissolution of SiCp in Ti-melt. The mixed layer origins at damages of a monocrystalline TiC layer. Only the monocrystalline TiC layer exhibits particular ORs with the lattice of SiCp, i.e. (111)TiC(0001)SiC and 具110典TiC具11¯ 00典SiC or (111)TiC(101¯ 2)SiC and 具11¯ 0典TiC具12¯ 10典SiC. 4. The thickness of the reaction layer is minimized by injecting the SiCp fully outside the laser beam in the stretched molten area behind the beam.
Acknowledgements The Netherlands Institute for Metals Research is acknowledged for its financial support. Drs A. Vreeling, P.M. Bronsveld and D.T.L. van Agterveld are acknowledged for their assistance and contributions during these examinations. S.A.J. Nijman and T. Silva are thanked for their assistance in the FEA modeling. H. Bron, K. Post and J. Harkema are acknowledged for their help with the analysis and laser experiments.
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References [1] Suresh S. Science 2447;292:2001. [2] De Hosson JThM, Kooi BJ. In: Nalwa H, editor. Handbook of surfaces and interfaces. New York: Academic Press; 2001. p. 1–114 [chapter 1]. [3] Kloosterman AB, Kooi BJ, De Hosson JThM. Acta mater 1998;46:6205. [4] Vreeling AJ, Ocelı´k V, Pei YT, Van Agterveld DTL, De Hosson JThM. Acta mater 2000;48:4225. [5] Howe JM. Int Mater Rev 1993;38:233. [6] Lloyd DJ. Acta mater 1991;39:59. [7] Pei YT, De Hosson JThM. Acta mater 2001;49:561. [8] Volz R, Reichelt U, Wolf S, Pei YT, Zuo TC. In: Roller D, editor. Proceedings of the 30th ISATA: Rapid Prototyping/laser Applications in the Automotive Industries. Croydon, England: Automotive Automation Ltd; 1997. p. 39-3. [9] Iwamoto C, Ichinose H, Tanaka SI. Philos Mag 1999;79:85. [10] Iwamoto C, Tanaka SI. Philos Mag 1998;78:835. [11] Kaptay G. Mater Sci Forum 1996;475:215–6. [12] Roth TA, Suppayak P. Mater Sci Eng 1978;35:187. [13] Andrade ENC. Philos Mag 1934;17:698. [14] Grosse AV. Science 1963;140:781. [15] Andrade ENC. Philos Mag 1934;17:497. [16] Agayev AD, Kostikov VI, Bobkovskiy VN. Russian Metall 1980;4:38. [17] Mooney M. J Colloid Sci 1951;6:162. [18] Toloui B. In: Harrigan WC, editor. ICCM-V. Warrendale, PA: TMS; 1985. p. 77-3. [19] Warrier SG, Lin RY. Metall Mater Trans 1995;26A:1885. [20] Massalski TB, editor. Binary alloy phase diagrams. Metals Park, OH: ASM; 1986. [p. 595 and 2056]. [21] Schneider S.J., editor. Engineering materials handbook, vol. 4: ceramics and glasses. ASM International; 1991. p. 805–10.
SiCp/Ti6Al4V functionally graded materials produced by laser melt injection Y.T. Pei, V. Ocelik, J.Th.M. De Hosson ∗ Department of Applied Physics, Materials Science Center and the Netherlands Institute for Metals Research, University of Groningen, Nijenborgh 4, 9747 AG Groningen, The Netherlands Received 2 January 2002; received in revised form 25 January 2002; accepted 25 January 2002
Abstract With a well-controlled laser melt injection (LMI) process, for the first time the feasibility is demonstrated to produce SiC particles (SiCp) reinforced Ti6Al4V functionally graded materials (FGMs). SiCp are injected just behind the laser beam into the extended part of the laser melt pool that is formed at relatively high beam scanning velocities. The process allows for the minimization of the decomposition reaction between SiCp and Ti6Al4V melt, and also leads to FGMs of SiCp/Ti6Al4V instead of a homogeneous composite layer on Ti6Al4V substrates. An injection model is designed based on the temperature/viscosity field of the laser pool for a deeper understanding of the mechanism of formation of the FGMs with LMI. The model is based on finite element calculations of the temperature field in the melt pool, physical considerations of the LMI process and it is supported by experimental observations. Three types of reaction layers are observed around SiCp, namely a thin monocrystalline TiC layer, a cellular polycrystalline TiC layer and a thick mixed layer of TiC with Ti5Si3. Among them, only the monocrystalline TiC layer exhibits particular orientation relationships (ORs) to the SiCp lattice, i.e. (111)TiC(0001)SiC and具110典TiC具11¯ 00典SiC or (111)TiC(101¯ 2)SiC and 具11¯ 0典TiC具12¯ 10典SiC. These two kinds of TiC reaction layers act as a barrier against the interfacial reaction and its swift formation during rapid cooling hinders the dissolution of SiCp in the Ti-melt. 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Functionally graded materials; Laser treatment; Microstructure; Metals; Carbides
1. Introduction The recent resurgence of interest in materials with spatial gradients in composition and structure has been driven by the fact that these materials exhibit characteristics not attainable by conven∗ Corresponding author. Tel.: +31-50-363-4898; fax: +3150-363-4881. E-mail address: [email protected] (J.Th.M. De Hosson).
