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Microstructure evolutions and nucleation mechanisms of dynamic recrystallization of a powder metallurgy Ni-based superalloy during hot compression
Author’s Accepted Manuscript Microstructure evolutions and nucleation mechanisms of dynamic recrystallization of a powder metallurgy Ni-based superalloy during hot compression Guoai He, Feng Liu, Lan Huang, Zaiwang Huang, Liang Jiang www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(16)31165-0 http://dx.doi.org/10.1016/j.msea.2016.09.083 MSA34168
To appear in: Materials Science & Engineering A Received date: 26 August 2016 Revised date: 19 September 2016 Accepted date: 20 September 2016 Cite this article as: Guoai He, Feng Liu, Lan Huang, Zaiwang Huang and Liang Jiang, Microstructure evolutions and nucleation mechanisms of dynamic recrystallization of a powder metallurgy Ni-based superalloy during hot c o m p r e s s i o n , Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.09.083 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Microstructure evolutions and nucleation mechanisms of dynamic recrystallization of a powder metallurgy Ni-based superalloy during hot compression Guoai Hea,b,c, Feng Liua,b,c, Lan Huanga,b,c, Zaiwang Huanga,b,c,*, Liang Jianga,b,c
a
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China
b
Powder Metallurgy Research Institute, Central South University, Changsha 410083, China
c
High Temperature Materials Research Institute, Central South University, Changsha 410083,
China *
Corresponding author at: Powder Metallurgy Research Institute, Central South
University, Changsha 410013, PR China. Tel. & Fax: +86 0731 8883 0938. E-mail: [email protected]
Abstract Dynamic recrystallization (DRX) has been of great concern throughout the manufacturing processes, and it deeply affects in-service performance of powder metallurgy Ni-based disk components. Understanding the underpinning mechanisms of DRX is vital to produce the desired microstructure and mechanical properties of the superalloys. In this article, microstructure evolutions and nucleation mechanisms of DRX of an advanced Ni-based superalloy during hot deformation were studied using high resolution EBSD and TEM. The experimental results show that low angle 1
grain boundaries were formed at low temperature and readily evolved to high angle grain boundaries with temperature increasing and/or strain rate decreasing. Effects of strain amount on DRX were examined and significant DRX was detected when the strain increased to 0.5 under conditions of 1050 °C/0.01 s-1. Three different nucleation mechanisms were found under different deformation parameters. Nucleation of DRX was strongly related to the bulged-original boundaries, which acted as the interaction barrier with mobile dislocations at relatively low temperature and strain rate of 0.1 s-1. In contrast, twining boundaries were identified as the nucleation sites at a higher temperature and strain rate of 0.01 s-1. At 1100 °C and low strain rate of 0.001 s-1, coalesced γ’ was encompassed by bulged-original boundary, where nucleation of DRX was detected.
Keywords: Powder metallurgy Ni-based superalloy; hot compression; grain boundary evolution; dynamic recrystallization; nucleation mechanism
2
1 Introduction Powder metallurgy Ni-based superalloys are the most widely used for hot section rotating components of gas turbine aero-engine owing to their excellent high temperature
strength,
superior
creep/fatigue
resistance
at
elevated
temperature[1-4]. Forged alloys with fine grain size and uniform distribution of gamma prime were a desired selection to achieve the optimized mechanical performance. Generally, the pre-alloyed powders atomized by Argon were initially consolidated into nearly fully dense compacts using hot isostatic pressing (HIP), which were subsequently transformed for hot extrusion (HEX) to obtain favorable microstructure. Finally, isothermal forging was conducted under controllable strain rates and precise temperatures to a pre-designed reduction. During forging, the processing variables such as strain rate, deformation temperature and strain occupied very important place in controlling microstructure of final parts. Abnormal deformation conditions would cause catastrophic defects, i.e. surface cracks, inhomogeneous distribution of grain size and flow localization, and whereas, the grains
would
be
further
refined
when
deformed
under
appropriate
thermomechanical parameters owing to the dynamic recrystallization[5, 6]. Effects of thermomechanical processing parameters and initial conditions on microstructure evolution of PM Ni-based superalloys were studied extensively in past decades primarily using hot compression methods. Immarigeon and Floyd[7] investigated the importance of different initial grain sizes on plastic flow behaviors and their corresponding mechanisms. Three different mechanical behaviors of 3
flow-hardening,
steady-state
and
flow-softening
were
observed
and
the
corresponding physics could be attributed to dynamic grain growth, classical superplasticity under constant grain size and dynamic recrystallization. Similar works [8, 9]concerning initial grain sizes and forming parameters on microstructure development were conducted by Shen. Physical tests combined with finite element method were used to develop an empirical model for microstructure evolution in superalloy forming, which could predict the percentage of recrystallization and grain size distribution. Park et al. [10] applied two-steps forging process to evaluate the effects of strain on flow behavior and microstructures using hot compression tests. The results of flow behaviors showed that the strain rate and temperature substantially affected the process of DRX, while the initial grain size and applied strain exhibited a difference to meta-dynamic recrystallization (MDRX).In the two step forging process, the larger second strains resulted in smaller grain sizes owing to prominent recrystallization. Through introducing strain gradients using double cone specimen, Weaver et al.[11] further investigated the effects of industry-scale processing conditions i.e. monotonic hit and multi-hit with dwell time of 30 s or 60 s on microstructure development of Ni-based superalloy. In addition, ahangiri[12] found in IN 939 superalloy that dynamic recovery (DRV) contributes alloy softening in the temperature range from 1000 °C to 1050 °C, while dynamic recovery (DRV) takes a leading role in triggering material softening upon the temperature of 1100 °C-1150 °C. It has been widely documented that there exists a competition mechanism 4
between DRV and DRX in the superalloy forging. Three types of recrystallization behaviors can be summarized as [13]: discontinuous dynamic recrystallization (DDRX) is featured by obvious nucleation and grain growth, continuous dynamic recrystallization (CDRX) is characteristic for nuclei grow by migration of grain boundaries, geometric dynamic recrystallization (GDRX) is recognized by the refined grains induced by the impingement of initial elongated grain boundaries during severe plastic deformation. It should be noted that DDRX mechanism dominated flow softening behavior during hot deformation of Inconel 625 superalloy, as presented by Guo et al. [14]. Ning and others[15] reported two different nucleation mechanisms, i.e., grain boundary bulging nucleation and dislocation induced phase nucleation for DRX in the hot deformation of FGH4096 superalloy. While new dynamic recrystallized grains (DRXGs) were observed to form in the vicinity of the initial grain boundaries and second phase particles after the evolution of subgrains in Incoloy 901 superalloy[16]. Microscopic mechanism of grain boundaries evolution related to aforementioned mechanism can be revealed by electron backscatter diffraction (EBSD) technique. The misorientation angle (θ) between the adjacent grains is employed to characterize the subgrains and DRXGs, which can be divided into low angle grain boundaries (LAGBs, 2° ≤ θ< 15°) and high angle grain boundaries (HAGBs, θ ≥ 15°). For the case of a recovered polycrystalline containing both grains and subgrains, those grains with low angle grain boundaries are defined as subgrains [17]. Previous reports [18, 19] showed that LAGBs were formed at the initial period of deformation, which would be 5
further developed into HAGBs after the initiation of DRX. Based on the previous results, the hot working parameters significantly affected the development of microstructure i.e. DRV, DRX and MDRX, and their mechanisms were intensively investigated in the past decades. However, the microstructure evolution during hot deformation is quite complex and the deformation mechanisms vary with different superalloys. The nucleation mechanisms of DRX are not conclusive and strongly depend on the deformation parameters. In this study, hot compression tests were performed to simulate the actual forging process under a temperature range from 1000 °C to 1100 °C and strain rate 0.001, 0.01, 0.1, 1.0 s-1.Techniques of scanning electron microscopy (SEM), EBSD together with transmission electron microscope (TEM) were utilized to observe the microstructure features of an advanced PM Ni based superalloy. Herein, we demonstrate that there different nucleation mechanisms for DRX corresponding to certain deformation conditions.
