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Experimental investigation on recrystallization mechanism of a Ni-base single crystal superalloy
Journal of Alloys and Compounds 672 (2016) 457e469
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Experimental investigation on recrystallization mechanism of a Ni-base single crystal superalloy Zhonglin Li, Qingyan Xu*, Baicheng Liu Key Laboratory for Advanced Materials Processing Technology (Ministry of Education), School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 July 2015 Received in revised form 14 February 2016 Accepted 16 February 2016 Available online 23 February 2016
Experimentation by TEM, EBSD and optical microscope is used to understand recrystallization in a Nibased single crystal superalloy. Hot compression is employed at different temperatures to provide driving force for recrystallization. Recrystallization sensitivity for this investigated alloy is provided. The results indicate that deformation temperature, as well as annealing conditions, has a great influence on recrystallization behavior. Samples deformed around 980 C have the highest propensity for recrystallization, and stacking faults can facilitate the recrystallization process by themal twinning nucleation. Microstructural observation shows that as-cast inhomogeneity plays a significant role in the microstructure evolution, especially below g0 phase solvus. Recrystallization nucleates first and grow rapidly in the dendritic arms. In the interdendritic regions, thermal twinning play the dominant role, and small grains remain after recrystallization is completed. Grain coarsening is rather difficult owing to abundant twinning grain boundaries. The eutectics in the IDRs can undergo recrystallization by themselves in the case of high plastic strains, and impede the grain boundaries at low plastic strains. © 2016 Elsevier B.V. All rights reserved.
Keywords: Recrystallization Single-crystal superalloys Deformation temperature As-cast microstructure Thermal twinning
1. Introduction Nickel-based single-crystal (SX) superalloys have been widely used due to excellent creep resistance for blades in aero-engines and land-based turbines, allowing higher turbine entry temperatures [1]. However, the geometry of turbine blades is becoming increasingly complicated, resulting in greater difficulty in casting the SX components [2e6]. Recrystallization (RX) can occur during solution treatment. Due to the complete removal of grain boundary strengthening elements such as C, B and Zr, the presence of recrystallized grains will be detrimental to the material's performance, particularly fatigue and creep resistance [7e11]. RX, which is intolerant in single crystal components, may arise during heat treatment. It can be ascribed to the plastic deformation during manufacturing process [12,13]. Plastic deformation comes from numerous sources, such as contraction stress during solidification, removal of shell and cores, grinding the airfoil, etc. RX was noticed both in the vicinity of a slight gibbosity on the surface and the rooteairfoil transition area, as shown in Fig. 1. The new grains can arise during solution heat treatment even after relatively small
* Corresponding author. E-mail address: [email protected] (Q. Xu). http://dx.doi.org/10.1016/j.jallcom.2016.02.149 0925-8388/© 2016 Elsevier B.V. All rights reserved.
degrees of plastic deformation [14]. RX in superalloys has been studied for decades. However, the main focus on RX has been put to wrought [15,16], powder metallurgy [17e19], and oxide dispersion strengthened superalloys [20]. During the last 10 years, much work [21e23] has been conducted to investigate RX in SX nickel-base superalloys in the open literature. The influence of annealing conditions [24,25], orientational dependence [26,27] and microstructural features [22,28e30] (g0 , carbides and eutectics) has been experimentally studied by a lot of researchers. In addition, predictions [12,14] and microstructural evolution [27,31] of RX have also been performed by modelling. The strategies to characterize primary RX in such materials require, at first, a well-defined deformation step to provide driving force. Mostly, indentation with different indenter geometries and shot peening has been employed for this purpose (see Table 1). RX microstructure after annealing was observed, and RX nucleation and grain growth in SX nickel-based superalloys were researched and discussed by these kinds of deformation methods. However, such deformation methods cause a high local dislocation density in the deformed regions, and it is difficult to determine the amount of plastic strains. Therefore, some conclusions are contrary as a result of the inhomogeneous distribution of the stored energy [22,28,32]. Additionally, the sensitivity of RX to annealing temperatures and
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Fig. 2. The schematic of test pieces.
2. Experimental 2.1. Materials
Fig. 1. Recrystallization Phenomena on turbine airfoils after solution heat treatment (a) in the vicinity of a slight gibbosity on the surface; (b) the rooteairfoil transition zone.
One Chinese second-generation SX superalloy is employed. The chemical composition is shown in Table 2. The master alloy is melted in a vacuum induction furnace, and cylindrical bars are cast in Bridgman furnaces with their longitudinal axes in the [001] di-
Table 1 Deformation methods used in earlier works on RX of SX nickel-base superalloys. Alloys
Deformation Methods
Reference
Ni8.5Cr5.0Co9.5W2.8Ta5.5Al2.2Ti (wt. %) Ni9Cr10Co7W2Mo4Al3Ti3Ta (wt. %) Ni13Cr4Co8(Al þ Ti)12(Ta þ W þ Mo) (wt. %) CMSX-4 CMSX-11B, PWA1483, SRR99,CMSX-6 Ni5.5Al8.5Cr5.0Co2.8Ta2.2Ti9.5W(atom %) CMSX-6 DD6 CMSX-4 CMSX-4 CMSX-2
Shot peening Shot peening Indentation (sphere) Indentation(sphere) Compression and indentation(pyramid) Indentation (V-shaped), grit blasting Indentation, grit blasting Shot peening Indentation(pyramid and sphere) Indentation(sphere), tensile Shot peening
[7] [21] [26] [27] [28] [33] [29,32] [30] [34] [35] [36,37]
Table 2 Nominal chemical composition of the investigated alloy [38]. Element
Cr
Co
Mo
W
Ta
Re
Nb
Al
Hf
Ni
wt.%
4.3
9
2
8
7.5
2
0.5
5.6
0.1
Balance
plastic strains can hardly be provided through indentation and shot peening. The present work aims to obtain a comprehensive understanding of RX in a Nickel-based SX superalloy. Hot compression is used to obtain a uniform distribution of plastic strains and stored energy. Deformation at different temperatures is induced since plastic deformation can occur at different temperatures during solidification and cooling of castings. Different annealing temperatures are employed. Deformation dislocations are examined using TEM. RX microstructure is observed by electron back-scatter diffraction (EBSD) and optical microscope. The influence of deformation temperature, annealing temperature and plastic amount, as well as the sensitivity of RX, is investigated and discussed. Furthermore, the effect of as-cast dendritic microstructure on RX microstructure evolution is also discussed in this research.
Table 3 Experimental details used in this research. Sample Group
Deformation Temperature/ C
Plastic Strains
Annealing Temperatures/ C
1# 2# 3# 4# 5# 6# 7# 8# 9# 10#
20 650 700 760 850 980 1070 1100 1150 1200
0e8% 5%, 22% 3%e8%, 22% 5%, 21% 3%e8%, 21% 0e12%, 21% 0~12%, 22% 4%e8%, 22% 4%e9%, 22% 4%e12%, 24%
1260e1310, STa 1260e1310 1240e1310, STa 1240e1310 12401310, STa 1250e1300, STa 12701310, STa 12601300 12801290 12601310
a
ST means standard treatment (1290ºC/1h þ 1300ºC/2h þ 1315ºC/4h).
rection with less than 15 deviation. (The crystallographic plane was determined by X-ray diffraction.) As-cast alloy exhibits dendritic characteristics, with a coherent two-phase microstructure (g matrix and g0 precipitate). Fine and regular cubic g0 phases are present in the dendritic arms (DAs), while coarse and irregular cubic g0 phases emerged in the interdendritic regions (IDRs). Bulky
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Fig. 3. Summary of recrystallization sensitivity at typical deformation temperatures (a) 20 C; (b) 700 C; (c) 850 C; (d) 980 C; (e) 1100 C; (f) 1200 C.