tional materials. Recent theoretical and experimental work has established that controlled gradients in mechanical properties offer attractive challenges for the design of surfaces with resistance to contact deformation and damage. In particular the damage and failure resistance of surfaces to normal and sliding contact or impact can be changed substantially through such gradients. For a recent review see Ref. [1]. SiC reinforced metals are appropriate candidates for enhanced tribological applications, e.g. the automotive industry, but
1359-6454/02/$22.00 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 2 ) 0 0 0 4 9 - 6
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the mechanical performance strongly depends on the degree of dissolution of SiC in the matrix and the type of reaction layer [2]. Abrasive wear and sliding wear tests show a good relative wear resistance of SiC particles (SiCp) embedded in Ti-6Al4V produced by laser treatments [2,3]. The wear rate has been found equal to 10⫺8 mm3/Nm. Injection of ceramic particles into a laser melt pool provides an advanced method to tailor a composite layer onto a surface of a metallic substrate [2–5]. The most important factors that dominate the success of the injection process and consequently the properties of the composite layer are the following: (a) particles dissolution and reaction with the matrix melt at high temperatures; (b) distribution and volume fraction of the injected particles; and (c) thermal stresses built up in the composite layer during cooling of the melt pool. Regarding the first aspect, heating of ceramic particles under direct irradiation of a laser beam promotes particle dissolution/reaction and therefore it is undesirable. Further, it has been shown that a thick reaction layer around the particles can be formed and the reaction products in the matrix result in a brittle behavior of composite layer that affects the transfer of load from the matrix to the particular reinforcements [2–5]. In a special case, e.g. laser melt injection (LMI) of SiCp into Al-alloys [4], the particles could be injected only near the center of the Al melt pool where the oxide skin of the Al melt was partially broken under the irradiation of the beam. In such a situation, heating up of the particles by the laser beam could not be avoided. Regarding the second factor, a homogeneous particle distribution along the injection track is a primary requirement for LMI, which can be obtained with a stable particle stream from an injection nozzle and a constant powder-feeding rate. Due to the fast solidification process of the laser pool, clustering of the particles in the melt does not appear under optimum injection conditions [6]. In order to reduce the residual stresses involved a gradual increase in volume fraction of the injected particles from the bottom toward the surface of the composite layer is highly desirable, which, to the best of our knowledge, has never been obtained by a single-step injection process. In our previous work in the field of LMI [2,4]
the dissolution of the SiCp and the amount of the reaction products in the SiCp/Ti6Al4V composite were clearly reduced by positioning the powder stream partially out of the beam. This idea is extended further in this paper not only allowing for a minimization of the reaction between SiCp and Ti melt, but also to obtain a functionally graded material (FGM) of SiCp/Ti6Al4V. The work is divided into two parts. The focus is on the laser injection process of the FGMs and the interface reaction between SiCp and Ti6Al4V melt. For a better physical understanding of the mechanism of formation of the FGMs a model is designed which is based on the temperature/viscosity field of the laser pool. The model is based on finite element calculations of the temperature field in the melt pool, physical considerations of the LMI process and it is supported by experimental observations. In addition, the microstructure and crystallography of the reaction layer around SiCp are scrutinized with orientation imaging microscopy (OIM) and scanning electron/scanning Auger microscopy (SEM/SAM). In a subsequent paper, the defect structure in the reaction layer and the interface bonding are emphasized.
2. Experimental details The experimental setup of LMI has been documented in detail in Ref. [2]. A Rofin Sinar 2 kW Nd:YAG laser is employed for the injection experiments. A numerically controlled X–Y table executed the specimen movement. The powder feeder used is a Metco 9MP-type commercial instrument. The processing parameters varied in the range of 600–1000 W laser power, 500–1200 mm/min beam scanning velocity and 1–3 g/m powder feeding rate. Selection of SiC/Ti6Al4V is based on the fact that Ti6Al4V has a very low thermal conductivity, which leads to the possibility of elongating the laser pool along the beam scanning direction. As a consequence it allows for the injection of SiCp outside the beam and also to obtain FGMs. Ti6Al4V substrates of dimensions 50 × 40 × 5mm3 are cut with electrical discharge
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erosion from a rectangular bar. The surface of the substrates did not experience any further post finishing after cutting and it was just cleaned by acetone prior to LMI. Monocrystalline 6H-SiC powder with a size distribution ranging between 50 and 90 µm is used. The particles are injected into the melt pool at an angle of 35° with respect to the surface normal. The injection velocity of the SiCp particles is measured by recording their flight out of the powder nozzle with a Kodak 4540 Ektapro highspeed camera at speed of 27 × 103frames / s and turns out to lie in the range between 2 and 4 m/s. No pre-heating of the substrate is necessary for the injection, which is different from the situation of SiCp injection into Al [4]. Transverse and longitudinal sections of the injected composite tracks are cut for microstructural studies and quantitative metallographic analysis. The polished specimens are etched with Keller’s reagent for 10–20 s at room temperature. A Philips XL-30s FEG scanning electron microscope (SEM) equipped with an electron backscattered pattern-orientation imaging microscope (EBSPOIM) and an energy dispersive spectrometer (EDS) is employed for the combined microstructure and crystallography analysis. For the operational principle of OIM as well as the definition of the sample reference frame reference is made to [7]. The composition of the reaction layers is analyzed by a dedicated small-spot and ultrahigh vacuum (UHV) scanning electron/scanning Auger microscope (SEM/SAM) derived from a JEOL JAMP 7800F.
