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Effect of severe plastic deformation on grain boundary liquation of a nickel–base superalloy
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Scripta Materialia 60 (2009) 244–247 www.elsevier.com/locate/scriptamat
Effect of severe plastic deformation on grain boundary liquation of a nickel–base superalloy S.M. Mousavizade,a F. Malek Ghaini,a,* M.J. Torkamany,b J. Sabbaghzadehb,c and A. Abdollah-zadeha a
Department of Materials Engineering, Tarbiat Modares University, P.O. Box 14115-143, Tehran, Iran b Paya Partov Laser Research Centre, P.O. Box 14665-576, Tehran, Iran c NSTR Institute, AEOI, Tehran, Iron Received 30 July 2008; revised 12 October 2008; accepted 14 October 2008 Available online 29 October 2008
The effect of severe plastic deformation associated with friction stir processing (FSP) on grain boundary liquation and the attendant cracking of Inconel 738 superalloy was examined. The intense plastic deformation and the associated heat generation during FSP result in a significant change in the relevant microstructural features, such as the size and distribution of secondary phases. These changes provide a liquation-resistant microstructure to the extent required for suppression of grain boundary liquation during laser surface melting. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Nickel alloys; Friction stir processing; Liquation cracking; Laser welding
Grain boundary (GB) liquation and the attendant cracking has been observed during welding of a variety of alloys, such as nickel–base superalloys, aluminum alloys, stainless steels, maraging steels and magnesium alloys [1–5]. GB liquation cracking is a major problem associated with fusion welding of nickel–base superalloys [3–5]. It is normally associated with local or partial melting of the GBs in the heat-affected zone (HAZ) [1]. The concurrent presence of liquid film and welding stresses and strains during weld cooling could result in HAZ liquation cracking by GB separation [6–8]. To eliminate or reduce the occurrence of GB liquation during the welding of nickel–base superalloys, especially those containing large amounts of (Ti + Al), such as IN738 alloy, several approaches have been used. These include preheating, welding above the homogenization temperature, development of a pre-weld heat treatment, using very low heat input welding processes, such as laser cladding, and utilizing low strength filler alloys [9,10]. Jahnke [11] suggested that electron beam welding with a preheating of the alloy to above the homogenization
* Corresponding author. Tel.: +98 9124401035; fax: +98 21 82883381; e-mail addresses: [email protected]; smmz_2831@yahoo. com
temperature of 1120 °C and the subsequent use of hot isostatic pressing can be used to weld IN738LC without the use of a filler metal. Ojo et al. [7–9] and Chaturvedi [12] reported several GB liquation-related phenomena in precipitation-hardened nickel–base superalloys. They indicated that these alloys are highly susceptible to GB liquation with various heat treatments, several filler metals and different welding processes. A more conventional method to prevent weld cracking in these alloys, however, is to employ lower strength and ductile solid solution strengthened fillers such as IN625. The welds produced by these fillers, even if relatively crack free, have inferior mechanical properties. Attempts are therefore being made to carry out welding with stronger agehardenable fillers that more closely match the base metal properties. However, the use of these fillers promotes the tendency to liquation cracking [10]. Pre-weld heat treatment can improve the weldability of precipitation-hardened nickel–base superalloys through a combination of improved ductility and a desirable microstructure, but it has been indicated that GB liquation is not completely eliminated through pre-weld heat treatment [8]. The microstructural characteristics of precipitationhardened nickel–base superalloys are potentially detrimental in terms of increasing the liquation of GBs and consequently decreasing material resistance to HAZ
1359-6462/$ - see front matter Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2008.10.021
S. M. Mousavizade et al. / Scripta Materialia 60 (2009) 244–247
liquation cracking [7]. A substantial volume of possible liquating particles, such as c0 precipitates, and secondary solidification products, such as carbides, borides, sulphocarbides and c–c0 eutectic phase, along the GBs coupled with a coarse grain structure and a high level of segregation in the cast alloy, contribute to the high susceptibility of these alloys to GB liquation [9]. Significant efforts have been made towards eliminating the GB liquation cracking of nickel–base superalloys. However, none of them have been very efficient and/or completely successful because they cannot alter the liquation-related microstructural features of the alloy completely. This work examines an alternative approach which can effectively change the detrimental microstructure features in terms of GB liquation and, therefore, can produce a liquation-resistant microstructure. With the objective of altering the microstructure, friction stir processing (FSP) was used. FSP is often considered as a generic tool for microstructural modification, such as microstructural refinement and homogenization of cast alloys [13–19]. Severe plastic deformation associated with FSP is expected to alter the liquation-related microstructural features. After FSP of the material’s surface, laser surface melting was used to evaluate its GB liquation susceptibility. Figure 1 shows schematic illustrations of the experimental sequence, sectioning and investigated regions. The cast alloy used in this study was Inconel 738, with the following composition (wt.