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Solidification microstructures induced by laser surface alloying: influence of the substrate
Materials Science and Engineering, A 134 ( 1991 ) 1283-1287
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Solidification microstructures induced by laser surface alloying: influence of the substrate J. M. Pelletier, L. Renaud and F. Fouquet GEMPPM-('ALFETMA 7-, Bat. 502, INSA, 69621 Villeur))anne Cedex (t=rance)
Abstract Surface-treated alloys have been realized on both aluminium-based and iron-based substrates, by laser melting of a nickel predeposited layer. Microstructures have been characterized by optical and scanning electron microscopies and microanalyses. Phases observed are in agreement with equilibrium phase diagrams. The microstructures are mainly dendritic and various phases have been identified. A major difference exists between the two kinds of alloys: a plane front growth occurred at the beginning of the solidification process (interface between the substrate and the melted zone) in iron-based surface alloys, but not in aluminium-based ones. This difference is explained using the constitutional supercooling criterion.
1. Introduction
2. Experimental procedure
Solidification microstructure depends on many parameters, especially on solidification rate and material composition [1]. During recent years considerable interest has occurred concerning rapid solidification phenomena. Laser surface melting offers an attractive means of achieving rapid solidification [2] because high cooling rates can be obtained due to the cooling of the melt oool by the substrate (self-quenching). From a netallurgical point of view, various microstruc,ures may be obtained, if the incident power and the interaction time are varied, thus leading to various melted depths, temperature gradients and quenching rates, and, consequently, to various solidification rates. In addition, laser melting of both a predeposited layer of a specific material and part of the substrate induces a particular composition in the treated layer after solidification [3]. This phenomenon, called laser surface alloying (LSA), is studied in the present work. The objective is to correlate the microstructure of the various laser surface alloys with the solidification conditions. Two substrates are considered: aluminium-based and iron-based ones. Nickel is the main alloying element, introduced either by electrolytic, electroless or plasma spraying deposition. These systems have been chosen owing to the large variety of possible microstructures.
2.1. Materials and specimens For aluminium substrates, aluminium cast alloys have the following composition: Al5wt.%Si-3wt.%Cu. Specimen ( 8 0 . 2 0 . 5 m m ~) were machined from untreated sheets and coated with nickel, either by plasma or by electrolytic deposition. Various thicknesses of the layers a r e used: 25 ~m, 50 ~tm or 100/~m. Concerning iron-based substrates, mild steels samples (0.1 wt.% C) of parallelepipedic shape are used ( 5 0 . 2 0 . 5 mm~). They were plated with a 75 ktm thick electroless nickel-phosphorus layer. The phosphorus content is about 8 wt.%.
0921-5093/91/$3.50
2.2. Laser processing Specimens were mounted on a numerically controlled X-Y table and irradiated with a CO~ continuous laser, the nominal power P of which is up to 3.5 kW. The X-Y table allowed scanning rates Vr of the sample under the laser beam from 1 mm s i to 300 mm s J. A variation of the interaction time, defined by T= d/lZr, where d is the equivalent beam diameter on the sample [4], was thus possible. During scanning of the sample under the laser beam, argon was blown through a gas nozzle to prevent any oxidation. This protection must be carefully controlled to achieve reproducible treatments. Multiple laser tracks were performed and the shift d between two © Elsevier Sequoia/Printed in The Netherlands
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successive tracks was kept constant. No absorbing coating was used to avoid any contamination of the sample during melting. SURFACE ALLOY
2.3. Characterization of the microstructure For optical microscopy, the samples were polished and etched in NaOH (10%) for aluminium-based alloys and in nital (3%) for ironbased ones. For scanning electron microscopy, images were carried out using backscattered electrons and thus etching was not necessary. Various microanalysis experiments were also carried out by SIMS, electron microprobe or X-ray imaging.