3. Results 3.1. Particle injection process The LMI process is schematically depicted in Fig. 1. An over-defocused beam is used and the focal spot of the beam lies 9 mm above the substrate surface, which represents a spot of about ⭋3mm. SiCp are injected into the laser pool just behind the beam in such a way that the powder stream is positioned close to the beam, but without touching it. This permits the particles to penetrate in the melt to certain depths and the method also
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Fig. 1. Sketch of the laser particle injection process. The white arrows schematically indicate the maximum injection depth of SiCp at different positions from the starting point A to the end B of the injection process.
avoids reaction of the particles with the melt at higher temperatures. A special powder nozzle developed for laser powder cladding and welding [8] satisfies all the demands of a stable injection process. The nozzle possesses a co-axial flow of helium shielding gas not only to protect the laser pool from the atmosphere but also to focus simultaneously the powder stream. A stable flow and straight wire-like shape of the powder stream from the nozzle are essential for controlling the injection process and obtaining a homogeneous distribution of the injected particles along the track. The key point for the realization of the injection conditions is to extend the laser pool backwards behind the beam as far as possible, which will provide the necessary space for particles to inject without touching to the beam. This is achieved with the aid of relatively high scanning velocities of the beam in combination with appropriate laser powers. For understanding the dependence of the size of the pool on the beam scanning velocity, finite element analyses (FEA) are executed with a three-dimensional model attaining an estimate of the shape and the temperature profile of the laser melt pool. The simulation is performed for a YAG laser beam with a top-hat energy distribution and a Ti6Al4V plate with the actual dimensions of the substrates used in the experiments. The model takes into account the latent heat but it does not account for convection and particles in-take. The results are summarized in Fig. 2 and in Table 1. It appears that both the length (L) and the extended part length (Le) of the pool behind the beam increase with increasing beam velocity at constant
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Fig. 2. Temperature field and dimensions of the laser pool obtained by FEA modeling at a constant power density (I0) and different beam velocity of: (a) 0.6 m/min; and (b) 0.9 m/min. Table 1 Lengths and surface temperature of laser melted Ti6Al4V pool obtained from the FEA modeling vb (m/min) 0.6 Le (mm) L (mm) Le /D L /D Temperature at X=0 (°C)
1.02 2.95 0.34 0.98 2418
0.9
1.2
2.45 4.53 0.82 1.51 3007
3.83 5.93 1.28 1.98 3717
laser energy density (I0), which is defined as the product of the laser power and the irradiation time on a local area divided by the area of the beam, according to: I0 ⫽
4P πDvb
(1)
where P is the laser power, D the spot size of the beam and vb the scanning velocity of the beam. For instance, the ratios of L/D and Le/D change from 0.98 to 1.98 and 0.34 to 1.28, respectively, by increasing the laser beam velocity from 0.6 to 1.2 m/min. Although the FEA result may differ slightly from the real temperature field due to the simplicities made in the model, it gives a clear indication that a powder stream of approximate the
same diameter as the beam spot can be injected in the extended part of the pool at high beam velocities without making contact with the beam. This set-up does not disturb the formation of a stable laser pool and favors a well-controlled injection process. Therefore, all the problems associated with the flight of SiCp through the beam, such as heating up the particles, shadowing the pool and increasing the temperature of the melt, will not arise anymore. It should be realized that the extension of the laser pool strongly relies on the low thermal conductivity 6.7 W/m K of the Ti6Al4V alloy with respect to other alloys or metals. 3.2. Particle distribution and FGM formation Fig. 3a shows a polished surface of a SiCp/Ti6Al4V composite track produced with LMI, which is prepared for tensile tests in such a way that the convex surface (part) of the track has been just polished away to make it flat and coplanar with the substrate surface. The picture gives an overview on the overall quality of the composite material in which SiCp are embedded over the whole molten area of the substrate. The particles are distributed homogeneously along the track. One of the striking findings is the fact that the particle distribution at different depths can be controlled by the laser beam scanning velocity and the injection position. As shown in Fig. 3b, SiCp are
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Fig. 3. (a) Plane view of a SiCp/Ti6Al4V FGM track produced with laser particle injection; (b) cross-section of a track injected at 0.5 m/min beam scanning velocity; (c) and (d) cross-section and longitudinal section, respectively, of an FGM track produced at 0.8 m/min beam scanning velocity.
distributed much more evenly over the whole cross-section of the composite track produced at 500 mm/min. In contrast, the track injected at 800 mm/min has more SiCp on the top area and fewer particles at the bottom as shown in Fig. 3c, that is to say, a FGM is formed. A typical longitudinal cross-section of the FGMs displayed in Fig. 3d clearly demonstrates the gradual distribution of SiCp along the depth direction of the composite
and the smooth transition between the composite layer and the substrate. To characterize the particle distribution as a function of depth, quantitative metallographic analysis is carried out with the aid of computer software that works on the principle of image analysis. A digitized picture of a longitudinal cross-section of the FGM tracks, e.g. Fig. 3d, is first divided into a series of horizontal strips with
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a width of 100 µm, which is just slightly larger than the maximum particle size. The software measures the pixels occupied by SiCp and subsequently the ratio of the particle-occupied pixels to the total pixels of a strip area is calculated. This ratio is then assigned as the equivalent volume fraction of SiCp to the depth, where the middle of the strip is located on the longitudinal cross-section. In this way, a profile of the particle distribution along the depth direction of a composite track can be drawn as displayed in Fig. 4. It is clear that there is a transition from a homogeneous composite material, produced at 500 mm/min beam velocity, to the graded composite, i.e. FGM obtained at beam velocities over 800 mm/min. The higher the beam scanning velocity the steeper is the slope of the profile. Because the powder-feeding rate (g/m) is fixed for different beam scanning velocities, however, the total amount of the injected particles in those tracks maintains nearly the same, which is equal to the corresponding area of volume fraction versus depth for the profiles in Fig. 4. The formation of the FGMs is most likely controlled by the laser beam scanning velocity, if the powder stream is strictly focused onto the extended part of the laser pool and the injection velocity of the particles is appropriate.