%): 0.10C, 15.50Cr, 9.8Co, 3.04W, 2.27Mo, 0.70Nb, 0.09Fe, 4.36Al, 3.15Ti, 1.81Ta, 0.04Zr, 0.01B, balance nickel. Single-pass FSP was performed on 100 50 5 mm rectangular plates machined from the as-received cast billets. FSP was performed using a tungsten carbide base alloy tool that consisted of a 14 mm shoulder diameter, 4 mm probe diameter and 1.4 mm probe length. After a number of test runs and producing a processed material with about 1.7 mm depth, the following parameters were selected: tool rotation speed of 800 rpm and tool travel speed of 50 mm min1. The tool was tilted by 3°. The FSP sample was then autogenously laser melted transverse to the direction of FSP, as shown in Figure 1. A pulsed Nd:YAG laser with a maximum mean laser power of 400 W was used. The focusing optical system was composed of three lenses, each with a 75 mm focal length. Pure argon gas at 160 106 m3 s1 flow rate was used for shielding. To achieve a laser fusion zone of adequate depth while keeping the HAZ width at levels lower than the FSP depth, the following process parameters were used: spot diameter of 1 mm; pulse frequency of 40 Hz; pulse duration of 10 ms; travel speed of 3 mm s1 at an average power of 170 W. The final specimen was sectioned along the laser melting direction and transverse to the FSP direction (Fig. 1) and was etched electrolytically in 12 ml H3PO4 + 40 ml HNO3 + 48 ml H2SO4 solution at 6 V for 5 s. The microstructure of the as-cast base metal (BM), the laser fusion zone (FZ), the friction stir processed zone (FSPZ), the HAZ of laser melting in the as-cast base metal and the HAZ of laser melting in the FSPZ (Fig. 1) were studied by optical and scanning electron microscope. Figure 2 shows an optical micrograph of the laser FZ and the as-cast BM. After laser melting of the as-cast plate with prescribed parameters, a shallow FZ of about 1 mm
245
width and 225 lm depth is obtained. No evidence of centerline solidification cracking in the FZ was observed. This could be related to the rapid solidification of the laser melting process and also the effect of the relatively high carbon content [8] of the investigated alloy (0.10 wt.%) in reducing the amount as well as the distribution of interdendritic liquid during the last stages of weld solidification. Microstructural examination of the HAZ of laser melting in the as-cast base metal showed that any GB intersecting the FZ had cracked, with resolidified products formed along them (Fig. 3a), or had liquated but not cracked (Fig. 3b). The HAZ liquation cracks extended into the BM to a distance varying from 20 to 110 lm. Some of the HAZ liquation cracks appeared to have extended a small distance into the FZ (Figs. 2 and 3a). This could be related to the stress concentration due to the formation of a liquation crack. As the liquid pool passes the liquation susceptible region, GB melting occurs in the HAZ. As the heat source passes on, the HAZ contracts on cooling, resulting in the development of tensile stresses which would cause separation of the liquid film at the HAZ grain boundaries, i.e. liquation cracks form in the HAZ. At this point, the fusion zone is in a mushy condition and solidification cracking does not occur since the shrinkage stresses are uniformly distributed. At a later time, however, the fusion zone reaches the liquid film stage and strain concentration develops at the tip of the HAZ crack at the fusion line, leading to hot tearing in the fusion zone [20]. Therefore, the observed hot tear in the fusion zone is essentially a continuation of the HAZ liquation crack. The fact that all hot tears in the fusion zone were associated with liquation cracks in the HAZ supports the proposed mechanism. It is interesting to note that the microstructural examination of the HAZ of laser melting in the FSPZ showed that no GB was cracked (Fig. 4). Complete suppression of liquation cracking in the HAZ of laser melting in the FSPZ and severe liquation cracking of the HAZ of laser melting in the as-cast base metal could be related to the distinct difference between the microstructural characteristics of the as-cast BM (Fig. 5a and d) and the FSP zone (Fig. 5b and c), as discussed below. The microstructure of the as-cast IN738 alloy was cored dendritic (Fig. 2), with enriched interdendritic regions which consisted of several secondary solidification constituents, such as carbides and c–c0 eutectic (Fig. 5a), due to the microsegregation that occurred during solidification. Most of the carbides that formed along the GBs are coarse and semi-continuous (Fig. 5d). Some blocky carbides are also observed within grains. Ojo et al. [7–9] discussed the second-phase particles capable of liquating, the microstructural features of the different liquation phenomena involved and the characteristics of the liquid film contributing to the high susceptibility of IN738 alloy to HAZ liquation cracking in various preweld heat treatments, such as solution or overage heat treatments. In brief, it has been recognized that the basic requirement for constitutional liquation of the secondphase particles is that they should exist at temperatures equal to or above their eutectic temperature on heating. Therefore, the susceptibility of a second-phase particle to constitutional liquation in the weld HAZ will depend upon its solid-state dissolution behavior, as complete
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Figure 1. Schematic illustrations showing the sequence of FSP and laser surface melting; sectioning and investigated regions.
