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Study of aluminium–copper–iron alloys: application for laser cladding
Surface and Coatings Technology 151 – 152 (2002) 329–332
Study of aluminium–copper–iron alloys: application for laser cladding L. Dubourga,b,*, F. Hlawkab, A. Cornetb a
ˆ API, 67400 Illkirch, France IREPA LASER, Parc d’Innovation, Pole b LISS, ENSAIS, 24 bd de la Victoire, 67000 Strasbourg, France
Abstract The modifications brought to aluminium by addition of copper and iron are investigated after rapid surface solidification by a CO2 laser. In order to carry out an exhaustive study, the laser surface alloys are obtained using a precoating process. Samples are characterised using microscopic examination, X-ray diffraction and hardness measurement. The optimum composition (42 wt.% Cu – 5 wt.% Fe) is selected for its fine and homogeneous microstructure (Al2 Cu and Al7 Cu2 Fe phases) and its high hardness (370 Hv0.2). 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Aluminium; Laser alloying; Laser cladding; Intermetallic compounds
1. Introduction Aluminium alloys offer advantages in automotive, aircraft and spacecraft industries, notably due to their lightness and corrosion resistance. However, these materials possess poor mechanical and tribological characteristics. Laser surface treatments can correct these drawbacks without affecting the global properties of the aluminium piece. For this reason, several authors have investigated the formation of intermetallic compounds with aluminium by laser (Fe w1x, Cu w2x, Mo w3x and Cr w4x). Previous studies indicated the great utility of copper in aluminium laser alloying w2x. The mechanical characteristics and the wear behaviour of such laser Al –Cu alloys make them ideal for laser alloying or cladding applications w5x. The present study shows the influence of Cu and Fe content in aluminium upon the microstructure and hardness. In order to do an exhaustive study, the laser surface alloys are obtained using a precoating process. A powder layer with a fixed Cu–Fe composition is placed on the aluminium specimen. A CO2 laser beam melts this layer together with the aluminium surface. Convection flows, generated by the temperature gradient between the beam centre and the sides of the melted pool, ensure the precoating incorporation and the homogeneity of this surface alloy w6x. * Corresponding author. Tel.: q33-3-88-65-54-00; fax: q33-3-8865-54-01. E-mail address: [email protected] (L. Dubourg).
The tracks are analysed using microscopic examination, X-ray diffraction and micro-hardness measurements. Subsequently, the optimum composition in relation to the microstructure, homogeneity and hardness is selected. 2. Experimental set-up Surface treatments are carried out with a continuous CO2 laser of 10.6-mm wavelength and energy spatial distribution of TEM20 type. The process parameters are as follows: laser power of 4000 W on the sample, scanning speed of 4.2=10y3 m sy1, beam diameter of 4 mm on the substrate and sample preheating of 150 8C. The interaction zone is protected from air oxidation with an inert gas (1 bar pressure, Argon flow rate of 8.3=10y2 l sy1 and helium flow rate of 16.7=10y2 l sy1). The precoating is composed of metallic powders and a binder. The powders have a granulometry of 45– 150 mm and their compositions are as follows: Fe (99.5 wt.%), Cu (90 wt.% Cu, 10 wt.% Al). The binder is composed of 94 wt.% of water (evaporated during the precoating drying at 80 8C for 30 min) and 6 wt.% of an organic compound (vaporised during laser scanning). The microstructure analysis is performed by optical microscopy after a Keller’s reagent etching and by Xray diffraction using a Siemens D5000 diffractometer (CuKa radiation). Vickers hardness measurements are performed in the cross-section with a load of 2 N.
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 1 . 0 1 5 9 1 - 2
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Fig. 1. Micrographs of coatings composed of (wt.%): (a) 15Fe; (b) 42Cu – 5Fe.