Fig. 4. Profile of SiCp volume fraction as a function of the depth in the FGM tracks produced at different velocities of beam scanning.
3.3. Microstructure of FGM and reaction layers One of the main effects by shifting the powder stream fully out of the beam is the minimization of the dissolution of SiCp in the melt. Nevertheless, as SiCp slightly dissolve during penetration in the melt the matrix composition changes depending on the temperature of the melt. The matrix consists mainly of α-Ti dendritic grains that exhibit a lenticular martensite structure and interdendritic eutectics. In addition, some faceted Ti5Si3 particles in a micrometers range as well as a few TiC freedendrites are observed in the top area of the composite track. The latter are rarely found near the bottom of the track. Indeed, the total amount of these reaction products is considerably reduced in comparison with the results described previously [2]. A reaction layer surrounds always the injected SiCp. Three types of reaction layers can be distinguished: (a) a thin monocrystalline TiC layer with a particular OR to SiCp; (b) a cellular polycrystalline TiC layer; and (c) a mixed layer of TiC and Ti5Si3 particles. The monocrystalline TiC layer shown in Fig. 5a appears relatively smooth, few hundred nanometers thick and it may grow on one or a few flat facets of a SiCp. A SiCp is not fully surrounded by this type of TiC layer due to the orientations of different facets of the particle. The SiCp that exhibit a monocrystalline TiC reaction layer are mostly located in the top region of the injection track. The cellular TiC layer depicted in Fig.5b forms the majority the reaction layers (also see Ref. [2]), and can be found on most of the SiCp that are distributed everywhere over the whole cross-section. The third albeit minor type of reaction layer is more complicated and its microstructure is revealed in Fig. 6. It is composed of more or less spherical TiC particles and faceted Ti5Si3 particles together with a broken TiC ‘skin’ at the outside. In particular, this complex structure does not form a dense layer of TiC and Ti5Si3 particles in the Ti matrix. The increase in size of the TiC particles from the inner interface towards the outer TiC skin clearly indicates a continuous nucleation and growth feature of the layer, as seen in Fig. 6a. Such a mixed reaction layer is often found on the SiCp near the surface of the composite tracks and
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Fig. 5. SEM micrographs showing: (a) monocrystalline TiC reaction layer formed at a flat facet of SiCp; and (b) cellular polycrystalline TiC layer surrounding SiCp.
may partially cover a SiCp. Moreover, it frequently shows a transition from the monocrystalline TiC layers, as shown in Fig. 6b and c. In another words, the mixed layer originates at the sites of the monocrystalline TiC layer previously formed. It is interesting to note that the mixed layer stimulates the SiCp to dissolve continuously, whereas the monocrystalline TiC layer prevents further dissolution of the SiCp as soon as it forms on the facets. As revealed in Fig. 6c and d, the protrusions on the SiCp facets indicated by arrows, are still covered by a monocrystalline TiC layer and the concave parts adjacent on both sides of the protrusions are exposed to the mixed layers. The step height of 1 µm between the protrusion and the adjacent concaves points to the stimulated dissol-
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ution of the SiCp. Obviously, the mixed reaction layer is much thicker than the other two types of reaction layers due to the different mechanisms of formation. Because TiC is not a phase that has a stoichiometric composition, the carbon content of the TiC reaction layers is analyzed with SEM/SAM in order to reveal its formation. Before taking Auger electron spectra, the surface of the specimens is cleaned in situ by Ar+-ion sputtering under UHV. The intensities of Si, C and Ti elemental peaks in the Auger electron spectra are analyzed. Both the area mapping and line profile of carbon intensity presented in Fig. 7 clearly show a rapid decrease of the carbon content in the monocrystalline TiC layer from the SiCp side to the matrix side. In contrast, the polycrystalline TiC layer exhibits less change in carbon content over the thickness, as seen in Fig. 8. This implies different conditions of carbon exhaustion/supplement during the growth of the two layers, which are most likely associated with a different growth of the layers. Temperature dependence of the compositional change during solidification is not expected to play a role in this case, because the monocrystalline TiC layer is even thinner than the polycrystalline TiC layer. Besides the morphology of the reaction layers, its thickness is another important factor that may affect the interfacial bond and load transfer from the matrix to the SiCp. The TiC reaction layers change in thickness depending on the beam velocity and the depth to which the SiCp penetrate in the melt as shown in Fig. 9. At the low beam velocity namely 500 mm/min, the layer thickness is relatively large and decreases as SiCp distribute from the top to the bottom of the track. In contrast, the layer thickness is in general much thinner at higher beam velocities but it increases as the particle locations change from the top towards the bottom of the tracks. Such a dependence of the layers thickness indicates a different reaction circumstance of the particles in the melt, as discussed in the next section. 3.4. OIM study of reaction layers In order to obtain crystallographic information of the reaction layers OIM has been employed. A
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Fig. 6. SEM micrographs showing: (a) microstructure of the mixed reaction layer, and the evolution of the mixed reaction layer formed at the damaged place of the monocrystalline TiC layer from a single site (b); multiple sites (c) to the formation of a continuous layer along the interface (d).
typical orientation map of the cellular TiC reaction layer is presented in Fig. 10a. The reaction layer on the left facet of the SiCp consists of more than 20 TiC grains that are displayed in different colors in Fig. 10b, which are assigned on the basis of the local details of the lattice orientation as indicated by the inserted unit triangle of the inverse pole figure (IPF), Fig. 10c. The TiC grains are connected to each other with high angle boundaries. In the pole figure shown in Fig. 10c, 具001典 poles of all the TiC grains are plotted on the projection plane of the sample coordinates TD and RD, with corresponding colors of the grains. Together with the [0001] pole of the SiCp and Fig. 10c it was clear that there is no specific OR between the polycrystalline TiC reaction layer and the SiCp.
