- Email: [email protected]
M superalloy using processing maps
Materials and Design 87 (2015) 256–265
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/jmad
Characterization of hot compression behavior of a new HIPed nickel-based P/M superalloy using processing maps Guoai He a,b,c, Feng Liu a,b,c,⁎, Jiayong Si c,d, Chuan Yang a,c, Liang Jiang a,b,c a
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China Powder Metallurgy Research Institute, Central South University, Changsha 410083, China High Temperature Materials Research Institute, Central South University, Changsha 410083, China d School of Mechanical &Electrical Engineering, Central South University of Forestry &Technology, Changsha 410004, China b c
a r t i c l e
i n f o
Article history: Received 13 April 2015 Received in revised form 6 August 2015 Accepted 7 August 2015 Available online 12 August 2015 Keywords: Powder metallurgy nickel-based superalloy Hot compression Constitutive equation Processing map Electron back scatter diffraction Dynamic recrystallization
a b s t r a c t Isothermal forging was a critical step process to fabricate the high-performance nickel-based superalloy. The temperature and strain rate served the most critical role in determining its microstructure and mechanical properties. In this article, we employed the hot compression to simulate the isothermal forging process upon the temperature ranging from 1000 °C to 1100 °C in combination with a strain rate of 0.001–1.0 s−1 for a new P/M nickelbased alloy. The activation energy was determined as 903.58 kJ/mol and the processing maps at a strain range of 0.4–0.7 were developed. The instability domains were more inclined to occur at strain rates higher than 0.1 s−1 and manifested in the form of adiabatic shear bands. The map further demonstrated that the regions with peak efficiency of 55% were located at 1080 °C/0.0015 s−1 and 1095 °C/0.014 s−1, respectively. Obvious dynamic recrystallization could be detected at the strain rate 0.01 s−1 leading to a significant flow stress drop and the grain growth was remarkably triggered under 1100 °C. The findings can shed light on the forging processing optimization of the new nickel-based superalloy. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Powder metallurgy (PM) nickel-based superalloy was widely used in the hot section turbine disk of the aeroengine owing to its superior high-temperature strength, creep resistance, and fatigue performance [1–3]. The process technique of superalloys involved atomization of the melt into micro-sized powders, which were subsequently consolidated into near fully dense compacts through hot isostatic pressing (HIP) and/or hot extrusion (HEX) and finally transferred for isothermal forging (IF). These extreme hot deformation processes would cause the occurrence of catastrophic defects, i.e., cracks, critical grain growth and flow localization. Previous researches [4–8] showed that the strain rate, deformation temperature and excursion strain served the primary roles in tuning the microstructures and mechanical properties of forged superalloys. Over the past 30 years, the processing map had been well established as an effective technique to understand hot working behavior of numerous materials [9–15] and nickel-based superalloys using hot compression and hot torsion, for example: Rene 95 [1], IN-625 [3], IN-718 [16], NIMONIC AP-1 [17], IN-100 [18], FGH96 [19], IN-939 [20], and others [21–23]. In this article, we focused on the deformation behaviors of a new designed P/M nickel-based superalloy. ⁎ Corresponding author at: Powder Metallurgy Research Institute, Central South University, Changsha 4100813, China. E-mail address: [email protected] (F. Liu).
http://dx.doi.org/10.1016/j.matdes.2015.08.035 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
The new alloy possessed a higher cobalt (Co) content compared to aforementioned counterparts [24], reaching as high as 20 wt.%. Previous results [25–27] showed the evidence that increasing Co content might improve phase stability, lower the gamma prime solvus, and reduce the quenching residual stress. Recently, a series of new-developed nickel-based superalloys containing up to 30% Co demonstrated an excellent combination of high-temperature strength [28,29], creep resistance [30], low cycle fatigue, and fatigue crack growth rate [31]. Meanwhile, the addition of hafnium (Hf) was favorable for improving grain-boundary stability [2], as well as promoting the formation of a cellular γ′ precipitation [32]. Therefore, it's pivotal to characterize the deformation behavior of this new P/M nickel-based superalloy considering its future application. Herein, we employed the hot compression to investigate the deformation mechanism of this new alloy. The activation energy, constitutive equation and the processing maps were established. The physics for the difference about the activation energy for this alloy and other similar alloy was discussed. Our results unambiguously showed that the notable dynamic recrystallization occurred at a strain rate of 0.01 s−1, resulting in the remarkable decrease of flow stress. The processing map at a strain of 0.7 demonstrated that the peak efficiency upon to 55% at 1095 °C and 0.014 s−1, wherein a lower flow stress could be detected in our alloy. We investigated the instability and stability regions with corresponding microstructures and proposed the optimum hot working parameters for this new alloy.
G. He et al. / Materials and Design 87 (2015) 256–265
2. Experimental work 2.1. Material The basic chemical composition of this new alloy PM-0001 contained Co, Cr, W, Mo, Al, Ti, Nb, Zr, Hf, C, B, and balance was Ni, wherein the sum of Cr + Mo + W was about 21.0 wt.%, Al + Ti about 7.0 wt.%, Co about 26 wt.%, Nb + Zr about 1.0 wt.% with minor amounts of C, B, and Hf. The processing history of the alloy consisted of vacuum induction melting (VIM) and a plasma rotating electrode process (PREP), by which the powder was produced. Then the powders were screened and blended before being loaded into a stainless steel container. Finally, hot isostatic pressing was conducted under the established conditions after the evacuated container was sealed. A section, with a thickness of 10 mm, was cut from an as-HIPed P/M nickel-based alloy billet (140 mm in height and 90 mm in diameter) in the direction of cross-section. Cylindrical specimens, with a diameter of 6 mm and a height of 9 mm, were machined from the center of the billet to ensure a uniform of microstructure across all samples. The microstructure of as-HIPed PM-0001 was shown in Fig. 1, which demonstrated clearly the existence of prior particle boundary (PPB), one of the three metallurgy defects in a P/M alloy. 2.2. Hot compression test and microstructure examination Isothermal compression tests were conducted under constant strain rates of 1.0, 0.1, 0.01 and 0.001 s− 1 at temperatures ranging from 1000 °C to 1100 °C in increments of 25 °C using Gleeble 3180D. Many temperature adjustment tests with three thermocouples welded on the specimen (shown in Fig. 2(a)) were conducted to ensure a uniform temperature along the length of the specimen prior to compression. The specimens were heated at 5 °C/s and soaked for 2 min at the deformation temperature before compressing to a final true strain of 0.7, and then quenched upon completion of their deformation. The deformed specimens were sectioned parallel to the compression axis and prepared for optical and electronic observation using standard metallographic techniques. An optical microscope (LEICA-DM4000M) and a field-emission gun SEM, equipped with an electron backscatter diffraction (EBSD) detector and Channel 5 software, were used to characterize the deformed microstructure. The parameters used were as follows: accelerating voltage 25 kV, step size 0.05, 0.1 μm, and a spot size of 6.0. The etchant used was a solution of 100 ml HCl, 100 ml ethanol, and 5 g CuCl2 [33].
257
accuracy in actual stress, as can be seen in Fig. 2(b). Fig. 3 demonstrated the true stress-true strain curves of the alloy deformed at different temperatures and strain rates. Typically, the trend of the curves exhibited a sharp increase with increasing strain before hitting a peak stress within a small strain range and then decreased gradually to a steady state at a strain greater than about 0.3. Generally, material deformation at elevated temperatures was a competing process of dynamic softening and work hardening. At the beginning of the deformation, the dislocation density increased and work hardening dominated the hot working process leading to the rapid increase of flow stress. As the deformation proceeded, the flow stress reached a peak meaning that work hardening, as the dominant process, came to an end. Then the softening due to dynamic recovery and recrystallization started to dominate the deformation process, leading to a gradual decrease in flow stress [34]. The flow stress, which remained stable with increasing strain, could be defined as a steady stress, indicating that work hardening and softening reached equilibrium. As Fig. 3 showed, for a particular strain rate, the flow stress decreased with increasing deformation temperature. At a given deformation temperature, the flow stress increased with increasing strain rate owing to an insufficiency of time for dislocation annihilation at higher strain rates [3,27]. At a strain rate of 0.01 s−1 and temperatures of 1000 °C/1025 °C, the flow stress decreased significantly after the strain reached 0.2 compared to higher strain rates (1.0/0.1 s−1), as shown in Fig. 3(a), (b) and (c). The mechanism could be attributed to the occurrence of significant dynamic recrystallization, which would be validated by further microstructural observation in the following sections. In addition, the flow stress of the specimen deformed at 1100 °C and 0.01 s−1 was approximately 60 MPa, which was lower than other asHIPed P/M nickel-based superalloys [17–19,35,36] under the similar deformation parameters, indicating that this new alloy would exhibit superior workability. Further work was necessary to verify the hypothesis. Table 1 listed the steady flow stress for samples deformed at 1100 °C and 0.01 s− 1 for other as-HIPed P/M nickel-based superalloys. 3.1. Kinetic analysis of the deformation During deformation, the peak stress could be related to strain rate and deformation temperature through the Arrhenius kinetic rate equation [37]: n expð−Q =RTÞ ε ¼ A sinh ασ p
ð1Þ
3. Results and discussion 3.0.1. Flow behavior under hot deformation The result of temperature calibration showed the superior uniformity of temperature along the whole specimen, which would improve the
where ε was the strain rate, A and α were material constants, σp was the peak stress, n was the stress exponent, Q was the activation energy for hot working, R was the gas constant, and T was the absolute
Fig. 1. Initial microstructure obtained for the as-HIPed nickel-based superalloy: (a) at low magnification; (b) at high magnification. Fig. 1 shows obviously the existence of prior particle boundary, one of three metallurgy defects in powder metallurgy alloy.