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Table 4 Critical annealing temperatures for RX of compressed samples with 5% plastic strain. Deformation Temperature/ C
Critical Annealing Temperature Range for RX/ C
20 650 700 760 850 980 1070 1100 1150 1200
1305e1310 1310e1315 1315e1320 1310e1315 1270e1280 1250e1260 1260e1270 1270e1280 1280e1285 1310e1315
Fig. 4. Deformed microstructure with 5.9% plastic strain at 980 C; red arrows show abundant stacking faults. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
using EDM, and tubed in silica glass under inert argon atmosphere to avoid oxidation. The tubed samples are annealed at elevated temperatures. The annealing temperatures are shown in Table 3. The annealing time is at least 5 min. Cooling in air was employed. In this research, optical microscope and electron back-scattered diffraction (EBSD) technique are used to examine whether RX occurs. RX microstructure on the middle section face (Fig. 2) is observed and analyzed. For optical observation, samples are mechanically ground and polished, and then chemically etched using a solution of 5 ml HCl þ 2g CuSO4 þ 2 ml HF þ 23.5 ml H2O for 20s. For EBSD observation, the samples are mechanically ground and then electropolished using perchloric acid (10% in volume) and dehedrated ethanol (90% in volume). An Oxford detector is employed within a MIRA3 LMH field emission gun scanning electron microscope. In this experiment, g and g0 are detected as the same phase due to their highly similar lattice parameters. Data are analyzed using the commercial HKL CHANNEL5 software and self-developed Matlab codes, assuming a FCC Nisuperalloy structure with a lattice parameter of 0.357 nm. In this research, the inverse pole figure (IPF), grain average misorientation (GAM) and kernel average misorientation (KAM) map are employed to analyze RX microstructure in SX superalloys. Each pixel in the IPF map shows the crystal direction that is aligned with the chosen direction of reference (the normal direction of sample is used as the reference direction in this research). GAM map represents the average misorientation between each neighboring pair of measurement points, assigning the same value to every point contained within the grain [47]. KAM map shows the average misorientation within a kernel which is a set of points of prescribed size surrounding the scan point of interest. The 1st nearestneighbor kernel is used in this research. It should be noted that GAM and KAM are draw by self-developed matlab codes, and they are both sensitive to the step size. To make different microstructure comparable, 2 mm is employed as step size in GAM and KAM map in this research. 3. Results and discussion 3.1. Influence of deformation temperature
g/g0
eutectic structure appeared in the IDRs, and approximately 6 vol% eutectics are present in the as-cast alloy. 2.2. Hot compression details As-cast samples are cut into smaller cylinders of 10 mm in length and 6 mm in diameter (as shown in Fig. 2) using electro discharge machining (EDM) to avoid additional harms [11]. Uniaxial compressive testing is carried out on the Gleeble 1500D (thermophysical simulator) at a strain rate of 3 103 s1 to provide driving force for RX. Graphite gaskets are placed between the test piece and the compression head for lubrication. The length of the compressed cylinders is measured again to obtain the final logarithmic plastic strain. The deformation temperatures and induced plastic strains are shown in Table 3. The total sample number is 143. The foil is cut perpendicular to the [001] direction from the middle cross section of sample deformed at 980 C. After mechanically grinding, the thin foils were electrochemically thinned in a solution of 10 ml perchloric acid and 90 ml ethanol at a current of 30 mA and temperature of 25 C. The deformed microstructure is subsequently examined by TEM. 2.3. Annealing and recrystallization observation Each compressed sample is then cut into two same cylinders
Fig. 3 shows the RX sensitivity at six typical deformation temperatures. It indicates whether the strained samples undergo recrystallization or not at different annealing temperatures. Logarithmic plastic strain is used in this research. Data at the annealing temperature of 1315 C (green vertical lines) represents the RX sensitivity at standard solution treatment. Plastic strain of 5% (purple horizontal line) is firstly considered. Table 4 summarizes critical annealing temperature range for RX at different deformation temperatures. Obviously, there exists a big difference between different deformation temperatures. Critical temperature for RX of deformed samples at 980 C is the lowest. This is different from previous work [14,35] which held that samples deformed at higher temperatures have a higher tendency for RX. However, limited deformation temperature range was chosen and critical annealing temperatures for RX were not given in their research. Higher driving force for RX can be induced at elevated temperatures, especially above 1000 C. This seems to mean that RX can nucleate easily at elevated temperatures. However, the experimental results in this research show that it becomes more difficult for RX to occur with increasing temperature above 980 C. There should be other factors influencing RX in SX superalloys. Our previous work shows that stacking faults can facilitate RX nucleation in SX superalloys at low plastic strains [39,40]. In this investigated alloys, a large amount of stacking faults can present when deformed at 980 C (Fig. 4). This is the reason why samples
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Fig. 5. Comparison of RX microstructure of samples annealed at different temperatures for 10min (deformed with around 5% plastic strain at 850 C) (a) 1280 C; (b) 1290 C; (c) 1295 C; (d) 1300 C; (e) 1305 C; (f) 1310 C. The color key map is inserted in the right bottom.
deformed at around 980 C have the highest propensity for recrystallization. Another interesting finding is that samples deformed at room temperature have a little higher propensity for recrystallization than those at around 700 C. This can be concluded not only from the critical annealing temperatures (Table 4), but also from the experimental results under standard solution treatment. For samples deformed at 700 C, RX does not occur even at 5% plastic strain (Fig. 3(b)). Nevertheless, RX occurs even below 2% plastic strain for samples deformed at room temperature (Fig. 3(a)). When checking the RX sensitivity at annealing temperature of 1310 C, we can also find that samples deformed at room temperature have a little higher propensity for recrystallization than those at between 650 C and 760 C. In addition, this finding can also be found from the critical temperatures for RX, as shown in Table 4.
3.2. Influence of annealing temperatures RX microstructure of deformed samples annealed at different temperatures for 10 min is shown in Fig. 5 and Fig. 6. All IPF maps in this paper employ the same color key map as that in Fig. 5. Dendrite-like recrystallized grains are present below 1300 C, showing clearly the RX grain boundary motion is retarded by the IDRs. Previous research indicates that dispersions of coherent particles will be at least four times more effective than incoherent particles, in restricting grain boundary motion [41]. It is evident that g0 particles dissolve slower in the IDRs than in the DAs [28]. This is mainly related to solutal microsegregation. The g0 forming elements, Al and Ta, are among the most strongly partitioning elements which are easily rich in the IDRs during casting [42e46]. Therefore, the g0 solvus is much higher in the IDRs. According to the
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Fig. 6. Comparison of RX microstructure of samples annealed at different temperatures for 10min (deformed with around 5% plastic strain at 980 C) (a) 1260 C; (b) 1270 C; (c) 1280 C; (d) 1285 C; (e) 1295 C; (f) 1300 C.
thermodynamic calculations by Pandat®, the g0 solvus in the DAs and the IDRs are around 1260 C and 1300 C respectively. In the g single phase, RX can nucleate and migrate rapidly. Thus, RX grains emerge as dendritic shape during the early stage under 1300 C, while almost all regions undergo RX and large grains are present above 1300 C (Fig. 5(e) and (f)). Evidently, the inhomogeneity in as-cast SX superalloys, particularly the dendritic microstructure, plays the significant role in the formation of RX. Additionally, dendrite-like RX grains become larger with increasing temperatures (Fig. 5(a)e(d), and Fig. 6) when annealing under 1300 C. This shows there are no obvious boundaries between the DAs and the IDRs. The dissolution of the g0 phase plays a great role in RX of SX nickel-based superalloys, especially in the case of low plastic strains. When the annealing temperature is below the g0 solvus (around 1260 C), RX is very difficult to occur,
even at medium plastic strains (above 20%) (Fig. 5(a), (c), (d)).
