- Email: [email protected]
Solidification path of single-crystal nickel-base superalloys with minor carbon additions under laser rapid directional solidification conditions
Scripta Materialia 127 (2017) 58–62
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular Article
Solidification path of single-crystal nickel-base superalloys with minor carbon additions under laser rapid directional solidification conditions Yao-Jian Liang, Jia Li ⁎, An Li, Xiao-Tong Pang, Hua-Ming Wang National Engineering Laboratory of Additive Manufacturing for Large Metallic Components and Engineering Research Center of Ministry of Education on Laser Direct Manufacturing for Large Metallic Components, School of Materials Science and Engineering, Beihang University, 37 Xueyuan Road, Beijing 100191, People's Republic of China
a r t i c l e
i n f o
Article history: Received 25 June 2016 Received in revised form 31 August 2016 Accepted 31 August 2016 Available online xxxx Keywords: Laser treatment Superalloy Rapid directional solidification Solidification microstructure Phase transformations
a b s t r a c t The solidification path of single-crystal nickel-base superalloys containing minor carbon was investigated under various laser rapid directional solidification (LRDS) conditions. By controlling the solidification rate, LRDS processing can provide the evidence whether some diffusion-controlled phase transformations occur because such transformations will be suppressed under high cooling rates. Results show that the solidification path and final solidification microstructure depend upon the cooling rate; the microstructure without γ-γ′ eutectic can be obtained as long as the cooling rate is high enough. A peritectic transformation in carbon-containing singlecrystal superalloys was first experimentally verified by controlling the cooling rate during LRDS processing. © 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Single-crystal (SX) nickel-base superalloys are widely used in turbine engines due to their excellent mechanical properties at elevated temperatures. Although carbon, a grain boundary strengthener, was fully eliminated from early SX superalloys, minor additions of carbon (≤0.05 wt%) have been reintroduced to a number of commercial SX alloys because of the frequent existence of low-angle grain boundaries [1, 2]. Moreover, recent research by Tin and Pollock et al. [3–6] has confirmed that under conventional DS conditions, minor carbon additions to high-refractory SX superalloys are beneficial to lower the tendency of grain-defect formation and refractory element segregation. Hence, it is desirable to understand the solidification behavior of these SX alloys containing minor carbon because it is critical for the control of their final solidification microstructure and mechanical properties. Recently, laser processing, e.g., laser additive manufacturing (LAM), has exhibited its potential for precise repair and rapid forming of SX components [7–16] with high solidification cooling rates [17–19]. It is well known that rapid solidification (RS) processing is an effective way to reduce microsegregation and refine microstructure. More importantly, by controlling solidification velocity, RS processing can provide the evidence whether some diffusion-controlled phase transformations occur because these transformations will be suppressed under high solidification rates [20]. It is therefore specifically appropriate to be used to investigate solidification phase transformations. With this in mind, of particular interest is the effect of minor carbon additions on the solidification sequence of SX superalloys under ⁎ Corresponding author. E-mail address: [email protected] (J. Li).
http://dx.doi.org/10.1016/j.scriptamat.2016.08.039 1359-6462/© 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
laser rapid directional solidification (LRDS) conditions. Unfortunately, although many researchers have focused on the morphology and growth mechanism of MC carbides [21,22] and on the solidification behaviors of early DS superalloys with relatively high carbon contents [23, 24], the solidification path of the SX superalloys with minor carbon additions has not yet been reported under LRDS conditions. Consequently, the main objective of this study is to reveal the solidification sequence of the SX nickel-base superalloys with minor carbon additions under RS conditions. For this purpose, a SX superalloy containing minor carbon was treated under various LRDS conditions to understand its RS path. To obtain various LRDS conditions, laser remelting experiments were conducted with different processing parameters (Table 1) by a 6 kW LAM system with an oxygen content b 50 ppm. The chemical composition of the SX Ni-base superalloy studied here is Ni–5.7Al–9.2Cr– 5.3Co–2.7Ta–2.3Ti–8.9W–0.015C (in wt%). All SX substrates were machined from a conventional DS SX cast ingot with the [001] orientation normal to the remelted surfaces. For all experiments, the substrate surfaces were ground with 600-grit SiC paper, cleaned in methanol, and maintained at room temperature prior to laser remelting. In order to relate a set of given laser processing parameters to resulting solidification behavior, the temperature gradient was estimated by the laser temperature field, and the solidification velocity was related to the laser scanning velocity and melt-pool geometry, where the temperature field and melt-pool geometry were given by a laser-remelted heat-transfer equation reported elsewhere [11,15]. Subsequently, mean temperature gradient, GM, and solidification velocity, VM, were calculated over the total melt-pool depth along the symmetrical centerline of liquidus isotherm. Corresponding mean cooling rates in melt pools, RC, listed in
Y.-J. Liang et al. / Scripta Materialia 127 (2017) 58–62 Table 1 Laser processing parameters (laser power, P, scanning velocity, Vb, and beam diameter, Db) and corresponding cooling rate, RC. No.