Average hardness is calculated using eight micro-hardness measurements (see Fig. 1a). 3. Results and discussion Different coatings are achieved with ratios of 0 up to 45 wt.% Cu and of 0 up to 30 wt.% Fe. For these concentrations, each coating is free of porosity (Fig. 1). Between 0 and 40 wt.% Cu, the Al–Cu alloys are homogeneous on the entire coating. Above 40 wt.% Cu, cracks appear in the coatings due to the low ductility of the alloy w7x. Al–Fe alloys are homogeneous up to approximately 15 wt.% Fe. Above 15 wt.% Fe, ironrich zones and small cracks appear in the coating bottom (cf. Fig. 1a). Above 30 wt.% Fe, alloy homogeneity becomes too poor to ensure an isotropic mechanical behaviour of the track. Cu addition improves the homogeneity of the Al–Fe alloys (Fig. 1b). Nevertheless, when the Fe ratio is higher than the Cu ratio (approx. wt.% Fe)1.5=wt.% Cu), the homogeneity decreases. As a result, the appearance of some non-melted particles in the coating bottom is observed. Fig. 2 shows the obtained alloys and their positions on the Al–Cu–Fe equilibrium diagram at various Cu and Fe ratios (aluminium-rich zone of the Al–Cu–Fe diagram w8x). Grey areas of the diagram illustrate the zones where the microstructure of the obtained alloys is similar. If only Cu is added, the phases identified by microscopic examinations and X-ray diffraction (Fig. 3a) fit with the equilibrium diagram. In this case, rapid quenching due to the laser process creates equilibrium phases. Up to 24 wt.% Cu, a hypoeutectic microstructure is observed: white primary grains (Al)qa dark eutectic (Fig. 2i). In the range of 27 and 37 wt.% Cu, the eutectic microstructure is always present, as shown in Fig. 2h. Theoretically, according to the equilibrium Al– Cu diagram, this structure exists for a specific Cu percentage, approximately 32 wt.%. The growth of primary phases w(Al) and u-Al2Cux is limited because of the rapid quenching w7x. Finally, in the range of 37–
40 wt.% Cu, micrographs (Fig. 2g) show a hypereutectic microstructure: angular primary grains u-Al2Cuqa dark eutectic. If only Fe is added, the alloys have the same microstructure (Fig. 2l) confirmed by X-ray diffraction analysis. This microstructure is composed of primary needle-like phases of l-Al3Fe, white dendrites (Al) and a dark eutectic. The microstructure does not fit with the equilibrium diagram, as primary phases of (Al) and lAl3Fe can be observed at the same time. For 24 wt.% Fe, the eutectic phase disappears except at grain boundaries, as shown in Fig. 2a. Moreover, X-ray diffraction analysis also indicates the presence of a metastable phase Al6Fe at 24 wt.% Fe. According to Richmond et al. w9x, this metastable phase is present in the eutectic. In the case of Cu–Fe combined addition, four different microstructures can be noted according to the composition. For low Cu and Fe ratios (0-wt.% Cu-15 and 0-wt.% Fe-6), the microstructure is composed of primary dendrites (Al) and eutectic (Fig. 2i). For low Cu (0-wt.% Cu-15) and high Fe ratios (8-wt.% Fe-22), the microstructure is composed of phases lAl3Fe, dendrites (Al) and eutectic (Fig. 2b). The lAl3Fe phase precipitates with a needle-like shape as in Al–Fe alloy (Fig. 2l) in spite of the presence of copper. The X-ray diffraction analysis does not show the presence of the phase a-Al82Cu4Fe14, according to the equilibrium diagram shown in Fig. 2. For high Cu (35wt.% Cu-42) and low Fe ratios (5-wt.% Fe-13), the microstructure is composed of phases u-Al2Cu, vAl7Cu2Fe and eutectic (see the X-ray diffraction analysis shown in Fig. 3b). The primary phase v-Al7Cu2Fe precipitates with a needle-like shape for 8-wt.% Fe13 (Fig. 2e) and with an equi-axis shape for 5-wt.% Fe-8 (Fig. 2f). This last composition (42Cu–5Fe) is particularly attractive for its fine and homogeneous microstructure. For high Cu (20-wt.% Cu-30) and high Fe ratios (10-wt.% Fe-20), the microstructure is composed of primary phases v-Al7Cu2Fe (needles) and eutectic (Fig. 2c,d). For these concentrations, the X-ray diffraction analysis also indicates the presence of the phase u-Al2Cu. This observation does not fit with
L. Dubourg et al. / Surface and Coatings Technology 151 – 152 (2002) 329–332
331
Fig. 2. Micrographs of the phases at various Cu and Fe contents.