INTERMEDIATE ZONE UNAFFECTED SUBSTRATE
SURFACE ALLOY HEAT-AFFECTED ZONE
3. Observation and characterization of surface alloys
After optimization of the laser processing conditions, a cross-section of the sample reveals the existence of three different zones (Fig. 1 ), i.e. from the core to the surface of the sample. 3.1. The unaffected substrate The hypoeutectic A1-Si-Cu substrate is made up from large aluminium grains and dispersed silicon and A12Cu particles (Fig. l(a)). The low carbon steel microstructure is classical and is composed of ferrite grains and pearlite (Fig. l(b)). 3.2. The intermediate zone Partial melting of the A1-Si matrix, which has a low melting temperature (650 °C), followed by rapid solidification, induces a large grain refinement. No solid state transformation may occur in this alloy. In iron-based substrate, microstructure modifications have occurred during laser treatment, through solid state phase transformations. Near the melt zone [5], the pearlite dissolution is complete and grain refinement has been induced during cooling with the formation of an acicular microstructure. In the vicinity of the unaffected substrate, this pearlite dissolution is only partial. 3.3. Laser surface alloy 3.3.1. Aluminium-based alloys A dendritic microstructure is shown in Fig. 2. Owing to the cross-section effect of a threedimensional growth network, the dendrites appear to be randomly oriented. From the value of the interdendritic spacing (about a few microns), the quenching rate may be estimated to be about 104 K s -1 during solidification [8].
UNAFFECTED SUBSTRATE
Fig. 1. Cross-sections of laser surface alloys; optical micrographs: (a) AI-5%Si-3%Cu, electrolytically coated with a 50 ktm thick nickel layer and laser treated; P = 2.4 kW; V~= 10 m m s 1; d = 1 mm; 6 = 0 . 3 ram; 10% NaOH etching. (b) Fe-0.1wt.%C, coated with a 75 ktm thick electroless nickel layer and laser treated; P = 1.1 kW; V~=20 mm s ~, d = 0 . 6 ram; 6 = 0.15 ram; 3% nital etching.
Fig. 2. Backscattered electron micrograph of the laser aluminium-based surface alloy, obtained with the same operating conditions as in Fig. 1 (without etching).
Chemical homogenization has been mainly achieved in the melt pool by convection movements. The solidification microstructure has been studied using various characterization methods [6, 7], and it appears that no metastable phases were observed, only classical phenomena occurred. Dendrites have been identified as composite intermetallic compounds: the inner part (bright zone in Fig. 2) is A13Ni2, while the
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outer part (grey zone) is AI~Ni. The interdendritic dark zone is essentially composed of aluminium, with very small nickel concentration (0.3%). No extension of the solid solubility was achieved during quenching from the liquid state [7]. The interface between substrate and surface alloy is shown in Fig. 3. There is no transition zone: the microstructure is wholly dendritic in the laser surface alloy.
homogeneous zone, as well as molten depth zv and interdendritic spacing, depend on laser processing conditions, especially on scanning velocity, i.e. on interaction time (Fig. 5). The higher the scanning velocity, the smaller the A value. The composition analyses [9] together with the Fe-Ni, Fe-P and Ni-P phase diagrams [ 10], make
3.3.2. Iron-based allovs A fine and dendritic microstructure was also observed in the greater part of this surface alloy. The segregating element is phosphorus. Figure 4 clearly shows that strong local composition heterogeneities exist in the surface alloy, mainly between dendrites and interdendritic areas. However, near the interface with the unmelted substrate a bright strip, corresponding to a plane front growth, is observed. The thickness zXof this
/a)
[]
Fig. 3. Backscattered electron micrograph of the interfacial region between laser surface alloy and substrate in aluminium-based alloy (same conditions as in Fig. 1 ).
.....
(b)
(ci Fig. 4. Backscattered electron micrograph of the laser ironbased surface alloy, obtained with the same operating conditions as in Fig. 1 (without etching).
Fig. 5. Backscattered electron micrographs of the interfacial region between laser surface alloy and substrate: iron-based alloy, coated with electroless nickel, laser treated: 1' = 1420 W : d - 0 . 6 m m ; 0 = 0 . 1 5 m m : ( a ) V j = 2 0 m m s ~;(b) V~=40 mms ';(c) V ~ = 8 0 m m s ~.