3.3. Influence of annealing time Fig. 7 shows the evolution of RX microstructure at the annealing temperature of 1280 C. In the early stage, recrystallization first nucleates and grows rapidly in the dendritic arms (Fig. 7(a)). The grain boundaries are inhibited or retarded by the g0 precipitate in the interdendritic regions. In the meanwhile, nucleation also occurs in the IDRs (Fig. 7(b)). However, it's difficult for the nucleation to grow into larger grains due to the retardation of g0 phase. Thus, many small grains remain in the IDRs after RX is completed (Fig. 7(c)). After that, small grains are merged by the grain boundary curved driving force (Fig. 7(d)). It should be noted that irregular grain boundaries remain due to the different kinetics between the
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Fig. 7. Evolution of recrystallization microstructure of samples with plastic strains of around 11% (deformed at 980 C and annealed at 1280 C): (a) 5min; (b) 10min; (c) 1h; (d) 4h.
Fig. 8. Comparison of recrystallization microstructures of samples with different plastic strains (deformed at 980 C and annealed at 1280 C for 10min) (a) 5.86%; (b) 11.03%.
DAs and the IDRs.
3.4. Influence of plastic strains Plastic deformation provides the driving force for RX. Higher plastic strains result in higher stored energy in deformed alloys. Fig. 8 show the comparison of RX microstructure of samples with different plastic strains at temperatures below solvus. For samples with a lower plastic strain, only undeveloped recrystallized grains emerge after 10 min. Recrystallized grains grow faster at higher plastic strains. In addition, a large amount of small grains start to present in the IDRs after 10 min. This indicates that higher plastic
strains not only provide higher driving force for the motion of grain boundaries, but also facilitate the recrystallization nucleation. Two compressed samples with different plastic strains are annealed at standard solution treatment, and their RX microstructure is shown in Fig. 9. Small grains merge after long time, and more grains remain in samples with higher plastic strain. This is the result of the higher nucleation rate in the early stage.
3.5. Influence of as-cast inhomogeneity To clarify the recrystallization behavior in the IDRs, two partially recrystallized samples are examined by observing the
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Fig. 9. Comparison of recrystallization microstructures of samples with different plastic strains (deformed at room temperature and annealed by standard solution treatment) (a) 1.57%; (b) 4.92%.
Fig. 10. Orientation index maps (OIM) from a partially recrystallized sample. (deformed at 650 C for 21.4% plastic strain and annealed at 1280 C for 10min) (a) IPF map; (b) GAM map; (c) FSE map; (d) KAM map.
microstructure in this area, as shown in Fig. 10 and Fig. 11. The samples are heavily and slightly strained before heat treatment, respectively. Silver lines represent small angle grain boundaries (3º15 ) in the IPF (inverse polar figure) map (Fig. 10(a) and Fig. 11(a)). Forward scattered electron (FSE) can greatly enhance diffraction contrast in imaging, and clearly reveal grains and grain boundaries. Heavily rugged areas in FSE map (Fig. 10(c) and Fig. 11(c)) represent the regions with the most serious elemental segregation in the IDRs. In addition, eutectics can also be identified in FSE map, as indicated by red arrows in Fig. 10(c) and Fig. 11(c).
Both GAM and KAM maps can provide an indication of the strain distribution in the material, and are sensitive to the step size of the measurement grid. We can find that local and average misorientation are extremely low in recrystallized regions, as shown in Fig. 10(b), (d) and Fig. 11(b), (d). In the unrecrystallized regions, these values are higher, especially in samples with higher plastic strains. Obviously, RX arises in the dendritic arms, as well as some small islands in the IDRs. In the heavily strained sample, all eutectics undergo RX completely or partially, as shown in Fig. 10. However,
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Fig. 11. OIM from a partially recrystallized sample. (deformed at 1070 C for 5.4% plastic strain and annealed at 1290 C for 10min) (a) IPF map; (b) GAM map; (c) FSE map; (d) KAM map.
RX just appears in few eutectics for the lightly strained sample (Fig. 11). This indicates that eutectics play a different role in RX of SX superalloys for samples with different plastic strains. As known, g/ g0 eutectics are soft particles, and plastic deformation can easily arise in this region. In heavily strained samples or regions, the dislocation density within g/g0 eutectics can be high enough, and the eutectics themselves undergo RX by nucleation or the motion of RX grain boundary. The dislocation density in the g/g0 eutectics is low for lightly strained samples, thus the RX grain boundary is retarded by the eutectic. This is consistent with Wang's conclusion about the influence of the eutectics [22,48]. Till now, it is still a question whether RX nucleation can occur in the IDRs except the eutectics, since grain boundary migration is difficult in this area. To clarify this issue, RX microstructure of one sample annealed at 1270 C for 2 h is observed in detail, as shown in Fig. 12. Unexpectedly, a large amount of green points appear in the IDRs (Fig. 12(a) using a step size of 10 mm). To identify these green points, the microstructure in the IDRs is carefully examined using a smaller step size, 3 mm (Fig. 12(b)) and 0.05 mm (Fig. 12(c), (d)). Clearly, twin grains appear as a lot of parallel thin strips with a width of only a few hundred nanometers. Since no deformation twinning was found in nickel-based SX superalloys [22,37,49,50], particularly at low plastic strains, RX nucleation can take place through thermal twinning in the IDRs. It has been demonstrated that annealing twinning can easily arise in materials with low stacking fault energy [51e55]. In this alloy, more twinning grains appear in the IDRs than in the DAs, due to the retarding of g0 phase in the IDRs. In addition, twinning grains are only found in the IDRs alongside developed dendritic RX grains. This indicates that RX in
the DAs can trigger the twinning grains in the IDRs at the front of RX grain boundaries, though the grain boundary motion is difficult. According to RX nucleation theory, nucleation through thermal fluctuation is impossible. So many twinning strips can't be formed by subgrain coarsening nucleation. Therefore, nucleation by thermal twinning in the front of the grain boundary takes place, and it is easier for twinning in the IDRs to move into the original matrix compared with the normal grains in the DAs. Fig. 13 shows another examples with a lot of thermal twinning in the IDRs at the annealing temperature of 1280 C. The twinning strips are wider, compared with the previous one annealed at 1270 C, though the number of twinning grains is significantly less. In general, parallel twinning strips are easier to arise when samples with higher plastic deformation are annealed at lower temperatures. To further clarify the complete RX process in the IDRs, two samples with high plastic strains after annealing at 1280 C for 4 h are examined by EBSD (Fig. 14) and optical microscope (Fig. 15). Fig. 14 shows the RX microstructure in the IDRs which can be easily observed in the FSE micrographs (the area enclosed by red lines in Fig. 14(c) (f)). As stated above, small grains remain in the IDRs after RX is completed, as shown in both Figs. 14 and 15. RX grains can be observed by optical micrographs though it is not so clear to distinguish different grains. The average grain size in the IDRs is approximately 20e30 mm, while it was larger than 100 mm in the DAs. As shown in Fig. 14(b) and (e), plentiful twinning grain boundaries remain in the final RX microstructures. The mechanism of twinning is still controversial, and it could result from a growth accident or partial dislocations from the migrating grain
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Fig. 12. OIM from a partially recrystallized sample. (deformed at 980 C for 10% plastic strain and annealed at 1270 C for 2h) (a) IPF map with step size of 10 mm; (b) enlarged area of the black rectangle in (a), IPF map with step size of 3 mm; (c) enlarged area of the partial recrystallization, IPF map with step size of 0.05 mm; (d) grain boundary map superimposed with kikuchi band contrast (twinning grain boundaries represented by red lines).
Fig. 13. OIM from a partially recrystallized sample. (deformed at 1100 C for 7.8%% plastic strain and annealed at 1280 C for 10min):(a1) IPF map with step size of 10 mm; (b1) enlarged area of the black rectangle in (a1), IPF map with step size of 2 mm; (c1) enlarged area of the black rectangle in (b1), IPF map with step size of 0.658 mm; (a2) (b2) (c2) corresponding FSE maps with (a1) (b1) (c1).