258
G. He et al. / Materials and Design 87 (2015) 256–265
Fig. 2. (a) A specimen with three thermocouples welded thereon; (b) the results of the temperature adjustment at 1100 °C. Fig. 2b shows the results of temperature adjustment at the deformation temperature of 1100 °C, which indicates the superior uniformity of temperature along the whole specimen.
deformation temperature. The activation energy Q could then be defined as: " Q ¼ R
∂ ln ε ∂ ln sinh ασ p
# ∂ ln sinh ασ p : ∂ð1=TÞ ε
this study, the stress exponent n was strain rate dependent and could be calculated from the lnσ-ln ε plot using linear regression. The plot of peak stress σp versus ln ε was shown in Fig. 4(b). The optimum value of constant α could be determined by linear regression to fit the relationships of ln ε-lnσp and σp-ln ε. These two plots together with the plots of ln[sinh(ασp)]-(1000/T) and ln[sinh(ασp)]-ln ε (see Fig. 4(c) and (d)) were adopted here to obtain the deformation activation energy. The deformation activation energy of the studied alloy and other as-HIPed P/M nickel-based superalloys [36,38–40] was listed for comparison in
ð2Þ
T
In this investigation, the variation of the natural logarithm of the peak stress with strain rate at various deformation temperatures was presented in Fig. 4(a). Over the entire range of strain rates covered in
Fig. 3. Typical flow stress–strain curves of nickel-based superalloy PM-0001 deformed at (a) 1.0 s−1; (b) 0.1 s−1; (c) 0.01 s−1; and (d) 0.001 s−1.
G. He et al. / Materials and Design 87 (2015) 256–265
259
Table 1 Steady stress of other as-HIPed P/M Ni-base superalloys deformed at 1100 °C/0.01 s−1. Author, year
Alloy/initial microstructure (as-HIPed)
HIP condition
Steady stress at 1100 °C, 0.01 s−1/MPa
Ref.
Somani M. 1992 Somani M. 1998 Ning Y., 2010 Xie X.H., 2002 Xu Wei, 2006 He G.A., 2015
NIMONIC-AP-1 IN100 FGH4096 FGH4096 FGH4096 PM-0001
1180 °C/108 MPa/3 h 1200 °C/115 MPa/3 h 1030 °C/120 MPa/1 h + 1170 °C/140 MPa/2 h / / 1100 °C/140 MPa/4 h
123.2 122 80–120 85 (roughly) 80 (roughly) 62
[17] [18] [19] [35] [36] This work
Table 2. The average activation energy for our alloy was somewhat higher than that of similar reported superalloys. The deformation activation energy of alloy was influenced by many factors, such as material chemistry, preparation method, as well as processing history and so on. Leverant and Kear [41] reported that the rate-controlling deformation process might involve self-diffusion in the Ni–Co based alloy, which was affected by interchange of Co and Ni atoms. The diffusion activation energy increased with increasing Co concentration, mechanism of which could be ascribed to a decrease of atomic order degree due to a substitution of Ni atoms by Co atoms [42]. The content of Co in our alloy was higher than that of other similar nickel-based superalloys [24], which served the dominated factor leading to a higher value of Q in our alloy. Moreover, additional energy was necessary for eliminating the inherent defects (i.e., PPB, as shown in Fig. 1) in the process of HIP, leading to a lift of deformation activation energy. The temperature compensated strain rate parameter, or the Zener– Hollomon parameter, could be defined as follows [43]: Q Z ¼ ε exp RT
ð3Þ
or, n Z ¼ A sinh ασ p :
ð4Þ
Fig. 5 showed the relationship between Zener–Hollomon parameter and [sinh(ασp)]n. This plot indicated that the superior linear relationship between ln Z and ln [sinh(ασp)], which was adopted to determined the constant A and exponent n. The value of n could be obtained by calculating the slope of the ln Z-ln [sinh(ασp)] plot and the constant A could be determined as 1.211 × 10 33 based on Eq. (4). Finally, substituted the values of α, A, n, and Q into Eq. (1), the constitutive equation of hot deformation for alloy could be expressed as:
ε ¼ 1:211 1033 ½ sinhð0:0062σ pÞ
3:011
903580 : exp − RT
ð5Þ
Fig. 4. Relationships between (a) ln σp and ln ε; (b) σp and ln ε; (c) ln [sinh(ασp)] and (1000/T); (d) ln [sinh(ασp)] and ln ε. These four plots are adopted here to obtain the constant α and deformation activation energy.
260
G. He et al. / Materials and Design 87 (2015) 256–265
transformation, and precipitation under dynamic conditions, and could be evaluated by integrating the following equation [18]:
Table 2 Activation energy for some other as-HIPed P/M Ni-Base superalloys. Alloy/initial microstructure
Temperature range (°C)
PM-0001/HIPed FGH96/HIPed FGH96/HIPed FGH96/HIPed FGH4096/HIPed
1000–1100 1000–1100 1000–1100 1000–1150 1080–1140
Strain rate range (s−1) 0.01–1.0 0.001–0.1 0.001–0.1 0.001–1.0 0.02–1.0
Total true strain 0.7 0.5 0.5 1.0 0.69
Activation energy (kJ/mol)
Ref.
903.58 754 740.4 693.21 595
This work [38] [39] [36] [40]
Z J¼
σ 0
mσε : 1þm
εdσ ¼
ð9Þ
For an ideal linear dissipater, m = 1 and J ¼ Jmax ¼ σ ε =2. For a non-linear dissipater, the efficiency of power dissipation could be expressed in terms of a dimensionless parameter [18]: η ¼ 2m=ðm þ 1Þ:
3.2. Processing maps Generally, the processing maps consisted of a power dissipation map and an instability map, and could be divided into safe domains and unsafe domains. In this model, the total power P might be divided into two complementary functions: G content and J co-content [44]: Z P¼Gþ J¼
ε 0
Z σdε þ
σ 0
εdσ:
ð6Þ
The strength of materials increased with increasing plastic strain rate [34]. The rate sensitive flow behavior was given [45] by:
ð10Þ
The efficiency of power dissipation η represented the proportion of the power dissipated by microstructure evolution over total power. Increase of η implied an increase in the power dissipated by microstructure evolution. The iso-efficiency contours of the power dissipation map, in which η varied with deformation temperature and strain rate, form the power dissipation map. This map exhibited different domains correlated with specific microstructures [46]. The continuum instability criterion based on the extremum principle of irreversible thermodynamic had been developed to identify the domains of flow instabilities, and given by another dimensionless parameter [18]: m ∂ ln m þ 1 ξ ε ¼ þ m b 0: ∂ ln ε
ð11Þ
m
σ ¼kε
ð7Þ
where σ was the flow stress, ε was the strain rate, m was the strain rate sensitivity, and k was a constant. The value of m was determined by a variant of Eq. (8) given at a constant strain and deformation temperature [17]:
∂ðlnσ Þ
: m¼
∂ ln ε
ε;T
ð8Þ
The G term mentioned in Eq. (6) referred to the power dissipated by plastic work, most of which would be converted into visco-plastic heat and the remaining power would be stored as lattice defects. However, the J term was related to metallurgical mechanisms such as dynamic recovery and recrystallization, deformation-induced phase
Fig. 5. Relationship between ln Z and ln [sinh(ασp)]. This plot indicates that the superior linear relationship between ln Z and ln [sinh(ασp)], which is adopted to determined the constant A.
The variation of instability parameter ξðεÞ might be evaluated as a function of deformation temperature and strain rate to obtain an instability map. The metallurgical instability during plastic flow occurred in domains where ξðεÞ was negative. Fig. 6 showed the processing maps of the alloy generated in the temperature range from 1000 to 1100 °C and strain rate range from 0.001 to 1.0 s−1 at different strains (from 0.4 to 0.7). As can be seen in Fig. 6, the stability domains were more likely to exhibit higher η, indicating more power dissipated by microstructure evolution. The instability regions were mainly located at the deformation conditions of high strain rates (N0.1 s−1) and/or low deformation temperature. Both of the two domains would be discussed in the following sections.
3.2.1. Instability domain In general, the maps at high strain rates were more inclined to exhibit instability domains regardless of strain and temperature. Fig. 6 showed that the material tended to exhibit a larger area instability domain with accumulation of strain. For the case of a strain of 0.7, the unsafe domain was located within the deformation temperature range from 1000 to 1100 °C and strain rate range from 0.1 s− 1 to 1.0 s−1. The mechanism of instability had been reported as being probably associated with cracking, located plastic flow, or adiabatic shear bands. The macroscopic features of specimens deformed at different temperatures and strain rates were presented in Fig. 7. As can be seen, there were no visible cracks in any of the specimen surfaces. The surface of all of the specimens was seen to be smooth and did not exhibit severe barrelling. However, the appearance of the specimens deformed at higher strain rates and lower temperatures tended to be more irregular when compared to other deformation conditions. Almost all specimens deformed at high strain rates (1.0 s−1) had an irregular shape regardless of their deformation temperature. The shape tended to be increasingly regular with decreasing strain rate, which matched the evidence of the processing maps. If the plasticity of the material was worse, shear deformation along an axis at about 45° to the direction of compression could be seen [34]. This shear deformation would result in the formation of shear bands and was consistent with the lower η in processing map. The reason might be that there was more adiabatic shear deformation heat dissipating in the shear deformation zone. Consequently, the processing
G. He et al. / Materials and Design 87 (2015) 256–265
261
Fig. 6. Processing maps at different strains: (a) 0.4; (b), 0.5; (c) 0.6; (d) 0.7. The shadows represent the instability domains.
map of a material was somehow related to its plasticity, and even workability: the higher the value of η, the better plasticity and workability of the material. Fig. 8 showed the microstructures of the specimen deformed at 1025 °C/1.0 s− 1 and at a strain of 0.7. The specimen exhibited local shear bands, which were characteristics of shear deformation, oriented at about 45° with respect to the deformation axis. A similar situation could be seen in other specimens deformed at this domain. The formation of shear bands during deformation was related to the generation of adiabatic heat associated with the lower thermal conductivity in comparison to other materials [47,48]. The mechanism might be attributed to the insufficient time available for releasing the deformation heat when the specimen was deformed at higher strain rates, leading to a temperature increase in regions with severe
Fig. 7. Macroscopic features of specimens deformed under different conditions, at a strain of 0.7. Specimens exhibit irregular appearance deformed at high strain rate and the shape become more and more regular with decreasing of strain rate.
deformation. The accumulation of deformation heat would result in the formation of bands of flow localization, which stimulated the occurrence of flow instability.