Figure 2. Optical micrograph showing the fusion zone and the HAZ of laser melting in the as-cast base metal with two liquation cracks.
Figure 3. SEM micrographs showing (a) a liquation crack and (b) a liquated but uncracked grain boundary in the HAZ of laser melting in the as-cast base metal.
Figure 4. (a) Optical micrograph showing that liquation cracking did not occur in the HAZ of laser melting in the FSPZ. (b) SEM micrograph showing the transition zone between the uncracked FSP zone and the cracked as-cast base metal.
dissolution prior to reaching the eutectic temperature will preclude the occurrence of liquation. Due to the ra-
pid heating during welding, the dissolution behavior of the second-phase is expected to deviate from equilibrium. Hence, the limited integration time available for homogenization by the diffusion process could cause relatively larger second-phase particles to survive well above their solvus to temperatures above the eutectic reaction temperature of the alloy [7–9]. The microstructural features of the as-cast BM used in this work exacerbate the GB liquation. The as-cast BM consisted of a high volume fraction of coarse secondphase particles capable of liquating, such as the coarse and semi-continuous carbides that formed along the GBs and also the low melting point c–c0 eutectic. As discussed above, these coarse second phases have a high tendency for constitutional liquation. The large grain size of the as-cast BM (400–700 lm) results in a smaller GB area, which promotes the existence of a large volume of coarse carbides and other liquating phases, with the attendant detrimental semi-continuous distribution along the GBs. The observed amount, size and distribution of second phases and the large grain size in the as-cast base metal could increase the intergranular liquid film thickness and therefore lead to higher GB liquation susceptibility. Ojo et al. [8] reported that an increase in the intergranular liquid film thickness would lower the stress required to cause GB decohesion. Also, the high level of GB segregation of melting point depressing elements like boron due to the cast structure with large grain size might have an additional detrimental effect. A high level of such elemental segregation usually tends to lower the terminal resolidification temperature of the intergranular liquid film [21,22] with an attendant build-up of higher welding stresses across the grain boundaries during cooling, and thus can enhance liquation cracking susceptibility. All of these microstructural characteristics, especially the coarse liquating phases, would contribute to the severe liquation cracking when the as-cast microstructure becomes the HAZ of laser melting. In contrast, a single FSP pass dramatically alters the microstructure from that of the as-cast metal. During this process, the intense plastic deformation and thermal exposure that the material undergoes results in significant evolution of the local microstructure. As shown in Figure 5b and the higher magnification SEM image in Figure 5c, FSP resulted in a break-up of both the semi-continuous carbides and the segregated dendritic structure of the as-cast BM. It can be seen that these finer carbides in the FSP zone have relatively less association with the GBs. The distribution and size of other secondary solidification constituents capable of liquating are also expected to significantly change. In addition, FSP results in the generation of very fine recrystallized grains of about 5–10 lm size, as shown in Figure 5c. It should noted that the as-cast BM grain size is about 400– 700 lm. It is reasonable to expect these new GBs to show less segregation as compared to the considerably coarse dendritic structure of the BM formed during the solidification process. Furthermore, the finer grains result in the availability of a larger GB area and hence a lower concentration of liquation-inducing material at the GBs of the FSP zone. As mentioned above, the obtained microstructural features due to the FSP, observed in this work, led to
S. M. Mousavizade et al. / Scripta Materialia 60 (2009) 244–247
247
Figure 5. SEM micrographs showing microstructure of (a) the as-cast base metal (b) and the FSP zone. Relatively higher magnification SEM micrographs of (c) FSP zone and (d) as-cast base metal.