2. Materials and Experimental Procedure 2.1 Materials The material used in current work was a newly designed superalloy with an nominal composition of (wt. %): Co was about 26%, Cr 13%, Ti 3.7%,W 4%, Mo 4%, Al 3.2%, Nb 0.95%, Hf 0.2%, and minor addition of B, C, Zr with balance nickel. The masteralloy was melted in a vacuum induction melting furnace and subsequently atomized into powder. The contained powder was then consolidated using hot 6
isostatic pressing (HIP) at a temperature and pressure of 1100°C/140 MPa. After HIP treatment, the billet was cut out and subjected to canned hot extrusion (HEX) to a reduction ratio of 10:1.
2.2 Hot Compression Tests Isothermal hot compression tests were carried out to simulate the real forging process and evaluate the microstructure evolution during hot working under the temperature range from 1000 °C to 1100 °C and strain rate 0.001, 0.01, 0.1, 1.0 s-1. These deformation procedures are chosen mainly to encourage DRX and superplastic deformation, and to avoid remarkable grain growth during hot forging. For this purpose, cylindrical specimens with dimension of 8 mm in diameter and 12 mm in height were machined from the extruded billet on the same circumference to ensure a uniform initial microstructure. Graphite foils with a thickness of 0.5 mm were applied at the end of the alloy specimen in order to reduce the friction with the dies. Subsequently, specimen was mounted in Gleeble 3180 test system with a thermocouple welded on the specimen surface to monitor and control the temperature. All the specimens were heated to preset temperature with a rate of 5°C s-1 and compressed up to a true strain of 0.7 () after temperature soaking for 3 minutes. The true strain was calculated by the following equation: ε=ln((L0+∆L)/L0 ), where, L0 is the initial length of the specimen, ∆L is the length of deformation. After compression, the specimens were quenched manually to room temperature within 2-3 s. 7
2.3 Microstructure analysis Deformed specimens were cut axially using a high-precision diamond wafer blade. The sections were prepared by virtue of standard metallographic techniques and finally polished with 0.05μm colloidal silica. For EBSD observation, vibration polishing was performed for 4 hours. A field-emission gun SEM (Quanta 650), equipped with an EBSD detector and Channel 5 software, was utilized to examine the microstructure evolution. Depending on specific grain size, measurements were performed over two types of areas using the following parameters: a spot size of 6.0, accelerating voltage 20 kV, and a step size between 0.3 and 1.0 μm. The average grain size was determined by the function of ‘detect grains’ in Channel 5. Thus, orientation information will be recorded automatically and utilized to calculate the misorientation angle among neighboring grains. The misorientation gik is calculated between the orientation gi of a grid point and its nearest neighbor gk (two per point, one to the right and one below) using the following equation: gik = gi-1·gk. Accordingly, misorientation angle within one degree will be classified as a unit. The software will count the number of each interval and calculate the corresponding frequency, according to which the frequency of misorientation angle under a specific value can be obtained and plotted. Herein, misorientation degree represents a certain range value. Misorientation angle of 15° marks a line to discriminate the LAGBs and HAGBs. For TEM observation, thin foils with a diameter of 3 mm were punched out at the geometry center of specimen after grinding to a thickness of 0.07-0.08 mm. After that, the discs were electro polished in a corrosive solution of ethanol and perchloric 8
acid (9:1 in volume fraction) at the temperature of -25 ° and voltage of 20 v. TEM examination was conducted on a field emission transmission electron microscopy (FEI Tecnai G2 F20) under an accelerating voltage of 200kV.
3. Results and discussions 3.1 Pristine microstructure before compression Figure 1 presents the typical orientation image microscopy (OIM) map (perpendicular to extrusion direction) prior to hot compression. Equiaxed grains occupy a dominated fraction and the average grain size was determined as 15.07 μm, as shown in Fig.1 (a). LAGBs and HAGBs are marked by gray lines and blue lines respectively. The statistical measurements in figure 1b show that most of grain boundaries are in the regime of HAGBs, corresponding to the state of recrystallization grains
3.2 Microstructure evolution of hot deformation Figure 2 demonstrates the microstructure evolution after imposing a true strain of 0.7 under various deformation conditions. The microscopic details reveal the 9
prominent microstructure sensibility that is related to deformation parameters. It could be seen that most of the initial equiaxed grains were highly deformed and remarkably developed to elongated counterparts under certain conditions and there fined grains were formed at the grain boundaries regardless of deformation conditions. Compared to the microstructure of as-HEXed alloy, new grains formed at initial grain boundaries after compressed to a true strain of 0.7 at all deformation conditions. For the case of specimen that deformed at 1000°C /1.0 s-1, a few fine and new grains are generated around the elongated grain boundaries, amount of which increases with the increase of temperature and the decrease of strain rate. With temperature increased up to 1100 °C, completed DRX microstructure with a relatively small grain size corresponds to the higher strain rate of 1.0 s-1, in contrast to the substantial growth of recrystallization grains under low strain rate.