3.2.2. Stability domain For a strain of 0.7, one domain was located at 1080 °C/0.0015 s−1 with a peak efficiency of 55%, and the other occurred at about 1095 °C/0.014 s− 1 with a peak efficiency of 55%. The higher η meant a better workability at this domain, wherein the optimum hot working parameters for this alloy may be located [49].
Fig. 8. Microstructure deformed within the instability domain of 1025 °C s−1, at a strain of 0.7. Inside the dash line shows the characteristics of shear deformation, and the shear band can be seen clearly.
262
G. He et al. / Materials and Design 87 (2015) 256–265
Fig. 9. OIM maps of specimens obtained by EBSD under different deformation conditions: (a) 1025 °C/1.0 s−1; (b) 1050 °C/1.0 s−1; (c) 1075 °C/1.0 s−1; (d) 1025 °C/0.1 s−1; (e) 1050 °C/ 0.1 s−1; (f) 1075 °C/0.1 s−1; (g) 1025 °C/0.01 s−1; (h) 1050 °C/0.01 s−1; and (i) 1075 °C/0.01 s−1, at a strain of 0.7. Fig. 9 showed the effects of strain rates and temperature on microstructure evolution. The degree of dynamic recrystallization increased with increasing of temperature and decreasing of strain rate.
The orientation imaging microscopy (OIM) of the deformed specimens obtained by EBSD was shown in Fig. 9, among which Figs (d) to (i) were deformed at stability domains. Fig. 9 demonstrated the effects of deformation parameters on the degree of dynamic recrystallization and microstructure evolution. It might be noted that during deformation, the boundaries would become serrated and bulged before developing into the familiar ‘necklace’ structure: these were indicative of dynamic recrystallization in nickel-based superalloys [46]. As shown in Fig. 9, the deformed grains occurred to recrystallize at 1050 °C for strain rate of 1.0 s−1 and the degree of dynamic recrystallization tended to increase with increasing deformation temperature and decreasing strain rate. The physics behind this could be ascribed to more energy and the sufficient time available for nucleation as well as growth of dynamic recrystallization. Fig. 10 demonstrated the detailed features of grain boundary caused by dynamic recrystallization and microstructure evolution. The EBSD observation at 1050 °C/1.0 s−1 (Fig. 10(a)) demonstrated that the grain boundaries bulged and became to be serrated, implying the occurrence of dynamic recrystallization. At the grain boundary, the necklace-like microstructure was ubiquitously observed, showing that the grain boundary acted as the nucleation sites of dynamic recrystallization. In accompanying with decrease of strain rate, the longer compression time could facilitate the dynamic recrystallization, as shown in Fig. 10(b), till to be completed. As the temperature
increased to 1075 °C (see Fig. 9(i)), almost all of the grains had undergone dynamic recrystallization to yield a fine-grained microstructure at a strain rate of 0.01 s−1. With strain rate further decreased, some of the recrystallized grains occurred to grow up, as shown in Fig. 10(c). Moreover, the higher temperature could accelerate the grain growth and reach a polygonal morphology (Fig. 10(d)). Fig. 11 revealed the deformation microstructure of dynamic recrystallization under the temperature of 1100 °C, indicating the effects of strain rate on grain growth. The EBSD micrographs showed that the deformed grains growth occurred and developed into equiaxed morphology. While the strain rate was decreased to be 0.001 s− 1, the grain growth could be obviously observed, which was not favorable for microstructure design requirement of nickel-based superalloy. Therefore, the parameter of 1100 °C/0.001 s− 1 was advisable to be avoided for the actual forging processing. The significant decrease in stress after reaching peak stress (shown in Fig. 3) might be attributed to the formation of new, fine, recrystallized grain boundaries (see Fig. 9(g)). Considerable stress concentrations formed at grain boundary triple junctions during deformation: these would be released by grain boundary sliding and would accelerate if there were a considerable number of grain boundaries present [50]. Consequently, both a lower strain rate and a higher deformation temperature made a remarkable contribution to accelerating the grain
G. He et al. / Materials and Design 87 (2015) 256–265
263
Fig. 10. High-magnification microstructure of specimen under different deformation conditions: (a) 1050 °C/1.0 s− 1 ; (b) 1025 °C/0.1 s − 1 ; (c) 1075 °C/0.001 s− 1 ; and (d) 1100 °C/0.001 s−1, at a strain of 0.7. Fig. 10 showed the High-magnification microstructural features and microstructure evolution.
Fig. 11. OIM maps of the alloy deformed at 1100 °C: (a) 1.0 s−1; (b) 0.1 s−1; (c) 0.01 s−1; and (d) 0.001 s−1, at a strain of 0.7. Fig. 11 showed the effects of strain rate on degree of grain growth for the alloy at 1100 °C.
264
G. He et al. / Materials and Design 87 (2015) 256–265
boundary sliding process, which was considered to be the main reason for the softening that partially offset the deformation-induced stress. From the aforementioned discussion, it might be concluded that both deformation temperature and strain rate exerted a significant effect on dynamic recrystallization and microstructure evolution. The material exhibited more dynamic recrystallization grains under low strain rate and high temperature deformation conditions. The optimum deformation conditions for this alloy might be at 1075 to 1095 °C and 0.01 to 0.04 s− 1, at which the material exhibits favorably complete recrystallizated microstructure and potential superior workability. 4. Conclusion 1. Obvious dynamic recrysallization could be readily detected at the strain rate of 0.01 s−1, resulting in the significant flow stress decrease compared to 0.1 and 1.0 s−1. 2. The activation energy was determined to be 903.58 kJ/mol and a constitutive equation for the new alloy had been developed as follows: ε ¼ 1:211 1033 ½ sinhð0:0062σ pÞ
3:011
903580 exp − RT
3. The flow stress instability occurred at strain rate higher than 0.1 s−1 and temperature range from 1000 °C to 1100 °C, the adiabatic shear bands acted on the leading failure mode. 4. The peak efficiency of 55% upon a strain of 0.7 could be found at 1080 °C/0.0015 s− 1 and 1095 °C/0.014 s− 1. The grain growth was remarkably triggered under 1100 °C. 5. For our new alloy, the forging conditions could be optimized to the temperature range of 1075 °C–1095 °C and strain rate range of 0.014–0.04 s−1. Acknowledgment This work was supported by The National High-Tech Research and Development Programme of China, Grant no. 2012AA03A514, the Doctoral-Independent Exploration and Innovation of Central South University, Grant no. 2015zzts031 and the outstanding graduate project of Advanced Non-ferrous Metal Structural Materials and Manufacturing Collaborative Innovation Center. The author, Liu Feng, thanks The Natural Science Foundation of China for their financial support under Grant nos. 51401242, 61271356, and 51301209, and The Postdoctoral Science Foundation of Central South University for their grant no. 74341016096. References [1] M.O. Alniak, F. Bedir, Modelling of deformation and microstructural changes in P/M Rene 95 under isothermal forging conditions, Mater. Sci. Eng. A 429 (2006) 295–303. [2] M. Donachie, S. Donachie, Superalloys—A Technical Guide, ASM International, Materials Park, OH, 2002. [3] D. Li, Q. Guo, S. Guo, H. Peng, Z. Wu, The microstructure evolution and nucleation mechanisms of dynamic recrystallization in hot-deformed Inconel 625 superalloy, Mater. Des. 32 (2011) 696–705. [4] M. Zhang, F. Li, Z. Yuan, J. Li, S. Wang, Effect of heat treatment on the microindentation behavior of powder metallurgy nickel based superalloy FGH96, Mater. Des. 49 (2013) 705–715. [5] X.-M. Chen, Y.C. Lin, D.-X. Wen, J.-L. Zhang, M. He, Dynamic recrystallization behavior of a typical nickel-based superalloy during hot deformation, Mater. Des. 57 (2014) 568–577. [6] H. Zhang, K. Zhang, H. Zhou, Z. Lu, C. Zhao, X. Yang, Effect of strain rate on microstructure evolution of a nickel-based superalloy during hot deformation, Mater. Des. 80 (2015) 51–62. [7] X.-M. Chen, Y.C. Lin, M.-S. Chen, H.-B. Li, D.-X. Wen, J.-L. Zhang, et al., Microstructural evolution of a nickel-based superalloy during hot deformation, Mater. Des. 77 (2015) 41–49. [8] N.K. Park, I.S. Kim, Y.S. Na, J.T. Yeom, Hot forging of a nickel-base superalloy, J. Mater. Process. Technol. 111 (2001) 98–102. [9] H. Sun, Y. Sun, R. Zhang, M. Wang, R. Tang, Z. Zhou, Study on hot workability and optimization of process parameters of a modified 310 austenitic stainless steel using processing maps, Mater. Des. 67 (2015) 165–172.