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it possible to interpret the solidification process as follows: solidification starts with a plane front growth of an austenitic Fe-Ni solid solution, simultaneously rejecting phosphorus in the melt. This segregation is due to the fact that there is no solubility of phosphorus in nickel or in iron, but a full solid solubility between nickel and iron. As the solidification rate increases, the plane front becomes unstable and the solid growth goes ahead with austenitic Fe-Ni solid solution dendrites. Phosphorus is still rejected in the interdendritic areas. At the end of the solidification sequence, the interdendritic areas, enriched with phosphorus, are composed of a eutectic mixture (Fe, Ni)-(Fe, Ni)3P [11]. The laser processing conditions have an influence on geometric features--the shape and thickness of the molten zone, the interdendritic spacing, and the plane front solidification zone thickness--but the solidification sequence and the phases formed remain the same for all treatments.
4. Discussion
Let us note that, from the experimental observations, a major difference appears between the two kinds of alloys: a plane front growth occurs at the beginning of the solidification process, at the interface between the substrate and melted zone, in iron-based alloys (Fig. 5), but not in aluminium-based alloys (Fig. 3). This difference can be explained using the constitutional supercooling theory [12], which introduces two pertinent parameters: the temperature gradient within the liquid at the liquid-solid interface G and the solidification rate V~. This theory states that if G/V~is higher than a critical value (G/V~)*, a plane front solidification takes place, and if G/V~ is lower than this critical value, a cellular or dendritic solidification occurs. Since after laser melting the solidification rate is initially zero at the interface between substrate and melt pool, the G/~ ratio is infinite and thus the solidification should occur, in all cases, with a plane front. After that, the solidification rate rapidly increases and the temperature gradient, which is maximum at the beginning, decreases, leading to a subcritical value of G/~ and hence to a dendritic solidification (Fig. 6(a)). In such conditions a transition between a plane front and dendritic growth appears. This sequence should always be observed in laser surface alloying; the absence of
G~
G~
Zc
ZF
Fig. 6. Schematic representation of the evolution of the temperature gradient at the solid-liquid interface G and the solidification rate V, w'. depth in the surface alloy (Z); Zt=melted depth; Zc=depth for which G / ~ becomes critical. (a) Substrate with a low thermal conductivity; (b) substrate with a high thermal conductivity.
this phenomenon in the present aluminium-based alloy is due to the fact that aluminium-based substrate has a high thermal conductivity. Therefore G is small and V~rapidly grows, leading to a very short supercritical stage (Fig. 6(b)). Thus, A is negligible and solidification appears fully dendritic, since a plane front growth cannot significantly occur. Furthermore, in the case of aluminium cast alloy, two phases coexist in the substrate, with a large difference of melting points; consequently a partial melting is achieved, which reduces the homogeneous nucleation ability from the melt. We can now explain the previous observations, concerning the decrease of A induced by an increase of the scanning velocity for iron-based alloys (Fig. 5). If we consider that an increase of scanning velocity produces an increase in the average solidification rate, and also in the average temperature gradient in the melt, it is not possible to explain the reduction of A. The probable reason is that the critical value (G/~)* is proportional to the concentration of the segregating element [12]. When the scanning velocity is higher, the melted depth is smaller and hence the concentration (percent of phosphorus in the present case) is larger, leading to an increase of (G/Vss)*.Therefore it is then more difficult to have (G/~)>(G/V~)*; in these conditions this ratio becomes more rapidly subcritical and a diminution of A is obtained. Experimental observations show clearly that (G/~)* depends on the phosphorus concentration, which may be estimated using a simple initial layer thickness-melted depth ratio, to be about 1.5% in Fig. 5(a), and about 8% in Fig. 5(c); the concentration ratio is in agreement with the observed ratio of the white strip thicknesses.