10
11
3.3 Effects of strain rate and temperature on microstructure and misorientation evolution
Figure 3 presents the typical OIMs when deformed at different conditions, indicating the evolution of LAGBs (marked as gray lines) and HAGBs (marked as black lines). As demonstrated in Fig.3 (a), a high density of LAGBs was observed along the deformed grain boundaries when compressed to a true strain of 0.7 under 1000 °C/1.0 s-1. For the specimen deformed at a strain rate of 1.0 s-1, the fraction of 12
LAGBs was a little higher than that of 0.1 s-1 at the same deformation temperature of 1000 °C, which was about 88.79% and 81.35%, respectively. The LAGBs are inherently associated with the substructures and dislocations, which would be further developed into HAGBs in accompany with the DRX initiation. After plastic deformation exceeding a critical value, mobile dislocations piled up on the grain boundary and readily triggered DRX process. Herein, new grains nucleate and grow in virtue of high temperature and strain rate. Figure 4 shows a high density of dislocations under1000 °C/1.0 s-1, where dislocations piled up inside the grains. This observation is consistent with the EBSD measurement that indicates LAGBs are ubiquitous.
While the strain rate is decreased to 0.1 s-1, discernable embryos nearby the grain boundary can be found and point to the occurrence of DRX, shown in Fig.3 (b). When the temperature increases to 1025 °C, the fraction of LAGBs continues to decrease and the fraction of HAGBs with about 60°starts to increase, compared to that of 1000 °C at the same strain rate. Misorientation about 60° is considered as a signal of twinning mechanism, which was believed to play a primary role in nucleation process of DRX for the alloys with low stacking fault energy [20, 21]. Gradually increasing recrystallized grains formed at the deformed grain boundaries and the percentageof 13
HAGBs starts to increase with the temperature increases. Thus, the higher temperatures can provide higher driving force for grain boundaries migration, and meanwhile, lower strain rates will offer sufficient time for the grain growth, both of which are ascribed to be the incentive of DRX behavior. Figure 3 (g)~(l) shows the point to point and point to origin misorientations along the initial grain boundaries and along the boundaries to the grain interior, respectively. Fig. 3 (g) presents the misorientations along the boundaries (line A1) when deformed at temperature of 1000 °C and strain rate of 1.0 s-1. The steady low-misorientations of point to point (marked as straight line) along boundaries indicate the high density of LAGBs. While the low misorientations (lower than 10°) of point to origin (marked as dash line) along the boundaries demonstrate the absence of HAGBs. Small angle grain boundaries lower than 10° were considered to be necessary for nucleation of DRX. The results imply that a significant fraction of subgrains were obtained when the specimen was deformed at temperature of 1000 °C and strain rate of 1.0 s-1, where DRX process is limited. The misorientation from grain boundary to the grain interior (line A2) is presented in Fig.3 (j). The point-to-point misorientation does not exceed 4° while those of point-to-origin gradually increase to an extent of 22°, indicating the continuous development of grain boundaries from the marginal to the center. These results also imply that the grain boundaries are the preferential sources for the initiation of new grains. As strain rate decreases to 0.1 s-1, obvious fluctuation in misorientations along line B1 is observed and presented in Fig.3 (h). The frequent changes in misorientations 14
further validate the lower density of LAGBs compared to those at 1.0 s-1, indicating that some of the developed subgrains have been transformed into higher angle grain boundaries. Fig. 3 (k) illustrates the misorientation from boundary to the center (line B2). As can be seen, a rapid rise within a distance of 1 μm is detected, indicating the occurrence of DRX. A few new grains frequently nucleate at initial grain boundaries, which agree with the indication of misorientation evolution. While the temperature is enhanced up to 1025 °C, the scenario is different that sharp peaks marked by dark arrows emerge, as illustrated in Fig.3 (i). The continual-sharp rise in misorientation is generally considered as a signal of significant DRX occurring, which is consistent with the microstructure observation in Fig.3 (c). The point-to-origin misorientation (along line C2) is higher than 15° with a larger distance of 10 μm, indicating a higher degree of DRX of the specimen when deformed under conditions of 1025 °C and 0.1 s-1, as shown in Fig.3 (l).
15
3.4 Effects of strain on microstructure and misorientation evolution
The OIMs of specimens that were deformed to different strains under conditions of 1050 °C/0.01 s-1 are illustrated in Figure 5 (a)~(d), showing the effects of strain on microstructure and misorientation evolution. The summary on the variations of the misorientation at different strains was presented in Fig.5 (e) and (f). As can be seen clearly, the fraction of LAGBs increases firstly to a peak with the increase of strain and starts to decrease with the further accumulation of strain. The physics behind this phenomenon could be ascribed to the adequate accumulation of strain for germinating DRX, of which after the occurrence the density of LAGBs starts to decrease. It can be found in Fig.5 (a) that LAGBs are formed near grain boundaries at the early deformation, while a certain amount of original HAGs still retains. Generally, grain boundaries with misorientation about 10~15o are treated as the transition zone 16
from LAGBs to HAGBs[22]. It should be noted in figure 5 (e) that the fraction of transition zone is fairly low, which indicates the rapid migration from LAGBs to HAGBs. As the deformation proceeding, the LAGBs continue to accumulate and a few new grains are generated due to the development of some LAGs, as shown in Fig.5 (b). The fraction of LAGBs is observed to decrease obviously with the strain increasing to 0.5, indicating the significant occurrence of DRX. After reaching a strain of 0.7, the amount of HAGBs increases to a high fraction of 68%, and average grain size becomes remarkably smaller owing to the DRX process.
3.5 Mechanism of dynamic recrystallization Discontinuous dynamic recrystallization (DDRX) and continuous dynamic recrystallization (CDRX) were widely recognized as the primary mechanisms of DRX in Ni based superalloy [22], wherein the deformation conditions, i.e. strain, temperature and strain rate, determine the nucleation mechanisms. In order to understand the specific nucleation mechanisms of DRX under different deformation background, TEM examinations were employed to investigate the deformation microstructure at the nanoscale.
