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Microscopic morphology and distribution of TiC phase in laser clad NiCrBSiC–TiC layer on titanium alloy substrate
Surface and Coatings Technology 155 (2002) 203–207
Microscopic morphology and distribution of TiC phase in laser clad NiCrBSiC–TiC layer on titanium alloy substrate R.L. Sun*, J.F. Mao, D.Z. Yang School of Material Science and Engineering, Harbin Institute of Technology, P.O. Box 433, Harbin 150001, PR China Received 4 September 2001; accepted in revised form 3 January 2002
Abstract The microscopic morphology and distribution of TiC phase, as well as its interfaces with a g–Ni matrix, in a laser clad NiCrBSiC–TiC layer on a Ti–6Al–4V alloy substrate were characterized using TEM and SEM. The experimental results showed that during laser irradiation heating, TiC particles were partially dissolved into the melted Ni-base alloy, and the dissolved Ti and C atoms were precipitated in the form of TiC dendrites during cooling. The liquid-precipitated TiC particles with smaller sizes are mainly distributed within the g–Ni dendrites, while the larger liquid-precipitated TiC and the undissolved TiC particles are located in the regions between the g–Ni dendrites. Also, a phenomenon of some TiC particles meeting with each other and merging together was observed in the clad zone. The interfaces of liquid-precipitated TiC phase with the g–Ni matrix possess a faceted feature, while the undissolved TiC particles have smooth curved interfaces of a non-faceted type. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Titanium alloy; Laser cladding; Metalyceramics composite coatings; TiC phase; Microstructure
1. Introduction It is well known that laser cladding is an effective technique to prevent metallic materials from wearing w1,2x, eroding w3,4x and oxidizing w5,6x. In recent years, laser cladding has been investigated widely. The substrate materials have been extended from carbon and alloy steels to aluminum w7–9x and titanium w10–16x alloys. The clad materials ranged from the Fe-base, Nibase and Co-base alloys to the ceramic powders of WC w17,18x, TiC w19,20x, SiC w21x, etc. The ceramic powders are used for incorporating with alloy powders to obtain the ceramicsymetal composite coatings. Liu and Mazumder et al. w7,8x have investigated the microstructure of laser clad Ni–Al bronze on an Al alloy substrate. Molin and Hualun w12x demonstrated that, after laser cladding of Ti–6Al–4V alloy with BNqNiCrCoAlY powders, the maximum hardness of the clad layer reached HV1600, and the sliding wear rates were two orders of magnitude less than that of the heat-treated *Corresponding author. Tel.: q86-0451-6414445; fax: q86-04516415168. E-mail address: [email protected] (R.L. Sun).
titanium alloy. Ayers w16x has examined the wear behavior of laser clad TiC andyor WC layers on Al-base and Ti-base alloys, indicating that the wear-resistance of the clad layers would depend on the type, size and amount of the carbides, and a wear-resistance 7;38 times that of the substrate was obtained. Previous work by the present authors has shown that the laser clad NiCrBSiC w22x and NiCrBSiC–TiC w23x layers on Ti–6Al–4V alloy exhibit very high microhardness (HV900;1200) and wear-resistance (a sliding wear rate an order of magnitude less). The aim of this study was to examine the TiC phase morphology and distribution in NiCrBSiC–TiC composite coatings using SEM and TEM, including the feature of interface between the Ni-base alloy matrix and TiC, as a function of laser processing energy. 2. Experimental Ti–6Al–4V alloy hot rolled in the aqb range and annealed at 780 8C was used as the substrate material in this study, and the size of samples was f30=20 mm. The clad binder material was NiCrBSiC pre-alloyed powders with the composition (wt.%) of: 17.0Cr, 3.5B,
0257-8972/02/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 2 . 0 0 0 0 6 - 3
R.L. Sun et al. / Surface and Coatings Technology 155 (2002) 203–207
204
Table 1 Effects of specific laser energy on the dimensions and dilution of the clad layer PyVD (kJycm2)
W (mm)
H (mm)
h (mm)
A (%)
12.1 8.3 6.2 4.4
6.38 6.16 5.90 5.86
0.88 0.87 0.69 0.58
0.54 0.26 0.19 0.10
38 23 21 14