P (W)
Vb (mm/s)
Db (mm)
RC (K/s)
A B
2000 2000
5 45
4 2
1101.1 17,606.4
Table 1, can be obtained by RC = dT/dt = GM · VM, where T is the temperature and t is the time [20]. The microstructures were characterized by a CamScan Apollo300 field emission scanning electron microscope (SEM) equipped with an Oxford INCA energy dispersive spectrometer (EDS). The thermal characteristics of the samples were investigated by a Netzsch STA-449F3 differential scanning calorimetry (DSC) instrument on heating. Fig. 1a and b shows the solidification microstructures of the LRDS SX samples. Both the two processing conditions can produce good rapid directional solidification microstructures. The fine dendrites grow epitaxially from the SX substrates to the top of laser tracks along the [001]/ 〈100〉 crystallographic orientation. For investigating the solidification path of the RS carbon-containing SX alloys, Fig. 1c and d shows the cross-section dendritic morphologies of the two samples. As shown in Fig. 1c, with minor carbon additions, some white, branching precipitates occur in the interdendritic regions of epitaxial γ dendrites in sample A. The DSC profiles show that there is a reaction peak at ~ 1612 K just below the liquidus (~1635 K), which was identified as a γ-MC eutectic peak in earlier research [3–6] and verifies these phases formed after the growth of γ dendrites to be MC carbides. Similar to earlier research [21,22], the EDS analysis shows that these carbides are mainly Ti- and Ta-rich. In comparison, owing to a great increase in cooling rate, the primary dendrites are finer and the carbides distributed in the interdendritic regions of sample B become non-branching, small rods (or layers) (see Fig. 1d). Additionally, the porosity in sample B is visibly higher than that in sample A. This appears to be because as solidification velocity increases, residual liquid has no enough time to fill up the interdendritic regions. Even so, their solidification times are rather
59
short compared with conventional DS process, which implies that there exist other reasons affecting the porosity and they will be discussed later. It is also noteworthy that no visible γ-γ′ eutectic is found in Fig. 1. However, the last stage solidification is generally considered to be completed by forming the γ-γ′ eutectic microstructure in earlier research on the solidification behaviors of welded superalloys with relatively high carbon contents by Ojo and Sidhu et al. [23,24]. Therefore, the understanding of the solidification path of the LRDS SX alloys with a low carbon content requires careful investigations into the interdendritic microstructures. In order to assess the solidification sequence, the LRDS microstructures at interdendritic regions of samples A and B are shown in Fig. 2a and b, respectively. The branching morphology of the MC carbides can be seen more clearly in Fig. 2a. It is interesting that a small number of γ-γ′ eutectic islands are found, as shown in the middle-top of the figure. Compared with sample A, the branches of the MC carbides in sample B are not visible and no γ-γ′ eutectic is observed. To better understand the solidification sequence, Fig. 2c illustrates the back-scattered electron (BSE) image of the interdendritic region same as Fig. 2a. As can be seen, the γ-γ′ eutectic is very close to (or even locates at) interdendritic shrinkage defects, which implies that the initiation of these reactions indeed lags behind the precipitation of the MC carbides. In particular, a few MC carbides are observed to be wrapped in these γ-γ′ eutectic microstructures and this characteristic can be seen more clearly in a highmagnification view (Fig. 2d). However, such a microstructure is not found in the conventional DS SX substrate with a low cooling rate though the MC carbides also adjoin the γ-γ′ eutectic islands (see Fig. 2e). According to above careful analyses of solidification microstructure, the solidification paths of the SX superalloy containing minor carbon under various LRDS conditions can be proposed by using a simplified pseudo-ternary composition triangle of γ, γ′, and MC similar to previous research [23,25] (Fig. 3a). Step I. Solidification commences with the well-known phase transformation, L → γ0. As a result, the solidification microstructure of the LRDS SX exhibits mostly γ dendrites epitaxially growing from the SX
Fig. 1. OM images showing the solidification microstructures of the LRDS SX samples: (a) A and (b) B. SEM images of the cross-section dendritic morphologies of samples (c) A and (d) B.
60
Y.-J. Liang et al. / Scripta Materialia 127 (2017) 58–62
Fig. 2. SEM images showing the solidification microstructures at interdendritic regions of the LRDS SX samples: (a) A and (b) B. (c) SEM BSE image of the interdendritic region same as (a). (d) The high-magnification view of the region of the red box in (a). (e) SEM image of the interdendritic microstructures of the SX substrate. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
substrate. As the solidification proceeds, the liquid will be enriched in the elements with partition coefficients, k, less than unity, including Ti, Ta, and C. In addition, the kC (~ 0.2 [26]) is much less than the kTi and kTa (~ 0.7 [26–29]), which indicates that the enrichment rate of C atoms in liquid is higher. Such enrichments will lead to the composition preferentially reaching the tri-phase equilibrium among liquid, γ, and MC, rather than liquid, γ, and γ′, as the path C0e schematically shown
in Fig. 3a. It should be noted that when the solidification velocity is high enough, partition coefficients should be considered to be velocity-dependent rather than constant [20], which can lead to the reduction in the difference between kC and kTa/Ti. Nevertheless, the enrichment rate of C atoms in liquid is still higher than that of Ti/Ta atoms and similar γ-MC eutectic reactions have also been reported under higher cooling rates (see the following).
Fig. 3. (a) The liquidus projection for a pseudo-ternary composition triangle of γ, γ′, and MC. Schematics of the proposed solidification sequence under different LRDS conditions: (b) a relatively high cooling rate (A), and (c) a rather high cooling rate (B). (d) A pseudo-binary peritectic phase diagram of γ-MC and γ-γ′.