Figure 6 shows the TEM observation when deformed to a true strain of 0.7 at 17
conditions of 1000 °C/0.1 s-1 and 1050 °C/0.1 s-1, respectively. A high density of dislocations is observed at a lower temperature of 1000 °C compared to that of 1050 °C, indicating faster dislocation annihilating at a higher temperature. Dark region in figure 6 (a) shows that a high density of dislocations piling up at the grain boundaries after plastic deformation. With a further deformation, as dislocations continuing to pile up at the boundaries, the dislocation walls are formed due to the limited time for annihilation. The deformation substructures are even more ubiquitous and can be attributed to the dislocation activities, eventually enable the formation of bulged grain boundaries, as shown in Fig.6 (a). For the DRX nuclei, the serrated boundaries provide favorable sites for dislocation annihilation. Grain growth from nuclei to DRX grains is commonly associated with the reduction of boundary curvature, as shown in figure 6. In contrast, very few dislocations and substructures can be found (Fig. 6b) at 1050 °C because most of grains have been well developed by DRX process. A faster atom diffusion will be expected at higher deformation temperature, i.e., 1050 °C/0.01s-1 (Fig. 6b), dislocation/substructure annihilation become easier and promote the formation of the dislocation-free grains.
When deformed at 1050 °C /0.01 s-1 and 1100 °C/0.01 s-1, the deformation microstructure exhibits more complex behavior. Deformation twinning crossing grain 18
interior becomes prevalent, and is believed to accelerate boundary bulging and the separation of bulged parts from original grains [19], as marked by dark arrows in Fig. 7. The new DRX nuclei with some dislocations inside are generated along the bulged twining boundary. The twining boundaries are more inclined to transform into random high angle boundaries probably due to the interaction between mobile dislocations and bulged twining boundaries, as shown in Fig .7 (a). At even higher temperature to 1100 °C, more deformation twinnings are observed in the new-formed grains and inside a lower dislocation density is obtained. This finding is in good agreement with previous reports[23], implying that the twinning boundary can serve as the sites for DRX nucleation.
Figure 8 presents the TEM examination of specimen which was deformed to a true strain of 0.7 under conditions of 1100 °C/0.001 s-1. Electron diffraction pattern measurement indicates the existence of two phases of gamma and gamma prime. As 19
can be seen, two or multiple γ’ interact with one another at elevated temperature, some of which will be further developed into integration. Coalesced γ’ near grain boundaries will come into contact with grain boundaries gradually with the deformation progressing. The boundaries will become serrated and bulged under deformation, by which the coalesced γ’ was surrounded, as shown in figure 8 (b). The bulged boundary and enclosed γ’ will be the preferable nucleation sites for DRX. DRX nucleus will be generated due to the interaction between the mobile dislocations and bulged boundaries. In addition, higher temperature accelerates the atom diffusion rate; and the bulged boundaries provide larger diffusion interface, both of which will promote the development of DRX nuclei, as shown in figure 8 (c). During atom diffusion, boundary formation elements will diffuse toward the interface to form the boundary while the formation elements of γ’ will diffuse into the γ matrix. Meanwhile, the elements in gamma matrix and γ’ path (matrix between two γ’) will diffuse toward the DRX nuclei, leading to the development of DRX nuclei, as shown in figure 8 (d). As a result, new dislocation-free DRX grains with a size range of 0.4-0.8 μm will be generated along the original boundary after sufficient diffusion.
4 Conclusion 1) Remarkable LAGBs form1000 °C/1.0 s-1 and further develop into HAGBs with the increase of temperature and the decrease of strain rate. DRX can be fully finished upon 1100 °C/1.0 s-1, while an obvious grain growth can be stimulated at the strain rate of 0.001 s-1. 20
2) The fraction of LAGBs increases firstly to a peak with the increase of strain and then decreases gradually with further deformation. Degree of DRX increases with the accumulation of straining and significant DRX was detected when the strain increases to 0.5 at temperature of 1050 °C / 0.01 s-1. 3) Three different nucleation mechanisms were observed at different deformation parameters. Nucleation of DRX was found to form at the bulged-original boundaries which were interacted with mobile dislocations at relatively low temperature and strain rate of 0.1 s-1, while twining boundaries were observed to be preferable sites for nucleation at higher temperature and strain rate of 0.01 s-1. At a temperature of 1100 °C and low strain rate of 0.001 s-1, coalesced γ’ was encompassed by bulged-original boundary, where nucleation of DRX was detected.
Conflicts of interest: The authors declare that there are no conflicts of interest regarding the publication of this paper.
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Acknowledgment Z. H. and L. J. appreciate the financial support from The National Key Research and Development Program of China (2016YFB0700300). G. H. is grateful for the support from the Fundamental Research Funds for the Central Universities of Central South University (2015zzts031) and the outstanding graduate project of Advanced Non-ferrous Metal Structural Materials and Manufacturing Collaborative Innovation Center. F. L. and L. H. are thankful for the funding sponsorship of the Natural Science Foundation of China (51401242, 51301209). We would like to acknowledge the help from Hong JIANG for the helpful technical discussion.
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Figure 1 Typical orientation image microscopy (OIM) map (along direction of HEX) of as-HEXed alloy before hot compression (a); distribution of misorientation of as HEXed alloy (b). Figure 2 Microstructure development of the alloy compressed to a true strain of 0.7 under various deformation conditions. 23
Figure 3 Typical OIMs and the corresponding distribution of misorientation of specimens when deformed at different temperatures and strain rates: (a) and (d) 1000 °C/1.0 s-1; (b) and (e) 1000 °C/0.1 s-1; (c) and (f) 1025 °C/0.01 s-1. Figures (g)~(l) show the misorientations of point to point and point to origin along the initial grain boundaries and from the boundaries to the interior of grains, as marked by lines: (g), A1; (h), B1; (i), C1; (j), A2; (k), B2; (l), C2. Figure 4 TEM observation of the specimen which was deformed under the condition of 1000 °C/1.0 s-1, showing the characteristics of dislocation. Figure 5 The OIMs of specimens deformed to a different of strains under conditions of 1050 °C and 0.01 s-1: (a), 0.1; (b), 0.3; (c), 0.5; (d), 0.7; (e), summary on the variations of misorientation; (f), illustration of variation tendency of misorientation. Figure 6 TEM observation of the specimen deformed under deformation conditions: (a) 1000 °C/0.1 s-1; (b) 1050 °C/0.1 s-1; showing the nucleation of DRX at bulged boundaries with high density of dislocation and substructure. Figure 7 TEM observation of the specimen when deformed under deformation conditions: (a) 1050 °C/0.01 s-1; (b) 1100 °C/0.01 s-1; showing the nucleation of DRX along the twining boundaries. Figure 8 TEM observation of the specimen which was deformed under deformation conditions of 1100 °C/0.001 s-1 ; showing the coalesced γ’ is encompassed by bulged-original boundary, where nucleation of DRX was detected.