and a Philips CM12 type transmission electron microscope with an EDAX-9100 system. The specimens for SEM observations were cut perpendicularly and transversely to the clad layer, mechanically polished and then etched using an etchant of HFyHNO3s1:15. The TEM samples were cut along the direction parallel to the clad layer, and the foils were prepared by ion-thinning. 3. Results and discussion 3.1. Morphology of TiC phase in clad zone
4.0Si, 1.0C, -12Fe, and balance Ni. The particle sizes of the powders were in the range of 50–100 mm. In order to obtain the laser clad NiCrBSiC–TiC layer, TiC and pre-alloyed NiCrBSiC powders were mechanically mixed according to the radio TiC:NiCrBSiCs1:2, i.e. 33 vol.%TiC. Clad powders were pre-placed on the surface of substrates using an organic binder, to form a layer of 1.0 mm thickness. Before the powders were pasted on the surface of the substrate was ground, using emery paper to remove the oxide scale. A 9-kW CO2 laser was used, providing a beam that was directed onto the sample surface using gold-coated water-cooled copper mirrors. The parameters of laser processing were selected as: the output power Ps4–6 kW, scanning speed Vs5–15 mmys, and beam diameter ds6.0 mm. The laser cladding was carried out in argon, using side jet shielding with an argon flux of 20 lymin. The effect of laser cladding parameters can be expressed by the ratio PyVD, which is referred to as specific laser energy. The laser clad layer can be divided into two regions: the clad zone (CZ) and the dilution zone (DZ) w23x. The dimension of the laser clad layer can be given by the width (w), height (H) of the clad zone and the depth (h) of melted substrate material. The dilution of the clad layer is calculated in term of A%shyHqh. The microstructure and phase constitution of the laser clad layer were examined using a Hitachi S-570 type scanning electron microscope, a S-530 type SEM equipped with a Link ISIS energy spectrum analyzer,
The dimensions of the clad and the dilution zones are related to the specific laser energy, as shown in Table 1. Fig. 1 shows the effect of specific laser energy on morphology of the TiC phase in the clad zone. Because the hardness of TiC phase (HVs3000–3200) is much higher than the Ni-base alloy (HVf650–800), TiC particles are exposed on the polished surface. When the specific laser energy is relatively low, most TiC particles have edge angles and irregular shapes (Fig. 1a). With the increase of specific laser energy, the size of the TiC particles is reduced and their edges become relatively smooth (Fig. 1b), implying that the TiC particles are partially dissolved during laser cladding. When the specific laser energy is further increased, TiC particles with irregular shapes are rarely observed and many fine dendrites appear instead (Fig. 1c). The SEM energy spectrum analysis showed that the dendrites were TiC phase. The solution of Ti and C from the TiC particles into the liquid leads to the formation of TiC phase (e.g. as dendrites) during solidification. Dilution from the Ti alloy substrate is a further source of Ti for the melt pool, while the Ni based alloy also provides a source of C. The TiC phase formed during solidification is referred to here as the liquid-precipitated phase. Therefore, with increasing the specific laser energy, the extent of enrichment of the melt in Ti and C is increased, leading to more liquid-precipitated TiC formed during solidification.
Fig. 1. SEM micrographs showing the effect of specific laser energy on TiC phase morphology on mechanically polished surface of the clad zone. PyVD: (a) 6.2; (b) 8.3; and (c) 12.1 kJycm2.
R.L. Sun et al. / Surface and Coatings Technology 155 (2002) 203–207
Fig. 2. SEM micrograph showing that the liquid-precipitated TiC phase nucleates at an undissolved TiC particle and grows into fine dendrites (PyVDs12.1 kJycm2).
Fig. 2 shows that the liquid-precipitated TiC phase nucleates at undissolved TiC particles, and grows into fine dendrites. Fig. 3 shows that the TiC phase can also nucleate independently in the liquid, and exhibits a less well-developed dendrite morphology with some faceting. When the specific laser energy is relatively low, e.g. as in Fig. 1a, epitaxial growth occurs on the undissolved TiC particles, as fine columnar grains (Fig. 4). This epitaxially grown layer is too narrow (approx. 5;20 nm in extent) to be observed by SEM (Fig. 1a). With the increase in the specific laser energy, the extent of solution of TiC particles in the liquid increases, leading to the differences in TiC morphology and distribution in the clad layer.