Although it is uncertain whether the left facet of the SiCp orients edge-on or parallel to the basal plan of its lattice, which is inclined at an angle of about 10° to the edge-on direction, it is found that most of the TiC grains grow in the [001] direction nearly perpendicular to the facet, within a maximum inclined angle of 15° assuming the facet is edge-on. This is evidenced in Fig. 10c where the [001] poles of the TiC grains are clustered mainly into two groups, close to the edge-on position on the projection plane marked by the arrow. These two groups of the [001] poles are actually located within a deviating cone of the maximum inclined angle around the facet normal in the stereo-projection. This is indicated by the fact that one group of the poles is orientated above the projection plane
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Fig. 7. SEM/SAM microscopy of the monocrystalline TiC reaction layer: (a) micrograph of the observation area; (b) carbon content mapping; and (c) intensity profiles of carbon and titanium elements along the scanning line indicated in (a).
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Fig. 8. SEM/SAM microscopy of the polycrystalline TiC reaction layer: (a) micrograph of the observation area; (b) carbon content mapping; and (c) intensity profiles of carbon and titanium elements along the scanning line indicated in (a).
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case, the OR is recognized as (111)TiC(101¯ 2)SiC and 具11¯ 0典TiC具12¯ 10典SiC. The ORs recognized by OIM are considered to be the same as those observed with HRTEM [9,10], although it is impossible for OIM to distinguish (111)TiC(0001)SiC and 具11¯ 0典TiC具112¯ 0典SiC from (111)TiC(0006)SiC and [12¯ 1]TiC[101¯ 0]SiC. 4. Discussions 4.1. Formation mechanism of FGM
Fig. 9. Average thickness of the TiC reaction layers as a function of the depth to which SiCp are injected in the FGM tracks, produced at different beam-scanning velocities.
and the other below the plane. Those TiC grains having their 001 pole tilted far away from the facet normal, e.g. grain No. 2, 7 and 25, are smaller in size and are probably suppressed by the neighboring grains during growth. Therefore, the preferred orientation of the TiC grains is a growth phenomenon related to the steepest thermal gradient, as the thermal conductivity of TiC is maximum in the [001] direction. Associated with this growth phenomenon, geometrically necessary tilts of the 001 growing direction of the TiC grains on the slightly curved facet can still be distinguished in Fig. 10b, where the TiC grain Nos. 1–14 nucleated on the top half of the facet slightly tilt towards topleft. In another words, their 001 poles are located above the projection line of the facet normal if the poles are on the left to ND-RD or below the line if the poles are on the right to ND-RD. The TiC grains No. 15–28 on the bottom half of the facet just tilt their 001 pole in the opposite direction. Another OIM map shown in Fig. 11 reveals not only the specific OR between the monocrystalline TiC layer and SiCp lattices, but also twinning in the layer. The monocrystalline TiC layer forms on a facet almost parallel to the basal plane of SiC lattice. All the TiC twins have the same OR with the lattice of SiCp, i.e. (111)TiC(0001)SiC and 具11¯ 0典TiC具112¯ 0典SiC. Other examples of the monocrystalline TiC layer are observed on (101¯ 2) facets of SiCp and show denser twining inside. In this
The FGMs of SiCp/Ti6Al4V are formed in situ during a one-step process of LMI. To understand the formation mechanism of the FGMs with LMI, a simple model will be designed. The particle injection process can be divided into two stages: penetration through the melt surface and movement of the particles inside the laser pool. To simplify the problem the particles are assumed to be of spherical shape, having a radius R and a vertical component of velocity v0 approaching the melt surface. If v0 is larger than the minimum velocity vmin that is necessary to overcome the surface barrier, the particles will penetrate into the melt with a reduced velocity vr [4] vr ⫽ 冑v02⫺vmin2.
(2)
The forces acting on a moving particle in the melt are Stokes’ force F(1): F(1) ⫽ ⫺6πRhv(t)
(3)
and a buoyant force F(2) according to Archimedes’ principle 4 F(2) ⫽ πR3(rSiC⫺rTi)g, 3
(4)
where h is the viscosity of the melt, rTi ⫽ 4150kg / m3 and rSiC ⫽ 3217kg / m3 are the densities of the melt and SiCp respectively, and g the acceleration due to gravitation. According to Newton’s second law of motion, one can write dv 4 3 πR rSiC ⫽ F(1) ⫹ F(2) 3 dt with
(5)
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Fig. 10. Combined microstructure and crystallographic analysis of a cellular polycrystalline TiC reaction layer: (a) SEM micrograph; (b) [001] IPF map showing TiC grains connected by high angle grain boundaries inside the layer; (c) 具001典 pole figure of the TiC grains. The poles with corresponding color and numbers to the grains are clustered into two groups c1 and c2 that are orientated close to the facet normal of the SiCp, represented by the arrow.