Y.-J. Liang et al. / Scripta Materialia 127 (2017) 58–62
Step II. When the composition reaches the point e in Fig. 3a, the solidification continues with a γ-MC eutectic reaction: L → γE + MC, the path eP, according to the DSC results. This eutectic-type solidification reaction can explain the formation of MC carbides with a branching morphology. Similar eutectic-type solidification reaction of MC carbides was also observed under higher cooling rates (~6.2 × 104 and 1.4 × 105 K/s) in earlier research on a superalloy with a lower C content (b0.01%) by Wang et al. [21], reconfirming that with a relatively high C content (0.015%) this reaction is indeed after the epitaxial directional growth of single-phase γ dendrites. Such γ-MC eutectic reactions have also been widely reported in the superalloys with higher C contents based on the analyses of DSC curves [3–6] or their morphology, location, and compositions [22–24]. Step III—Absent under rather high cooling rates. On further solidification, the point P in Fig. 3a represents that the solidification appears to continue with a peritectic reaction: L + MC → γ + γ′. Although Ojo et al. [23] have suggested that this solidification step is not a ternary eutectic reaction and may be a peritectic transformation by theoretical (phase rule) analyses in their previous study on the solidification behavior of IN738LC superalloy, they do not provide any evidence to confirm this process. Fortunately, peritectic transformation requires the diffusion of solute atoms in peritectic (solid-state) phase, meaning that it will be suppressed under very high solidification rates [30]. As shown in Fig. 2a, c, and d, a few residual MC carbides are observed to be wrapped in γ-γ′ eutectic microstructures under a high cooling rate (N 1000 K/s). Interestingly, this phenomenon does not exist in both the LRDS sample B with a very high cooling rate (N 17,000 K/s) (Fig. 2b) and the conventional DS substrate with a rather low cooling rate (Fig. 2e). All these observations have adequately verified that this transformation is peritectic-type. Because of very high diffusivity of C atoms, this peritectic process is complete under low cooling rates (substrate). As a result, only the MC carbides adjoining the γ-γ′ eutectic islands can be seen (see Fig. 2e and Ref. [23]). As the cooling rate increases, solidification time is shortened and the peritectic process becomes difficult to be completed (sample A). Accordingly, a few residual MC carbides will be wrapped in γ-γ′ microstructures (Fig. 2a, c, and d). Under very high cooling rates, the solidification time is rather short; therefore the diffusion of C atoms in solid is quite limited and the peritectic process is very difficult to conduct (sample B). Consequently, only the primary MC carbides exist and no γ-γ′ microstructure can be observed (Fig. 2b). Step IV—Absent under rather high cooling rates. The solidification is finally completed by forming γ-γ′ eutectic: L → γ + γ′. Strictly speaking, this process may also be peritectic-type [31–33]. However, understanding the last stage phase transformation requires further in-depth and careful investigations and this is not the emphasis of this work. Therefore, we retain to this well-known term ‘γ-γ′ eutectic’ to address such a process and resulting microstructure consistent with previous research [33–35]. Fig. 3b and c schematically summarizes the solidification sequences under different LRDS conditions. As can be seen, solidification initiates at TL with the epitaxial directional growth of columnar γ dendrites, L → γ0, and continues with a γ-MC eutectic reaction resulting in the formation of MC carbides, L → γE + MC, when the temperature declines below TMC. Subsequently, if the cooling rate is low enough to allow the sufficient diffusion of C atoms in solid and γ′ formers in liquid, a peritectic transformation will occur at TP by consuming the primary MC carbides, L + MC → γ + γ′, and solidification will be finally completed by forming the γ-γ′ eutectic microstructure: L → γ + γ′ (Fig. 3b and Fig. 2a). Otherwise, only Step I and II phase transformations occur under rather high cooling rates, as shown in Fig. 3c and Fig. 2b. The difference in solidification paths is easy to be explained by a pseudo-binary phase diagram of γ-MC and γ-γ′ (Fig. 3d). If cooling rates are not high enough, solidification will continue with the equilibrium liquidus after point P (solid line) whereas for very high cooling rates it will continue with a nonequilibrium transformation resulting in continuous formation of
61
γ-MC eutectic because of very low diffusivities of C atoms in solid and γ′ formers in liquid (dash line extrapolating from eP). Such nonequilibrium transformations have been widely reported in earlier research on peritectic alloys under RS conditions [36]. More importantly, these findings demonstrate a feasible method for preparing the SX alloys without γ-γ′ eutectic by controlling the solidification rate. Furthermore, MC carbides become non-branching, small rods (or layers) within interdendritic regions under very high cooling rates (Fig. 1d and Fig. 2b). This is likely due to a divorced eutectic reaction caused by limited diffusion of C atoms in liquid. Interestingly, these subsequent diffusion phase transformations occurring at the last stage solidification appear to be able to fill up the shrinkage defects (see Fig. 2c), implying that a rather high cooling rate that will suppress them is not always beneficial to the solidification