205
Fig. 3. TEM micrograph showing a TiC dendrite with faceted interfaces, nucleated in the liquid independently of undissolved TiC particles (PyVDs12.1 kJycm2).
(Fig. 6a). The density of TiC phase is less than that of the Ni-base alloy, leading to a tendency for the TiC particles to float up in the melt pool. According to the Stokes law w24x, the floatation velocity is proportional to the square of the particle radius. The larger TiC particles would float up more quickly than the smaller ones, thus allowing TiC particles with various dimensions to come into contact. Also, the intense convection movement in the melt pool might promote the mutual collision between TiC particles. In the powder mixture used in this study, the volume fraction of TiC is 33 vol.%, and the sizes of TiC powder particles are in the range of 1;10 mm. The particle ‘density’ is
3.2. Distribution of TiC phase in clad zone The distribution of TiC phase in the clad zone might be related to the convection movement of the melt, the interaction between TiC particles, and the interaction of TiC phase with liquidysolid interfaces in the melt pool during solidification. The smaller liquid-precipitated TiC particles are mainly distributed within the g–Ni dendrites (Fig. 5a). The larger liquid-precipitated TiC and the undissolved TiC particles exist between the g–Ni dendrites (Fig. 5b,c). During solidification, the smaller TiC particles might easily be captured by the moving liquidysolid interfaces of g–Ni dendrites and incorporated into the g–Ni phase. In contrast, the larger TiC particles would be rejected by the moving liquidysolid interfaces and are thus located in the regions of final solidification between g–Ni dendrites. In the clad zone, some TiC particles, usually having different dimensions, came into contact with each other
Fig. 4. TEM micrograph showing epitaxial growth of TiC phase in Fig. 1a (PyVDs6.2 kJycm2).
206
R.L. Sun et al. / Surface and Coatings Technology 155 (2002) 203–207
Fig. 7. TEM micrographs showing the interfaces of liquid-precipitated (a) and undissolved (b) TiC particles with g–Ni matrix in the clad zone (PyVDs8.3 kJycm2).
matrix in the clad zone, respectively. The interfaces of the liquid-precipitated TiC particles with the g–Ni matrix are faceted, while the undissolved TiC particles have the smooth curved interfaces of non-faceted type. 3.3. Morphology of TiC phase in dilution zone Fig. 5. TEM micrographs showing distribution of TiC phase in the clad zone: (a) fine liquid-precipitated TiC particles located within g– Ni dendrite; (b) larger liquid-precipitated TiC particles existing between g–Ni dendrites; (c) undissolved TiC particles existing between g–Ni dendrites.
106;107TiCymm3, and the probability of particle collision in the melt would be high (e.g. Fig. 6b). Furthermore, the epitaxially solidified TiC layer on the undissolved TiC particles would reduce the space between collided particles (Fig. 6b). The extent of TiC particles merging together will depend on the time period for which the melt pool exists prior to completion of the rapid solidification process. Fig. 7a,b show interfaces of the liquid-precipitated TiC and the undissolved TiC particles with the g–Ni
Fig. 6. TEM micrographs showing some TiC particles meeting with each other (a) and merging together (b) in clad zone.
In the dilution zone, most TiC phase is in the form of dendrites (Fig. 8), and undissolved TiC particles with irregular shapes are rarely observed. Fig. 9 shows the cross-section morphology of the TiC dendrite arms in the b–Ti matrix. The composition (excluding carbon and boron) of the b–Ti matrix as determined by TEM EDX analysis is (wt.%): 73.7Ti, 11.3Ni, 4.6Fe, 5.3Cr, 0.8Si and 4.4Al. Regions of ‘dark contrast’ with a thickness of 0.1;0.2mm exist between the TiC dendrite arms and the b–Ti matrix (see the arrow in Fig. 9a). The composition of the ‘dark contrast’ bands is (wt.%): 63.5Ti, 25.5Ni, 5.3Fe, 2.6Cr, 0.9Si and 2.2Al, showing lower Ti and higher Ni contents near the b–Ti matrix. This feature is associated with the diffusion process, as
Fig. 8. SEM micrograph showing the morphology of TiC phase in the form of dendrites in the dilution zone (PyVDs12.1 kJycm2).
R.L. Sun et al. / Surface and Coatings Technology 155 (2002) 203–207
207
References
Fig. 9. TEM micrographs showing the cross-section morphology of TiC dendrite arms in the dilution zone (PyVDs12.1 kJycm2): (a) bright field image; (b) SADP of TiC; (c) SADP of b–Ti.