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Fig. 14. OIM from two completely recrystallized samples.(a) (b) (c) deformed at 1100 C for 22.6% plastic strain and annealed at 1280 C for 4h; (d) (e) (f) deformed at 1200 C for 22.7% plastic strain and annealed at 1280 C for 4h; (a) (d) Euler angle map superimposed with grain boundaries; (b) (e) grain boundaries map with red lines representing twinning boundaries; (c) (f) FSE maps with red lines representing interdendritic regions (IDRs). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
Fig. 15. Optical microstructure from two completely recrystallized samples. These two samples correspond to those in Fig. 14. (a) (b) (c) deformed at 1100 C for 22.6% plastic strain and annealed at 1280 C for 4h; (d) (e) (f) deformed at 1200 C for 22.7% plastic strain and annealed at 1280 C for 4h.
boundaries. Both a gain in grain-boundary energy and a decrease in deformation stored energy play the dominant role in thermal twinning. Furthermore, thermal twinning would typically lead to the formation of rather immobile boundaries. Thus, grain coarsening is difficult owing to the presence of abundant twinning grain boundaries. In materials with FCC structure, a random twinning sequence using randomly one out of the possible 12 {111}<112>variants available would lead to nearly any orientation possible. In addition, it is assumed that there must be the preference of certain twin sequences. Thus, different orientations can arise though a lot of twinning grains with only one orientation are present at the early stage in the IDRs [56].
4. Conclusions RX in a nickel-based single crystal superalloy is investigated systematically using TEM, EBSD and optical microscope. Hot compression is employed to provide driving force for RX. The following conclusions can be drawn. 1) Deformation temperature has a great influence on RX behavior. For DD6 alloy, samples deformed around 980 C have the highest propensity for RX, because the presence of stacking faults can facilitate the RX nucleation. 2) The higher the deformation amount, the faster the grain boundary moves and the higher the RX nucleation rate.
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3) The g0 phase retards the movement of RX grain boundaries. When the annealing temperature is below the g0 solvus, as-cast dendritic microstructure plays a significant role in RX as a result of the different dissolution of g0 phase in the DAs and IDRs. 4) RX nucleates first and grow rapidly in the DAs. In the IDRs, thermal twinning play the dominant role, and small grains remain after RX is completed. Grain coarsening is rather difficult owing to abundant twinning grain boundaries. Acknowledgments This research was funded by the National Basic Research Program of China (No. 2011CB706801) and National Natural Science Foundation of China (Nos. 51171089 and 51374137). The authors are grateful for this support. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jallcom.2016.02.149. References [1] T. Pollock, S. Tin, Nickel-based superalloys for advanced turbine engines: chemistry, microstructure, and properties, J. Propul Power 22 (2006) 361e371. [2] T.M. Pollock, W.H. Murphy, The breakdown of single-crystal solidification in high refractory nickel-base alloys, Metall. Mater. Trans. A 27 (1996) 1081e1094. [3] T.M. Pollock, W.H. Murphy, E.H. Goldman, D.L. Uram, J.S. Tu, Grain defect formation during directional solidification of nickel base single crystals, in: S.D. Antolovich, R.W. Stusrud, R.A. MacKay, D.L. Anton, T. Khan, R.D. Kissinger, D.L. Klarstrom (Eds.), Superalloys 1992, The Minerals, Metals and Materials Society, Warrendale, PA, 1992, pp. 125e134. [4] D. Pan, Q. Xu, B. Liu, J. Li, H. Yuan, H. Jin, Modeling of grain selection during directional solidification of single crystal superalloy turbine blade castings, JOM 62 (2010) 30e34. [5] D. Ma, A. Bührig-Polaczek, The influence of surface roughness on freckle formation in directionally solidified superalloy samples, Metall. Mater. Trans. B 43B (2012) 671e677. [6] D. Ma, Q. Wu, A. Bührig-Polaczek, Some new observations on freckle formation in directionally solidified superalloy components, Metall. Mater. Trans. B 43 (2012) 344e353. [7] B. Zhang, D. Liu, Influence of recrystallization on high-temperature stress rupture property and fracture behavior of single crystal superalloy, Mat. Sci. Eng. A 551 (2012) 149e153. [8] G. Xie, L.H. Lou, Influence of the characteristic of recrystallization grain boundary on the formation of creep cracks in a directionally solidified Ni-base superalloy, Mat. Sci. Eng. A 532 (2012) 579e584. [9] G. Xie, L. Wang, J. Zhang, L.H. Lou, Intermediate temperature creep of directionally solidified Ni-based superalloy containing local recrystallization, Mat. Sci. Eng. A 528 (2011) 3062e3068. [10] G. Xie, L. Wang, J. Zhang, L.H. Lou, Influence of recrystallization on the hightemperature properties of a directionally solidified Ni-base superalloy, Metall. Mater Trans. A 39A (2008) 206e210. [11] T. Khan, P. Caron, Y.G. Nakagawa, Mechanical behavior and processing of DS and single crystal superalloys, J. Metals (1986) 16e19. [12] Z. Li, J. Xiong, Q. Xu, J. Li, B. Liu, Deformation and recrystallization of single crystal nickel-based superalloys during investment casting, J. Mater Process Tech. 217 (2015) 1e12. [13] Z. Li, Q. Xu, J. Xiong, J. Li, B. Liu, Simulation and experiments on plasticity and recrystallisation in SX superalloys by investment casting, Mater Res. Innov. 18 (2014) S4eS331. [14] C. Panwisawas, H. Mathur, J. Gebelin, D. Putman, C.M.F. Rae, R.C. Reed, Prediction of recrystallization in investment cast single-crystal superalloys, Acta Mater 61 (2013) 51e66. [15] A. Porter, B. Ralph, The Recrystallization of nickel-based superalloys, J. Mater Sci. 16 (1981) 707e713. [16] D.S. Weaver, Recrystallization and grain-growth behavior of a nickel-base superalloy during multi-hit deformation, Scr. Mater 57 (2007) 1044e1047. [17] M. Dahlen, L. Winberg, Influence of gamma'-precipitation on the recrystallization of a nickel-base superalloy, Acta Metall. 28 (1980) 41e50. [18] A.J. Porter, B. Ralph, Recrystallization of a nickel-base superalloy e kinetics and microstructural development, Mater. Sci. Eng. 59 (1983) 69e78. [19] D.D. Whitis, Recovery and recrystallization after critical strain in the nickelbased superalloy rene 88DT., Superalloys 2004, Minerals Metals Mater. Soc. Warrendale (2004) 391e400. [20] M. Kouichi, Directional Recrystallisation in Mechanically Alloyed ODS Nickel
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Liu, Microstructure simulation on recrystallization of an as-cast nickel based single crystal superalloy, Comp. Mater Sci. 107 (2015) 122e133. [32] D. Goldschmidt, U. Paul, P.R. Sahm, Porosity clusters and recrystallization in single-crystal components, in: S.D. Antolovich, R.W. Stusrud, R.A. MacKay, D.L. Anton, T. Khan, R.D. Kissinger, D.L. Klarstrom (Eds.), Superalloys 1992, TMS, Warrendale, PA, 1992, pp. 155e164. [33] S.D. Bond, J.W. Martin, Surface recrystallization in a single-crystal nickelbased superalloy, J. Mater Sci. 19 (1984) 3867e3872. [34] M. Okazaki, T. Hiura, T. Suzuki, Effect of local cellular transformation on fatigue small crack growth in CMSX-4 and CMSX-2 at high temperature. e For refurbishment technology, in: T.M. Pollock, R.D. Kissinger, R.R. Bowman, K.A. Green, M. McLean, S. Olson, J.J. Schirra (Eds.), Superalloys 2000, TMS, Warrendale, PA, 2000, pp. 505e514. [35] D.C. Cox, B. Roebuck, C.M.F. Rae, R.C. Reed, Recrystallisation of single crystal superalloy CMSXe4, Mater. Sci. Technol. 19 (2003) 440e446. [36] C.Y. Jo, H.M. Kim, Effect of recrystallisation on microstructural evolution and mechanical properties of single crystal nickel based superalloy CMSX-2-Part 2-Creep behaviour of surface recrystallised single crystal, Mater. Sci. Technol. 19 (2003) 1671e1676. [37] C.Y. Jo, H.Y. Cho, H.M. Kim, Effect of recrystallisation on microstructural evolution and mechanical properties of single crystal nickel base superalloy CMSX-2-Part 1-Microstructural evolution during recrystallisation of single crystal, Mater. Sci. Technol. 