the equilibrium diagram (Fig. 2), as the phase aAl82Cu4Fe14 is not present while the phase u-Al2Cu is present. These differences between the analysed alloys and the equilibrium diagram are probably due to the
rapid quenching which characterises the laser processes. As a result, the observed phenomena are the widening or the shift of phase zones on the equilibrium diagram w7x, the precipitation of metastable phases (Al6Fe in our
Fig. 3. X-Ray diffraction spectra (wt.%): (a) 37Cu alloy; (b) 35Cu – 13Fe alloy; (c) average hardness of the tracks vs. Cu and Fe ratios (isohardness contours are calculated with a polynomial regression of degree 3).
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L. Dubourg et al. / Surface and Coatings Technology 151 – 152 (2002) 329–332
case) or the absence of equilibrium phases (aAl82Cu4Fe14 in our case). The micro-hardness average of each coating is shown in Fig. 3c with iso-hardness contours (fitting with a polynomial regression of degree 3). If only Cu is added, alloy hardness increases in a linear way with the Cu ratio w5x. For 40 wt.% Cu, the micro-hardness average is 250 Hv0.2. If only Fe is added, micro-hardness improvement is higher than Cu addition. For 15 wt.% Fe, the micro-hardness average is equal to 63 Hv0.2 and it is approximately 380 Hv0.2 at 28 wt.% Fe. An exponential increase is noticed between the Fe ratio and the alloy micro-hardness. In case of a combined addition of Cu and Fe, the hardness increases quickly in comparison with the non-treated sample (hardness of 20 Hv0.2). Maximum hardness without cracks has been observed for 38Cu–11Fe (wt.%): 550 Hv0.2 and for 28Cu–14Fe (wt.%): 515 Hv0.2. The structural refinement, the presence of intermetallic compounds and the structural hardening of the Al matrix may explain the hardness improvement. The structural refinement materialises in a grain size reduction due to the fast solidification w10x. The hardness of the intermetallic compounds depends on their composition: 400–600 Hv0.2 for uAl2Cu w11x, 600–750 Hv0.2 for l-Al3Fe w12x and 600– 700 Hv0.2 for v-Al7Cu2Fe. The hardness of the vAl7Cu2Fe phase is deduced from its position on Fig. 3c. The structural hardening of the Al matrix is due to the supersaturation of (Al) primary phases and especially the ageing reactions in the case of copper w13x. During a rapid quenching, the maximum solubility in aluminium is 5.5 wt.% for copper w14x and 9 wt.% for iron w15x. As a result, the hardness improvement thanks to Fe addition is higher than with Cu addition due to two reasons: the hardness of the intermetallic phase l-Al3Fe is higher (600–750 Hv0.2) than that of the phase uAl2Cu (400–600 Hv0.2) and the solubility limit of iron in aluminium (9 wt.%) is superior to the one of copper (5.5 wt.%).