Ti is incorporated into the TiC with ‘rejection’ of Ni from the carbide. 4. Conclusions During laser irradiation heating, TiC powder particles were dissolved into the melted Ni-base alloy, to extents depending on specific energy of the laser processing. During solidification of the laser generated melt pool, TiC dendrites formed. The liquid-precipitated TiC particles with relatively small sizes are mainly distributed within g–Ni dendrites, while the larger liquid-precipitated TiC and the undissolved TiC particles are located in the regions between g–Ni dendrites. The interfaces of liquid-precipitated TiC phase with the g–Ni matrix show a faceting, and the undissolved TiC particles have smooth curved interface of nonfaceted type.
w1x M. Corchia, P. Delogu, F. Nenci, A. Belmondo, S. Corcorulo, W. Stabielli, Wear 119 (1987) 137. w2x T.C. Lei, J.H. Ouyang, Y.T. Pei, Y. Zhou, Mater. Sci. Technol. 11 (1995) 520. w3x V.M. Weerasinghe, W.M. Steen, D.R.F. West, Surf. Eng. 3 (1987) 147. w4x R. Li, M.G.S. Ferreira, M.A. Anjos, R. Vilar, Surf. Coat. Technol. 88 (1996) 96. w5x H.M. Tawancy, N.M. Abbas, A. Bennett, Surf. Coat. Technol. 68y69 (1994) 10. w6x C.A. Liu, M.J. Humphries, R.C. Krutenat, Thin Solid Films 107 (1983) 269. w7x Y. Liu, J. Mazumder, K. Shibata, Metall. Mater.Trans. A 25A (1994) 37. w8x Y. Liu, J. Mazumder, K. Shibata, Metall. Mater. Trans. A 26A (1995) 1519. w9x G.Y. Liang, T.T. Wong, J.M.K. MacAlpine, J.Y. Su, Surf. Coat. Technol. 127 (2000) 233. w10x J.A. Folks, K. Shibata, J. Laser Appl. 6 (1994) 88. w11x J.H. Abboud, D.R.F. West, R.H. Hibberd, Surf. Eng. 9 (1993) 221. w12x P.A. Molian, L. Hualum, Wear 130 (1989) 337. w13x J.H. Abboud, D.R.F. West, R.D. Rawlings, Mater. Sci. Technol. 10 (1994) 848. w14x S. Mridha, T.N. Baker, Mater. Sci. Technol. 12 (1996) 595. w15x A. Mehlmann, S.F. Dirnfeld, I. Minkoff, Surf. Coat. Technol. 42 (1990) 275. w16x J.D. Ayers, Wear 97 (1984) 249. w17x W. Cerri, R. Martinella, G.P. Mor, Surf. Coat. Technol. 49 (1991) 40. w18x J. Mateos, J.M. Cuetos, E. Fernandez, R. Vijande, Wear 239 (2000) 274. w19x J.H. Ouyang, Y.T. Pei, T.C. Lei, Y. Zhou, Wear 185 (1995) 167. w20x Q. Li, T.C. Lei, W.Z. Chen, Surf. Coat. Technol. 114 (1999) 278. w21x G. Abbas, D.R.F. West, Wear 143 (1991) 353. w22x R.L. Sun, D.Z. Yang, L.X. Guo, et al., Surf. Coat. Technol. 132 (2000) 251. w23x R.L. Sun, D.Z. Yang, L.X. Guo, et al., Surf. Coat. Technol. 135 (2001) 307. w24x M.C. Flemings, Solidification Processing, McGraw-Hill, New York, 1974, p. 194.