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冉
冊
rTi b ⫽ g 1⫺ rSiC and c⫽
9 h . 2 R2rSiC
From Eq. (7) the injection depth of SiCp depends on the viscosity (h) of the melt, the injection velocity (v0) of SiCp, the minimum velocity (vmin) of SiCp necessary for penetration through the melt surface and the time t. The value of vmin can be estimated according to [2,4]: vmin ⫽
冪2s Rr 3
lv
Fig. 11. OIM of a monocrystalline TiC layer showing the twin boundaries inside the layer.
v(t) ⫽
dz(t) , dt
(6)
where z(t) is the penetration depth of the particle in the melt pool at time t. Integration of Eqs. (5) and (6) and substitution of the initial conditions v ⫽ vp and z(t ⫽ 0) ⫽ 0 yields: z(t) ⫽ where
1 [bct⫺(b⫺vrc)(1⫺e⫺ct)] c2
(7)
.(slv ⫹ slp⫺spv),
(8)
SiC
where s is the interface tension between the phases indicated by the subscripts l=liquid, p=solid SiCp and v=vapor. For slv ⫽ 1650mJ / m2, spv ⫽ 1920mJ / m2, slp ⫽ 1000mJ / m2 [4,11–12] and R ⫽ 40µm, vmin becomes 1.94 m/s, which lies in the same range of the experimental value of vi, that is 1.5–2 m/s. In the following calculations, vmin is assumed to be constant in the melt. Reliable data are scarcely available in literature on the temperature effects on the surface tensions. According to Eq. (7) the time dependence of the injection depth of a moving particle in the melt is shown in Fig. 12 for various values of the viscosity. The particle reaches the maximum depth independent of this viscosity range within a time lapse of 5 ms, which is much shorter than the cooling time of the melt (80–150 ms at the used beam velocities). Afterwards, the particle may be pushed back by the buoyant force over a few micrometers, which is, however, so small that it can be neglected with respect to the maximum injection depth in the range of millimeter. Therefore, for simplicity the time for SiCp to penetrate in the melt is fixed at 5 ms. The most sensitive parameters that dominate the final injection depth of SiCp are likely the viscosity of the melt and the injection velocity of SiCp. To correlate the viscosity of the melt to the temperature field of the laser pool, as demonstrated in Fig. 2, it is necessary to find a relationship between the viscosity of the melt and the temperature. Although up to now little work is published on the
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and hm into Eq. (9) C is found to be 0.168 cP. Therefore, the correlation between h0 and T is fixed and depicted in Fig. 13a. On the other hand, from the rheological point of view the actual injection conditions are rather different over the whole extended part of the pool. In another words, particles are injected at the starting point A (see Fig. 1) into an infinitely dilute Timelt at a temperature around 3000 °C, and at the end B injected into a dense suspension containing
Fig. 12.
Injection depth of SiCp as a function of time.
thermophysical properties of Ti6Al4V, the following equation [13] relates the viscosity h0 of a metal to the temperature T: h ⫽ Cexp(E / RT),
(9)
where C is a constant, R the molar gas constant and E the activation energy of the viscosity. It has been observed in about 20 different liquid metals [14] that the activation energy (in kcal/mol) exhibits a power-like behavior as a function of the melting point Tm, according to: E ⫽ 0.431T1.348 m
(10)
Further, at Tm the viscosity of a liquid metal reads (in cP) [15]: hm ⫽ 5.7 × 10⫺6(AwTm)1/2 / V2/3 m ,
(11)
where Aw is the atomic weight and Vm is the atomic volume at Tm. Combining Eqs. (9)–(11) provides a relationship between viscosity and temperature. For Ti6Al4V melt, E and hm are found to be 11.585 kcal/mol and 3.41 cP, respectively, by substituting the melting temperature 1935 K, atomic weight 46.75 g/mol and atomic volume 11.265 10⫺6 m3/mol in Eqs. (10) and (11). The calculated value of hm is consistent with the measured viscosity of Ti–C alloys [16]. Next, by substituting E
Fig. 13. (a) Calculated viscosity (h0) of liquid Ti6Al4V versus the temperature profile of the laser pool according to the FEM results. The origin of the X-coordinate coincides with the starting point of injection process. heff ⫽ h0exp(2.5f/ 1⫺1.5f) is also plotted assuming SiCp concentration increases linearly from 0 at X ⫽ 0mm to 35% at X ⫽ ⫺2.3mm; (b) maximum injection depth (d⬘max) of SiCp at different injection positions determined by heff. dmax is an approximate estimation of the maximum injection depth that takes into account the viscosity change of the suspension in depth direction.