4. Conclusions Laser surface alloys are produced with ratios of 0 up to 45 wt.% Cu and of 0 up to 30 wt.% Fe. All tracks are free of porosity; cracks and non-melted particles do not appear up to 40 wt.% Cu or 15 wt.% Fe. Depending on the Cu and Fe ratios, the track is composed of (Al), u-Al2Cu, v-Al7Cu2Fe or l-Al3Fe phases and the hardness evolves from 20 to 550 Hv0.2. For a cladding application, results indicate the optimum composition to be: Al 42Cu–5Fe (wt.%). Its microstructure is fine, homogeneous, hard (370 Hv0.2) and composed of uAl2Cu, v-Al7Cu2Fe and eutectic phases with compact shapes. Needle-like phases are not selected in order to avoid the stress concentrations leading to the appearance of cracks during wear tests. References w1x W.J. Tomlinson, A.S. Bransden, J. Mater. Sci. Lett. 13 (1994) 1086–1088. w2x L. Dubourg, F. Hlawka, A. Cornet, J. Phys. IV 10 (2000) 137– 144. w3x Y.Y. Qui, A. Almeida, R. Vilar, J. Mater. Sci. 33 (1998) 2639– 2651. w4x Y.Y. Qui, A. Almeida, R. Vilar, Scr. Metal. Mater. 33 (6) (1995) 863–870. w5x L. Dubourg, H. Pelletier, D. Vaissiere, F. Hlawka, A. Cornet, Wear (submitted). w6x H.J. Hegge, J. Th, M. De Hosson, J. Mater. Sci. 26 (1991) 711–714. w7x Y. Liu, J. Mazumder, K. Shibata, Metal. Mater. Trans. A 26A (1995) 1519–1533. w8x F. Faudot, Ann. Chim. Fr. 18 (1993) 445–456. w9x J.J. Richmond, S.E. Lebeau, K.P. Cooper, Rapid Solidification Processing: Principles and Technologies III, NBS, 1982. w10x J. Th, M. De Hosson, J. Noordhuis, J. Phys. IV 3 (1993) 927– 932. w11x D.R. Lide, Handbook of Aluminium and Aluminium Alloys, CRC Press, 1996. w12x L.F. Mondolfo, Aluminium Alloys, Butterworth and Co ltd (1976). w13x E. Cerri, E. Evangelista, N. Ryum, Metal. Mater. Trans. A 28A (1997) 257–263. w14x J.L. Murray, Phase Diagrams of Binary Copper Alloys, ASM International, Metals Park, OH, 1986, pp. 103–108. w15x A. Tonejc, A. Bonefacic, J. Appl. Phys. 40 (1969) 419–420.
Study of aluminium–copper–iron alloys: application for laser cladding L. Dubourga,b,*, F. Hlawkab, A. Cornetb a
ˆ API, 67400 Illkirch, France IREPA LASER, Parc d’Innovation, Pole b LISS, ENSAIS, 24 bd de la Victoire, 67000 Strasbourg, France
Abstract The modifications brought to aluminium by addition of copper and iron are investigated after rapid surface solidification by a CO2 laser. In order to carry out an exhaustive study, the laser surface alloys are obtained using a precoating process. Samples are characterised using microscopic examination, X-ray diffraction and hardness measurement. The optimum composition (42 wt.% Cu – 5 wt.% Fe) is selected for its fine and homogeneous microstructure (Al2 Cu and Al7 Cu2 Fe phases) and its high hardness (370 Hv0.2). 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Aluminium; Laser alloying; Laser cladding; Intermetallic compounds
1. Introduction Aluminium alloys offer advantages in automotive, aircraft and spacecraft industries, notably due to their lightness and corrosion resistance. However, these materials possess poor mechanical and tribological characteristics. Laser surface treatments can correct these drawbacks without affecting the global properties of the aluminium piece. For this reason, several authors have investigated the formation of intermetallic compounds with aluminium by laser (Fe w1x, Cu w2x, Mo w3x and Cr w4x). Previous studies indicated the great utility of copper in aluminium laser alloying w2x. The mechanical characteristics and the wear behaviour of such laser Al –Cu alloys make them ideal for laser alloying or cladding applications w5x. The present study shows the influence of Cu and Fe content in aluminium upon the microstructure and hardness. In order to do an exhaustive study, the laser surface alloys are obtained using a precoating process. A powder layer with a fixed Cu–Fe composition is placed on the aluminium specimen. A CO2 laser beam melts this layer together with the aluminium surface. Convection flows, generated by the temperature gradient between the beam centre and the sides of the melted pool, ensure the precoating incorporation and the homogeneity of this surface alloy w6x. * Corresponding author. Tel.: q33-3-88-65-54-00; fax: q33-3-8865-54-01. E-mail address: [email protected] (L. Dubourg).