the increased resistance of the alloy to HAZ liquation cracking. During laser melting, smaller second phases, such as the fine carbides observed in the FSP zone, can dissolve completely prior to reaching the eutectic temperature and consequently preclude the occurrence of liquation. The resistance to GB liquation is greatly influenced by fine and smaller second-phase particles, such as carbides and especially the c0 precipitates that are produced during the FSP. Constitutional liquation of the main strengthening phase of precipitation-hardened nickel–base superalloys, the c0 precipitates, was reported recently by Ojo et al. [23]. The size of c0 in the interdendritic regions of the as-cast BM is about 0.45 lm, as compared to the FSP zone, where the average size of the fine c0 particles is about 0.05 lm. The particle size has a significant effect on its solid-state dissolution behavior and consequently its susceptibility to constitutional liquation. For a better understanding of this, an approximate estimation of the dissolution time of the c0 particles is presented here. Eq. (1) [8] gives a simplified description of the dissolution kinetics of spherical precipitates under isothermal conditions: r2 ¼ r20 2kDm t
ð1Þ
where r is the radius of the particle, r0 is the initial particle radius, t is time, Dm is the solute bulk diffusivity and k is given by k¼2
Ci Co Cp Co
ð2Þ
where Ci, Co and Cp are the solute concentration at the particle/matrix interface, in the particle and at infinity, respectively. The time required for a complete isothermal dissolution of the particle can be obtained from Eq. (1) by setting r = 0, so that tf ¼
r20 kDm
ð3Þ
By substituting approximate values for the c0 size, the ratio of the time required for complete isothermal dissolution of c0 in the interdendritic regions of the as-cast BM to the time required for complete isothermal dissolution of fine c0 in the FSP zone is about 80. Therefore, as a first approximation, the complete dissolution of the fine c0 in the FSP zone is 80 times faster than the c0 particles of the as-cast BM, and hence it can be reasonably expected that the fine c0 of the FSP zone is dissolved prior to reaching the eutectic temperature and consequently the occurrence of its constitutional liquation is precluded. In conclusion, the combination of very large deformations and heating associated with FSP of the as-cast
IN738 superalloy can change its GB liquation crack susceptible microstructure to a crack-resistant microstructure to that extent required for the complete suppression of liquation cracking in the laser surface melting. The microstructural features of FSP zone observed in this work, especially the small secondary phases and the fine grain structure, have increased the resistance of IN738 to GB liquation and the attendant cracking. The authors acknowledge the useful discussions with Dr. H. Asadi and T. Saeid of the Tarbiat Modares University. [1] S. Kou, JOM 55 (6) (2003) 37. [2] J.J. Pepe, W.F. Savage, Weld. J. 46 (1967) 411-s. [3] Y.G. Liu, R. Nakkalil, N.L. Richards, M.C. Chaturvedi, Mater. Sci. Eng. A 202 (1995) 179. [4] P. Ferro, A. Zambon, F. Bonollo, Mater. Sci. Eng. A 392 (2005) 94. [5] M. Qian, J.C. Lippold, Mater. Sci. Eng. A 340 (2003) 225. [6] M.B. Henderson, D. Arrell, R. Larsson, M. Heobel, G. Marchant, Sci. Technol. Weld. Join. 9 (2004) 13. [7] O.A. Ojo, N. Richards, M. Chaturvedi, Metall. Trans. A 37 (2006) 421. [8] O.A. Ojo, M.C. Chaturvedi, Metall. Trans. A 38 (2007) 356. [9] O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Scr. Mater. 51 (2004) 141. [10] R.K. Sidhu, N.L. Richards, M.C. Chaturvedi, Mater. Sci. Technol. 21 (2005) 1119. [11] B. Jahnke, Weld. J. 61 (1982) 343s. [12] M.C. Chaturvedi, Mater. Sci. Forum 546–549 (2007) 1163. [13] R.S. Mishra, M.W. Mahoney, S.X. McFadden, N.A. Mara, A.K. Mukherjee, Scr. Mater. 42 (1999) 163. [14] R.S. Mishra, Z.Y. Ma, Mater. Sci. Eng. R 50 (2005) 1. [15] Z.Y. Ma, A.L. Pilchak, M.C. Juhas, J.C. Williams, Scr. Mater. 58 (2008) 361. [16] M.L. Santella, T. Engstrom, D. Storjohann, T.Y. Pan, Scr. Mater. 53 (2005) 201. [17] A.L. Pilchak, M.C. Juhas, J.C. Williams, Metall. Trans. A 38 (2007) 401. [18] K. Oh-Ishi, T.R. McNelley, Metall. Trans. A 35 (2004) 2951. [19] Z.Y. Ma, S.R. Sharma, R.S. Mishra, Mater. Sci. Eng. A 433 (2006) 269. [20] O.A. Ojo, On liquation cracking of Inconel 738LC superalloy welds, PhD Thesis, University of Manitoba, 2004. [21] R.K. Sidhu, O.A. Ojo, M.C. Chaturvedi, Metall. Trans. A 38 (2007) 858. [22] H. Guo, M.C. Chaturvedi, N.L. Richards, G.S. McMahon, Scr. Mater. 40 (1999) 383. [23] O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Scr. Mater. 50 (2004) 641.