microstructure. This may be another reason why the porosity in sample B is visibly higher than that in sample A in addition to the effect of solidification time (Fig. 1). In view of this, a good solidification microstructure requires an appropriate control of solidification rate. In summary, the solidification path of the SX nickel-base superalloys containing minor carbon was investigated under various LRDS conditions. It was found that solidification initiates with the epitaxial directional growth of γ dendrites, L → γ0, and continues with a γ-MC eutectic reaction resulting in the formation of MC carbides, L → γE + MC. Subsequently, only when the cooling rate is low enough to allow the diffusion of C atoms in solid and γ′ formers in liquid, will a peritectic transformation occur by consuming the MC carbides, L + MC → γ + γ′, and solidification be finally completed by forming the well-known γ-γ′ eutectic microstructure: L → γ + γ′. By controlling the cooling rate during LRDS processing, above peritectic transformation was first experimentally verified and a feasible method for preparing the SX alloys without γ-γ′ eutectic was demonstrated. Nevertheless, these subsequent diffusion phase transformations are able to fill up the shrinkage defects, implying that a rather high cooling rate that will suppress them is not always beneficial to the solidification microstructure. A good solidification microstructure requires an appropriate control of solidification rate. The understanding of the solidification sequence of these LRDS carbon-containing SX superalloys contributes to the prediction and control of the morphology and quantity of carbides and further understanding of the solidification behavior in the future. Acknowledgements This work was supported by the National High Technology Research and Development Program of China (Grant No. 2014AA041701) and the Joint Funds of NSFC-Liaoning (Grant No. U1508213). The authors thank Drs. ZHANG Shu-Quan and TIAN Xiang-Jun for laser processing. References [1] R.C. Reed, The Superalloys: Fundamentals and Applications, Cambridge university Press, 2006. [2] M.J. Donachie, S.J. Donachie, Superalloys: A Technical Guide, Second Edition ASM International, Materials Park, Ohio, 2002. [3] S. Tin, T.M. Pollock, W.T. King, in: T.M. Pollock, R.D. Kissinger, R.R. Bowman, K.A. Green, M. McLean, S. Olson, J.J. Scbirra (Eds.), Superalloys 2000, TMS (The Minerals, Metals &Materials Society) 2000, pp. 201–210. [4] S. Tin, T.M. Pollock, Metall. Mater. Trans. A 34A (2003) 1953–1967. [5] S. Tin, T.M. Pollock, W. Murphy, Metall. Mater. Trans. A 32A (2001) 1743–1753. [6] S. Tin, T.M. Pollock, J. Mater. Sci. 39 (2004) 7199–7205. [7] R. Acharya, R. Bansal, J.J. Gambone, S. Das, Metall. Mater. Trans. B Process Metall. Mater. Process. Sci. 45 (2014) 2247–2261. [8] R. Acharya, R. Bansal, J.J. Gambone, S. Das, Metall. Mater. Trans. B Process Metall. Mater. Process. Sci. 45 (2014) 2279–2290. [9] T.D. Anderson, J.N. DuPont, T. DebRoy, Metall. Mater. Trans. A 41 (2009) 181–193. [10] T.D. Anderson, J.N. DuPont, T. DebRoy, Acta Mater. 58 (2010) 1441–1454. [11] M. Gäumann, C. Bezençon, P. Canalis, W. Kurz, Acta Mater. 49 (2001) 1051–1062. [12] W. Liu, J.N. DuPont, Acta Mater. 52 (2004) 4833–4847. [13] W. Liu, J.N. DuPont, Acta Mater. 53 (2005) 1545–1558. [14] J.M. Vitek, Acta Mater. 53 (2005) 53–67. [15] Y.-J. Liang, H.-M. Wang, Mater. Des. 102 (2016) 297–302.
62
Y.-J. Liang et al. / Scripta Materialia 127 (2017) 58–62
[16] Y.-J. Liang, A. Li, X. Cheng, X.-T. Pang, H.-M. Wang, J. Alloys Compd. 688 (2016) 133–142. [17] Y. Xiong, W.H. Hofmeister, Z. Cheng, J.E. Smugeresky, E.J. Lavernia, J.M. Schoenung, Acta Mater. 57 (2009) 5419–5429. [18] M. Gremaud, J.-D. Wagnière, A. Zryd, W. Kurz, Surf. Eng. 12 (1996) 251–259. [19] Y.-J. Liang, X. Cheng, H.-M. Wang, Acta Mater. 118 (2016) 17–27. [20] W. Kurz, D.J. Fisher, Fundamentals of Solidification, Trans Tech Publications, Switzerland, 1998. [21] H.M. Wang, J.H. Zhang, Y.J. Tang, Z.Q. Hu, N. Yukawa, M. Morinaga, Y. Murata, Mater. Sci. Eng. A 156 (1992) 109–116. [22] L. Liu, F. Sommer, Scr. Metall. Mater. 30 (1994) 587–591. [23] O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Metall. Mater. Trans. A 37A (2006) 421–433. [24] R.K. Sidhu, O.A. Ojo, M.C. Chaturvedi, Metall. Mater. Trans. A 38A (2007) 858–870. [25] B. Radhakrishnan, R.G. Thompson, Metall. Trans. A 22A (1991) 887–902. [26] J.N. DuPont, J.C. Lippold, S.D. Kiser, Welding Metallurgy and Weldability of Nickelbase Alloys, John Wiley & Sons, Hoboken, New Jersey, 2009.
[27] S.-M. Seo, H.-W. Jeong, Y.-K. Ahn, D.W. Yun, J.-H. Lee, Y.-S. Yoo, Mater. Charact. 89 (2014) 43–55. [28] A. Thirumalai, A. Akhtar, R.C. Reed, Mater. Sci. Technol. 22 (2006) 1–13. [29] R.A. Hobbs, S. Tin, C.M.F. Rae, R.W. Broomfield, C.J. Humphreys, in: K.A. Green, T.M. Pollock, H. Harada, T.E. Howson, R.C. Reed, J.J. Schirra, S. Walston (Eds.), Superalloys 2004, TMS (The Minerals, Metals & Materials Society) 2004, pp. 819–825. [30] D.E. Laughlin, K. Hono, Physical Metallurgy, Elsevier, Fifth ed, 2014. [31] J.H. Lee, J.D. Verhoeven, J. Phase Equilib. 15 (1994) 136–146. [32] J.H. Lee, J.D. Verhoeven, J. Cryst. Growth 143 (1994) 86–102. [33] N. Warnken, J. Phase Equilib. Diffus. 37 (2016) 100–107. [34] A. Heckl, R. Rettig, S. Cenanovic, M. Göken, R.F. Singer, J. Cryst. Growth 312 (2010) 2137–2144. [35] N. Warnken, D. Ma, M. Mathes, I. Steinbach, Mater. Sci. Eng. A 413-414 (2005) 267–271. [36] H.W. Kerr, W. Kurz, Int. Mater. Rev. 41 (1996) 129–164.