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PII: DOI: Reference:
S0921-5093(16)31165-0 http://dx.doi.org/10.1016/j.msea.2016.09.083 MSA34168
To appear in: Materials Science & Engineering A Received date: 26 August 2016 Revised date: 19 September 2016 Accepted date: 20 September 2016 Cite this article as: Guoai He, Feng Liu, Lan Huang, Zaiwang Huang and Liang Jiang, Microstructure evolutions and nucleation mechanisms of dynamic recrystallization of a powder metallurgy Ni-based superalloy during hot c o m p r e s s i o n , Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.09.083 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Microstructure evolutions and nucleation mechanisms of dynamic recrystallization of a powder metallurgy Ni-based superalloy during hot compression Guoai Hea,b,c, Feng Liua,b,c, Lan Huanga,b,c, Zaiwang Huanga,b,c,*, Liang Jianga,b,c
a
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China
b
Powder Metallurgy Research Institute, Central South University, Changsha 410083, China
c
High Temperature Materials Research Institute, Central South University, Changsha 410083,
China *
Corresponding author at: Powder Metallurgy Research Institute, Central South
University, Changsha 410013, PR China. Tel. & Fax: +86 0731 8883 0938. E-mail: [email protected]
Abstract Dynamic recrystallization (DRX) has been of great concern throughout the manufacturing processes, and it deeply affects in-service performance of powder metallurgy Ni-based disk components. Understanding the underpinning mechanisms of DRX is vital to produce the desired microstructure and mechanical properties of the superalloys. In this article, microstructure evolutions and nucleation mechanisms of DRX of an advanced Ni-based superalloy during hot deformation were studied using high resolution EBSD and TEM. The experimental results show that low angle 1
grain boundaries were formed at low temperature and readily evolved to high angle grain boundaries with temperature increasing and/or strain rate decreasing. Effects of strain amount on DRX were examined and significant DRX was detected when the strain increased to 0.5 under conditions of 1050 °C/0.01 s-1. Three different nucleation mechanisms were found under different deformation parameters. Nucleation of DRX was strongly related to the bulged-original boundaries, which acted as the interaction barrier with mobile dislocations at relatively low temperature and strain rate of 0.1 s-1. In contrast, twining boundaries were identified as the nucleation sites at a higher temperature and strain rate of 0.01 s-1. At 1100 °C and low strain rate of 0.001 s-1, coalesced γ’ was encompassed by bulged-original boundary, where nucleation of DRX was detected.
Keywords: Powder metallurgy Ni-based superalloy; hot compression; grain boundary evolution; dynamic recrystallization; nucleation mechanism
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1 Introduction Powder metallurgy Ni-based superalloys are the most widely used for hot section rotating components of gas turbine aero-engine owing to their excellent high temperature
strength,
superior
creep/fatigue
resistance
at
elevated
temperature[1-4]. Forged alloys with fine grain size and uniform distribution of gamma prime were a desired selection to achieve the optimized mechanical performance. Generally, the pre-alloyed powders atomized by Argon were initially consolidated into nearly fully dense compacts using hot isostatic pressing (HIP), which were subsequently transformed for hot extrusion (HEX) to obtain favorable microstructure. Finally, isothermal forging was conducted under controllable strain rates and precise temperatures to a pre-designed reduction. During forging, the processing variables such as strain rate, deformation temperature and strain occupied very important place in controlling microstructure of final parts. Abnormal deformation conditions would cause catastrophic defects, i.e. surface cracks, inhomogeneous distribution of grain size and flow localization, and whereas, the grains
would
be
further
refined
when
deformed
under
appropriate
thermomechanical parameters owing to the dynamic recrystallization[5, 6]. Effects of thermomechanical processing parameters and initial conditions on microstructure evolution of PM Ni-based superalloys were studied extensively in past decades primarily using hot compression methods. Immarigeon and Floyd[7] investigated the importance of different initial grain sizes on plastic flow behaviors and their corresponding mechanisms. Three different mechanical behaviors of 3
flow-hardening,
steady-state
and
flow-softening
were
observed
and
the
corresponding physics could be attributed to dynamic grain growth, classical superplasticity under constant grain size and dynamic recrystallization. Similar works [8, 9]concerning initial grain sizes and forming parameters on microstructure development were conducted by Shen. Physical tests combined with finite element method were used to develop an empirical model for microstructure evolution in superalloy forming, which could predict the percentage of recrystallization and grain size distribution. Park et al. [10] applied two-steps forging process to evaluate the effects of strain on flow behavior and microstructures using hot compression tests. The results of flow behaviors showed that the strain rate and temperature substantially affected the process of DRX, while the initial grain size and applied strain exhibited a difference to meta-dynamic recrystallization (MDRX).In the two step forging process, the larger second strains resulted in smaller grain sizes owing to prominent recrystallization. Through introducing strain gradients using double cone specimen, Weaver et al.[11] further investigated the effects of industry-scale processing conditions i.e. monotonic hit and multi-hit with dwell time of 30 s or 60 s on microstructure development of Ni-based superalloy. In addition, ahangiri[12] found in IN 939 superalloy that dynamic recovery (DRV) contributes alloy softening in the temperature range from 1000 °C to 1050 °C, while dynamic recovery (DRV) takes a leading role in triggering material softening upon the temperature of 1100 °C-1150 °C. It has been widely documented that there exists a competition mechanism 4
between DRV and DRX in the superalloy forging. Three types of recrystallization behaviors can be summarized as [13]: discontinuous dynamic recrystallization (DDRX) is featured by obvious nucleation and grain growth, continuous dynamic recrystallization (CDRX) is characteristic for nuclei grow by migration of grain boundaries, geometric dynamic recrystallization (GDRX) is recognized by the refined grains induced by the impingement of initial elongated grain boundaries during severe plastic deformation. It should be noted that DDRX mechanism dominated flow softening behavior during hot deformation of Inconel 625 superalloy, as presented by Guo et al. [14]. Ning and others[15] reported two different nucleation mechanisms, i.e., grain boundary bulging nucleation and dislocation induced phase nucleation for DRX in the hot deformation of FGH4096 superalloy. While new dynamic recrystallized grains (DRXGs) were observed to form in the vicinity of the initial grain boundaries and second phase particles after the evolution of subgrains in Incoloy 901 superalloy[16]. Microscopic mechanism of grain boundaries evolution related to aforementioned mechanism can be revealed by electron backscatter diffraction (EBSD) technique. The misorientation angle (θ) between the adjacent grains is employed to characterize the subgrains and DRXGs, which can be divided into low angle grain boundaries (LAGBs, 2° ≤ θ< 15°) and high angle grain boundaries (HAGBs, θ ≥ 15°). For the case of a recovered polycrystalline containing both grains and subgrains, those grains with low angle grain boundaries are defined as subgrains [17]. Previous reports [18, 19] showed that LAGBs were formed at the initial period of deformation, which would be 5
further developed into HAGBs after the initiation of DRX. Based on the previous results, the hot working parameters significantly affected the development of microstructure i.e. DRV, DRX and MDRX, and their mechanisms were intensively investigated in the past decades. However, the microstructure evolution during hot deformation is quite complex and the deformation mechanisms vary with different superalloys. The nucleation mechanisms of DRX are not conclusive and strongly depend on the deformation parameters. In this study, hot compression tests were performed to simulate the actual forging process under a temperature range from 1000 °C to 1100 °C and strain rate 0.001, 0.01, 0.1, 1.0 s-1.Techniques of scanning electron microscopy (SEM), EBSD together with transmission electron microscope (TEM) were utilized to observe the microstructure features of an advanced PM Ni based superalloy. Herein, we demonstrate that there different nucleation mechanisms for DRX corresponding to certain deformation conditions.