Microscopic morphology and distribution of TiC phase in laser clad NiCrBSiC–TiC layer on titanium alloy substrate R.L. Sun*, J.F. Mao, D.Z. Yang School of Material Science and Engineering, Harbin Institute of Technology, P.O. Box 433, Harbin 150001, PR China Received 4 September 2001; accepted in revised form 3 January 2002
Abstract The microscopic morphology and distribution of TiC phase, as well as its interfaces with a g–Ni matrix, in a laser clad NiCrBSiC–TiC layer on a Ti–6Al–4V alloy substrate were characterized using TEM and SEM. The experimental results showed that during laser irradiation heating, TiC particles were partially dissolved into the melted Ni-base alloy, and the dissolved Ti and C atoms were precipitated in the form of TiC dendrites during cooling. The liquid-precipitated TiC particles with smaller sizes are mainly distributed within the g–Ni dendrites, while the larger liquid-precipitated TiC and the undissolved TiC particles are located in the regions between the g–Ni dendrites. Also, a phenomenon of some TiC particles meeting with each other and merging together was observed in the clad zone. The interfaces of liquid-precipitated TiC phase with the g–Ni matrix possess a faceted feature, while the undissolved TiC particles have smooth curved interfaces of a non-faceted type. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Titanium alloy; Laser cladding; Metalyceramics composite coatings; TiC phase; Microstructure
1. Introduction It is well known that laser cladding is an effective technique to prevent metallic materials from wearing w1,2x, eroding w3,4x and oxidizing w5,6x. In recent years, laser cladding has been investigated widely. The substrate materials have been extended from carbon and alloy steels to aluminum w7–9x and titanium w10–16x alloys. The clad materials ranged from the Fe-base, Nibase and Co-base alloys to the ceramic powders of WC w17,18x, TiC w19,20x, SiC w21x, etc. The ceramic powders are used for incorporating with alloy powders to obtain the ceramicsymetal composite coatings. Liu and Mazumder et al. w7,8x have investigated the microstructure of laser clad Ni–Al bronze on an Al alloy substrate. Molin and Hualun w12x demonstrated that, after laser cladding of Ti–6Al–4V alloy with BNqNiCrCoAlY powders, the maximum hardness of the clad layer reached HV1600, and the sliding wear rates were two orders of magnitude less than that of the heat-treated *Corresponding author. Tel.: q86-0451-6414445; fax: q86-04516415168. E-mail address: [email protected] (R.L. Sun).
titanium alloy. Ayers w16x has examined the wear behavior of laser clad TiC andyor WC layers on Al-base and Ti-base alloys, indicating that the wear-resistance of the clad layers would depend on the type, size and amount of the carbides, and a wear-resistance 7;38 times that of the substrate was obtained. Previous work by the present authors has shown that the laser clad NiCrBSiC w22x and NiCrBSiC–TiC w23x layers on Ti–6Al–4V alloy exhibit very high microhardness (HV900;1200) and wear-resistance (a sliding wear rate an order of magnitude less). The aim of this study was to examine the TiC phase morphology and distribution in NiCrBSiC–TiC composite coatings using SEM and TEM, including the feature of interface between the Ni-base alloy matrix and TiC, as a function of laser processing energy. 2. Experimental Ti–6Al–4V alloy hot rolled in the aqb range and annealed at 780 8C was used as the substrate material in this study, and the size of samples was f30=20 mm. The clad binder material was NiCrBSiC pre-alloyed powders with the composition (wt.%) of: 17.0Cr, 3.5B,
0257-8972/02/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 2 . 0 0 0 0 6 - 3
R.L. Sun et al. / Surface and Coatings Technology 155 (2002) 203–207
204
Table 1 Effects of specific laser energy on the dimensions and dilution of the clad layer PyVD (kJycm2)
W (mm)
H (mm)
h (mm)
A (%)
12.1 8.3 6.2 4.4
6.38 6.16 5.90 5.86
0.88 0.87 0.69 0.58
0.54 0.26 0.19 0.10
38 23 21 14
and a Philips CM12 type transmission electron microscope with an EDAX-9100 system. The specimens for SEM observations were cut perpendicularly and transversely to the clad layer, mechanically polished and then etched using an etchant of HFyHNO3s1:15. The TEM samples were cut along the direction parallel to the clad layer, and the foils were prepared by ion-thinning. 3. Results and discussion 3.1. Morphology of TiC phase in clad zone