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular Article
Solidification path of single-crystal nickel-base superalloys with minor carbon additions under laser rapid directional solidification conditions Yao-Jian Liang, Jia Li ⁎, An Li, Xiao-Tong Pang, Hua-Ming Wang National Engineering Laboratory of Additive Manufacturing for Large Metallic Components and Engineering Research Center of Ministry of Education on Laser Direct Manufacturing for Large Metallic Components, School of Materials Science and Engineering, Beihang University, 37 Xueyuan Road, Beijing 100191, People's Republic of China
a r t i c l e
i n f o
Article history: Received 25 June 2016 Received in revised form 31 August 2016 Accepted 31 August 2016 Available online xxxx Keywords: Laser treatment Superalloy Rapid directional solidification Solidification microstructure Phase transformations
a b s t r a c t The solidification path of single-crystal nickel-base superalloys containing minor carbon was investigated under various laser rapid directional solidification (LRDS) conditions. By controlling the solidification rate, LRDS processing can provide the evidence whether some diffusion-controlled phase transformations occur because such transformations will be suppressed under high cooling rates. Results show that the solidification path and final solidification microstructure depend upon the cooling rate; the microstructure without γ-γ′ eutectic can be obtained as long as the cooling rate is high enough. A peritectic transformation in carbon-containing singlecrystal superalloys was first experimentally verified by controlling the cooling rate during LRDS processing. © 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Single-crystal (SX) nickel-base superalloys are widely used in turbine engines due to their excellent mechanical properties at elevated temperatures. Although carbon, a grain boundary strengthener, was fully eliminated from early SX superalloys, minor additions of carbon (≤0.05 wt%) have been reintroduced to a number of commercial SX alloys because of the frequent existence of low-angle grain boundaries [1, 2]. Moreover, recent research by Tin and Pollock et al. [3–6] has confirmed that under conventional DS conditions, minor carbon additions to high-refractory SX superalloys are beneficial to lower the tendency of grain-defect formation and refractory element segregation. Hence, it is desirable to understand the solidification behavior of these SX alloys containing minor carbon because it is critical for the control of their final solidification microstructure and mechanical properties. Recently, laser processing, e.g., laser additive manufacturing (LAM), has exhibited its potential for precise repair and rapid forming of SX components [7–16] with high solidification cooling rates [17–19]. It is well known that rapid solidification (RS) processing is an effective way to reduce microsegregation and refine microstructure. More importantly, by controlling solidification velocity, RS processing can provide the evidence whether some diffusion-controlled phase transformations occur because these transformations will be suppressed under high solidification rates [20]. It is therefore specifically appropriate to be used to investigate solidification phase transformations. With this in mind, of particular interest is the effect of minor carbon additions on the solidification sequence of SX superalloys under ⁎ Corresponding author. E-mail address: [email protected] (J. Li).
http://dx.doi.org/10.1016/j.scriptamat.2016.08.039 1359-6462/© 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
laser rapid directional solidification (LRDS) conditions. Unfortunately, although many researchers have focused on the morphology and growth mechanism of MC carbides [21,22] and on the solidification behaviors of early DS superalloys with relatively high carbon contents [23, 24], the solidification path of the SX superalloys with minor carbon additions has not yet been reported under LRDS conditions. Consequently, the main objective of this study is to reveal the solidification sequence of the SX nickel-base superalloys with minor carbon additions under RS conditions. For this purpose, a SX superalloy containing minor carbon was treated under various LRDS conditions to understand its RS path. To obtain various LRDS conditions, laser remelting experiments were conducted with different processing parameters (Table 1) by a 6 kW LAM system with an oxygen content b 50 ppm. The chemical composition of the SX Ni-base superalloy studied here is Ni–5.7Al–9.2Cr– 5.3Co–2.7Ta–2.3Ti–8.9W–0.015C (in wt%). All SX substrates were machined from a conventional DS SX cast ingot with the [001] orientation normal to the remelted surfaces. For all experiments, the substrate surfaces were ground with 600-grit SiC paper, cleaned in methanol, and maintained at room temperature prior to laser remelting. In order to relate a set of given laser processing parameters to resulting solidification behavior, the temperature gradient was estimated by the laser temperature field, and the solidification velocity was related to the laser scanning velocity and melt-pool geometry, where the temperature field and melt-pool geometry were given by a laser-remelted heat-transfer equation reported elsewhere [11,15]. Subsequently, mean temperature gradient, GM, and solidification velocity, VM, were calculated over the total melt-pool depth along the symmetrical centerline of liquidus isotherm. Corresponding mean cooling rates in melt pools, RC, listed in
Y.-J. Liang et al. / Scripta Materialia 127 (2017) 58–62 Table 1 Laser processing parameters (laser power, P, scanning velocity, Vb, and beam diameter, Db) and corresponding cooling rate, RC. No.