19 (2003) 1665e1670. [38] J.R. Li, J.Q. Zhao, S.Z. Liu, M. Han, Effects of low angle boundaries on the mechanical properties of single crystal superalloy DD6, in: R.C. Reed, K.A. Green, P. Caron, T.P. Gabb, M.G. Fahrmann, E.S. Huron (Eds.), Superalloys 2008, Champion, PA, 2008, pp. 443e451. [39] S. Dash, N. Brown, An investigation of the origin and growth of annealing twins, Acta Metall. 11 (1963) 1067e1075. [40] S. Mahajan, Critique of mechanisms of formation of deformation, annealing and growth twins: face-centered cubic metals and alloys, Scr. Mater 68 (2013) 95e99. [41] R.D. Doherty, Role of interfaces in kinetics of internal shape changes, Metal Sci. 16 (1982) 1e14. [42] G.E. Fuchs, Solution heat treatment response of a third generation single crystal Ni-base superalloy, Mat. Sci. Eng. A 300 (2001) 52e60. [43] H.T. Pang, L. Zhang, R.A. Hobbs, H.J. Stone, C.M.F. Rae, Solution heat treatment optimization of fourth-generation single-crystal nickel-base superalloys, Metall. Mater. Trans. A 43 (2012) 3264e3282. [44] R.M. Kearsey, J.C. Beddoes, P. Jones, P. Au, Compositional design considerations for microsegregation in single crystal superalloy systems, Intermetallics 12 (2004) 903e910. [45] E.C. Caldwell, F.J. Fela, G.E. 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Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Experimental investigation on recrystallization mechanism of a Ni-base single crystal superalloy Zhonglin Li, Qingyan Xu*, Baicheng Liu Key Laboratory for Advanced Materials Processing Technology (Ministry of Education), School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 July 2015 Received in revised form 14 February 2016 Accepted 16 February 2016 Available online 23 February 2016
Experimentation by TEM, EBSD and optical microscope is used to understand recrystallization in a Nibased single crystal superalloy. Hot compression is employed at different temperatures to provide driving force for recrystallization. Recrystallization sensitivity for this investigated alloy is provided. The results indicate that deformation temperature, as well as annealing conditions, has a great influence on recrystallization behavior. Samples deformed around 980 C have the highest propensity for recrystallization, and stacking faults can facilitate the recrystallization process by themal twinning nucleation. Microstructural observation shows that as-cast inhomogeneity plays a significant role in the microstructure evolution, especially below g0 phase solvus. Recrystallization nucleates first and grow rapidly in the dendritic arms. In the interdendritic regions, thermal twinning play the dominant role, and small grains remain after recrystallization is completed. Grain coarsening is rather difficult owing to abundant twinning grain boundaries. The eutectics in the IDRs can undergo recrystallization by themselves in the case of high plastic strains, and impede the grain boundaries at low plastic strains. © 2016 Elsevier B.V. All rights reserved.
Keywords: Recrystallization Single-crystal superalloys Deformation temperature As-cast microstructure Thermal twinning
1. Introduction Nickel-based single-crystal (SX) superalloys have been widely used due to excellent creep resistance for blades in aero-engines and land-based turbines, allowing higher turbine entry temperatures [1]. However, the geometry of turbine blades is becoming increasingly complicated, resulting in greater difficulty in casting the SX components [2e6]. Recrystallization (RX) can occur during solution treatment. Due to the complete removal of grain boundary strengthening elements such as C, B and Zr, the presence of recrystallized grains will be detrimental to the material's performance, particularly fatigue and creep resistance [7e11]. RX, which is intolerant in single crystal components, may arise during heat treatment. It can be ascribed to the plastic deformation during manufacturing process [12,13]. Plastic deformation comes from numerous sources, such as contraction stress during solidification, removal of shell and cores, grinding the airfoil, etc. RX was noticed both in the vicinity of a slight gibbosity on the surface and the rooteairfoil transition area, as shown in Fig. 1. The new grains can arise during solution heat treatment even after relatively small
* Corresponding author. E-mail address: [email protected] (Q. Xu). http://dx.doi.org/10.1016/j.jallcom.2016.02.149 0925-8388/© 2016 Elsevier B.V. All rights reserved.
degrees of plastic deformation [14]. RX in superalloys has been studied for decades. However, the main focus on RX has been put to wrought [15,16], powder metallurgy [17e19], and oxide dispersion strengthened superalloys [20]. During the last 10 years, much work [21e23] has been conducted to investigate RX in SX nickel-base superalloys in the open literature. The influence of annealing conditions [24,25], orientational dependence [26,27] and microstructural features [22,28e30] (g0 , carbides and eutectics) has been experimentally studied by a lot of researchers. In addition, predictions [12,14] and microstructural evolution [27,31] of RX have also been performed by modelling. The strategies to characterize primary RX in such materials require, at first, a well-defined deformation step to provide driving force. Mostly, indentation with different indenter geometries and shot peening has been employed for this purpose (see Table 1). RX microstructure after annealing was observed, and RX nucleation and grain growth in SX nickel-based superalloys were researched and discussed by these kinds of deformation methods. However, such deformation methods cause a high local dislocation density in the deformed regions, and it is difficult to determine the amount of plastic strains. Therefore, some conclusions are contrary as a result of the inhomogeneous distribution of the stored energy [22,28,32]. Additionally, the sensitivity of RX to annealing temperatures and
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Fig. 2. The schematic of test pieces.
2. Experimental 2.1. Materials
Fig. 1. Recrystallization Phenomena on turbine airfoils after solution heat treatment (a) in the vicinity of a slight gibbosity on the surface; (b) the rooteairfoil transition zone.
One Chinese second-generation SX superalloy is employed. The chemical composition is shown in Table 2. The master alloy is melted in a vacuum induction furnace, and cylindrical bars are cast in Bridgman furnaces with their longitudinal axes in the [001] di-
Table 1 Deformation methods used in earlier works on RX of SX nickel-base superalloys. Alloys
Deformation Methods
Reference
Ni8.5Cr5.0Co9.5W2.8Ta5.5Al2.2Ti (wt. %) Ni9Cr10Co7W2Mo4Al3Ti3Ta (wt. %) Ni13Cr4Co8(Al þ Ti)12(Ta þ W þ Mo) (wt. %) CMSX-4 CMSX-11B, PWA1483, SRR99,CMSX-6 Ni5.5Al8.5Cr5.0Co2.8Ta2.2Ti9.5W(atom %) CMSX-6 DD6 CMSX-4 CMSX-4 CMSX-2
Shot peening Shot peening Indentation (sphere) Indentation(sphere) Compression and indentation(pyramid) Indentation (V-shaped), grit blasting Indentation, grit blasting Shot peening Indentation(pyramid and sphere) Indentation(sphere), tensile Shot peening
[7] [21] [26] [27] [28] [33] [29,32] [30] [34] [35] [36,37]
Table 2 Nominal chemical composition of the investigated alloy [38]. Element
Cr
Co
Mo
W
Ta
Re
Nb
Al
Hf
Ni
wt.%
4.3
9
2
8
7.5
2
0.5
5.6
0.1
Balance
plastic strains can hardly be provided through indentation and shot peening. The present work aims to obtain a comprehensive understanding of RX in a Nickel-based SX superalloy. Hot compression is used to obtain a uniform distribution of plastic strains and stored energy. Deformation at different temperatures is induced since plastic deformation can occur at different temperatures during solidification and cooling of castings. Different annealing temperatures are employed. Deformation dislocations are examined using TEM. RX microstructure is observed by electron back-scatter diffraction (EBSD) and optical microscope. The influence of deformation temperature, annealing temperature and plastic amount, as well as the sensitivity of RX, is investigated and discussed. Furthermore, the effect of as-cast dendritic microstructure on RX microstructure evolution is also discussed in this research.