2. Materials and Experimental Procedure 2.1 Materials The material used in current work was a newly designed superalloy with an nominal composition of (wt. %): Co was about 26%, Cr 13%, Ti 3.7%,W 4%, Mo 4%, Al 3.2%, Nb 0.95%, Hf 0.2%, and minor addition of B, C, Zr with balance nickel. The masteralloy was melted in a vacuum induction melting furnace and subsequently atomized into powder. The contained powder was then consolidated using hot 6
isostatic pressing (HIP) at a temperature and pressure of 1100°C/140 MPa. After HIP treatment, the billet was cut out and subjected to canned hot extrusion (HEX) to a reduction ratio of 10:1.
2.2 Hot Compression Tests Isothermal hot compression tests were carried out to simulate the real forging process and evaluate the microstructure evolution during hot working under the temperature range from 1000 °C to 1100 °C and strain rate 0.001, 0.01, 0.1, 1.0 s-1. These deformation procedures are chosen mainly to encourage DRX and superplastic deformation, and to avoid remarkable grain growth during hot forging. For this purpose, cylindrical specimens with dimension of 8 mm in diameter and 12 mm in height were machined from the extruded billet on the same circumference to ensure a uniform initial microstructure. Graphite foils with a thickness of 0.5 mm were applied at the end of the alloy specimen in order to reduce the friction with the dies. Subsequently, specimen was mounted in Gleeble 3180 test system with a thermocouple welded on the specimen surface to monitor and control the temperature. All the specimens were heated to preset temperature with a rate of 5°C s-1 and compressed up to a true strain of 0.7 () after temperature soaking for 3 minutes. The true strain was calculated by the following equation: ε=ln((L0+∆L)/L0 ), where, L0 is the initial length of the specimen, ∆L is the length of deformation. After compression, the specimens were quenched manually to room temperature within 2-3 s. 7
2.3 Microstructure analysis Deformed specimens were cut axially using a high-precision diamond wafer blade. The sections were prepared by virtue of standard metallographic techniques and finally polished with 0.05μm colloidal silica. For EBSD observation, vibration polishing was performed for 4 hours. A field-emission gun SEM (Quanta 650), equipped with an EBSD detector and Channel 5 software, was utilized to examine the microstructure evolution. Depending on specific grain size, measurements were performed over two types of areas using the following parameters: a spot size of 6.0, accelerating voltage 20 kV, and a step size between 0.3 and 1.0 μm. The average grain size was determined by the function of ‘detect grains’ in Channel 5. Thus, orientation information will be recorded automatically and utilized to calculate the misorientation angle among neighboring grains. The misorientation gik is calculated between the orientation gi of a grid point and its nearest neighbor gk (two per point, one to the right and one below) using the following equation: gik = gi-1·gk. Accordingly, misorientation angle within one degree will be classified as a unit. The software will count the number of each interval and calculate the corresponding frequency, according to which the frequency of misorientation angle under a specific value can be obtained and plotted. Herein, misorientation degree represents a certain range value. Misorientation angle of 15° marks a line to discriminate the LAGBs and HAGBs. For TEM observation, thin foils with a diameter of 3 mm were punched out at the geometry center of specimen after grinding to a thickness of 0.07-0.08 mm. After that, the discs were electro polished in a corrosive solution of ethanol and perchloric 8
acid (9:1 in volume fraction) at the temperature of -25 ° and voltage of 20 v. TEM examination was conducted on a field emission transmission electron microscopy (FEI Tecnai G2 F20) under an accelerating voltage of 200kV.
3. Results and discussions 3.1 Pristine microstructure before compression Figure 1 presents the typical orientation image microscopy (OIM) map (perpendicular to extrusion direction) prior to hot compression. Equiaxed grains occupy a dominated fraction and the average grain size was determined as 15.07 μm, as shown in Fig.1 (a). LAGBs and HAGBs are marked by gray lines and blue lines respectively. The statistical measurements in figure 1b show that most of grain boundaries are in the regime of HAGBs, corresponding to the state of recrystallization grains
3.2 Microstructure evolution of hot deformation Figure 2 demonstrates the microstructure evolution after imposing a true strain of 0.7 under various deformation conditions. The microscopic details reveal the 9
prominent microstructure sensibility that is related to deformation parameters. It could be seen that most of the initial equiaxed grains were highly deformed and remarkably developed to elongated counterparts under certain conditions and there fined grains were formed at the grain boundaries regardless of deformation conditions. Compared to the microstructure of as-HEXed alloy, new grains formed at initial grain boundaries after compressed to a true strain of 0.7 at all deformation conditions. For the case of specimen that deformed at 1000°C /1.0 s-1, a few fine and new grains are generated around the elongated grain boundaries, amount of which increases with the increase of temperature and the decrease of strain rate. With temperature increased up to 1100 °C, completed DRX microstructure with a relatively small grain size corresponds to the higher strain rate of 1.0 s-1, in contrast to the substantial growth of recrystallization grains under low strain rate.