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about 35 vol.% SiCp (measured value, see Fig. 4) at a temperature close to the solidification temperature of the melt. Therefore, our considerations should focus not only on the temperature dependence of the viscosity, but also on the interactions with the injected particles inside the melt. That is to say, rather the effective viscosity (heff) of a SiCp/Ti suspension is relevant for the injection processing mechanism. Since Einstein proposed a viscosity prediction theory for the well-dispersed dilute suspension with uniform rigid spheres, numerous equations have been developed for suspensions at higher concentrations. Among others, Mooney’s model is recommended because of its applicability. Especially, the crowding effects considered in his model fit in generally with the situation of particles injection. The following relationship between heff and particle concentration f is proposed [17]: heff ⫽ h0exp
冉 冊 2.5f 1⫺kf
(12)
where h0 is the viscosity of the melt at different temperatures, and k is a constant, the crowding factor, estimated to be in the range of 1.35 ⬍ k ⬍ 1.91. Here, k is set to 1.5 according to Mooney’s suggestion. Assuming that SiCp are homogeneously distributed over the cross-section of the powder stream, the SiCp concentration in the melt increases linearly from 0 at the starting point to 35 vol.% at the end of injection. According to Eq. (12), heff is plotted in Fig. 13a as a function of relative position (X-coordinate) over the elongated part of the laser pool. By replacing h with heff, one is now able to calculated with Eq. (7) the maximum injection depth (d⬘max) of SiCp, the results of which are illustrated in Fig. 13b. It should be realized that in this way the calculated value of the maximum injection depth does not take into account any influence of the viscosity change due to the suspension in the vertical direction, i.e. as a function of depth. To include these influences, the actual depth (dmax) can be approximately estimated by using the middle-value theorem of a definite integral, as shown in Fig. 13b. The calculated dmax is somewhat larger than the measured maximum injection depth on the cross-
sections. It may be attributed to the fact that the injected SiCp particles actually also cool the melt. The last point of consideration is the effect of the injection velocity of SiCp on the injection depth. Fig. 14 presents the injection depth as a function of the injection velocity of SiCp. The steepness of the curve reveals the fact that the particles may be distributed over the whole distance of penetration, i.e. from the molten surface to the maximum injection depth due to the wide interval (2–4 m/s) of v0. Provided the particles are evenly distributed along the injection depth, the total amount of the injected SiCp at different depths of a composite layer is simply equal to the sum of the particles injected at different positions from the starting point to the end of injection process. In another words, the total amount of SiCp at different depths is proportional to the lengths of the gray bars shown in Fig. 13b. Therefore, it is understandable that the SiCp concentration at the top of a FGM layer is much higher than the value at the bottom, and an FGM is thus formed with a gradual profile of SiCp volume fraction along the depth as shown in Fig. 4. 4.2. Interfacial reaction and formation of reaction layers Titanium matrix composites are commonly fabricated by solid-state diffusion-bonding processes,
Fig. 14. Calculated injection depth of SiCp versus injection velocity.
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because severe reactions exist between liquid titanium and reinforcing fibers/particles. Scant work [18,19] on liquid infiltration process suggests that methods of minimizing the reaction at the interface during fabrication must be developed in order to ensure the success of the molten metal route. These may include shortening the processing time, providing protective coatings on the reinforcements, or incorporating alloying elements that retard the reaction. The LMI process as shown in this work is one of the potential solutions for minimizing the reaction, with which no any other process can compete in shortening the processing/reaction time. The processing time is typically shorter than 150 ms in this case. The observed interfacial microstructures suggest the following elementary steps of the interaction between the injected SiCp and the molten Ti-alloy referring to the phase diagrams [20]: SiC→Si∗ ⫹ C∗(T ⬎ 1650°C)
(13)
Ti ⫹ C →TiC(1650°C ⬍ T ⬍ 3067°C)
(14)
∗
∗
and 5Ti∗ ⫹ 3Si∗→Ti5Si3(1650°C ⬍ T
(15)
⬍ 2130°C) where ∗ refers to the element in the molten Ti6Al4V material. First, when SiCp comes into contact with the liquid Ti-alloy, SiCp starts to dissolve in the liquid metal into carbon and silicon atoms according to reaction (13). As the concentrations of C and Si atoms in the adjacent melt increase and the melt rapidly cools down, TiC will form first by reaction (14), because its melting point is higher than that of Ti5Si3. The surface of SiCp serves as a preferential heterogeneous nucleation site of TiC reaction layer. Reaction (15) may occur only when Si content in the melt is high enough. The essential factor that governs the interfacial reaction is the formation of a TiC reaction layer surrounding the particles, which separates SiCp from a direct contact with the liquid metal. Dissolution of SiCp in Ti-melt at temperatures higher than 1650 °C is a fast process, indicated by the exothermal character according to simple thermal dynamic analysis. What should be kept in mind, however,
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are the facts that the LMI process is a non-equilibrium process and that the reaction happens at a wide temperature interval from approximately 3000–1650 °C. Due to rapid cooling of the melt behind the laser beam, the dissolution of SiCp makes the adjacent melt very quickly constitutionally supercooled by the carbon content. Therefore, TiC phases will precipitate and nucleate heterogeneously on SiCp. Depending on the character of the interfaces/facets of SiCp, the TiC nuclei may exhibit a definite OR to the SiC lattice on the basal plane or on the (101¯ 2) plane, or have no specific ORs on randomly oriented and curved interfaces. Iwamoto et al. [9,10] showed that both the basal planes and the (101¯ 2) planes of 6H-SiC possess a polar character. The exposed carbon-atom plane is suitable for a carbon bridge between TiC and SiC lattices that determines the ORs. They observed in situ during reactive wetting of SiC in a Ti-containing AgCu-alloy that SiC tends to dissolve along the basal plane and to produce an atomically flat plane. The nucleation sites of TiC nuclei are steps or defects on the basal