Table 3 Experimental details used in this research. Sample Group
Deformation Temperature/ C
Plastic Strains
Annealing Temperatures/ C
1# 2# 3# 4# 5# 6# 7# 8# 9# 10#
20 650 700 760 850 980 1070 1100 1150 1200
0e8% 5%, 22% 3%e8%, 22% 5%, 21% 3%e8%, 21% 0e12%, 21% 0~12%, 22% 4%e8%, 22% 4%e9%, 22% 4%e12%, 24%
1260e1310, STa 1260e1310 1240e1310, STa 1240e1310 12401310, STa 1250e1300, STa 12701310, STa 12601300 12801290 12601310
a
ST means standard treatment (1290ºC/1h þ 1300ºC/2h þ 1315ºC/4h).
rection with less than 15 deviation. (The crystallographic plane was determined by X-ray diffraction.) As-cast alloy exhibits dendritic characteristics, with a coherent two-phase microstructure (g matrix and g0 precipitate). Fine and regular cubic g0 phases are present in the dendritic arms (DAs), while coarse and irregular cubic g0 phases emerged in the interdendritic regions (IDRs). Bulky
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Fig. 3. Summary of recrystallization sensitivity at typical deformation temperatures (a) 20 C; (b) 700 C; (c) 850 C; (d) 980 C; (e) 1100 C; (f) 1200 C.
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Table 4 Critical annealing temperatures for RX of compressed samples with 5% plastic strain. Deformation Temperature/ C
Critical Annealing Temperature Range for RX/ C
20 650 700 760 850 980 1070 1100 1150 1200
1305e1310 1310e1315 1315e1320 1310e1315 1270e1280 1250e1260 1260e1270 1270e1280 1280e1285 1310e1315
Fig. 4. Deformed microstructure with 5.9% plastic strain at 980 C; red arrows show abundant stacking faults. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
using EDM, and tubed in silica glass under inert argon atmosphere to avoid oxidation. The tubed samples are annealed at elevated temperatures. The annealing temperatures are shown in Table 3. The annealing time is at least 5 min. Cooling in air was employed. In this research, optical microscope and electron back-scattered diffraction (EBSD) technique are used to examine whether RX occurs. RX microstructure on the middle section face (Fig. 2) is observed and analyzed. For optical observation, samples are mechanically ground and polished, and then chemically etched using a solution of 5 ml HCl þ 2g CuSO4 þ 2 ml HF þ 23.5 ml H2O for 20s. For EBSD observation, the samples are mechanically ground and then electropolished using perchloric acid (10% in volume) and dehedrated ethanol (90% in volume). An Oxford detector is employed within a MIRA3 LMH field emission gun scanning electron microscope. In this experiment, g and g0 are detected as the same phase due to their highly similar lattice parameters. Data are analyzed using the commercial HKL CHANNEL5 software and self-developed Matlab codes, assuming a FCC Nisuperalloy structure with a lattice parameter of 0.357 nm. In this research, the inverse pole figure (IPF), grain average misorientation (GAM) and kernel average misorientation (KAM) map are employed to analyze RX microstructure in SX superalloys. Each pixel in the IPF map shows the crystal direction that is aligned with the chosen direction of reference (the normal direction of sample is used as the reference direction in this research). GAM map represents the average misorientation between each neighboring pair of measurement points, assigning the same value to every point contained within the grain [47]. KAM map shows the average misorientation within a kernel which is a set of points of prescribed size surrounding the scan point of interest. The 1st nearestneighbor kernel is used in this research. It should be noted that GAM and KAM are draw by self-developed matlab codes, and they are both sensitive to the step size. To make different microstructure comparable, 2 mm is employed as step size in GAM and KAM map in this research. 3. Results and discussion 3.1. Influence of deformation temperature
g/g0
eutectic structure appeared in the IDRs, and approximately 6 vol% eutectics are present in the as-cast alloy. 2.2. Hot compression details As-cast samples are cut into smaller cylinders of 10 mm in length and 6 mm in diameter (as shown in Fig. 2) using electro discharge machining (EDM) to avoid additional harms [11]. Uniaxial compressive testing is carried out on the Gleeble 1500D (thermophysical simulator) at a strain rate of 3 103 s1 to provide driving force for RX. Graphite gaskets are placed between the test piece and the compression head for lubrication. The length of the compressed cylinders is measured again to obtain the final logarithmic plastic strain. The deformation temperatures and induced plastic strains are shown in Table 3. The total sample number is 143. The foil is cut perpendicular to the [001] direction from the middle cross section of sample deformed at 980 C. After mechanically grinding, the thin foils were electrochemically thinned in a solution of 10 ml perchloric acid and 90 ml ethanol at a current of 30 mA and temperature of 25 C. The deformed microstructure is subsequently examined by TEM. 2.3. Annealing and recrystallization observation Each compressed sample is then cut into two same cylinders
Fig. 3 shows the RX sensitivity at six typical deformation temperatures. It indicates whether the strained samples undergo recrystallization or not at different annealing temperatures. Logarithmic plastic strain is used in this research. Data at the annealing temperature of 1315 C (green vertical lines) represents the RX sensitivity at standard solution treatment. Plastic strain of 5% (purple horizontal line) is firstly considered. Table 4 summarizes critical annealing temperature range for RX at different deformation temperatures. Obviously, there exists a big difference between different deformation temperatures. Critical temperature for RX of deformed samples at 980 C is the lowest. This is different from previous work [14,35] which held that samples deformed at higher temperatures have a higher tendency for RX. However, limited deformation temperature range was chosen and critical annealing temperatures for RX were not given in their research. Higher driving force for RX can be induced at elevated temperatures, especially above 1000 C. This seems to mean that RX can nucleate easily at elevated temperatures. However, the experimental results in this research show that it becomes more difficult for RX to occur with increasing temperature above 980 C. There should be other factors influencing RX in SX superalloys. Our previous work shows that stacking faults can facilitate RX nucleation in SX superalloys at low plastic strains [39,40]. In this investigated alloys, a large amount of stacking faults can present when deformed at 980 C (Fig. 4). This is the reason why samples
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Fig. 5. Comparison of RX microstructure of samples annealed at different temperatures for 10min (deformed with around 5% plastic strain at 850 C) (a) 1280 C; (b) 1290 C; (c) 1295 C; (d) 1300 C; (e) 1305 C; (f) 1310 C. The color key map is inserted in the right bottom.
deformed at around 980 C have the highest propensity for recrystallization. Another interesting finding is that samples deformed at room temperature have a little higher propensity for recrystallization than those at around 700 C. This can be concluded not only from the critical annealing temperatures (Table 4), but also from the experimental results under standard solution treatment. For samples deformed at 700 C, RX does not occur even at 5% plastic strain (Fig. 3(b)). Nevertheless, RX occurs even below 2% plastic strain for samples deformed at room temperature (Fig. 3(a)). When checking the RX sensitivity at annealing temperature of 1310 C, we can also find that samples deformed at room temperature have a little higher propensity for recrystallization than those at between 650 C and 760 C. In addition, this finding can also be found from the critical temperatures for RX, as shown in Table 4.