4.0Si, 1.0C, -12Fe, and balance Ni. The particle sizes of the powders were in the range of 50–100 mm. In order to obtain the laser clad NiCrBSiC–TiC layer, TiC and pre-alloyed NiCrBSiC powders were mechanically mixed according to the radio TiC:NiCrBSiCs1:2, i.e. 33 vol.%TiC. Clad powders were pre-placed on the surface of substrates using an organic binder, to form a layer of 1.0 mm thickness. Before the powders were pasted on the surface of the substrate was ground, using emery paper to remove the oxide scale. A 9-kW CO2 laser was used, providing a beam that was directed onto the sample surface using gold-coated water-cooled copper mirrors. The parameters of laser processing were selected as: the output power Ps4–6 kW, scanning speed Vs5–15 mmys, and beam diameter ds6.0 mm. The laser cladding was carried out in argon, using side jet shielding with an argon flux of 20 lymin. The effect of laser cladding parameters can be expressed by the ratio PyVD, which is referred to as specific laser energy. The laser clad layer can be divided into two regions: the clad zone (CZ) and the dilution zone (DZ) w23x. The dimension of the laser clad layer can be given by the width (w), height (H) of the clad zone and the depth (h) of melted substrate material. The dilution of the clad layer is calculated in term of A%shyHqh. The microstructure and phase constitution of the laser clad layer were examined using a Hitachi S-570 type scanning electron microscope, a S-530 type SEM equipped with a Link ISIS energy spectrum analyzer,
The dimensions of the clad and the dilution zones are related to the specific laser energy, as shown in Table 1. Fig. 1 shows the effect of specific laser energy on morphology of the TiC phase in the clad zone. Because the hardness of TiC phase (HVs3000–3200) is much higher than the Ni-base alloy (HVf650–800), TiC particles are exposed on the polished surface. When the specific laser energy is relatively low, most TiC particles have edge angles and irregular shapes (Fig. 1a). With the increase of specific laser energy, the size of the TiC particles is reduced and their edges become relatively smooth (Fig. 1b), implying that the TiC particles are partially dissolved during laser cladding. When the specific laser energy is further increased, TiC particles with irregular shapes are rarely observed and many fine dendrites appear instead (Fig. 1c). The SEM energy spectrum analysis showed that the dendrites were TiC phase. The solution of Ti and C from the TiC particles into the liquid leads to the formation of TiC phase (e.g. as dendrites) during solidification. Dilution from the Ti alloy substrate is a further source of Ti for the melt pool, while the Ni based alloy also provides a source of C. The TiC phase formed during solidification is referred to here as the liquid-precipitated phase. Therefore, with increasing the specific laser energy, the extent of enrichment of the melt in Ti and C is increased, leading to more liquid-precipitated TiC formed during solidification.
Fig. 1. SEM micrographs showing the effect of specific laser energy on TiC phase morphology on mechanically polished surface of the clad zone. PyVD: (a) 6.2; (b) 8.3; and (c) 12.1 kJycm2.
R.L. Sun et al. / Surface and Coatings Technology 155 (2002) 203–207
Fig. 2. SEM micrograph showing that the liquid-precipitated TiC phase nucleates at an undissolved TiC particle and grows into fine dendrites (PyVDs12.1 kJycm2).
Fig. 2 shows that the liquid-precipitated TiC phase nucleates at undissolved TiC particles, and grows into fine dendrites. Fig. 3 shows that the TiC phase can also nucleate independently in the liquid, and exhibits a less well-developed dendrite morphology with some faceting. When the specific laser energy is relatively low, e.g. as in Fig. 1a, epitaxial growth occurs on the undissolved TiC particles, as fine columnar grains (Fig. 4). This epitaxially grown layer is too narrow (approx. 5;20 nm in extent) to be observed by SEM (Fig. 1a). With the increase in the specific laser energy, the extent of solution of TiC particles in the liquid increases, leading to the differences in TiC morphology and distribution in the clad layer.
205
Fig. 3. TEM micrograph showing a TiC dendrite with faceted interfaces, nucleated in the liquid independently of undissolved TiC particles (PyVDs12.1 kJycm2).