[10] Z. Yang, F. Zhang, C. Zheng, M. Zhang, B. Lv, L. Qu, Study on hot deformation behaviour and processing maps of low carbon bainitic steel, Mater. Des. 66 (Part A) (2015) 258–266. [11] A. Łukaszek-Sołek, J. Krawczyk, The analysis of the hot deformation behaviour of the Ti–3Al–8 V–6Cr–4Zr–4Mo alloy, using processing maps, a map of microstructure and of hardness, Mater. Des. 65 (2015) 165–173. [12] H. Liao, Y. Wu, K. Zhou, J. Yang, Hot deformation behavior and processing map of Al– Si–Mg alloys containing different amount of silicon based on Gleebe-3500 hot compression simulation, Mater. Des. 65 (2015) 1091–1099. [13] M. Karami, R. Mahmudi, Hot shear deformation constitutive analysis and processing map of extruded Mg–12Li–1Zn bcc alloy, Mater. Des. 53 (2014) 534–539. [14] G. Liu, Y. Han, Z. Shi, J. Sun, D. Zou, G. Qiao, Hot deformation and optimization of process parameters of an as-cast 6Mo superaustenitic stainless steel: a study with processing map, Mater. Des. 53 (2014) 662–672. [15] K.-k. Deng, J.-c. Li, F.-j. Xu, K.-b. Nie, W. Liang, Hot deformation behavior and processing maps of fine-grained SiCp/AZ91 composite, Mater. Des. 67 (2015) 72–81. [16] F.-L. Sui, L.-X. Xu, L.-Q. Chen, X.-H. Liu, Processing map for hot working of Inconel 718 alloy, J. Mater. Process. Technol. 211 (2011) 433–440. [17] M. Somani, E.B. Rao, N. Birla, M. Bhatia, V. Singh, Y. Prasad, Processing map for controlling microstructure in hot working of hot isostatically pressed powder metallurgy NIMONIC AP-1 superalloy, Metall. Trans. A 23 (1992) 2849–2857. [18] M. Somani, K. Muraleedharan, Y. Prasad, V. Singh, Mechanical processing and microstructural control in hot working of hot isostatically pressed P/M IN-100 superalloy, Mater. Sci. Eng. A 245 (1998) 88–99. [19] Y. Ning, Z. Yao, H. Guo, M.W. Fu, H. li, X. Xie, Investigation on hot deformation behavior of P/M Ni-base superalloy FGH96 by using processing maps, Mater. Sci. Eng. A 527 (2010) 6794–6799. [20] M. Jahangiri, H. Arabi, S. Boutorabi, High-temperature compression behavior of cast and homogenized IN939 superalloy, Metall. Mater. Trans. A 44 (2013) 1827–1841. [21] K. Wu, G. Liu, B. Hu, F. Li, Y. Zhang, Y. Tao, et al., Hot compressive deformation behavior of a new hot isostatically pressed Ni–Cr–Co based powder metallurgy superalloy, Mater. Des. 32 (2011) 1872–1879. [22] P. Zhang, C. Hu, C.-g. Ding, Q. Zhu, H.-y. Qin, Plastic deformation behavior and processing maps of a Ni-based superalloy, Mater. Des. 65 (2015) 575–584. [23] Y. Wang, Q. Pan, Y. Song, C. Li, Z. Li, Hot deformation and processing maps of X-750 nickel-based superalloy, Mater. Des. 51 (2013) 154–160. [24] Y. Gu, H. Harada, C. Cui, D. Ping, A. Sato, J. Fujioka, New Ni–Co-base disk superalloys with higher strength and creep resistance, Scr. Mater. 55 (2006) 815–818. [25] J. Heslop, Wrought Nickel–Chromium Heat-Resisting Alloys Containing Cobalt, Cobblestone 24 (1964) 128. [26] J. Tien, T. Howson, G. Chen, X. Xie, Cobalt availability and superalloys, J. Meteorol. 32 (1980) 12–20. [27] R.N. Jarrett, J.K. Tien, Effects of cobalt on structure, microchemistry and properties of a wrought nickel-base superalloy, Metall. Trans. A 13 (1982) 1021–1032. [28] C.Y. Cui, Y.F. Gu, Y. Yuan, T. Osada, H. Harada, Enhanced mechanical properties in a new Ni–Co base superalloy by controlling microstructures, Mater. Sci. Eng. A 528 (2011) 5465–5469. [29] Y. Yuan, Y.F. Gu, Z.H. Zhong, T. Yokokawa, H. Harada, Enhanced strength at intermediate temperatures in a Ni-base disk superalloy with high Co addition, Mater. Sci. Eng. A 556 (2012) 595–600. [30] Y. Yuan, Y.F. Gu, Z.H. Zhong, T. Osada, C.Y. Cui, T. Tetsui, et al., Creep mechanisms of a new Ni–Co-base disc superalloy at an intermediate temperature, J. Microsc. 248 (2012) 34–41. [31] Z. Zhong, Y. Gu, Y. Yuan, C. Cui, T. Yokokawa, H. Harada, Fatigue crack growth behavior of a newly developed Ni–Co-base superalloy TMW-2 at elevated temperatures, Mater. Sci. Eng. A 552 (2012) 464–471. [32] R. Miner, Effects of C and Hf concentration on phase relations and microstructure of a wrought powder-metallurgy superalloy, Metall. Trans. A 8 (1977) 259–263. [33] Voort G. Vander, Metallography of superalloys, Ind. Heat. 70 (2003) 40–43. [34] Q.L. Pan, B. Li, Y. Wang, Y.W. Zhang, Z.M. Yin, Characterization of hot deformation behavior of Ni-base superalloy Rene'41 using processing map, Mater. Sci. Eng. A 585 (2013) 371–378. [35] X. Xie, Z. Yao, Y. Ning, H. Guo, Y. Tao, Y. Zhang, Investigation on hot deformation behavior of P/M superalloy FGH4096, Rare Metal Mater. Eng. 41 (2012) 82–86. [36] W. Xu, L.-w. Zhang, S.-d. Gu, J.-l. Zhang, Hot compressive deformation behavior and microstructure evolution of HIPed FGH96 superalloy, Trans. Nonferrous Metals Soc. 22 (2012) 66–71. [37] C. Bruni, A. Forcellese, F. Gabrielli, Hot workability and models for flow stress of NIMONIC 115 Ni-base superalloy, J. Mater. Process. Technol. 125–126 (2002) 242–247. [38] Y. Liu, F. Li, H. Yu, Deformation behavior of hot isostatic pressing FGH96 superalloy, Trans. Tianjin Univ. 12 (2006) 281–285. [39] R. Zhang, F. Li, X. Wang, Determining processing maps of FGH96 superalloy, J. Norhtwestern Polytechnical Univ. 25 (2007) 652–656. [40] Y. Ning, Z. Yao, M.W. Fu, H. Guo, Dynamic recrystallization of the hot isostatically pressed P/M superalloy FGH4096 in hot working process, Mater. Sci. Eng. A 527 (2010) 6968–6974. [41] G. Leverant, B. Kear, The mechanism of creep in gamma prime precipitationhardened nickel-base alloys at intermediate temperatures, Metall. Mater. Trans. B 1 (1970) 491–498. [42] A. Gazda, J. Čermák, Diffusion in Ni3Al-Co intermetallic, Metal 5 (2002) 14–16. [43] N.-K. Park, J.-T. Yeom, Y.-S. Na, Characterization of deformation stability in hot forging of conventional Ti–6Al–4 V using processing maps, J. Mater. Process. Technol. 130 (2002) 540–545. [44] H.R. Ezatpour, S.A. Sajjadi, M. Haddad-Sabzevar, G.R. Ebrahimi, Hot deformation and processing maps of K310 cold work tool steel, Mater. Sci. Eng. A 550 (2012) 152–159.
G. He et al. / Materials and Design 87 (2015) 256–265 [45] D. Cai, L. Xiong, W. Liu, G. Sun, M. Yao, Characterization of hot deformation behavior of a Ni-base superalloy using processing map, Mater. Des. 30 (2009) 921–925. [46] Y. Prasad, S. Sasidhara, Hot Working Guide: A Compendium of Processing Maps, ASM International, 1997. [47] Y. Han, G. Liu, D. Zou, R. Liu, G. Qiao, Deformation behavior and microstructural evolution of as-cast 904 L austenitic stainless steel during hot compression, Mater. Sci. Eng. A 565 (2013) 342–350.
265
[48] W. Peng, W. Zeng, Q. Wang, H. Yu, Characterization of high-temperature deformation behavior of as-cast Ti60 titanium alloy using processing map, Mater. Sci. Eng. A 571 (2013) 116–122. [49] J.K. Chakravartty, G.K. Dey, S. Banerjee, Y.V.R.K. Prasad, Characterization of hot deformation behaviour of Zr–2.5Nb–0.5Cu using processing maps, J. Nucl. Mater. 218 (1995) 247–255. [50] S. Medeiros, W. Frazier, Y. Prasad, Hot deformation mechanisms in a powder metallurgy nickel-base superalloy IN 625, Metall. Mater. Trans. A 31 (2000) 2317–2325.