3.2. Influence of annealing temperatures RX microstructure of deformed samples annealed at different temperatures for 10 min is shown in Fig. 5 and Fig. 6. All IPF maps in this paper employ the same color key map as that in Fig. 5. Dendrite-like recrystallized grains are present below 1300 C, showing clearly the RX grain boundary motion is retarded by the IDRs. Previous research indicates that dispersions of coherent particles will be at least four times more effective than incoherent particles, in restricting grain boundary motion [41]. It is evident that g0 particles dissolve slower in the IDRs than in the DAs [28]. This is mainly related to solutal microsegregation. The g0 forming elements, Al and Ta, are among the most strongly partitioning elements which are easily rich in the IDRs during casting [42e46]. Therefore, the g0 solvus is much higher in the IDRs. According to the
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Fig. 6. Comparison of RX microstructure of samples annealed at different temperatures for 10min (deformed with around 5% plastic strain at 980 C) (a) 1260 C; (b) 1270 C; (c) 1280 C; (d) 1285 C; (e) 1295 C; (f) 1300 C.
thermodynamic calculations by Pandat®, the g0 solvus in the DAs and the IDRs are around 1260 C and 1300 C respectively. In the g single phase, RX can nucleate and migrate rapidly. Thus, RX grains emerge as dendritic shape during the early stage under 1300 C, while almost all regions undergo RX and large grains are present above 1300 C (Fig. 5(e) and (f)). Evidently, the inhomogeneity in as-cast SX superalloys, particularly the dendritic microstructure, plays the significant role in the formation of RX. Additionally, dendrite-like RX grains become larger with increasing temperatures (Fig. 5(a)e(d), and Fig. 6) when annealing under 1300 C. This shows there are no obvious boundaries between the DAs and the IDRs. The dissolution of the g0 phase plays a great role in RX of SX nickel-based superalloys, especially in the case of low plastic strains. When the annealing temperature is below the g0 solvus (around 1260 C), RX is very difficult to occur,
even at medium plastic strains (above 20%) (Fig. 5(a), (c), (d)).
3.3. Influence of annealing time Fig. 7 shows the evolution of RX microstructure at the annealing temperature of 1280 C. In the early stage, recrystallization first nucleates and grows rapidly in the dendritic arms (Fig. 7(a)). The grain boundaries are inhibited or retarded by the g0 precipitate in the interdendritic regions. In the meanwhile, nucleation also occurs in the IDRs (Fig. 7(b)). However, it's difficult for the nucleation to grow into larger grains due to the retardation of g0 phase. Thus, many small grains remain in the IDRs after RX is completed (Fig. 7(c)). After that, small grains are merged by the grain boundary curved driving force (Fig. 7(d)). It should be noted that irregular grain boundaries remain due to the different kinetics between the
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Fig. 7. Evolution of recrystallization microstructure of samples with plastic strains of around 11% (deformed at 980 C and annealed at 1280 C): (a) 5min; (b) 10min; (c) 1h; (d) 4h.
Fig. 8. Comparison of recrystallization microstructures of samples with different plastic strains (deformed at 980 C and annealed at 1280 C for 10min) (a) 5.86%; (b) 11.03%.
DAs and the IDRs.
3.4. Influence of plastic strains Plastic deformation provides the driving force for RX. Higher plastic strains result in higher stored energy in deformed alloys. Fig. 8 show the comparison of RX microstructure of samples with different plastic strains at temperatures below solvus. For samples with a lower plastic strain, only undeveloped recrystallized grains emerge after 10 min. Recrystallized grains grow faster at higher plastic strains. In addition, a large amount of small grains start to present in the IDRs after 10 min. This indicates that higher plastic
strains not only provide higher driving force for the motion of grain boundaries, but also facilitate the recrystallization nucleation. Two compressed samples with different plastic strains are annealed at standard solution treatment, and their RX microstructure is shown in Fig. 9. Small grains merge after long time, and more grains remain in samples with higher plastic strain. This is the result of the higher nucleation rate in the early stage.
3.5. Influence of as-cast inhomogeneity To clarify the recrystallization behavior in the IDRs, two partially recrystallized samples are examined by observing the
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Fig. 9. Comparison of recrystallization microstructures of samples with different plastic strains (deformed at room temperature and annealed by standard solution treatment) (a) 1.57%; (b) 4.92%.
Fig. 10. Orientation index maps (OIM) from a partially recrystallized sample. (deformed at 650 C for 21.4% plastic strain and annealed at 1280 C for 10min) (a) IPF map; (b) GAM map; (c) FSE map; (d) KAM map.
microstructure in this area, as shown in Fig. 10 and Fig. 11. The samples are heavily and slightly strained before heat treatment, respectively. Silver lines represent small angle grain boundaries (3º15 ) in the IPF (inverse polar figure) map (Fig. 10(a) and Fig. 11(a)). Forward scattered electron (FSE) can greatly enhance diffraction contrast in imaging, and clearly reveal grains and grain boundaries. Heavily rugged areas in FSE map (Fig. 10(c) and Fig. 11(c)) represent the regions with the most serious elemental segregation in the IDRs. In addition, eutectics can also be identified in FSE map, as indicated by red arrows in Fig. 10(c) and Fig. 11(c).
Both GAM and KAM maps can provide an indication of the strain distribution in the material, and are sensitive to the step size of the measurement grid. We can find that local and average misorientation are extremely low in recrystallized regions, as shown in Fig. 10(b), (d) and Fig. 11(b), (d). In the unrecrystallized regions, these values are higher, especially in samples with higher plastic strains. Obviously, RX arises in the dendritic arms, as well as some small islands in the IDRs. In the heavily strained sample, all eutectics undergo RX completely or partially, as shown in Fig. 10. However,
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Fig. 11. OIM from a partially recrystallized sample. (deformed at 1070 C for 5.4% plastic strain and annealed at 1290 C for 10min) (a) IPF map; (b) GAM map; (c) FSE map; (d) KAM map.
RX just appears in few eutectics for the lightly strained sample (Fig. 11). This indicates that eutectics play a different role in RX of SX superalloys for samples with different plastic strains. As known, g/ g0 eutectics are soft particles, and plastic deformation can easily arise in this region. In heavily strained samples or regions, the dislocation density within g/g0 eutectics can be high enough, and the eutectics themselves undergo RX by nucleation or the motion of RX grain boundary. The dislocation density in the g/g0 eutectics is low for lightly strained samples, thus the RX grain boundary is retarded by the eutectic. This is consistent with Wang's conclusion about the influence of the eutectics [22,48]. Till now, it is still a question whether RX nucleation can occur in the IDRs except the eutectics, since grain boundary migration is difficult in this area. To clarify this issue, RX microstructure of one sample annealed at 1270 C for 2 h is observed in detail, as shown in Fig. 12. Unexpectedly, a large amount of green points appear in the IDRs (Fig. 12(a) using a step size of 10 mm). To identify these green points, the microstructure in the IDRs is carefully examined using a smaller step size, 3 mm (Fig. 12(b)) and 0.05 mm (Fig. 12(c), (d)). Clearly, twin grains appear as a lot of parallel thin strips with a width of only a few hundred nanometers. Since no deformation twinning was found in nickel-based SX superalloys [22,37,49,50], particularly at low plastic strains, RX nucleation can take place through thermal twinning in the IDRs. It has been demonstrated that annealing twinning can easily arise in materials with low stacking fault energy [51e55]. In this alloy, more twinning grains appear in the IDRs than in the DAs, due to the retarding of g0 phase in the IDRs. In addition, twinning grains are only found in the IDRs alongside developed dendritic RX grains. This indicates that RX in
the DAs can trigger the twinning grains in the IDRs at the front of RX grain boundaries, though the grain boundary motion is difficult. According to RX nucleation theory, nucleation through thermal fluctuation is impossible. So many twinning strips can't be formed by subgrain coarsening nucleation. Therefore, nucleation by thermal twinning in the front of the grain boundary takes place, and it is easier for twinning in the IDRs to move into the original matrix compared with the normal grains in the DAs. Fig. 13 shows another examples with a lot of thermal twinning in the IDRs at the annealing temperature of 1280 C. The twinning strips are wider, compared with the previous one annealed at 1270 C, though the number of twinning grains is significantly less. In general, parallel twinning strips are easier to arise when samples with higher plastic deformation are annealed at lower temperatures. To further clarify the complete RX process in the IDRs, two samples with high plastic strains after annealing at 1280 C for 4 h are examined by EBSD (Fig. 14) and optical microscope (Fig. 15). Fig. 14 shows the RX microstructure in the IDRs which can be easily observed in the FSE micrographs (the area enclosed by red lines in Fig. 14(c) (f)). As stated above, small grains remain in the IDRs after RX is completed, as shown in both Figs. 14 and 15. RX grains can be observed by optical micrographs though it is not so clear to distinguish different grains. The average grain size in the IDRs is approximately 20e30 mm, while it was larger than 100 mm in the DAs. As shown in Fig. 14(b) and (e), plentiful twinning grain boundaries remain in the final RX microstructures. The mechanism of twinning is still controversial, and it could result from a growth accident or partial dislocations from the migrating grain
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Fig. 12. OIM from a partially recrystallized sample. (deformed at 980 C for 10% plastic strain and annealed at 1270 C for 2h) (a) IPF map with step size of 10 mm; (b) enlarged area of the black rectangle in (a), IPF map with step size of 3 mm; (c) enlarged area of the partial recrystallization, IPF map with step size of 0.05 mm; (d) grain boundary map superimposed with kikuchi band contrast (twinning grain boundaries represented by red lines).