plane. Such nucleated TiC islands are actually surrounded by the most favorable fivefold coordinated sites and will quickly spread over the whole facets of SiCp [7]. The subsequent stacking of TiC atomic layers mainly exhausts carbon atoms from the adjacent supercooled melts and, therefore, ends by the depletion of carbon. This is evidenced in Fig. 7 where a gradual profile of carbon intensity clearly indicates the quick decrease of carbon content in the monocrystalline TiC layer. Solid-state diffusion of carbon atoms through the formed monocrystalline TiC layer is negligible with respect to the rapid growth of the layer. In contrast, the polycrystalline TiC layers grow from randomly oriented nuclei. The grain boundaries between TiC grains, which are the last solidifying part of the layer, provide channels for fast transfer of carbon atoms from SiCp during the later period of growth. Therefore, it yields a longer period of growth that corresponds to a thicker layer and the depletion of carbon in the polycrystalline layer is also less severe, as confirmed in Fig. 8. From Fig. 6 and the preceding discussion, it is clear that the TiC reaction layers actually serve as a barrier against the interfacial reaction and eventu-
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ally stops the dissolution of SiCp. Consider a monocrystalline TiC layer formed at a temperature of 2500 °C, for instance, in the top region of a FGM track. A drop-down of temperature exposes the monocrystalline TiC layer under tensile stresses because of the different thermal expansion coefficients between TiC and SiC phases, 7.7 × 10⫺6 / K and 4.4 × 10⫺6 / K [21], respectively. A simple estimate predicts a tensile stress of 590 MPa in the layer with a temperature drop to 2100 °C, which is much higher than the tensile strength of a TiC layer. As a result, direct fracture or mechanical twinning may be induced in the monocrystalline TiC layers depending on the real stress levels, and both of them have been observed as can be seen in Figs. 6 and 11. Fracture of the monocrystalline TiC layer brings SiCp partially into a direct contact again with Ti-melt and an interfacial reaction starts again. However, the reaction at temperatures below 2130 °C (the melting point of Ti5Si3 phase in Ti-melt) favors the simultaneous precipitation of both TiC phases and Ti5Si3 phases, which produces a mixed reaction layer as is shown in Fig. 6. The broken TiC layer may still play a critical role in the formation of Ti5Si3 phases, which prevents silicon atoms to diffuse freely away into the melt. Higher concentration of silicon in the melt is needed to form Ti5Si3 phases, in comparison with the necessary carbon concentration for the formation of TiC phases at the low temperatures. Probably due to different growth behavior, i.e. faceted and non-faceted characters between the two phases, the mixed layer is not a dense layer and cannot halt the reaction anymore. The reaction will last till the surrounding melt resolidifies. Formation of monocrystalline TiC layers at a lower temperature, e.g. injecting SiCp at lower temperatures or into a deeper depth, will reduce their fracture probability and thus restrains the formation of the mixed reaction layer. Another way to avoid fracture of the monocrystalline TiC layer is to choose other carbide ceramic particles that have a similar thermal expansion coefficient as TiC, e.g. WC as reinforcement. On the other hand, polycrystalline TiC layers may relax via plastic deformation and fracture in polycrystalline TiC layers is hardly observed. Based on the model for the formation of FGMs
and the interfacial reaction mechanism, it is possible to interpret the thickness dependence of the TiC reaction layers on the beam velocity and the injection depth of SiCp as seen in Fig. 9. All the SiCp at the bottom of an FGM track are injected at the beginning of the injection process through a hotter melt, thus, longer reaction/growth time yields thicker TiC layers on these SiCp. In contrast, most of SiCp in the top region of an FGM track are injected at a later stage of injection and therefore the shorter reaction time and lower temperatures make the TiC ration layers thinner. In addition, the appearance of both the mono- and polycrystalline TiC layers in this region leads to a large spread in layer thickness. To conclude, the linear increase in the thickness of the TiC layers with the injection depth of SiCp in the FGM tracks produced at high beam velocities is determined by the particle injection sequence that determines reaction temperatures and growth time. In the case of a homogeneous composite track produced at the beam velocity of 8.3 m/min, a direct interaction with the beam makes SiCp at the top region overheated and heavily dissolved in the melt. The decrease in thickness of the TiC layers with the injection depth of SiCp is mainly attributed to the large temperature gradient between the top and the bottom of the laser pool.
5. Conclusions The main conclusions of this work are the following: 1. For the first time particle-reinforced FGMs are produced with a well-controlled LMI process. 2. A clear physical picture of the particle injection process is obtained on the base of FEA modeling and theoretical calculations. Temperature/ viscosity field of the laser melt pool explains the formation of the FGMs by injecting SiCp into the elongated part of the laser pool behind the beam. The effective viscosity of the SiCp/Ti6Al4V suspension dominates the maximum injection depth of SiCp. 3. Three types of reaction layers are observed around SiCp: a thin monocrystalline TiC layer,
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a cellular polycrystalline TiC layer and a thick mixed layer of TiC with Ti5Si3 particular phases. Both kind of TiC reaction layers act as a barrier against the interfacial reaction and its swift formation during rapid cooling stops the dissolution of SiCp in Ti-melt. The mixed layer origins at damages of a monocrystalline TiC layer. Only the monocrystalline TiC layer exhibits particular ORs with the lattice of SiCp, i.e. (111)TiC(0001)SiC and 具110典TiC具11¯ 00典SiC or (111)TiC(101¯ 2)SiC and 具11¯ 0典TiC具12¯ 10典SiC. 4. The thickness of the reaction layer is minimized by injecting the SiCp fully outside the laser beam in the stretched molten area behind the beam.
Acknowledgements The Netherlands Institute for Metals Research is acknowledged for its financial support. Drs A. Vreeling, P.M. Bronsveld and D.T.L. van Agterveld are acknowledged for their assistance and contributions during these examinations. S.A.J. Nijman and T. Silva are thanked for their assistance in the FEA modeling. H. Bron, K. Post and J. Harkema are acknowledged for their help with the analysis and laser experiments.
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