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/jmad
Characterization of hot compression behavior of a new HIPed nickel-based P/M superalloy using processing maps Guoai He a,b,c, Feng Liu a,b,c,⁎, Jiayong Si c,d, Chuan Yang a,c, Liang Jiang a,b,c a
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China Powder Metallurgy Research Institute, Central South University, Changsha 410083, China High Temperature Materials Research Institute, Central South University, Changsha 410083, China d School of Mechanical &Electrical Engineering, Central South University of Forestry &Technology, Changsha 410004, China b c
a r t i c l e
i n f o
Article history: Received 13 April 2015 Received in revised form 6 August 2015 Accepted 7 August 2015 Available online 12 August 2015 Keywords: Powder metallurgy nickel-based superalloy Hot compression Constitutive equation Processing map Electron back scatter diffraction Dynamic recrystallization
a b s t r a c t Isothermal forging was a critical step process to fabricate the high-performance nickel-based superalloy. The temperature and strain rate served the most critical role in determining its microstructure and mechanical properties. In this article, we employed the hot compression to simulate the isothermal forging process upon the temperature ranging from 1000 °C to 1100 °C in combination with a strain rate of 0.001–1.0 s−1 for a new P/M nickelbased alloy. The activation energy was determined as 903.58 kJ/mol and the processing maps at a strain range of 0.4–0.7 were developed. The instability domains were more inclined to occur at strain rates higher than 0.1 s−1 and manifested in the form of adiabatic shear bands. The map further demonstrated that the regions with peak efficiency of 55% were located at 1080 °C/0.0015 s−1 and 1095 °C/0.014 s−1, respectively. Obvious dynamic recrystallization could be detected at the strain rate 0.01 s−1 leading to a significant flow stress drop and the grain growth was remarkably triggered under 1100 °C. The findings can shed light on the forging processing optimization of the new nickel-based superalloy. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Powder metallurgy (PM) nickel-based superalloy was widely used in the hot section turbine disk of the aeroengine owing to its superior high-temperature strength, creep resistance, and fatigue performance [1–3]. The process technique of superalloys involved atomization of the melt into micro-sized powders, which were subsequently consolidated into near fully dense compacts through hot isostatic pressing (HIP) and/or hot extrusion (HEX) and finally transferred for isothermal forging (IF). These extreme hot deformation processes would cause the occurrence of catastrophic defects, i.e., cracks, critical grain growth and flow localization. Previous researches [4–8] showed that the strain rate, deformation temperature and excursion strain served the primary roles in tuning the microstructures and mechanical properties of forged superalloys. Over the past 30 years, the processing map had been well established as an effective technique to understand hot working behavior of numerous materials [9–15] and nickel-based superalloys using hot compression and hot torsion, for example: Rene 95 [1], IN-625 [3], IN-718 [16], NIMONIC AP-1 [17], IN-100 [18], FGH96 [19], IN-939 [20], and others [21–23]. In this article, we focused on the deformation behaviors of a new designed P/M nickel-based superalloy. ⁎ Corresponding author at: Powder Metallurgy Research Institute, Central South University, Changsha 4100813, China. E-mail address: [email protected] (F. Liu).
http://dx.doi.org/10.1016/j.matdes.2015.08.035 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
The new alloy possessed a higher cobalt (Co) content compared to aforementioned counterparts [24], reaching as high as 20 wt.%. Previous results [25–27] showed the evidence that increasing Co content might improve phase stability, lower the gamma prime solvus, and reduce the quenching residual stress. Recently, a series of new-developed nickel-based superalloys containing up to 30% Co demonstrated an excellent combination of high-temperature strength [28,29], creep resistance [30], low cycle fatigue, and fatigue crack growth rate [31]. Meanwhile, the addition of hafnium (Hf) was favorable for improving grain-boundary stability [2], as well as promoting the formation of a cellular γ′ precipitation [32]. Therefore, it's pivotal to characterize the deformation behavior of this new P/M nickel-based superalloy considering its future application. Herein, we employed the hot compression to investigate the deformation mechanism of this new alloy. The activation energy, constitutive equation and the processing maps were established. The physics for the difference about the activation energy for this alloy and other similar alloy was discussed. Our results unambiguously showed that the notable dynamic recrystallization occurred at a strain rate of 0.01 s−1, resulting in the remarkable decrease of flow stress. The processing map at a strain of 0.7 demonstrated that the peak efficiency upon to 55% at 1095 °C and 0.014 s−1, wherein a lower flow stress could be detected in our alloy. We investigated the instability and stability regions with corresponding microstructures and proposed the optimum hot working parameters for this new alloy.
G. He et al. / Materials and Design 87 (2015) 256–265
2. Experimental work 2.1. Material The basic chemical composition of this new alloy PM-0001 contained Co, Cr, W, Mo, Al, Ti, Nb, Zr, Hf, C, B, and balance was Ni, wherein the sum of Cr + Mo + W was about 21.0 wt.%, Al + Ti about 7.0 wt.%, Co about 26 wt.%, Nb + Zr about 1.0 wt.% with minor amounts of C, B, and Hf. The processing history of the alloy consisted of vacuum induction melting (VIM) and a plasma rotating electrode process (PREP), by which the powder was produced. Then the powders were screened and blended before being loaded into a stainless steel container. Finally, hot isostatic pressing was conducted under the established conditions after the evacuated container was sealed. A section, with a thickness of 10 mm, was cut from an as-HIPed P/M nickel-based alloy billet (140 mm in height and 90 mm in diameter) in the direction of cross-section. Cylindrical specimens, with a diameter of 6 mm and a height of 9 mm, were machined from the center of the billet to ensure a uniform of microstructure across all samples. The microstructure of as-HIPed PM-0001 was shown in Fig. 1, which demonstrated clearly the existence of prior particle boundary (PPB), one of the three metallurgy defects in a P/M alloy. 2.2. Hot compression test and microstructure examination Isothermal compression tests were conducted under constant strain rates of 1.0, 0.1, 0.01 and 0.001 s− 1 at temperatures ranging from 1000 °C to 1100 °C in increments of 25 °C using Gleeble 3180D. Many temperature adjustment tests with three thermocouples welded on the specimen (shown in Fig. 2(a)) were conducted to ensure a uniform temperature along the length of the specimen prior to compression. The specimens were heated at 5 °C/s and soaked for 2 min at the deformation temperature before compressing to a final true strain of 0.7, and then quenched upon completion of their deformation. The deformed specimens were sectioned parallel to the compression axis and prepared for optical and electronic observation using standard metallographic techniques. An optical microscope (LEICA-DM4000M) and a field-emission gun SEM, equipped with an electron backscatter diffraction (EBSD) detector and Channel 5 software, were used to characterize the deformed microstructure. The parameters used were as follows: accelerating voltage 25 kV, step size 0.05, 0.1 μm, and a spot size of 6.0. The etchant used was a solution of 100 ml HCl, 100 ml ethanol, and 5 g CuCl2 [33].
257
accuracy in actual stress, as can be seen in Fig. 2(b). Fig. 3 demonstrated the true stress-true strain curves of the alloy deformed at different temperatures and strain rates. Typically, the trend of the curves exhibited a sharp increase with increasing strain before hitting a peak stress within a small strain range and then decreased gradually to a steady state at a strain greater than about 0.3. Generally, material deformation at elevated temperatures was a competing process of dynamic softening and work hardening. At the beginning of the deformation, the dislocation density increased and work hardening dominated the hot working process leading to the rapid increase of flow stress. As the deformation proceeded, the flow stress reached a peak meaning that work hardening, as the dominant process, came to an end. Then the softening due to dynamic recovery and recrystallization started to dominate the deformation process, leading to a gradual decrease in flow stress [34]. The flow stress, which remained stable with increasing strain, could be defined as a steady stress, indicating that work hardening and softening reached equilibrium. As Fig. 3 showed, for a particular strain rate, the flow stress decreased with increasing deformation temperature. At a given deformation temperature, the flow stress increased with increasing strain rate owing to an insufficiency of time for dislocation annihilation at higher strain rates [3,27]. At a strain rate of 0.01 s−1 and temperatures of 1000 °C/1025 °C, the flow stress decreased significantly after the strain reached 0.2 compared to higher strain rates (1.0/0.1 s−1), as shown in Fig. 3(a), (b) and (c). The mechanism could be attributed to the occurrence of significant dynamic recrystallization, which would be validated by further microstructural observation in the following sections. In addition, the flow stress of the specimen deformed at 1100 °C and 0.01 s−1 was approximately 60 MPa, which was lower than other asHIPed P/M nickel-based superalloys [17–19,35,36] under the similar deformation parameters, indicating that this new alloy would exhibit superior workability. Further work was necessary to verify the hypothesis. Table 1 listed the steady flow stress for samples deformed at 1100 °C and 0.01 s− 1 for other as-HIPed P/M nickel-based superalloys. 3.1. Kinetic analysis of the deformation During deformation, the peak stress could be related to strain rate and deformation temperature through the Arrhenius kinetic rate equation [37]: n expð−Q =RTÞ ε ¼ A sinh ασ p
ð1Þ
3. Results and discussion 3.0.1. Flow behavior under hot deformation The result of temperature calibration showed the superior uniformity of temperature along the whole specimen, which would improve the
where ε was the strain rate, A and α were material constants, σp was the peak stress, n was the stress exponent, Q was the activation energy for hot working, R was the gas constant, and T was the absolute
Fig. 1. Initial microstructure obtained for the as-HIPed nickel-based superalloy: (a) at low magnification; (b) at high magnification. Fig. 1 shows obviously the existence of prior particle boundary, one of three metallurgy defects in powder metallurgy alloy.