Fig. 13. OIM from a partially recrystallized sample. (deformed at 1100 C for 7.8%% plastic strain and annealed at 1280 C for 10min):(a1) IPF map with step size of 10 mm; (b1) enlarged area of the black rectangle in (a1), IPF map with step size of 2 mm; (c1) enlarged area of the black rectangle in (b1), IPF map with step size of 0.658 mm; (a2) (b2) (c2) corresponding FSE maps with (a1) (b1) (c1).
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Fig. 14. OIM from two completely recrystallized samples.(a) (b) (c) deformed at 1100 C for 22.6% plastic strain and annealed at 1280 C for 4h; (d) (e) (f) deformed at 1200 C for 22.7% plastic strain and annealed at 1280 C for 4h; (a) (d) Euler angle map superimposed with grain boundaries; (b) (e) grain boundaries map with red lines representing twinning boundaries; (c) (f) FSE maps with red lines representing interdendritic regions (IDRs). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
Fig. 15. Optical microstructure from two completely recrystallized samples. These two samples correspond to those in Fig. 14. (a) (b) (c) deformed at 1100 C for 22.6% plastic strain and annealed at 1280 C for 4h; (d) (e) (f) deformed at 1200 C for 22.7% plastic strain and annealed at 1280 C for 4h.
boundaries. Both a gain in grain-boundary energy and a decrease in deformation stored energy play the dominant role in thermal twinning. Furthermore, thermal twinning would typically lead to the formation of rather immobile boundaries. Thus, grain coarsening is difficult owing to the presence of abundant twinning grain boundaries. In materials with FCC structure, a random twinning sequence using randomly one out of the possible 12 {111}<112>variants available would lead to nearly any orientation possible. In addition, it is assumed that there must be the preference of certain twin sequences. Thus, different orientations can arise though a lot of twinning grains with only one orientation are present at the early stage in the IDRs [56].
4. Conclusions RX in a nickel-based single crystal superalloy is investigated systematically using TEM, EBSD and optical microscope. Hot compression is employed to provide driving force for RX. The following conclusions can be drawn. 1) Deformation temperature has a great influence on RX behavior. For DD6 alloy, samples deformed around 980 C have the highest propensity for RX, because the presence of stacking faults can facilitate the RX nucleation. 2) The higher the deformation amount, the faster the grain boundary moves and the higher the RX nucleation rate.
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3) The g0 phase retards the movement of RX grain boundaries. When the annealing temperature is below the g0 solvus, as-cast dendritic microstructure plays a significant role in RX as a result of the different dissolution of g0 phase in the DAs and IDRs. 4) RX nucleates first and grow rapidly in the DAs. In the IDRs, thermal twinning play the dominant role, and small grains remain after RX is completed. Grain coarsening is rather difficult owing to abundant twinning grain boundaries. Acknowledgments This research was funded by the National Basic Research Program of China (No. 2011CB706801) and National Natural Science Foundation of China (Nos. 51171089 and 51374137). The authors are grateful for this support. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jallcom.2016.02.149. References [1] T. Pollock, S. Tin, Nickel-based superalloys for advanced turbine engines: chemistry, microstructure, and properties, J. Propul Power 22 (2006) 361e371. [2] T.M. Pollock, W.H. Murphy, The breakdown of single-crystal solidification in high refractory nickel-base alloys, Metall. Mater. Trans. A 27 (1996) 1081e1094. [3] T.M. Pollock, W.H. Murphy, E.H. Goldman, D.L. Uram, J.S. Tu, Grain defect formation during directional solidification of nickel base single crystals, in: S.D. Antolovich, R.W. Stusrud, R.A. MacKay, D.L. Anton, T. Khan, R.D. Kissinger, D.L. Klarstrom (Eds.), Superalloys 1992, The Minerals, Metals and Materials Society, Warrendale, PA, 1992, pp. 125e134. [4] D. Pan, Q. Xu, B. Liu, J. Li, H. Yuan, H. Jin, Modeling of grain selection during directional solidification of single crystal superalloy turbine blade castings, JOM 62 (2010) 30e34. [5] D. Ma, A. Bührig-Polaczek, The influence of surface roughness on freckle formation in directionally solidified superalloy samples, Metall. Mater. Trans. B 43B (2012) 671e677. [6] D. Ma, Q. Wu, A. Bührig-Polaczek, Some new observations on freckle formation in directionally solidified superalloy components, Metall. Mater. Trans. B 43 (2012) 344e353. [7] B. Zhang, D. Liu, Influence of recrystallization on high-temperature stress rupture property and fracture behavior of single crystal superalloy, Mat. Sci. Eng. A 551 (2012) 149e153. [8] G. Xie, L.H. Lou, Influence of the characteristic of recrystallization grain boundary on the formation of creep cracks in a directionally solidified Ni-base superalloy, Mat. Sci. Eng. A 532 (2012) 579e584. [9] G. Xie, L. Wang, J. Zhang, L.H. Lou, Intermediate temperature creep of directionally solidified Ni-based superalloy containing local recrystallization, Mat. Sci. Eng. A 528 (2011) 3062e3068. [10] G. Xie, L. Wang, J. Zhang, L.H. Lou, Influence of recrystallization on the hightemperature properties of a directionally solidified Ni-base superalloy, Metall. Mater Trans. A 39A (2008) 206e210. [11] T. Khan, P. Caron, Y.G. Nakagawa, Mechanical behavior and processing of DS and single crystal superalloys, J. Metals (1986) 16e19. [12] Z. Li, J. Xiong, Q. Xu, J. Li, B. Liu, Deformation and recrystallization of single crystal nickel-based superalloys during investment casting, J. Mater Process Tech. 217 (2015) 1e12. [13] Z. Li, Q. Xu, J. Xiong, J. Li, B. Liu, Simulation and experiments on plasticity and recrystallisation in SX superalloys by investment casting, Mater Res. Innov. 18 (2014) S4eS331. [14] C. Panwisawas, H. Mathur, J. Gebelin, D. Putman, C.M.F. Rae, R.C. Reed, Prediction of recrystallization in investment cast single-crystal superalloys, Acta Mater 61 (2013) 51e66. [15] A. Porter, B. Ralph, The Recrystallization of nickel-based superalloys, J. Mater Sci. 16 (1981) 707e713. [16] D.S. Weaver, Recrystallization and grain-growth behavior of a nickel-base superalloy during multi-hit deformation, Scr. Mater 57 (2007) 1044e1047. [17] M. Dahlen, L. Winberg, Influence of gamma'-precipitation on the recrystallization of a nickel-base superalloy, Acta Metall. 28 (1980) 41e50. [18] A.J. Porter, B. Ralph, Recrystallization of a nickel-base superalloy e kinetics and microstructural development, Mater. Sci. Eng. 59 (1983) 69e78. [19] D.D. Whitis, Recovery and recrystallization after critical strain in the nickelbased superalloy rene 88DT., Superalloys 2004, Minerals Metals Mater. Soc. Warrendale (2004) 391e400. [20] M. Kouichi, Directional Recrystallisation in Mechanically Alloyed ODS Nickel
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