Scripta Materialia 60 (2009) 244–247 www.elsevier.com/locate/scriptamat
Effect of severe plastic deformation on grain boundary liquation of a nickel–base superalloy S.M. Mousavizade,a F. Malek Ghaini,a,* M.J. Torkamany,b J. Sabbaghzadehb,c and A. Abdollah-zadeha a
Department of Materials Engineering, Tarbiat Modares University, P.O. Box 14115-143, Tehran, Iran b Paya Partov Laser Research Centre, P.O. Box 14665-576, Tehran, Iran c NSTR Institute, AEOI, Tehran, Iron Received 30 July 2008; revised 12 October 2008; accepted 14 October 2008 Available online 29 October 2008
The effect of severe plastic deformation associated with friction stir processing (FSP) on grain boundary liquation and the attendant cracking of Inconel 738 superalloy was examined. The intense plastic deformation and the associated heat generation during FSP result in a significant change in the relevant microstructural features, such as the size and distribution of secondary phases. These changes provide a liquation-resistant microstructure to the extent required for suppression of grain boundary liquation during laser surface melting. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Nickel alloys; Friction stir processing; Liquation cracking; Laser welding
Grain boundary (GB) liquation and the attendant cracking has been observed during welding of a variety of alloys, such as nickel–base superalloys, aluminum alloys, stainless steels, maraging steels and magnesium alloys [1–5]. GB liquation cracking is a major problem associated with fusion welding of nickel–base superalloys [3–5]. It is normally associated with local or partial melting of the GBs in the heat-affected zone (HAZ) [1]. The concurrent presence of liquid film and welding stresses and strains during weld cooling could result in HAZ liquation cracking by GB separation [6–8]. To eliminate or reduce the occurrence of GB liquation during the welding of nickel–base superalloys, especially those containing large amounts of (Ti + Al), such as IN738 alloy, several approaches have been used. These include preheating, welding above the homogenization temperature, development of a pre-weld heat treatment, using very low heat input welding processes, such as laser cladding, and utilizing low strength filler alloys [9,10]. Jahnke [11] suggested that electron beam welding with a preheating of the alloy to above the homogenization
* Corresponding author. Tel.: +98 9124401035; fax: +98 21 82883381; e-mail addresses: [email protected]; smmz_2831@yahoo. com
temperature of 1120 °C and the subsequent use of hot isostatic pressing can be used to weld IN738LC without the use of a filler metal. Ojo et al. [7–9] and Chaturvedi [12] reported several GB liquation-related phenomena in precipitation-hardened nickel–base superalloys. They indicated that these alloys are highly susceptible to GB liquation with various heat treatments, several filler metals and different welding processes. A more conventional method to prevent weld cracking in these alloys, however, is to employ lower strength and ductile solid solution strengthened fillers such as IN625. The welds produced by these fillers, even if relatively crack free, have inferior mechanical properties. Attempts are therefore being made to carry out welding with stronger agehardenable fillers that more closely match the base metal properties. However, the use of these fillers promotes the tendency to liquation cracking [10]. Pre-weld heat treatment can improve the weldability of precipitation-hardened nickel–base superalloys through a combination of improved ductility and a desirable microstructure, but it has been indicated that GB liquation is not completely eliminated through pre-weld heat treatment [8]. The microstructural characteristics of precipitationhardened nickel–base superalloys are potentially detrimental in terms of increasing the liquation of GBs and consequently decreasing material resistance to HAZ
1359-6462/$ - see front matter Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2008.10.021
S. M. Mousavizade et al. / Scripta Materialia 60 (2009) 244–247
liquation cracking [7]. A substantial volume of possible liquating particles, such as c0 precipitates, and secondary solidification products, such as carbides, borides, sulphocarbides and c–c0 eutectic phase, along the GBs coupled with a coarse grain structure and a high level of segregation in the cast alloy, contribute to the high susceptibility of these alloys to GB liquation [9]. Significant efforts have been made towards eliminating the GB liquation cracking of nickel–base superalloys. However, none of them have been very efficient and/or completely successful because they cannot alter the liquation-related microstructural features of the alloy completely. This work examines an alternative approach which can effectively change the detrimental microstructure features in terms of GB liquation and, therefore, can produce a liquation-resistant microstructure. With the objective of altering the microstructure, friction stir processing (FSP) was used. FSP is often considered as a generic tool for microstructural modification, such as microstructural refinement and homogenization of cast alloys [13–19]. Severe plastic deformation associated with FSP is expected to alter the liquation-related microstructural features. After FSP of the material’s surface, laser surface melting was used to evaluate its GB liquation susceptibility. Figure 1 shows schematic illustrations of the experimental sequence, sectioning and investigated regions. The cast alloy used in this study was Inconel 738, with the following composition (wt.