The tracks are analysed using microscopic examination, X-ray diffraction and micro-hardness measurements. Subsequently, the optimum composition in relation to the microstructure, homogeneity and hardness is selected. 2. Experimental set-up Surface treatments are carried out with a continuous CO2 laser of 10.6-mm wavelength and energy spatial distribution of TEM20 type. The process parameters are as follows: laser power of 4000 W on the sample, scanning speed of 4.2=10y3 m sy1, beam diameter of 4 mm on the substrate and sample preheating of 150 8C. The interaction zone is protected from air oxidation with an inert gas (1 bar pressure, Argon flow rate of 8.3=10y2 l sy1 and helium flow rate of 16.7=10y2 l sy1). The precoating is composed of metallic powders and a binder. The powders have a granulometry of 45– 150 mm and their compositions are as follows: Fe (99.5 wt.%), Cu (90 wt.% Cu, 10 wt.% Al). The binder is composed of 94 wt.% of water (evaporated during the precoating drying at 80 8C for 30 min) and 6 wt.% of an organic compound (vaporised during laser scanning). The microstructure analysis is performed by optical microscopy after a Keller’s reagent etching and by Xray diffraction using a Siemens D5000 diffractometer (CuKa radiation). Vickers hardness measurements are performed in the cross-section with a load of 2 N.
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 1 . 0 1 5 9 1 - 2
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Fig. 1. Micrographs of coatings composed of (wt.%): (a) 15Fe; (b) 42Cu – 5Fe.
Average hardness is calculated using eight micro-hardness measurements (see Fig. 1a). 3. Results and discussion Different coatings are achieved with ratios of 0 up to 45 wt.% Cu and of 0 up to 30 wt.% Fe. For these concentrations, each coating is free of porosity (Fig. 1). Between 0 and 40 wt.% Cu, the Al–Cu alloys are homogeneous on the entire coating. Above 40 wt.% Cu, cracks appear in the coatings due to the low ductility of the alloy w7x. Al–Fe alloys are homogeneous up to approximately 15 wt.% Fe. Above 15 wt.% Fe, ironrich zones and small cracks appear in the coating bottom (cf. Fig. 1a). Above 30 wt.% Fe, alloy homogeneity becomes too poor to ensure an isotropic mechanical behaviour of the track. Cu addition improves the homogeneity of the Al–Fe alloys (Fig. 1b). Nevertheless, when the Fe ratio is higher than the Cu ratio (approx. wt.% Fe)1.5=wt.% Cu), the homogeneity decreases. As a result, the appearance of some non-melted particles in the coating bottom is observed. Fig. 2 shows the obtained alloys and their positions on the Al–Cu–Fe equilibrium diagram at various Cu and Fe ratios (aluminium-rich zone of the Al–Cu–Fe diagram w8x). Grey areas of the diagram illustrate the zones where the microstructure of the obtained alloys is similar. If only Cu is added, the phases identified by microscopic examinations and X-ray diffraction (Fig. 3a) fit with the equilibrium diagram. In this case, rapid quenching due to the laser process creates equilibrium phases. Up to 24 wt.% Cu, a hypoeutectic microstructure is observed: white primary grains (Al)qa dark eutectic (Fig. 2i). In the range of 27 and 37 wt.% Cu, the eutectic microstructure is always present, as shown in Fig. 2h. Theoretically, according to the equilibrium Al– Cu diagram, this structure exists for a specific Cu percentage, approximately 32 wt.%. The growth of primary phases w(Al) and u-Al2Cux is limited because of the rapid quenching w7x. Finally, in the range of 37–