258
G. He et al. / Materials and Design 87 (2015) 256–265
Fig. 2. (a) A specimen with three thermocouples welded thereon; (b) the results of the temperature adjustment at 1100 °C. Fig. 2b shows the results of temperature adjustment at the deformation temperature of 1100 °C, which indicates the superior uniformity of temperature along the whole specimen.
deformation temperature. The activation energy Q could then be defined as: " Q ¼ R
∂ ln ε ∂ ln sinh ασ p
# ∂ ln sinh ασ p : ∂ð1=TÞ ε
this study, the stress exponent n was strain rate dependent and could be calculated from the lnσ-ln ε plot using linear regression. The plot of peak stress σp versus ln ε was shown in Fig. 4(b). The optimum value of constant α could be determined by linear regression to fit the relationships of ln ε-lnσp and σp-ln ε. These two plots together with the plots of ln[sinh(ασp)]-(1000/T) and ln[sinh(ασp)]-ln ε (see Fig. 4(c) and (d)) were adopted here to obtain the deformation activation energy. The deformation activation energy of the studied alloy and other as-HIPed P/M nickel-based superalloys [36,38–40] was listed for comparison in
ð2Þ
T
In this investigation, the variation of the natural logarithm of the peak stress with strain rate at various deformation temperatures was presented in Fig. 4(a). Over the entire range of strain rates covered in
Fig. 3. Typical flow stress–strain curves of nickel-based superalloy PM-0001 deformed at (a) 1.0 s−1; (b) 0.1 s−1; (c) 0.01 s−1; and (d) 0.001 s−1.
G. He et al. / Materials and Design 87 (2015) 256–265
259
Table 1 Steady stress of other as-HIPed P/M Ni-base superalloys deformed at 1100 °C/0.01 s−1. Author, year
Alloy/initial microstructure (as-HIPed)
HIP condition
Steady stress at 1100 °C, 0.01 s−1/MPa
Ref.
Somani M. 1992 Somani M. 1998 Ning Y., 2010 Xie X.H., 2002 Xu Wei, 2006 He G.A., 2015
NIMONIC-AP-1 IN100 FGH4096 FGH4096 FGH4096 PM-0001
1180 °C/108 MPa/3 h 1200 °C/115 MPa/3 h 1030 °C/120 MPa/1 h + 1170 °C/140 MPa/2 h / / 1100 °C/140 MPa/4 h
123.2 122 80–120 85 (roughly) 80 (roughly) 62
[17] [18] [19] [35] [36] This work
Table 2. The average activation energy for our alloy was somewhat higher than that of similar reported superalloys. The deformation activation energy of alloy was influenced by many factors, such as material chemistry, preparation method, as well as processing history and so on. Leverant and Kear [41] reported that the rate-controlling deformation process might involve self-diffusion in the Ni–Co based alloy, which was affected by interchange of Co and Ni atoms. The diffusion activation energy increased with increasing Co concentration, mechanism of which could be ascribed to a decrease of atomic order degree due to a substitution of Ni atoms by Co atoms [42]. The content of Co in our alloy was higher than that of other similar nickel-based superalloys [24], which served the dominated factor leading to a higher value of Q in our alloy. Moreover, additional energy was necessary for eliminating the inherent defects (i.e., PPB, as shown in Fig. 1) in the process of HIP, leading to a lift of deformation activation energy. The temperature compensated strain rate parameter, or the Zener– Hollomon parameter, could be defined as follows [43]: Q Z ¼ ε exp RT
ð3Þ
or, n Z ¼ A sinh ασ p :
ð4Þ
Fig. 5 showed the relationship between Zener–Hollomon parameter and [sinh(ασp)]n. This plot indicated that the superior linear relationship between ln Z and ln [sinh(ασp)], which was adopted to determined the constant A and exponent n. The value of n could be obtained by calculating the slope of the ln Z-ln [sinh(ασp)] plot and the constant A could be determined as 1.211 × 10 33 based on Eq. (4). Finally, substituted the values of α, A, n, and Q into Eq. (1), the constitutive equation of hot deformation for alloy could be expressed as:
ε ¼ 1:211 1033 ½ sinhð0:0062σ pÞ
3:011
903580 : exp − RT
ð5Þ
Fig. 4. Relationships between (a) ln σp and ln ε; (b) σp and ln ε; (c) ln [sinh(ασp)] and (1000/T); (d) ln [sinh(ασp)] and ln ε. These four plots are adopted here to obtain the constant α and deformation activation energy.
260
G. He et al. / Materials and Design 87 (2015) 256–265
transformation, and precipitation under dynamic conditions, and could be evaluated by integrating the following equation [18]:
Table 2 Activation energy for some other as-HIPed P/M Ni-Base superalloys. Alloy/initial microstructure
Temperature range (°C)
PM-0001/HIPed FGH96/HIPed FGH96/HIPed FGH96/HIPed FGH4096/HIPed
1000–1100 1000–1100 1000–1100 1000–1150 1080–1140
Strain rate range (s−1) 0.01–1.0 0.001–0.1 0.001–0.1 0.001–1.0 0.02–1.0
Total true strain 0.7 0.5 0.5 1.0 0.69
Activation energy (kJ/mol)
Ref.
903.58 754 740.4 693.21 595
This work [38] [39] [36] [40]
Z J¼
σ 0
mσε : 1þm
εdσ ¼
ð9Þ
For an ideal linear dissipater, m = 1 and J ¼ Jmax ¼ σ ε =2. For a non-linear dissipater, the efficiency of power dissipation could be expressed in terms of a dimensionless parameter [18]: η ¼ 2m=ðm þ 1Þ:
3.2. Processing maps Generally, the processing maps consisted of a power dissipation map and an instability map, and could be divided into safe domains and unsafe domains. In this model, the total power P might be divided into two complementary functions: G content and J co-content [44]: Z P¼Gþ J¼
ε 0
Z σdε þ
σ 0
εdσ:
ð6Þ
The strength of materials increased with increasing plastic strain rate [34]. The rate sensitive flow behavior was given [45] by:
ð10Þ
The efficiency of power dissipation η represented the proportion of the power dissipated by microstructure evolution over total power. Increase of η implied an increase in the power dissipated by microstructure evolution. The iso-efficiency contours of the power dissipation map, in which η varied with deformation temperature and strain rate, form the power dissipation map. This map exhibited different domains correlated with specific microstructures [46]. The continuum instability criterion based on the extremum principle of irreversible thermodynamic had been developed to identify the domains of flow instabilities, and given by another dimensionless parameter [18]: m ∂ ln m þ 1 ξ ε ¼ þ m b 0: ∂ ln ε
ð11Þ
m
σ ¼kε
ð7Þ
where σ was the flow stress, ε was the strain rate, m was the strain rate sensitivity, and k was a constant. The value of m was determined by a variant of Eq. (8) given at a constant strain and deformation temperature [17]:
∂ðlnσ Þ
: m¼
∂ ln ε
ε;T
ð8Þ
The G term mentioned in Eq. (6) referred to the power dissipated by plastic work, most of which would be converted into visco-plastic heat and the remaining power would be stored as lattice defects. However, the J term was related to metallurgical mechanisms such as dynamic recovery and recrystallization, deformation-induced phase
Fig. 5. Relationship between ln Z and ln [sinh(ασp)]. This plot indicates that the superior linear relationship between ln Z and ln [sinh(ασp)], which is adopted to determined the constant A.
The variation of instability parameter ξðεÞ might be evaluated as a function of deformation temperature and strain rate to obtain an instability map. The metallurgical instability during plastic flow occurred in domains where ξðεÞ was negative. Fig. 6 showed the processing maps of the alloy generated in the temperature range from 1000 to 1100 °C and strain rate range from 0.001 to 1.0 s−1 at different strains (from 0.4 to 0.7). As can be seen in Fig. 6, the stability domains were more likely to exhibit higher η, indicating more power dissipated by microstructure evolution. The instability regions were mainly located at the deformation conditions of high strain rates (N0.1 s−1) and/or low deformation temperature. Both of the two domains would be discussed in the following sections.
3.2.1. Instability domain In general, the maps at high strain rates were more inclined to exhibit instability domains regardless of strain and temperature. Fig. 6 showed that the material tended to exhibit a larger area instability domain with accumulation of strain. For the case of a strain of 0.7, the unsafe domain was located within the deformation temperature range from 1000 to 1100 °C and strain rate range from 0.1 s− 1 to 1.0 s−1. The mechanism of instability had been reported as being probably associated with cracking, located plastic flow, or adiabatic shear bands. The macroscopic features of specimens deformed at different temperatures and strain rates were presented in Fig. 7. As can be seen, there were no visible cracks in any of the specimen surfaces. The surface of all of the specimens was seen to be smooth and did not exhibit severe barrelling. However, the appearance of the specimens deformed at higher strain rates and lower temperatures tended to be more irregular when compared to other deformation conditions. Almost all specimens deformed at high strain rates (1.0 s−1) had an irregular shape regardless of their deformation temperature. The shape tended to be increasingly regular with decreasing strain rate, which matched the evidence of the processing maps. If the plasticity of the material was worse, shear deformation along an axis at about 45° to the direction of compression could be seen [34]. This shear deformation would result in the formation of shear bands and was consistent with the lower η in processing map. The reason might be that there was more adiabatic shear deformation heat dissipating in the shear deformation zone. Consequently, the processing
G. He et al. / Materials and Design 87 (2015) 256–265
261
Fig. 6. Processing maps at different strains: (a) 0.4; (b), 0.5; (c) 0.6; (d) 0.7. The shadows represent the instability domains.
map of a material was somehow related to its plasticity, and even workability: the higher the value of η, the better plasticity and workability of the material. Fig. 8 showed the microstructures of the specimen deformed at 1025 °C/1.0 s− 1 and at a strain of 0.7. The specimen exhibited local shear bands, which were characteristics of shear deformation, oriented at about 45° with respect to the deformation axis. A similar situation could be seen in other specimens deformed at this domain. The formation of shear bands during deformation was related to the generation of adiabatic heat associated with the lower thermal conductivity in comparison to other materials [47,48]. The mechanism might be attributed to the insufficient time available for releasing the deformation heat when the specimen was deformed at higher strain rates, leading to a temperature increase in regions with severe
Fig. 7. Macroscopic features of specimens deformed under different conditions, at a strain of 0.7. Specimens exhibit irregular appearance deformed at high strain rate and the shape become more and more regular with decreasing of strain rate.
deformation. The accumulation of deformation heat would result in the formation of bands of flow localization, which stimulated the occurrence of flow instability.