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5. Conclusions
Laser surface alloying, performed on aluminlure-based and iron-based substrates with predeposited nickel layers, induces a dendritic microstructure in the chosen laser processing conditions. The phases observed are in agreement with equilibrium phase diagrams, in spite of the fairly high cooling rates--compared with those obtained by conventional casting methods--and are as follows: (1) In aluminium-based surface alloys, dendrites are constituted by A13Ni and AI~Ni 2 phases, while interdendritic areas are aluminiumrich. (2) In iron-based surface alloys, a Fe-Ni f.c.c. solid solution is found in dendrites and a eutectic mixture (Fe, Ni)-(Fe, Ni)~P is observed between these dendrites. However a main difference is found between the two cases: solidification starts with a plane front growth in iron-based surface alloys, owing to their lower thermal conductivity. In aluminium-based surface alloy, the plane front growth is not detectable and solidification occurs directly i n the dendritic regime. Furthermore, for iron-
based alloys, the thickness of the plane front growth zone is directly connected to the concentration of the segregating element in the melt. References I W. Kurz and D. J. Fisher, Fundamentals of A'olidification, Trans. Tech. Publications, Aedermansdorf. Swilzerland, 1986. 2 M, Lamb, W. M. Steen and l). R. F. West, 1(;4l, EO, 1/ol. 1, Laser Institute of America, 1984, p. 133. 3 A. Galerie, M. Pons and M, Caillet, Mater. Sei. Eng., ,%' (1987) 127. 4 J. Dietz and J. Merlin, Rev. Phys. Appl., 33 (1988) 1787. 5 M.F. Ashby and K. E. Easterling, Aeta Metall., 32 (1984) 1935. 6 S. Bonnet-Jobez, Thesis, INSA, Lyon, France, 1989. 7 E. Gaffet, J. M. Pelletier and S. Bonnel-Jobez. Acta Metall., 37 (1989) 32(/5. 8 H. Matija, B. C. Giessen and N. J. Grant, J. Inst. Metals, 96 (1968) 30. 9 L. Renaud, F. Fouquet, A. Elhamdaoui, J. P. Millet, H. Mazille and J. L. Crolet, Acta Memll. Mater. 38, 8 (1990) 1547. 10 T. B. Massalski, Binary All~:s Phase l)iagrarns, Vol. 1, ASM, Metals Park, Ohio, 1986. 11 L. Renaud, F. Fouquet and C. Esnouf, to be published. 12 W. A. Tiller, J. W. Rutter, K. A. Jackson and B. Chalmers, Acta Metall., I (1953) 428.
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Solidification microstructures induced by laser surface alloying: influence of the substrate J. M. Pelletier, L. Renaud and F. Fouquet GEMPPM-('ALFETMA 7-, Bat. 502, INSA, 69621 Villeur))anne Cedex (t=rance)
Abstract Surface-treated alloys have been realized on both aluminium-based and iron-based substrates, by laser melting of a nickel predeposited layer. Microstructures have been characterized by optical and scanning electron microscopies and microanalyses. Phases observed are in agreement with equilibrium phase diagrams. The microstructures are mainly dendritic and various phases have been identified. A major difference exists between the two kinds of alloys: a plane front growth occurred at the beginning of the solidification process (interface between the substrate and the melted zone) in iron-based surface alloys, but not in aluminium-based ones. This difference is explained using the constitutional supercooling criterion.
1. Introduction
2. Experimental procedure
Solidification microstructure depends on many parameters, especially on solidification rate and material composition [1]. During recent years considerable interest has occurred concerning rapid solidification phenomena. Laser surface melting offers an attractive means of achieving rapid solidification [2] because high cooling rates can be obtained due to the cooling of the melt oool by the substrate (self-quenching). From a netallurgical point of view, various microstruc,ures may be obtained, if the incident power and the interaction time are varied, thus leading to various melted depths, temperature gradients and quenching rates, and, consequently, to various solidification rates. In addition, laser melting of both a predeposited layer of a specific material and part of the substrate induces a particular composition in the treated layer after solidification [3]. This phenomenon, called laser surface alloying (LSA), is studied in the present work. The objective is to correlate the microstructure of the various laser surface alloys with the solidification conditions. Two substrates are considered: aluminium-based and iron-based ones. Nickel is the main alloying element, introduced either by electrolytic, electroless or plasma spraying deposition. These systems have been chosen owing to the large variety of possible microstructures.