%): 0.10C, 15.50Cr, 9.8Co, 3.04W, 2.27Mo, 0.70Nb, 0.09Fe, 4.36Al, 3.15Ti, 1.81Ta, 0.04Zr, 0.01B, balance nickel. Single-pass FSP was performed on 100 50 5 mm rectangular plates machined from the as-received cast billets. FSP was performed using a tungsten carbide base alloy tool that consisted of a 14 mm shoulder diameter, 4 mm probe diameter and 1.4 mm probe length. After a number of test runs and producing a processed material with about 1.7 mm depth, the following parameters were selected: tool rotation speed of 800 rpm and tool travel speed of 50 mm min1. The tool was tilted by 3°. The FSP sample was then autogenously laser melted transverse to the direction of FSP, as shown in Figure 1. A pulsed Nd:YAG laser with a maximum mean laser power of 400 W was used. The focusing optical system was composed of three lenses, each with a 75 mm focal length. Pure argon gas at 160 106 m3 s1 flow rate was used for shielding. To achieve a laser fusion zone of adequate depth while keeping the HAZ width at levels lower than the FSP depth, the following process parameters were used: spot diameter of 1 mm; pulse frequency of 40 Hz; pulse duration of 10 ms; travel speed of 3 mm s1 at an average power of 170 W. The final specimen was sectioned along the laser melting direction and transverse to the FSP direction (Fig. 1) and was etched electrolytically in 12 ml H3PO4 + 40 ml HNO3 + 48 ml H2SO4 solution at 6 V for 5 s. The microstructure of the as-cast base metal (BM), the laser fusion zone (FZ), the friction stir processed zone (FSPZ), the HAZ of laser melting in the as-cast base metal and the HAZ of laser melting in the FSPZ (Fig. 1) were studied by optical and scanning electron microscope. Figure 2 shows an optical micrograph of the laser FZ and the as-cast BM. After laser melting of the as-cast plate with prescribed parameters, a shallow FZ of about 1 mm
245
width and 225 lm depth is obtained. No evidence of centerline solidification cracking in the FZ was observed. This could be related to the rapid solidification of the laser melting process and also the effect of the relatively high carbon content [8] of the investigated alloy (0.10 wt.%) in reducing the amount as well as the distribution of interdendritic liquid during the last stages of weld solidification. Microstructural examination of the HAZ of laser melting in the as-cast base metal showed that any GB intersecting the FZ had cracked, with resolidified products formed along them (Fig. 3a), or had liquated but not cracked (Fig. 3b). The HAZ liquation cracks extended into the BM to a distance varying from 20 to 110 lm. Some of the HAZ liquation cracks appeared to have extended a small distance into the FZ (Figs. 2 and 3a). This could be related to the stress concentration due to the formation of a liquation crack. As the liquid pool passes the liquation susceptible region, GB melting occurs in the HAZ. As the heat source passes on, the HAZ contracts on cooling, resulting in the development of tensile stresses which would cause separation of the liquid film at the HAZ grain boundaries, i.e. liquation cracks form in the HAZ. At this point, the fusion zone is in a mushy condition and solidification cracking does not occur since the shrinkage stresses are uniformly distributed. At a later time, however, the fusion zone reaches the liquid film stage and strain concentration develops at the tip of the HAZ crack at the fusion line, leading to hot tearing in the fusion zone [20]. Therefore, the observed hot tear in the fusion zone is essentially a continuation of the HAZ liquation crack. The fact that all hot tears in the fusion zone were associated with liquation cracks in the HAZ supports the proposed mechanism. It is interesting to note that the microstructural examination of the HAZ of laser melting in the FSPZ showed that no GB was cracked (Fig. 4). Complete suppression of liquation cracking in the HAZ of laser melting in the FSPZ and severe liquation cracking of the HAZ of laser melting in the as-cast base metal could be related to the distinct difference between the microstructural characteristics of the as-cast BM (Fig. 5a and d) and the FSP zone (Fig. 5b and c), as discussed below. The microstructure of the as-cast IN738 alloy was cored dendritic (Fig. 2), with enriched interdendritic regions which consisted of several secondary solidification constituents, such as carbides and c–c0 eutectic (Fig. 5a), due to the microsegregation that occurred during solidification. Most of the carbides that formed along the GBs are coarse and semi-continuous (Fig. 5d). Some blocky carbides are also observed within grains. Ojo et al. [7–9] discussed the second-phase particles capable of liquating, the microstructural features of the different liquation phenomena involved and the characteristics of the liquid film contributing to the high susceptibility of IN738 alloy to HAZ liquation cracking in various preweld heat treatments, such as solution or overage heat treatments. In brief, it has been recognized that the basic requirement for constitutional liquation of the secondphase particles is that they should exist at temperatures equal to or above their eutectic temperature on heating. Therefore, the susceptibility of a second-phase particle to constitutional liquation in the weld HAZ will depend upon its solid-state dissolution behavior, as complete
246
S. M. Mousavizade et al. / Scripta Materialia 60 (2009) 244–247
Figure 1. Schematic illustrations showing the sequence of FSP and laser surface melting; sectioning and investigated regions.