P (W)
Vb (mm/s)
Db (mm)
RC (K/s)
A B
2000 2000
5 45
4 2
1101.1 17,606.4
Table 1, can be obtained by RC = dT/dt = GM · VM, where T is the temperature and t is the time [20]. The microstructures were characterized by a CamScan Apollo300 field emission scanning electron microscope (SEM) equipped with an Oxford INCA energy dispersive spectrometer (EDS). The thermal characteristics of the samples were investigated by a Netzsch STA-449F3 differential scanning calorimetry (DSC) instrument on heating. Fig. 1a and b shows the solidification microstructures of the LRDS SX samples. Both the two processing conditions can produce good rapid directional solidification microstructures. The fine dendrites grow epitaxially from the SX substrates to the top of laser tracks along the [001]/ 〈100〉 crystallographic orientation. For investigating the solidification path of the RS carbon-containing SX alloys, Fig. 1c and d shows the cross-section dendritic morphologies of the two samples. As shown in Fig. 1c, with minor carbon additions, some white, branching precipitates occur in the interdendritic regions of epitaxial γ dendrites in sample A. The DSC profiles show that there is a reaction peak at ~ 1612 K just below the liquidus (~1635 K), which was identified as a γ-MC eutectic peak in earlier research [3–6] and verifies these phases formed after the growth of γ dendrites to be MC carbides. Similar to earlier research [21,22], the EDS analysis shows that these carbides are mainly Ti- and Ta-rich. In comparison, owing to a great increase in cooling rate, the primary dendrites are finer and the carbides distributed in the interdendritic regions of sample B become non-branching, small rods (or layers) (see Fig. 1d). Additionally, the porosity in sample B is visibly higher than that in sample A. This appears to be because as solidification velocity increases, residual liquid has no enough time to fill up the interdendritic regions. Even so, their solidification times are rather
59
short compared with conventional DS process, which implies that there exist other reasons affecting the porosity and they will be discussed later. It is also noteworthy that no visible γ-γ′ eutectic is found in Fig. 1. However, the last stage solidification is generally considered to be completed by forming the γ-γ′ eutectic microstructure in earlier research on the solidification behaviors of welded superalloys with relatively high carbon contents by Ojo and Sidhu et al. [23,24]. Therefore, the understanding of the solidification path of the LRDS SX alloys with a low carbon content requires careful investigations into the interdendritic microstructures. In order to assess the solidification sequence, the LRDS microstructures at interdendritic regions of samples A and B are shown in Fig. 2a and b, respectively. The branching morphology of the MC carbides can be seen more clearly in Fig. 2a. It is interesting that a small number of γ-γ′ eutectic islands are found, as shown in the middle-top of the figure. Compared with sample A, the branches of the MC carbides in sample B are not visible and no γ-γ′ eutectic is observed. To better understand the solidification sequence, Fig. 2c illustrates the back-scattered electron (BSE) image of the interdendritic region same as Fig. 2a. As can be seen, the γ-γ′ eutectic is very close to (or even locates at) interdendritic shrinkage defects, which implies that the initiation of these reactions indeed lags behind the precipitation of the MC carbides. In particular, a few MC carbides are observed to be wrapped in these γ-γ′ eutectic microstructures and this characteristic can be seen more clearly in a highmagnification view (Fig. 2d). However, such a microstructure is not found in the conventional DS SX substrate with a low cooling rate though the MC carbides also adjoin the γ-γ′ eutectic islands (see Fig. 2e). According to above careful analyses of solidification microstructure, the solidification paths of the SX superalloy containing minor carbon under various LRDS conditions can be proposed by using a simplified pseudo-ternary composition triangle of γ, γ′, and MC similar to previous research [23,25] (Fig. 3a). Step I. Solidification commences with the well-known phase transformation, L → γ0. As a result, the solidification microstructure of the LRDS SX exhibits mostly γ dendrites epitaxially growing from the SX
Fig. 1. OM images showing the solidification microstructures of the LRDS SX samples: (a) A and (b) B. SEM images of the cross-section dendritic morphologies of samples (c) A and (d) B.
60
Y.-J. Liang et al. / Scripta Materialia 127 (2017) 58–62
Fig. 2. SEM images showing the solidification microstructures at interdendritic regions of the LRDS SX samples: (a) A and (b) B. (c) SEM BSE image of the interdendritic region same as (a). (d) The high-magnification view of the region of the red box in (a). (e) SEM image of the interdendritic microstructures of the SX substrate. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
substrate. As the solidification proceeds, the liquid will be enriched in the elements with partition coefficients, k, less than unity, including Ti, Ta, and C. In addition, the kC (~ 0.2 [26]) is much less than the kTi and kTa (~ 0.7 [26–29]), which indicates that the enrichment rate of C atoms in liquid is higher. Such enrichments will lead to the composition preferentially reaching the tri-phase equilibrium among liquid, γ, and MC, rather than liquid, γ, and γ′, as the path C0e schematically shown
in Fig. 3a. It should be noted that when the solidification velocity is high enough, partition coefficients should be considered to be velocity-dependent rather than constant [20], which can lead to the reduction in the difference between kC and kTa/Ti. Nevertheless, the enrichment rate of C atoms in liquid is still higher than that of Ti/Ta atoms and similar γ-MC eutectic reactions have also been reported under higher cooling rates (see the following).
Fig. 3. (a) The liquidus projection for a pseudo-ternary composition triangle of γ, γ′, and MC. Schematics of the proposed solidification sequence under different LRDS conditions: (b) a relatively high cooling rate (A), and (c) a rather high cooling rate (B). (d) A pseudo-binary peritectic phase diagram of γ-MC and γ-γ′.