40 wt.% Cu, micrographs (Fig. 2g) show a hypereutectic microstructure: angular primary grains u-Al2Cuqa dark eutectic. If only Fe is added, the alloys have the same microstructure (Fig. 2l) confirmed by X-ray diffraction analysis. This microstructure is composed of primary needle-like phases of l-Al3Fe, white dendrites (Al) and a dark eutectic. The microstructure does not fit with the equilibrium diagram, as primary phases of (Al) and lAl3Fe can be observed at the same time. For 24 wt.% Fe, the eutectic phase disappears except at grain boundaries, as shown in Fig. 2a. Moreover, X-ray diffraction analysis also indicates the presence of a metastable phase Al6Fe at 24 wt.% Fe. According to Richmond et al. w9x, this metastable phase is present in the eutectic. In the case of Cu–Fe combined addition, four different microstructures can be noted according to the composition. For low Cu and Fe ratios (0-wt.% Cu-15 and 0-wt.% Fe-6), the microstructure is composed of primary dendrites (Al) and eutectic (Fig. 2i). For low Cu (0-wt.% Cu-15) and high Fe ratios (8-wt.% Fe-22), the microstructure is composed of phases lAl3Fe, dendrites (Al) and eutectic (Fig. 2b). The lAl3Fe phase precipitates with a needle-like shape as in Al–Fe alloy (Fig. 2l) in spite of the presence of copper. The X-ray diffraction analysis does not show the presence of the phase a-Al82Cu4Fe14, according to the equilibrium diagram shown in Fig. 2. For high Cu (35wt.% Cu-42) and low Fe ratios (5-wt.% Fe-13), the microstructure is composed of phases u-Al2Cu, vAl7Cu2Fe and eutectic (see the X-ray diffraction analysis shown in Fig. 3b). The primary phase v-Al7Cu2Fe precipitates with a needle-like shape for 8-wt.% Fe13 (Fig. 2e) and with an equi-axis shape for 5-wt.% Fe-8 (Fig. 2f). This last composition (42Cu–5Fe) is particularly attractive for its fine and homogeneous microstructure. For high Cu (20-wt.% Cu-30) and high Fe ratios (10-wt.% Fe-20), the microstructure is composed of primary phases v-Al7Cu2Fe (needles) and eutectic (Fig. 2c,d). For these concentrations, the X-ray diffraction analysis also indicates the presence of the phase u-Al2Cu. This observation does not fit with
L. Dubourg et al. / Surface and Coatings Technology 151 – 152 (2002) 329–332
331
Fig. 2. Micrographs of the phases at various Cu and Fe contents.
the equilibrium diagram (Fig. 2), as the phase aAl82Cu4Fe14 is not present while the phase u-Al2Cu is present. These differences between the analysed alloys and the equilibrium diagram are probably due to the
rapid quenching which characterises the laser processes. As a result, the observed phenomena are the widening or the shift of phase zones on the equilibrium diagram w7x, the precipitation of metastable phases (Al6Fe in our
Fig. 3. X-Ray diffraction spectra (wt.%): (a) 37Cu alloy; (b) 35Cu – 13Fe alloy; (c) average hardness of the tracks vs. Cu and Fe ratios (isohardness contours are calculated with a polynomial regression of degree 3).