3.2.2. Stability domain For a strain of 0.7, one domain was located at 1080 °C/0.0015 s−1 with a peak efficiency of 55%, and the other occurred at about 1095 °C/0.014 s− 1 with a peak efficiency of 55%. The higher η meant a better workability at this domain, wherein the optimum hot working parameters for this alloy may be located [49].
Fig. 8. Microstructure deformed within the instability domain of 1025 °C s−1, at a strain of 0.7. Inside the dash line shows the characteristics of shear deformation, and the shear band can be seen clearly.
262
G. He et al. / Materials and Design 87 (2015) 256–265
Fig. 9. OIM maps of specimens obtained by EBSD under different deformation conditions: (a) 1025 °C/1.0 s−1; (b) 1050 °C/1.0 s−1; (c) 1075 °C/1.0 s−1; (d) 1025 °C/0.1 s−1; (e) 1050 °C/ 0.1 s−1; (f) 1075 °C/0.1 s−1; (g) 1025 °C/0.01 s−1; (h) 1050 °C/0.01 s−1; and (i) 1075 °C/0.01 s−1, at a strain of 0.7. Fig. 9 showed the effects of strain rates and temperature on microstructure evolution. The degree of dynamic recrystallization increased with increasing of temperature and decreasing of strain rate.
The orientation imaging microscopy (OIM) of the deformed specimens obtained by EBSD was shown in Fig. 9, among which Figs (d) to (i) were deformed at stability domains. Fig. 9 demonstrated the effects of deformation parameters on the degree of dynamic recrystallization and microstructure evolution. It might be noted that during deformation, the boundaries would become serrated and bulged before developing into the familiar ‘necklace’ structure: these were indicative of dynamic recrystallization in nickel-based superalloys [46]. As shown in Fig. 9, the deformed grains occurred to recrystallize at 1050 °C for strain rate of 1.0 s−1 and the degree of dynamic recrystallization tended to increase with increasing deformation temperature and decreasing strain rate. The physics behind this could be ascribed to more energy and the sufficient time available for nucleation as well as growth of dynamic recrystallization. Fig. 10 demonstrated the detailed features of grain boundary caused by dynamic recrystallization and microstructure evolution. The EBSD observation at 1050 °C/1.0 s−1 (Fig. 10(a)) demonstrated that the grain boundaries bulged and became to be serrated, implying the occurrence of dynamic recrystallization. At the grain boundary, the necklace-like microstructure was ubiquitously observed, showing that the grain boundary acted as the nucleation sites of dynamic recrystallization. In accompanying with decrease of strain rate, the longer compression time could facilitate the dynamic recrystallization, as shown in Fig. 10(b), till to be completed. As the temperature
increased to 1075 °C (see Fig. 9(i)), almost all of the grains had undergone dynamic recrystallization to yield a fine-grained microstructure at a strain rate of 0.01 s−1. With strain rate further decreased, some of the recrystallized grains occurred to grow up, as shown in Fig. 10(c). Moreover, the higher temperature could accelerate the grain growth and reach a polygonal morphology (Fig. 10(d)). Fig. 11 revealed the deformation microstructure of dynamic recrystallization under the temperature of 1100 °C, indicating the effects of strain rate on grain growth. The EBSD micrographs showed that the deformed grains growth occurred and developed into equiaxed morphology. While the strain rate was decreased to be 0.001 s− 1, the grain growth could be obviously observed, which was not favorable for microstructure design requirement of nickel-based superalloy. Therefore, the parameter of 1100 °C/0.001 s− 1 was advisable to be avoided for the actual forging processing. The significant decrease in stress after reaching peak stress (shown in Fig. 3) might be attributed to the formation of new, fine, recrystallized grain boundaries (see Fig. 9(g)). Considerable stress concentrations formed at grain boundary triple junctions during deformation: these would be released by grain boundary sliding and would accelerate if there were a considerable number of grain boundaries present [50]. Consequently, both a lower strain rate and a higher deformation temperature made a remarkable contribution to accelerating the grain
G. He et al. / Materials and Design 87 (2015) 256–265
263
Fig. 10. High-magnification microstructure of specimen under different deformation conditions: (a) 1050 °C/1.0 s− 1 ; (b) 1025 °C/0.1 s − 1 ; (c) 1075 °C/0.001 s− 1 ; and (d) 1100 °C/0.001 s−1, at a strain of 0.7. Fig. 10 showed the High-magnification microstructural features and microstructure evolution.
Fig. 11. OIM maps of the alloy deformed at 1100 °C: (a) 1.0 s−1; (b) 0.1 s−1; (c) 0.01 s−1; and (d) 0.001 s−1, at a strain of 0.7. Fig. 11 showed the effects of strain rate on degree of grain growth for the alloy at 1100 °C.
264
G. He et al. / Materials and Design 87 (2015) 256–265
boundary sliding process, which was considered to be the main reason for the softening that partially offset the deformation-induced stress. From the aforementioned discussion, it might be concluded that both deformation temperature and strain rate exerted a significant effect on dynamic recrystallization and microstructure evolution. The material exhibited more dynamic recrystallization grains under low strain rate and high temperature deformation conditions. The optimum deformation conditions for this alloy might be at 1075 to 1095 °C and 0.01 to 0.04 s− 1, at which the material exhibits favorably complete recrystallizated microstructure and potential superior workability. 4. Conclusion 1. Obvious dynamic recrysallization could be readily detected at the strain rate of 0.01 s−1, resulting in the significant flow stress decrease compared to 0.1 and 1.0 s−1. 2. The activation energy was determined to be 903.58 kJ/mol and a constitutive equation for the new alloy had been developed as follows: ε ¼ 1:211 1033 ½ sinhð0:0062σ pÞ
3:011
903580 exp − RT
3. The flow stress instability occurred at strain rate higher than 0.1 s−1 and temperature range from 1000 °C to 1100 °C, the adiabatic shear bands acted on the leading failure mode. 4. The peak efficiency of 55% upon a strain of 0.7 could be found at 1080 °C/0.0015 s− 1 and 1095 °C/0.014 s− 1. The grain growth was remarkably triggered under 1100 °C. 5. For our new alloy, the forging conditions could be optimized to the temperature range of 1075 °C–1095 °C and strain rate range of 0.014–0.04 s−1. Acknowledgment This work was supported by The National High-Tech Research and Development Programme of China, Grant no. 2012AA03A514, the Doctoral-Independent Exploration and Innovation of Central South University, Grant no. 2015zzts031 and the outstanding graduate project of Advanced Non-ferrous Metal Structural Materials and Manufacturing Collaborative Innovation Center. The author, Liu Feng, thanks The Natural Science Foundation of China for their financial support under Grant nos. 51401242, 61271356, and 51301209, and The Postdoctoral Science Foundation of Central South University for their grant no. 74341016096. References [1] M.O. Alniak, F. Bedir, Modelling of deformation and microstructural changes in P/M Rene 95 under isothermal forging conditions, Mater. Sci. Eng. A 429 (2006) 295–303. [2] M. Donachie, S. Donachie, Superalloys—A Technical Guide, ASM International, Materials Park, OH, 2002. [3] D. Li, Q. Guo, S. Guo, H. Peng, Z. Wu, The microstructure evolution and nucleation mechanisms of dynamic recrystallization in hot-deformed Inconel 625 superalloy, Mater. Des. 32 (2011) 696–705. [4] M. Zhang, F. Li, Z. Yuan, J. Li, S. Wang, Effect of heat treatment on the microindentation behavior of powder metallurgy nickel based superalloy FGH96, Mater. Des. 49 (2013) 705–715. [5] X.-M. Chen, Y.C. Lin, D.-X. Wen, J.-L. Zhang, M. He, Dynamic recrystallization behavior of a typical nickel-based superalloy during hot deformation, Mater. Des. 57 (2014) 568–577. [6] H. Zhang, K. Zhang, H. Zhou, Z. Lu, C. Zhao, X. Yang, Effect of strain rate on microstructure evolution of a nickel-based superalloy during hot deformation, Mater. Des. 80 (2015) 51–62. [7] X.-M. Chen, Y.C. Lin, M.-S. Chen, H.-B. Li, D.-X. Wen, J.-L. Zhang, et al., Microstructural evolution of a nickel-based superalloy during hot deformation, Mater. Des. 77 (2015) 41–49. [8] N.K. Park, I.S. Kim, Y.S. Na, J.T. Yeom, Hot forging of a nickel-base superalloy, J. Mater. Process. Technol. 111 (2001) 98–102. [9] H. Sun, Y. Sun, R. Zhang, M. Wang, R. Tang, Z. Zhou, Study on hot workability and optimization of process parameters of a modified 310 austenitic stainless steel using processing maps, Mater. Des. 67 (2015) 165–172.