Y.-J. Liang et al. / Scripta Materialia 127 (2017) 58–62
Step II. When the composition reaches the point e in Fig. 3a, the solidification continues with a γ-MC eutectic reaction: L → γE + MC, the path eP, according to the DSC results. This eutectic-type solidification reaction can explain the formation of MC carbides with a branching morphology. Similar eutectic-type solidification reaction of MC carbides was also observed under higher cooling rates (~6.2 × 104 and 1.4 × 105 K/s) in earlier research on a superalloy with a lower C content (b0.01%) by Wang et al. [21], reconfirming that with a relatively high C content (0.015%) this reaction is indeed after the epitaxial directional growth of single-phase γ dendrites. Such γ-MC eutectic reactions have also been widely reported in the superalloys with higher C contents based on the analyses of DSC curves [3–6] or their morphology, location, and compositions [22–24]. Step III—Absent under rather high cooling rates. On further solidification, the point P in Fig. 3a represents that the solidification appears to continue with a peritectic reaction: L + MC → γ + γ′. Although Ojo et al. [23] have suggested that this solidification step is not a ternary eutectic reaction and may be a peritectic transformation by theoretical (phase rule) analyses in their previous study on the solidification behavior of IN738LC superalloy, they do not provide any evidence to confirm this process. Fortunately, peritectic transformation requires the diffusion of solute atoms in peritectic (solid-state) phase, meaning that it will be suppressed under very high solidification rates [30]. As shown in Fig. 2a, c, and d, a few residual MC carbides are observed to be wrapped in γ-γ′ eutectic microstructures under a high cooling rate (N 1000 K/s). Interestingly, this phenomenon does not exist in both the LRDS sample B with a very high cooling rate (N 17,000 K/s) (Fig. 2b) and the conventional DS substrate with a rather low cooling rate (Fig. 2e). All these observations have adequately verified that this transformation is peritectic-type. Because of very high diffusivity of C atoms, this peritectic process is complete under low cooling rates (substrate). As a result, only the MC carbides adjoining the γ-γ′ eutectic islands can be seen (see Fig. 2e and Ref. [23]). As the cooling rate increases, solidification time is shortened and the peritectic process becomes difficult to be completed (sample A). Accordingly, a few residual MC carbides will be wrapped in γ-γ′ microstructures (Fig. 2a, c, and d). Under very high cooling rates, the solidification time is rather short; therefore the diffusion of C atoms in solid is quite limited and the peritectic process is very difficult to conduct (sample B). Consequently, only the primary MC carbides exist and no γ-γ′ microstructure can be observed (Fig. 2b). Step IV—Absent under rather high cooling rates. The solidification is finally completed by forming γ-γ′ eutectic: L → γ + γ′. Strictly speaking, this process may also be peritectic-type [31–33]. However, understanding the last stage phase transformation requires further in-depth and careful investigations and this is not the emphasis of this work. Therefore, we retain to this well-known term ‘γ-γ′ eutectic’ to address such a process and resulting microstructure consistent with previous research [33–35]. Fig. 3b and c schematically summarizes the solidification sequences under different LRDS conditions. As can be seen, solidification initiates at TL with the epitaxial directional growth of columnar γ dendrites, L → γ0, and continues with a γ-MC eutectic reaction resulting in the formation of MC carbides, L → γE + MC, when the temperature declines below TMC. Subsequently, if the cooling rate is low enough to allow the sufficient diffusion of C atoms in solid and γ′ formers in liquid, a peritectic transformation will occur at TP by consuming the primary MC carbides, L + MC → γ + γ′, and solidification will be finally completed by forming the γ-γ′ eutectic microstructure: L → γ + γ′ (Fig. 3b and Fig. 2a). Otherwise, only Step I and II phase transformations occur under rather high cooling rates, as shown in Fig. 3c and Fig. 2b. The difference in solidification paths is easy to be explained by a pseudo-binary phase diagram of γ-MC and γ-γ′ (Fig. 3d). If cooling rates are not high enough, solidification will continue with the equilibrium liquidus after point P (solid line) whereas for very high cooling rates it will continue with a nonequilibrium transformation resulting in continuous formation of
61
γ-MC eutectic because of very low diffusivities of C atoms in solid and γ′ formers in liquid (dash line extrapolating from eP). Such nonequilibrium transformations have been widely reported in earlier research on peritectic alloys under RS conditions [36]. More importantly, these findings demonstrate a feasible method for preparing the SX alloys without γ-γ′ eutectic by controlling the solidification rate. Furthermore, MC carbides become non-branching, small rods (or layers) within interdendritic regions under very high cooling rates (Fig. 1d and Fig. 2b). This is likely due to a divorced eutectic reaction caused by limited diffusion of C atoms in liquid. Interestingly, these subsequent diffusion phase transformations occurring at the last stage solidification appear to be able to fill up the shrinkage defects (see Fig. 2c), implying that a rather high cooling rate that will suppress them is not always beneficial to the solidification