2.1. Materials and specimens For aluminium substrates, aluminium cast alloys have the following composition: Al5wt.%Si-3wt.%Cu. Specimen ( 8 0 . 2 0 . 5 m m ~) were machined from untreated sheets and coated with nickel, either by plasma or by electrolytic deposition. Various thicknesses of the layers a r e used: 25 ~m, 50 ~tm or 100/~m. Concerning iron-based substrates, mild steels samples (0.1 wt.% C) of parallelepipedic shape are used ( 5 0 . 2 0 . 5 mm~). They were plated with a 75 ktm thick electroless nickel-phosphorus layer. The phosphorus content is about 8 wt.%.
0921-5093/91/$3.50
2.2. Laser processing Specimens were mounted on a numerically controlled X-Y table and irradiated with a CO~ continuous laser, the nominal power P of which is up to 3.5 kW. The X-Y table allowed scanning rates Vr of the sample under the laser beam from 1 mm s i to 300 mm s J. A variation of the interaction time, defined by T= d/lZr, where d is the equivalent beam diameter on the sample [4], was thus possible. During scanning of the sample under the laser beam, argon was blown through a gas nozzle to prevent any oxidation. This protection must be carefully controlled to achieve reproducible treatments. Multiple laser tracks were performed and the shift d between two © Elsevier Sequoia/Printed in The Netherlands
1284
successive tracks was kept constant. No absorbing coating was used to avoid any contamination of the sample during melting. SURFACE ALLOY
2.3. Characterization of the microstructure For optical microscopy, the samples were polished and etched in NaOH (10%) for aluminium-based alloys and in nital (3%) for ironbased ones. For scanning electron microscopy, images were carried out using backscattered electrons and thus etching was not necessary. Various microanalysis experiments were also carried out by SIMS, electron microprobe or X-ray imaging.
INTERMEDIATE ZONE UNAFFECTED SUBSTRATE
SURFACE ALLOY HEAT-AFFECTED ZONE
3. Observation and characterization of surface alloys
After optimization of the laser processing conditions, a cross-section of the sample reveals the existence of three different zones (Fig. 1 ), i.e. from the core to the surface of the sample. 3.1. The unaffected substrate The hypoeutectic A1-Si-Cu substrate is made up from large aluminium grains and dispersed silicon and A12Cu particles (Fig. l(a)). The low carbon steel microstructure is classical and is composed of ferrite grains and pearlite (Fig. l(b)). 3.2. The intermediate zone Partial melting of the A1-Si matrix, which has a low melting temperature (650 °C), followed by rapid solidification, induces a large grain refinement. No solid state transformation may occur in this alloy. In iron-based substrate, microstructure modifications have occurred during laser treatment, through solid state phase transformations. Near the melt zone [5], the pearlite dissolution is complete and grain refinement has been induced during cooling with the formation of an acicular microstructure. In the vicinity of the unaffected substrate, this pearlite dissolution is only partial. 3.3. Laser surface alloy 3.3.1. Aluminium-based alloys A dendritic microstructure is shown in Fig. 2. Owing to the cross-section effect of a threedimensional growth network, the dendrites appear to be randomly oriented. From the value of the interdendritic spacing (about a few microns), the quenching rate may be estimated to be about 104 K s -1 during solidification [8].
UNAFFECTED SUBSTRATE
Fig. 1. Cross-sections of laser surface alloys; optical micrographs: (a) AI-5%Si-3%Cu, electrolytically coated with a 50 ktm thick nickel layer and laser treated; P = 2.4 kW; V~= 10 m m s 1; d = 1 mm; 6 = 0 . 3 ram; 10% NaOH etching. (b) Fe-0.1wt.%C, coated with a 75 ktm thick electroless nickel layer and laser treated; P = 1.1 kW; V~=20 mm s ~, d = 0 . 6 ram; 6 = 0.15 ram; 3% nital etching.
Fig. 2. Backscattered electron micrograph of the laser aluminium-based surface alloy, obtained with the same operating conditions as in Fig. 1 (without etching).