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3.3 Effects of strain rate and temperature on microstructure and misorientation evolution
Figure 3 presents the typical OIMs when deformed at different conditions, indicating the evolution of LAGBs (marked as gray lines) and HAGBs (marked as black lines). As demonstrated in Fig.3 (a), a high density of LAGBs was observed along the deformed grain boundaries when compressed to a true strain of 0.7 under 1000 °C/1.0 s-1. For the specimen deformed at a strain rate of 1.0 s-1, the fraction of 12
LAGBs was a little higher than that of 0.1 s-1 at the same deformation temperature of 1000 °C, which was about 88.79% and 81.35%, respectively. The LAGBs are inherently associated with the substructures and dislocations, which would be further developed into HAGBs in accompany with the DRX initiation. After plastic deformation exceeding a critical value, mobile dislocations piled up on the grain boundary and readily triggered DRX process. Herein, new grains nucleate and grow in virtue of high temperature and strain rate. Figure 4 shows a high density of dislocations under1000 °C/1.0 s-1, where dislocations piled up inside the grains. This observation is consistent with the EBSD measurement that indicates LAGBs are ubiquitous.
While the strain rate is decreased to 0.1 s-1, discernable embryos nearby the grain boundary can be found and point to the occurrence of DRX, shown in Fig.3 (b). When the temperature increases to 1025 °C, the fraction of LAGBs continues to decrease and the fraction of HAGBs with about 60°starts to increase, compared to that of 1000 °C at the same strain rate. Misorientation about 60° is considered as a signal of twinning mechanism, which was believed to play a primary role in nucleation process of DRX for the alloys with low stacking fault energy [20, 21]. Gradually increasing recrystallized grains formed at the deformed grain boundaries and the percentageof 13
HAGBs starts to increase with the temperature increases. Thus, the higher temperatures can provide higher driving force for grain boundaries migration, and meanwhile, lower strain rates will offer sufficient time for the grain growth, both of which are ascribed to be the incentive of DRX behavior. Figure 3 (g)~(l) shows the point to point and point to origin misorientations along the initial grain boundaries and along the boundaries to the grain interior, respectively. Fig. 3 (g) presents the misorientations along the boundaries (line A1) when deformed at temperature of 1000 °C and strain rate of 1.0 s-1. The steady low-misorientations of point to point (marked as straight line) along boundaries indicate the high density of LAGBs. While the low misorientations (lower than 10°) of point to origin (marked as dash line) along the boundaries demonstrate the absence of HAGBs. Small angle grain boundaries lower than 10° were considered to be necessary for nucleation of DRX. The results imply that a significant fraction of subgrains were obtained when the specimen was deformed at temperature of 1000 °C and strain rate of 1.0 s-1, where DRX process is limited. The misorientation from grain boundary to the grain interior (line A2) is presented in Fig.3 (j). The point-to-point misorientation does not exceed 4° while those of point-to-origin gradually increase to an extent of 22°, indicating the continuous development of grain boundaries from the marginal to the center. These results also imply that the grain boundaries are the preferential sources for the initiation of new grains. As strain rate decreases to 0.1 s-1, obvious fluctuation in misorientations along line B1 is observed and presented in Fig.3 (h). The frequent changes in misorientations 14
further validate the lower density of LAGBs compared to those at 1.0 s-1, indicating that some of the developed subgrains have been transformed into higher angle grain boundaries. Fig. 3 (k) illustrates the misorientation from boundary to the center (line B2). As can be seen, a rapid rise within a distance of 1 μm is detected, indicating the occurrence of DRX. A few new grains frequently nucleate at initial grain boundaries, which agree with the indication of misorientation evolution. While the temperature is enhanced up to 1025 °C, the scenario is different that sharp peaks marked by dark arrows emerge, as illustrated in Fig.3 (i). The continual-sharp rise in misorientation is generally considered as a signal of significant DRX occurring, which is consistent with the microstructure observation in Fig.3 (c). The point-to-origin misorientation (along line C2) is higher than 15° with a larger distance of 10 μm, indicating a higher degree of DRX of the specimen when deformed under conditions of 1025 °C and 0.1 s-1, as shown in Fig.3 (l).
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3.4 Effects of strain on microstructure and misorientation evolution
The OIMs of specimens that were deformed to different strains under conditions of 1050 °C/0.01 s-1 are illustrated in Figure 5 (a)~(d), showing the effects of strain on microstructure and misorientation evolution. The summary on the variations of the misorientation at different strains was presented in Fig.5 (e) and (f). As can be seen clearly, the fraction of LAGBs increases firstly to a peak with the increase of strain and starts to decrease with the further accumulation of strain. The physics behind this phenomenon could be ascribed to the adequate accumulation of strain for germinating DRX, of which after the occurrence the density of LAGBs starts to decrease. It can be found in Fig.5 (a) that LAGBs are formed near grain boundaries at the early deformation, while a certain amount of original HAGs still retains. Generally, grain boundaries with misorientation about 10~15o are treated as the transition zone 16
from LAGBs to HAGBs[22]. It should be noted in figure 5 (e) that the fraction of transition zone is fairly low, which indicates the rapid migration from LAGBs to HAGBs. As the deformation proceeding, the LAGBs continue to accumulate and a few new grains are generated due to the development of some LAGs, as shown in Fig.5 (b). The fraction of LAGBs is observed to decrease obviously with the strain increasing to 0.5, indicating the significant occurrence of DRX. After reaching a strain of 0.7, the amount of HAGBs increases to a high fraction of 68%, and average grain size becomes remarkably smaller owing to the DRX process.
3.5 Mechanism of dynamic recrystallization Discontinuous dynamic recrystallization (DDRX) and continuous dynamic recrystallization (CDRX) were widely recognized as the primary mechanisms of DRX in Ni based superalloy [22], wherein the deformation conditions, i.e. strain, temperature and strain rate, determine the nucleation mechanisms. In order to understand the specific nucleation mechanisms of DRX under different deformation background, TEM examinations were employed to investigate the deformation microstructure at the nanoscale.