microstructure. This may be another reason why the porosity in sample B is visibly higher than that in sample A in addition to the effect of solidification time (Fig. 1). In view of this, a good solidification microstructure requires an appropriate control of solidification rate. In summary, the solidification path of the SX nickel-base superalloys containing minor carbon was investigated under various LRDS conditions. It was found that solidification initiates with the epitaxial directional growth of γ dendrites, L → γ0, and continues with a γ-MC eutectic reaction resulting in the formation of MC carbides, L → γE + MC. Subsequently, only when the cooling rate is low enough to allow the diffusion of C atoms in solid and γ′ formers in liquid, will a peritectic transformation occur by consuming the MC carbides, L + MC → γ + γ′, and solidification be finally completed by forming the well-known γ-γ′ eutectic microstructure: L → γ + γ′. By controlling the cooling rate during LRDS processing, above peritectic transformation was first experimentally verified and a feasible method for preparing the SX alloys without γ-γ′ eutectic was demonstrated. Nevertheless, these subsequent diffusion phase transformations are able to fill up the shrinkage defects, implying that a rather high cooling rate that will suppress them is not always beneficial to the solidification microstructure. A good solidification microstructure requires an appropriate control of solidification rate. The understanding of the solidification sequence of these LRDS carbon-containing SX superalloys contributes to the prediction and control of the morphology and quantity of carbides and further understanding of the solidification behavior in the future. Acknowledgements This work was supported by the National High Technology Research and Development Program of China (Grant No. 2014AA041701) and the Joint Funds of NSFC-Liaoning (Grant No. U1508213). The authors thank Drs. ZHANG Shu-Quan and TIAN Xiang-Jun for laser processing. References [1] R.C. Reed, The Superalloys: Fundamentals and Applications, Cambridge university Press, 2006. [2] M.J. Donachie, S.J. Donachie, Superalloys: A Technical Guide, Second Edition ASM International, Materials Park, Ohio, 2002. [3] S. Tin, T.M. Pollock, W.T. King, in: T.M. Pollock, R.D. Kissinger, R.R. Bowman, K.A. Green, M. McLean, S. Olson, J.J. Scbirra (Eds.), Superalloys 2000, TMS (The Minerals, Metals &Materials Society) 2000, pp. 201–210. [4] S. Tin, T.M. Pollock, Metall. Mater. Trans. A 34A (2003) 1953–1967. [5] S. Tin, T.M. Pollock, W. Murphy, Metall. Mater. Trans. A 32A (2001) 1743–1753. [6] S. Tin, T.M. Pollock, J. Mater. Sci. 39 (2004) 7199–7205. [7] R. Acharya, R. Bansal, J.J. Gambone, S. Das, Metall. Mater. Trans. B Process Metall. Mater. Process. Sci. 45 (2014) 2247–2261. [8] R. Acharya, R. Bansal, J.J. Gambone, S. Das, Metall. Mater. Trans. B Process Metall. Mater. Process. Sci. 45 (2014) 2279–2290. [9] T.D. Anderson, J.N. DuPont, T. DebRoy, Metall. Mater. Trans. A 41 (2009) 181–193. [10] T.D. Anderson, J.N. DuPont, T. DebRoy, Acta Mater. 58 (2010) 1441–1454. [11] M. Gäumann, C. Bezençon, P. Canalis, W. Kurz, Acta Mater. 49 (2001) 1051–1062. [12] W. Liu, J.N. DuPont, Acta Mater. 52 (2004) 4833–4847. [13] W. Liu, J.N. DuPont, Acta Mater. 53 (2005) 1545–1558. [14] J.M. Vitek, Acta Mater. 53 (2005) 53–67. [15] Y.-J. Liang, H.-M. Wang, Mater. Des. 102 (2016) 297–302.
62
Y.-J. Liang et al. / Scripta Materialia 127 (2017) 58–62
[16] Y.-J. Liang, A. Li, X. Cheng, X.-T. Pang, H.-M. Wang, J. Alloys Compd. 688 (2016) 133–142. [17] Y. Xiong, W.H. Hofmeister, Z. Cheng, J.E. Smugeresky, E.J. Lavernia, J.M. Schoenung, Acta Mater. 57 (2009) 5419–5429. [18] M. Gremaud, J.-D. Wagnière, A. Zryd, W. Kurz, Surf. Eng. 12 (1996) 251–259. [19] Y.-J. Liang, X. Cheng, H.-M. Wang, Acta Mater. 118 (2016) 17–27. [20] W. Kurz, D.J. Fisher, Fundamentals of Solidification, Trans Tech Publications, Switzerland, 1998. [21] H.M. Wang, J.H. Zhang, Y.J. Tang, Z.Q. Hu, N. Yukawa, M. Morinaga, Y. Murata, Mater. Sci. Eng. A 156 (1992) 109–116. [22] L. Liu, F. Sommer, Scr. Metall. Mater. 30 (1994) 587–591. [23] O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Metall. Mater. Trans. A 37A (2006) 421–433. [24] R.K. Sidhu, O.A. Ojo, M.C. Chaturvedi, Metall. Mater. Trans. A 38A (2007) 858–870. [25] B. Radhakrishnan, R.G. Thompson, Metall. Trans. A 22A (1991) 887–902. [26] J.N. DuPont, J.C. Lippold, S.D. Kiser, Welding Metallurgy and Weldability of Nickelbase Alloys, John Wiley & Sons, Hoboken, New Jersey, 2009.
[27] S.-M. Seo, H.-W. Jeong, Y.-K. Ahn, D.W. Yun, J.-H. Lee, Y.-S. Yoo, Mater. Charact. 89 (2014) 43–55. [28] A. Thirumalai, A. Akhtar, R.C. Reed, Mater. Sci. Technol. 22 (2006) 1–13. [29] R.A. Hobbs, S. Tin, C.M.F. Rae, R.W. Broomfield, C.J. Humphreys, in: K.A. Green, T.M. Pollock, H. Harada, T.E. Howson, R.C. Reed, J.J. Schirra, S. Walston (Eds.), Superalloys 2004, TMS (The Minerals, Metals & Materials Society) 2004, pp. 819–825. [30] D.E. Laughlin, K. Hono, Physical Metallurgy, Elsevier, Fifth ed, 2014. [31] J.H. Lee, J.D. Verhoeven, J. Phase Equilib. 15 (1994) 136–146. [32] J.H. Lee, J.D. Verhoeven, J. Cryst. Growth 143 (1994) 86–102. [33] N. Warnken, J. Phase Equilib. Diffus. 37 (2016) 100–107. [34] A. Heckl, R. Rettig, S. Cenanovic, M. Göken, R.F. Singer, J. Cryst. Growth 312 (2010) 2137–2144. [35] N. Warnken, D. Ma, M. Mathes, I. Steinbach, Mater. Sci. Eng. A 413-414 (2005) 267–271. [36] H.W. Kerr, W. Kurz, Int. Mater. Rev. 41 (1996) 129–164.