Figure 2. Optical micrograph showing the fusion zone and the HAZ of laser melting in the as-cast base metal with two liquation cracks.
Figure 3. SEM micrographs showing (a) a liquation crack and (b) a liquated but uncracked grain boundary in the HAZ of laser melting in the as-cast base metal.
Figure 4. (a) Optical micrograph showing that liquation cracking did not occur in the HAZ of laser melting in the FSPZ. (b) SEM micrograph showing the transition zone between the uncracked FSP zone and the cracked as-cast base metal.
dissolution prior to reaching the eutectic temperature will preclude the occurrence of liquation. Due to the ra-
pid heating during welding, the dissolution behavior of the second-phase is expected to deviate from equilibrium. Hence, the limited integration time available for homogenization by the diffusion process could cause relatively larger second-phase particles to survive well above their solvus to temperatures above the eutectic reaction temperature of the alloy [7–9]. The microstructural features of the as-cast BM used in this work exacerbate the GB liquation. The as-cast BM consisted of a high volume fraction of coarse secondphase particles capable of liquating, such as the coarse and semi-continuous carbides that formed along the GBs and also the low melting point c–c0 eutectic. As discussed above, these coarse second phases have a high tendency for constitutional liquation. The large grain size of the as-cast BM (400–700 lm) results in a smaller GB area, which promotes the existence of a large volume of coarse carbides and other liquating phases, with the attendant detrimental semi-continuous distribution along the GBs. The observed amount, size and distribution of second phases and the large grain size in the as-cast base metal could increase the intergranular liquid film thickness and therefore lead to higher GB liquation susceptibility. Ojo et al. [8] reported that an increase in the intergranular liquid film thickness would lower the stress required to cause GB decohesion. Also, the high level of GB segregation of melting point depressing elements like boron due to the cast structure with large grain size might have an additional detrimental effect. A high level of such elemental segregation usually tends to lower the terminal resolidification temperature of the intergranular liquid film [21,22] with an attendant build-up of higher welding stresses across the grain boundaries during cooling, and thus can enhance liquation cracking susceptibility. All of these microstructural characteristics, especially the coarse liquating phases, would contribute to the severe liquation cracking when the as-cast microstructure becomes the HAZ of laser melting. In contrast, a single FSP pass dramatically alters the microstructure from that of the as-cast metal. During this process, the intense plastic deformation and thermal exposure that the material undergoes results in significant evolution of the local microstructure. As shown in Figure 5b and the higher magnification SEM image in Figure 5c, FSP resulted in a break-up of both the semi-continuous carbides and the segregated dendritic structure of the as-cast BM. It can be seen that these finer carbides in the FSP zone have relatively less association with the GBs. The distribution and size of other secondary solidification constituents capable of liquating are also expected to significantly change. In addition, FSP results in the generation of very fine recrystallized grains of about 5–10 lm size, as shown in Figure 5c. It should noted that the as-cast BM grain size is about 400– 700 lm. It is reasonable to expect these new GBs to show less segregation as compared to the considerably coarse dendritic structure of the BM formed during the solidification process. Furthermore, the finer grains result in the availability of a larger GB area and hence a lower concentration of liquation-inducing material at the GBs of the FSP zone. As mentioned above, the obtained microstructural features due to the FSP, observed in this work, led to
S. M. Mousavizade et al. / Scripta Materialia 60 (2009) 244–247
247
Figure 5. SEM micrographs showing microstructure of (a) the as-cast base metal (b) and the FSP zone. Relatively higher magnification SEM micrographs of (c) FSP zone and (d) as-cast base metal.