Chemical homogenization has been mainly achieved in the melt pool by convection movements. The solidification microstructure has been studied using various characterization methods [6, 7], and it appears that no metastable phases were observed, only classical phenomena occurred. Dendrites have been identified as composite intermetallic compounds: the inner part (bright zone in Fig. 2) is A13Ni2, while the
1285
outer part (grey zone) is AI~Ni. The interdendritic dark zone is essentially composed of aluminium, with very small nickel concentration (0.3%). No extension of the solid solubility was achieved during quenching from the liquid state [7]. The interface between substrate and surface alloy is shown in Fig. 3. There is no transition zone: the microstructure is wholly dendritic in the laser surface alloy.
homogeneous zone, as well as molten depth zv and interdendritic spacing, depend on laser processing conditions, especially on scanning velocity, i.e. on interaction time (Fig. 5). The higher the scanning velocity, the smaller the A value. The composition analyses [9] together with the Fe-Ni, Fe-P and Ni-P phase diagrams [ 10], make
3.3.2. Iron-based allovs A fine and dendritic microstructure was also observed in the greater part of this surface alloy. The segregating element is phosphorus. Figure 4 clearly shows that strong local composition heterogeneities exist in the surface alloy, mainly between dendrites and interdendritic areas. However, near the interface with the unmelted substrate a bright strip, corresponding to a plane front growth, is observed. The thickness zXof this
/a)
[]
Fig. 3. Backscattered electron micrograph of the interfacial region between laser surface alloy and substrate in aluminium-based alloy (same conditions as in Fig. 1 ).
.....
(b)
(ci Fig. 4. Backscattered electron micrograph of the laser ironbased surface alloy, obtained with the same operating conditions as in Fig. 1 (without etching).
Fig. 5. Backscattered electron micrographs of the interfacial region between laser surface alloy and substrate: iron-based alloy, coated with electroless nickel, laser treated: 1' = 1420 W : d - 0 . 6 m m ; 0 = 0 . 1 5 m m : ( a ) V j = 2 0 m m s ~;(b) V~=40 mms ';(c) V ~ = 8 0 m m s ~.
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it possible to interpret the solidification process as follows: solidification starts with a plane front growth of an austenitic Fe-Ni solid solution, simultaneously rejecting phosphorus in the melt. This segregation is due to the fact that there is no solubility of phosphorus in nickel or in iron, but a full solid solubility between nickel and iron. As the solidification rate increases, the plane front becomes unstable and the solid growth goes ahead with austenitic Fe-Ni solid solution dendrites. Phosphorus is still rejected in the interdendritic areas. At the end of the solidification sequence, the interdendritic areas, enriched with phosphorus, are composed of a eutectic mixture (Fe, Ni)-(Fe, Ni)3P [11]. The laser processing conditions have an influence on geometric features--the shape and thickness of the molten zone, the interdendritic spacing, and the plane front solidification zone thickness--but the solidification sequence and the phases formed remain the same for all treatments.
4. Discussion
Let us note that, from the experimental observations, a major difference appears between the two kinds of alloys: a plane front growth occurs at the beginning of the solidification process, at the interface between the substrate and melted zone, in iron-based alloys (Fig. 5), but not in aluminium-based alloys (Fig. 3). This difference can be explained using the constitutional supercooling theory [12], which introduces two pertinent parameters: the temperature gradient within the liquid at the liquid-solid interface G and the solidification rate V~. This theory states that if G/V~is higher than a critical value (G/V~)*, a plane front solidification takes place, and if G/V~ is lower than this critical value, a cellular or dendritic solidification occurs. Since after laser melting the solidification rate is initially zero at the interface between substrate and melt pool, the G/~ ratio is infinite and thus the solidification should occur, in all cases, with a plane front. After that, the solidification rate rapidly increases and the temperature gradient, which is maximum at the beginning, decreases, leading to a subcritical value of G/~ and hence to a dendritic solidification (Fig. 6(a)). In such conditions a transition between a plane front and dendritic growth appears. This sequence should always be observed in laser surface alloying; the absence of
G~
G~
Zc
ZF
Fig. 6. Schematic representation of the evolution of the temperature gradient at the solid-liquid interface G and the solidification rate V, w'. depth in the surface alloy (Z); Zt=melted depth; Zc=depth for which G / ~ becomes critical. (a) Substrate with a low thermal conductivity; (b) substrate with a high thermal conductivity.