[10] Z. Yang, F. Zhang, C. Zheng, M. Zhang, B. Lv, L. Qu, Study on hot deformation behaviour and processing maps of low carbon bainitic steel, Mater. Des. 66 (Part A) (2015) 258–266. [11] A. Łukaszek-Sołek, J. Krawczyk, The analysis of the hot deformation behaviour of the Ti–3Al–8 V–6Cr–4Zr–4Mo alloy, using processing maps, a map of microstructure and of hardness, Mater. Des. 65 (2015) 165–173. [12] H. Liao, Y. Wu, K. Zhou, J. Yang, Hot deformation behavior and processing map of Al– Si–Mg alloys containing different amount of silicon based on Gleebe-3500 hot compression simulation, Mater. Des. 65 (2015) 1091–1099. [13] M. Karami, R. Mahmudi, Hot shear deformation constitutive analysis and processing map of extruded Mg–12Li–1Zn bcc alloy, Mater. Des. 53 (2014) 534–539. [14] G. Liu, Y. Han, Z. Shi, J. Sun, D. Zou, G. Qiao, Hot deformation and optimization of process parameters of an as-cast 6Mo superaustenitic stainless steel: a study with processing map, Mater. Des. 53 (2014) 662–672. [15] K.-k. Deng, J.-c. Li, F.-j. Xu, K.-b. Nie, W. Liang, Hot deformation behavior and processing maps of fine-grained SiCp/AZ91 composite, Mater. Des. 67 (2015) 72–81. [16] F.-L. Sui, L.-X. Xu, L.-Q. Chen, X.-H. Liu, Processing map for hot working of Inconel 718 alloy, J. Mater. Process. Technol. 211 (2011) 433–440. [17] M. Somani, E.B. Rao, N. Birla, M. Bhatia, V. Singh, Y. Prasad, Processing map for controlling microstructure in hot working of hot isostatically pressed powder metallurgy NIMONIC AP-1 superalloy, Metall. Trans. A 23 (1992) 2849–2857. [18] M. Somani, K. Muraleedharan, Y. Prasad, V. Singh, Mechanical processing and microstructural control in hot working of hot isostatically pressed P/M IN-100 superalloy, Mater. Sci. Eng. A 245 (1998) 88–99. [19] Y. Ning, Z. Yao, H. Guo, M.W. Fu, H. li, X. Xie, Investigation on hot deformation behavior of P/M Ni-base superalloy FGH96 by using processing maps, Mater. Sci. Eng. A 527 (2010) 6794–6799. [20] M. Jahangiri, H. Arabi, S. Boutorabi, High-temperature compression behavior of cast and homogenized IN939 superalloy, Metall. Mater. Trans. A 44 (2013) 1827–1841. [21] K. Wu, G. Liu, B. Hu, F. Li, Y. Zhang, Y. Tao, et al., Hot compressive deformation behavior of a new hot isostatically pressed Ni–Cr–Co based powder metallurgy superalloy, Mater. Des. 32 (2011) 1872–1879. [22] P. Zhang, C. Hu, C.-g. Ding, Q. Zhu, H.-y. Qin, Plastic deformation behavior and processing maps of a Ni-based superalloy, Mater. Des. 65 (2015) 575–584. [23] Y. Wang, Q. Pan, Y. Song, C. Li, Z. Li, Hot deformation and processing maps of X-750 nickel-based superalloy, Mater. Des. 51 (2013) 154–160. [24] Y. Gu, H. Harada, C. Cui, D. Ping, A. Sato, J. Fujioka, New Ni–Co-base disk superalloys with higher strength and creep resistance, Scr. Mater. 55 (2006) 815–818. [25] J. Heslop, Wrought Nickel–Chromium Heat-Resisting Alloys Containing Cobalt, Cobblestone 24 (1964) 128. [26] J. Tien, T. Howson, G. Chen, X. Xie, Cobalt availability and superalloys, J. Meteorol. 32 (1980) 12–20. [27] R.N. Jarrett, J.K. Tien, Effects of cobalt on structure, microchemistry and properties of a wrought nickel-base superalloy, Metall. Trans. A 13 (1982) 1021–1032. [28] C.Y. Cui, Y.F. Gu, Y. Yuan, T. Osada, H. Harada, Enhanced mechanical properties in a new Ni–Co base superalloy by controlling microstructures, Mater. Sci. Eng. A 528 (2011) 5465–5469. [29] Y. Yuan, Y.F. Gu, Z.H. Zhong, T. Yokokawa, H. Harada, Enhanced strength at intermediate temperatures in a Ni-base disk superalloy with high Co addition, Mater. Sci. Eng. A 556 (2012) 595–600. [30] Y. Yuan, Y.F. Gu, Z.H. Zhong, T. Osada, C.Y. Cui, T. Tetsui, et al., Creep mechanisms of a new Ni–Co-base disc superalloy at an intermediate temperature, J. Microsc. 248 (2012) 34–41. [31] Z. Zhong, Y. Gu, Y. Yuan, C. Cui, T. Yokokawa, H. Harada, Fatigue crack growth behavior of a newly developed Ni–Co-base superalloy TMW-2 at elevated temperatures, Mater. Sci. Eng. A 552 (2012) 464–471. [32] R. Miner, Effects of C and Hf concentration on phase relations and microstructure of a wrought powder-metallurgy superalloy, Metall. Trans. A 8 (1977) 259–263. [33] Voort G. Vander, Metallography of superalloys, Ind. Heat. 70 (2003) 40–43. [34] Q.L. Pan, B. Li, Y. Wang, Y.W. Zhang, Z.M. Yin, Characterization of hot deformation behavior of Ni-base superalloy Rene'41 using processing map, Mater. Sci. Eng. A 585 (2013) 371–378. [35] X. Xie, Z. Yao, Y. Ning, H. Guo, Y. Tao, Y. Zhang, Investigation on hot deformation behavior of P/M superalloy FGH4096, Rare Metal Mater. Eng. 41 (2012) 82–86. [36] W. Xu, L.-w. Zhang, S.-d. Gu, J.-l. Zhang, Hot compressive deformation behavior and microstructure evolution of HIPed FGH96 superalloy, Trans. Nonferrous Metals Soc. 22 (2012) 66–71. [37] C. Bruni, A. Forcellese, F. Gabrielli, Hot workability and models for flow stress of NIMONIC 115 Ni-base superalloy, J. Mater. Process. Technol. 125–126 (2002) 242–247. [38] Y. Liu, F. Li, H. Yu, Deformation behavior of hot isostatic pressing FGH96 superalloy, Trans. Tianjin Univ. 12 (2006) 281–285. [39] R. Zhang, F. Li, X. Wang, Determining processing maps of FGH96 superalloy, J. Norhtwestern Polytechnical Univ. 25 (2007) 652–656. [40] Y. Ning, Z. Yao, M.W. Fu, H. Guo, Dynamic recrystallization of the hot isostatically pressed P/M superalloy FGH4096 in hot working process, Mater. Sci. Eng. A 527 (2010) 6968–6974. [41] G. Leverant, B. Kear, The mechanism of creep in gamma prime precipitationhardened nickel-base alloys at intermediate temperatures, Metall. Mater. Trans. B 1 (1970) 491–498. [42] A. Gazda, J. Čermák, Diffusion in Ni3Al-Co intermetallic, Metal 5 (2002) 14–16. [43] N.-K. Park, J.-T. Yeom, Y.-S. Na, Characterization of deformation stability in hot forging of conventional Ti–6Al–4 V using processing maps, J. Mater. Process. Technol. 130 (2002) 540–545. [44] H.R. Ezatpour, S.A. Sajjadi, M. Haddad-Sabzevar, G.R. Ebrahimi, Hot deformation and processing maps of K310 cold work tool steel, Mater. Sci. Eng. A 550 (2012) 152–159.
G. He et al. / Materials and Design 87 (2015) 256–265 [45] D. Cai, L. Xiong, W. Liu, G. Sun, M. Yao, Characterization of hot deformation behavior of a Ni-base superalloy using processing map, Mater. Des. 30 (2009) 921–925. [46] Y. Prasad, S. Sasidhara, Hot Working Guide: A Compendium of Processing Maps, ASM International, 1997. [47] Y. Han, G. Liu, D. Zou, R. Liu, G. Qiao, Deformation behavior and microstructural evolution of as-cast 904 L austenitic stainless steel during hot compression, Mater. Sci. Eng. A 565 (2013) 342–350.
265
[48] W. Peng, W. Zeng, Q. Wang, H. Yu, Characterization of high-temperature deformation behavior of as-cast Ti60 titanium alloy using processing map, Mater. Sci. Eng. A 571 (2013) 116–122. [49] J.K. Chakravartty, G.K. Dey, S. Banerjee, Y.V.R.K. Prasad, Characterization of hot deformation behaviour of Zr–2.5Nb–0.5Cu using processing maps, J. Nucl. Mater. 218 (1995) 247–255. [50] S. Medeiros, W. Frazier, Y. Prasad, Hot deformation mechanisms in a powder metallurgy nickel-base superalloy IN 625, Metall. Mater. Trans. A 31 (2000) 2317–2325.