the increased resistance of the alloy to HAZ liquation cracking. During laser melting, smaller second phases, such as the fine carbides observed in the FSP zone, can dissolve completely prior to reaching the eutectic temperature and consequently preclude the occurrence of liquation. The resistance to GB liquation is greatly influenced by fine and smaller second-phase particles, such as carbides and especially the c0 precipitates that are produced during the FSP. Constitutional liquation of the main strengthening phase of precipitation-hardened nickel–base superalloys, the c0 precipitates, was reported recently by Ojo et al. [23]. The size of c0 in the interdendritic regions of the as-cast BM is about 0.45 lm, as compared to the FSP zone, where the average size of the fine c0 particles is about 0.05 lm. The particle size has a significant effect on its solid-state dissolution behavior and consequently its susceptibility to constitutional liquation. For a better understanding of this, an approximate estimation of the dissolution time of the c0 particles is presented here. Eq. (1) [8] gives a simplified description of the dissolution kinetics of spherical precipitates under isothermal conditions: r2 ¼ r20 2kDm t
ð1Þ
where r is the radius of the particle, r0 is the initial particle radius, t is time, Dm is the solute bulk diffusivity and k is given by k¼2
Ci Co Cp Co
ð2Þ
where Ci, Co and Cp are the solute concentration at the particle/matrix interface, in the particle and at infinity, respectively. The time required for a complete isothermal dissolution of the particle can be obtained from Eq. (1) by setting r = 0, so that tf ¼
r20 kDm
ð3Þ
By substituting approximate values for the c0 size, the ratio of the time required for complete isothermal dissolution of c0 in the interdendritic regions of the as-cast BM to the time required for complete isothermal dissolution of fine c0 in the FSP zone is about 80. Therefore, as a first approximation, the complete dissolution of the fine c0 in the FSP zone is 80 times faster than the c0 particles of the as-cast BM, and hence it can be reasonably expected that the fine c0 of the FSP zone is dissolved prior to reaching the eutectic temperature and consequently the occurrence of its constitutional liquation is precluded. In conclusion, the combination of very large deformations and heating associated with FSP of the as-cast
IN738 superalloy can change its GB liquation crack susceptible microstructure to a crack-resistant microstructure to that extent required for the complete suppression of liquation cracking in the laser surface melting. The microstructural features of FSP zone observed in this work, especially the small secondary phases and the fine grain structure, have increased the resistance of IN738 to GB liquation and the attendant cracking. The authors acknowledge the useful discussions with Dr. H. Asadi and T. Saeid of the Tarbiat Modares University. [1] S. Kou, JOM 55 (6) (2003) 37. [2] J.J. Pepe, W.F. Savage, Weld. J. 46 (1967) 411-s. [3] Y.G. Liu, R. Nakkalil, N.L. Richards, M.C. Chaturvedi, Mater. Sci. Eng. A 202 (1995) 179. [4] P. Ferro, A. Zambon, F. Bonollo, Mater. Sci. Eng. A 392 (2005) 94. [5] M. Qian, J.C. Lippold, Mater. Sci. Eng. A 340 (2003) 225. [6] M.B. Henderson, D. Arrell, R. Larsson, M. Heobel, G. Marchant, Sci. Technol. Weld. Join. 9 (2004) 13. [7] O.A. Ojo, N. Richards, M. Chaturvedi, Metall. Trans. A 37 (2006) 421. [8] O.A. Ojo, M.C. Chaturvedi, Metall. Trans. A 38 (2007) 356. [9] O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Scr. Mater. 51 (2004) 141. [10] R.K. Sidhu, N.L. Richards, M.C. Chaturvedi, Mater. Sci. Technol. 21 (2005) 1119. [11] B. Jahnke, Weld. J. 61 (1982) 343s. [12] M.C. Chaturvedi, Mater. Sci. Forum 546–549 (2007) 1163. [13] R.S. Mishra, M.W. Mahoney, S.X. McFadden, N.A. Mara, A.K. Mukherjee, Scr. Mater. 42 (1999) 163. [14] R.S. Mishra, Z.Y. Ma, Mater. Sci. Eng. R 50 (2005) 1. [15] Z.Y. Ma, A.L. Pilchak, M.C. Juhas, J.C. Williams, Scr. Mater. 58 (2008) 361. [16] M.L. Santella, T. Engstrom, D. Storjohann, T.Y. Pan, Scr. Mater. 53 (2005) 201. [17] A.L. Pilchak, M.C. Juhas, J.C. Williams, Metall. Trans. A 38 (2007) 401. [18] K. Oh-Ishi, T.R. McNelley, Metall. Trans. A 35 (2004) 2951. [19] Z.Y. Ma, S.R. Sharma, R.S. Mishra, Mater. Sci. Eng. A 433 (2006) 269. [20] O.A. Ojo, On liquation cracking of Inconel 738LC superalloy welds, PhD Thesis, University of Manitoba, 2004. [21] R.K. Sidhu, O.A. Ojo, M.C. Chaturvedi, Metall. Trans. A 38 (2007) 858. [22] H. Guo, M.C. Chaturvedi, N.L. Richards, G.S. McMahon, Scr. Mater. 40 (1999) 383. [23] O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Scr. Mater. 50 (2004) 641.