Figure 6 shows the TEM observation when deformed to a true strain of 0.7 at 17
conditions of 1000 °C/0.1 s-1 and 1050 °C/0.1 s-1, respectively. A high density of dislocations is observed at a lower temperature of 1000 °C compared to that of 1050 °C, indicating faster dislocation annihilating at a higher temperature. Dark region in figure 6 (a) shows that a high density of dislocations piling up at the grain boundaries after plastic deformation. With a further deformation, as dislocations continuing to pile up at the boundaries, the dislocation walls are formed due to the limited time for annihilation. The deformation substructures are even more ubiquitous and can be attributed to the dislocation activities, eventually enable the formation of bulged grain boundaries, as shown in Fig.6 (a). For the DRX nuclei, the serrated boundaries provide favorable sites for dislocation annihilation. Grain growth from nuclei to DRX grains is commonly associated with the reduction of boundary curvature, as shown in figure 6. In contrast, very few dislocations and substructures can be found (Fig. 6b) at 1050 °C because most of grains have been well developed by DRX process. A faster atom diffusion will be expected at higher deformation temperature, i.e., 1050 °C/0.01s-1 (Fig. 6b), dislocation/substructure annihilation become easier and promote the formation of the dislocation-free grains.
When deformed at 1050 °C /0.01 s-1 and 1100 °C/0.01 s-1, the deformation microstructure exhibits more complex behavior. Deformation twinning crossing grain 18
interior becomes prevalent, and is believed to accelerate boundary bulging and the separation of bulged parts from original grains [19], as marked by dark arrows in Fig. 7. The new DRX nuclei with some dislocations inside are generated along the bulged twining boundary. The twining boundaries are more inclined to transform into random high angle boundaries probably due to the interaction between mobile dislocations and bulged twining boundaries, as shown in Fig .7 (a). At even higher temperature to 1100 °C, more deformation twinnings are observed in the new-formed grains and inside a lower dislocation density is obtained. This finding is in good agreement with previous reports[23], implying that the twinning boundary can serve as the sites for DRX nucleation.
Figure 8 presents the TEM examination of specimen which was deformed to a true strain of 0.7 under conditions of 1100 °C/0.001 s-1. Electron diffraction pattern measurement indicates the existence of two phases of gamma and gamma prime. As 19
can be seen, two or multiple γ’ interact with one another at elevated temperature, some of which will be further developed into integration. Coalesced γ’ near grain boundaries will come into contact with grain boundaries gradually with the deformation progressing. The boundaries will become serrated and bulged under deformation, by which the coalesced γ’ was surrounded, as shown in figure 8 (b). The bulged boundary and enclosed γ’ will be the preferable nucleation sites for DRX. DRX nucleus will be generated due to the interaction between the mobile dislocations and bulged boundaries. In addition, higher temperature accelerates the atom diffusion rate; and the bulged boundaries provide larger diffusion interface, both of which will promote the development of DRX nuclei, as shown in figure 8 (c). During atom diffusion, boundary formation elements will diffuse toward the interface to form the boundary while the formation elements of γ’ will diffuse into the γ matrix. Meanwhile, the elements in gamma matrix and γ’ path (matrix between two γ’) will diffuse toward the DRX nuclei, leading to the development of DRX nuclei, as shown in figure 8 (d). As a result, new dislocation-free DRX grains with a size range of 0.4-0.8 μm will be generated along the original boundary after sufficient diffusion.
4 Conclusion 1) Remarkable LAGBs form1000 °C/1.0 s-1 and further develop into HAGBs with the increase of temperature and the decrease of strain rate. DRX can be fully finished upon 1100 °C/1.0 s-1, while an obvious grain growth can be stimulated at the strain rate of 0.001 s-1. 20
2) The fraction of LAGBs increases firstly to a peak with the increase of strain and then decreases gradually with further deformation. Degree of DRX increases with the accumulation of straining and significant DRX was detected when the strain increases to 0.5 at temperature of 1050 °C / 0.01 s-1. 3) Three different nucleation mechanisms were observed at different deformation parameters. Nucleation of DRX was found to form at the bulged-original boundaries which were interacted with mobile dislocations at relatively low temperature and strain rate of 0.1 s-1, while twining boundaries were observed to be preferable sites for nucleation at higher temperature and strain rate of 0.01 s-1. At a temperature of 1100 °C and low strain rate of 0.001 s-1, coalesced γ’ was encompassed by bulged-original boundary, where nucleation of DRX was detected.
Conflicts of interest: The authors declare that there are no conflicts of interest regarding the publication of this paper.
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Acknowledgment Z. H. and L. J. appreciate the financial support from The National Key Research and Development Program of China (2016YFB0700300). G. H. is grateful for the support from the Fundamental Research Funds for the Central Universities of Central South University (2015zzts031) and the outstanding graduate project of Advanced Non-ferrous Metal Structural Materials and Manufacturing Collaborative Innovation Center. F. L. and L. H. are thankful for the funding sponsorship of the Natural Science Foundation of China (51401242, 51301209). We would like to acknowledge the help from Hong JIANG for the helpful technical discussion.
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Figure 1 Typical orientation image microscopy (OIM) map (along direction of HEX) of as-HEXed alloy before hot compression (a); distribution of misorientation of as HEXed alloy (b). Figure 2 Microstructure development of the alloy compressed to a true strain of 0.7 under various deformation conditions. 23
Figure 3 Typical OIMs and the corresponding distribution of misorientation of specimens when deformed at different temperatures and strain rates: (a) and (d) 1000 °C/1.0 s-1; (b) and (e) 1000 °C/0.1 s-1; (c) and (f) 1025 °C/0.01 s-1. Figures (g)~(l) show the misorientations of point to point and point to origin along the initial grain boundaries and from the boundaries to the interior of grains, as marked by lines: (g), A1; (h), B1; (i), C1; (j), A2; (k), B2; (l), C2. Figure 4 TEM observation of the specimen which was deformed under the condition of 1000 °C/1.0 s-1, showing the characteristics of dislocation. Figure 5 The OIMs of specimens deformed to a different of strains under conditions of 1050 °C and 0.01 s-1: (a), 0.1; (b), 0.3; (c), 0.5; (d), 0.7; (e), summary on the variations of misorientation; (f), illustration of variation tendency of misorientation. Figure 6 TEM observation of the specimen deformed under deformation conditions: (a) 1000 °C/0.1 s-1; (b) 1050 °C/0.1 s-1; showing the nucleation of DRX at bulged boundaries with high density of dislocation and substructure. Figure 7 TEM observation of the specimen when deformed under deformation conditions: (a) 1050 °C/0.01 s-1; (b) 1100 °C/0.01 s-1; showing the nucleation of DRX along the twining boundaries. Figure 8 TEM observation of the specimen which was deformed under deformation conditions of 1100 °C/0.001 s-1 ; showing the coalesced γ’ is encompassed by bulged-original boundary, where nucleation of DRX was detected.
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