this phenomenon in the present aluminium-based alloy is due to the fact that aluminium-based substrate has a high thermal conductivity. Therefore G is small and V~rapidly grows, leading to a very short supercritical stage (Fig. 6(b)). Thus, A is negligible and solidification appears fully dendritic, since a plane front growth cannot significantly occur. Furthermore, in the case of aluminium cast alloy, two phases coexist in the substrate, with a large difference of melting points; consequently a partial melting is achieved, which reduces the homogeneous nucleation ability from the melt. We can now explain the previous observations, concerning the decrease of A induced by an increase of the scanning velocity for iron-based alloys (Fig. 5). If we consider that an increase of scanning velocity produces an increase in the average solidification rate, and also in the average temperature gradient in the melt, it is not possible to explain the reduction of A. The probable reason is that the critical value (G/~)* is proportional to the concentration of the segregating element [12]. When the scanning velocity is higher, the melted depth is smaller and hence the concentration (percent of phosphorus in the present case) is larger, leading to an increase of (G/Vss)*.Therefore it is then more difficult to have (G/~)>(G/V~)*; in these conditions this ratio becomes more rapidly subcritical and a diminution of A is obtained. Experimental observations show clearly that (G/~)* depends on the phosphorus concentration, which may be estimated using a simple initial layer thickness-melted depth ratio, to be about 1.5% in Fig. 5(a), and about 8% in Fig. 5(c); the concentration ratio is in agreement with the observed ratio of the white strip thicknesses.
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5. Conclusions
Laser surface alloying, performed on aluminlure-based and iron-based substrates with predeposited nickel layers, induces a dendritic microstructure in the chosen laser processing conditions. The phases observed are in agreement with equilibrium phase diagrams, in spite of the fairly high cooling rates--compared with those obtained by conventional casting methods--and are as follows: (1) In aluminium-based surface alloys, dendrites are constituted by A13Ni and AI~Ni 2 phases, while interdendritic areas are aluminiumrich. (2) In iron-based surface alloys, a Fe-Ni f.c.c. solid solution is found in dendrites and a eutectic mixture (Fe, Ni)-(Fe, Ni)~P is observed between these dendrites. However a main difference is found between the two cases: solidification starts with a plane front growth in iron-based surface alloys, owing to their lower thermal conductivity. In aluminium-based surface alloy, the plane front growth is not detectable and solidification occurs directly i n the dendritic regime. Furthermore, for iron-
based alloys, the thickness of the plane front growth zone is directly connected to the concentration of the segregating element in the melt. References I W. Kurz and D. J. Fisher, Fundamentals of A'olidification, Trans. Tech. Publications, Aedermansdorf. Swilzerland, 1986. 2 M, Lamb, W. M. Steen and l). R. F. West, 1(;4l, EO, 1/ol. 1, Laser Institute of America, 1984, p. 133. 3 A. Galerie, M. Pons and M, Caillet, Mater. Sei. Eng., ,%' (1987) 127. 4 J. Dietz and J. Merlin, Rev. Phys. Appl., 33 (1988) 1787. 5 M.F. Ashby and K. E. Easterling, Aeta Metall., 32 (1984) 1935. 6 S. Bonnet-Jobez, Thesis, INSA, Lyon, France, 1989. 7 E. Gaffet, J. M. Pelletier and S. Bonnel-Jobez. Acta Metall., 37 (1989) 32(/5. 8 H. Matija, B. C. Giessen and N. J. Grant, J. Inst. Metals, 96 (1968) 30. 9 L. Renaud, F. Fouquet, A. Elhamdaoui, J. P. Millet, H. Mazille and J. L. Crolet, Acta Memll. Mater. 38, 8 (1990) 1547. 10 T. B. Massalski, Binary All~:s Phase l)iagrarns, Vol. 1, ASM, Metals Park, Ohio, 1986. 11 L. Renaud, F. Fouquet and C. Esnouf, to be published. 12 W. A. Tiller, J. W. Rutter, K. A. Jackson and B. Chalmers, Acta Metall., I (1953) 428.