(Fig. 6a). The density of TiC phase is less than that of the Ni-base alloy, leading to a tendency for the TiC particles to float up in the melt pool. According to the Stokes law w24x, the floatation velocity is proportional to the square of the particle radius. The larger TiC particles would float up more quickly than the smaller ones, thus allowing TiC particles with various dimensions to come into contact. Also, the intense convection movement in the melt pool might promote the mutual collision between TiC particles. In the powder mixture used in this study, the volume fraction of TiC is 33 vol.%, and the sizes of TiC powder particles are in the range of 1;10 mm. The particle ‘density’ is
3.2. Distribution of TiC phase in clad zone The distribution of TiC phase in the clad zone might be related to the convection movement of the melt, the interaction between TiC particles, and the interaction of TiC phase with liquidysolid interfaces in the melt pool during solidification. The smaller liquid-precipitated TiC particles are mainly distributed within the g–Ni dendrites (Fig. 5a). The larger liquid-precipitated TiC and the undissolved TiC particles exist between the g–Ni dendrites (Fig. 5b,c). During solidification, the smaller TiC particles might easily be captured by the moving liquidysolid interfaces of g–Ni dendrites and incorporated into the g–Ni phase. In contrast, the larger TiC particles would be rejected by the moving liquidysolid interfaces and are thus located in the regions of final solidification between g–Ni dendrites. In the clad zone, some TiC particles, usually having different dimensions, came into contact with each other
Fig. 4. TEM micrograph showing epitaxial growth of TiC phase in Fig. 1a (PyVDs6.2 kJycm2).
206
R.L. Sun et al. / Surface and Coatings Technology 155 (2002) 203–207
Fig. 7. TEM micrographs showing the interfaces of liquid-precipitated (a) and undissolved (b) TiC particles with g–Ni matrix in the clad zone (PyVDs8.3 kJycm2).
matrix in the clad zone, respectively. The interfaces of the liquid-precipitated TiC particles with the g–Ni matrix are faceted, while the undissolved TiC particles have the smooth curved interfaces of non-faceted type. 3.3. Morphology of TiC phase in dilution zone Fig. 5. TEM micrographs showing distribution of TiC phase in the clad zone: (a) fine liquid-precipitated TiC particles located within g– Ni dendrite; (b) larger liquid-precipitated TiC particles existing between g–Ni dendrites; (c) undissolved TiC particles existing between g–Ni dendrites.
106;107TiCymm3, and the probability of particle collision in the melt would be high (e.g. Fig. 6b). Furthermore, the epitaxially solidified TiC layer on the undissolved TiC particles would reduce the space between collided particles (Fig. 6b). The extent of TiC particles merging together will depend on the time period for which the melt pool exists prior to completion of the rapid solidification process. Fig. 7a,b show interfaces of the liquid-precipitated TiC and the undissolved TiC particles with the g–Ni
Fig. 6. TEM micrographs showing some TiC particles meeting with each other (a) and merging together (b) in clad zone.
In the dilution zone, most TiC phase is in the form of dendrites (Fig. 8), and undissolved TiC particles with irregular shapes are rarely observed. Fig. 9 shows the cross-section morphology of the TiC dendrite arms in the b–Ti matrix. The composition (excluding carbon and boron) of the b–Ti matrix as determined by TEM EDX analysis is (wt.%): 73.7Ti, 11.3Ni, 4.6Fe, 5.3Cr, 0.8Si and 4.4Al. Regions of ‘dark contrast’ with a thickness of 0.1;0.2mm exist between the TiC dendrite arms and the b–Ti matrix (see the arrow in Fig. 9a). The composition of the ‘dark contrast’ bands is (wt.%): 63.5Ti, 25.5Ni, 5.3Fe, 2.6Cr, 0.9Si and 2.2Al, showing lower Ti and higher Ni contents near the b–Ti matrix. This feature is associated with the diffusion process, as
Fig. 8. SEM micrograph showing the morphology of TiC phase in the form of dendrites in the dilution zone (PyVDs12.1 kJycm2).
R.L. Sun et al. / Surface and Coatings Technology 155 (2002) 203–207
207
References
Fig. 9. TEM micrographs showing the cross-section morphology of TiC dendrite arms in the dilution zone (PyVDs12.1 kJycm2): (a) bright field image; (b) SADP of TiC; (c) SADP of b–Ti.
Ti is incorporated into the TiC with ‘rejection’ of Ni from the carbide. 4. Conclusions During laser irradiation heating, TiC powder particles were dissolved into the melted Ni-base alloy, to extents depending on specific energy of the laser processing. During solidification of the laser generated melt pool, TiC dendrites formed. The liquid-precipitated TiC particles with relatively small sizes are mainly distributed within g–Ni dendrites, while the larger liquid-precipitated TiC and the undissolved TiC particles are located in the regions between g–Ni dendrites. The interfaces of liquid-precipitated TiC phase with the g–Ni matrix show a faceting, and the undissolved TiC particles have smooth curved interface of nonfaceted type.
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