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L. Dubourg et al. / Surface and Coatings Technology 151 – 152 (2002) 329–332
case) or the absence of equilibrium phases (aAl82Cu4Fe14 in our case). The micro-hardness average of each coating is shown in Fig. 3c with iso-hardness contours (fitting with a polynomial regression of degree 3). If only Cu is added, alloy hardness increases in a linear way with the Cu ratio w5x. For 40 wt.% Cu, the micro-hardness average is 250 Hv0.2. If only Fe is added, micro-hardness improvement is higher than Cu addition. For 15 wt.% Fe, the micro-hardness average is equal to 63 Hv0.2 and it is approximately 380 Hv0.2 at 28 wt.% Fe. An exponential increase is noticed between the Fe ratio and the alloy micro-hardness. In case of a combined addition of Cu and Fe, the hardness increases quickly in comparison with the non-treated sample (hardness of 20 Hv0.2). Maximum hardness without cracks has been observed for 38Cu–11Fe (wt.%): 550 Hv0.2 and for 28Cu–14Fe (wt.%): 515 Hv0.2. The structural refinement, the presence of intermetallic compounds and the structural hardening of the Al matrix may explain the hardness improvement. The structural refinement materialises in a grain size reduction due to the fast solidification w10x. The hardness of the intermetallic compounds depends on their composition: 400–600 Hv0.2 for uAl2Cu w11x, 600–750 Hv0.2 for l-Al3Fe w12x and 600– 700 Hv0.2 for v-Al7Cu2Fe. The hardness of the vAl7Cu2Fe phase is deduced from its position on Fig. 3c. The structural hardening of the Al matrix is due to the supersaturation of (Al) primary phases and especially the ageing reactions in the case of copper w13x. During a rapid quenching, the maximum solubility in aluminium is 5.5 wt.% for copper w14x and 9 wt.% for iron w15x. As a result, the hardness improvement thanks to Fe addition is higher than with Cu addition due to two reasons: the hardness of the intermetallic phase l-Al3Fe is higher (600–750 Hv0.2) than that of the phase uAl2Cu (400–600 Hv0.2) and the solubility limit of iron in aluminium (9 wt.%) is superior to the one of copper (5.5 wt.%).
4. Conclusions Laser surface alloys are produced with ratios of 0 up to 45 wt.% Cu and of 0 up to 30 wt.% Fe. All tracks are free of porosity; cracks and non-melted particles do not appear up to 40 wt.% Cu or 15 wt.% Fe. Depending on the Cu and Fe ratios, the track is composed of (Al), u-Al2Cu, v-Al7Cu2Fe or l-Al3Fe phases and the hardness evolves from 20 to 550 Hv0.2. For a cladding application, results indicate the optimum composition to be: Al 42Cu–5Fe (wt.%). Its microstructure is fine, homogeneous, hard (370 Hv0.2) and composed of uAl2Cu, v-Al7Cu2Fe and eutectic phases with compact shapes. Needle-like phases are not selected in order to avoid the stress concentrations leading to the appearance of cracks during wear tests. References w1x W.J. Tomlinson, A.S. Bransden, J. Mater. Sci. Lett. 13 (1994) 1086–1088. w2x L. Dubourg, F. Hlawka, A. Cornet, J. Phys. IV 10 (2000) 137– 144. w3x Y.Y. Qui, A. Almeida, R. Vilar, J. Mater. Sci. 33 (1998) 2639– 2651. w4x Y.Y. Qui, A. Almeida, R. Vilar, Scr. Metal. Mater. 33 (6) (1995) 863–870. w5x L. Dubourg, H. Pelletier, D. Vaissiere, F. Hlawka, A. Cornet, Wear (submitted). w6x H.J. Hegge, J. Th, M. De Hosson, J. Mater. Sci. 26 (1991) 711–714. w7x Y. Liu, J. Mazumder, K. Shibata, Metal. Mater. Trans. A 26A (1995) 1519–1533. w8x F. Faudot, Ann. Chim. Fr. 18 (1993) 445–456. w9x J.J. Richmond, S.E. Lebeau, K.P. Cooper, Rapid Solidification Processing: Principles and Technologies III, NBS, 1982. w10x J. Th, M. De Hosson, J. Noordhuis, J. Phys. IV 3 (1993) 927– 932. w11x D.R. Lide, Handbook of Aluminium and Aluminium Alloys, CRC Press, 1996. w12x L.F. Mondolfo, Aluminium Alloys, Butterworth and Co ltd (1976). w13x E. Cerri, E. Evangelista, N. Ryum, Metal. Mater. Trans. A 28A (1997) 257–263. w14x J.L. Murray, Phase Diagrams of Binary Copper Alloys, ASM International, Metals Park, OH, 1986, pp. 103–108. w15x A. Tonejc, A. Bonefacic, J. Appl. Phys. 40 (1969) 419–420.