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Fig. 14 — SEM image of Gamma prime precipitation precipitated y’ of within the Inconel X-750 superalloy (a) deformed at temperature of 1000 °C and strain rate of 0.001 s~* to strain of 0.65 (b) soaked in 1000 °C for 3 min without any straining. potential softening DRX progress at these extended deforma- tion times (Fig. 16c and d). For higher strain rates of 1and0.1s"7, the size and number of DRX grains substantially decreased (Table 5) compared with the same strain rate at 1000 °C (Fig. 16a and b). The underlying reason for the less softening through DRX could be the emergence dynamically precipitation of some

Figure 14 — SEM image of Gamma prime precipitation precipitated y’ of within the Inconel X-750 superalloy (a) deformed at temperature of 1000 °C and strain rate of 0.001 s~* to strain of 0.65 (b) soaked in 1000 °C for 3 min without any straining. potential softening DRX progress at these extended deforma- tion times (Fig. 16c and d). For higher strain rates of 1and0.1s"7, the size and number of DRX grains substantially decreased (Table 5) compared with the same strain rate at 1000 °C (Fig. 16a and b). The underlying reason for the less softening through DRX could be the emergence dynamically precipitation of some

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temperature properties are shall meet the required specifica- tions to ensure about their performance at severe thermal situations. To attain good high-temperature properties and hot forging technology of superalloys, it is essential to meet the required specifications that ensure their performance in severe thermal situations [5]. Correlating flow behaviour with microstructure through hot deformation is meaningfully affected by temperature and strain rate variation of the pro- cess route [6]. Consequently, it is necessary to study the hot deformation behaviour of superalloys by fully understanding the damaging/restoration phenomena and their influence of on the flow behaviour and microstructure evolution of Inconel X-750. In this study, the Arrhenius [7], Zerilli-Armstrong [8] and Johnson—Cook [9] constitutive equations have been suc- cessfully utilized for prediction of hot flow curves through hot deformation. The Arrhenius-type equation, also known as stress compensated temperature-dependent model, can be stated in terms of Zener—Hollomon (Z) parameter. In this equation, superimposed recompense of working temperature, strain rate at any strain level is considered for flow stress description at elevated temperatures, accurately. The micro- structure is analyzed by different restoration mechanisms like dynamic recovery (DRV) and dynamic recrystallization (DRX) [10]. The ability of the other phases rather than matrix is related to the pinning effect and particle-stimulated nucle- ation depending upon the size and volume fraction of such phases [10]. Lately, hot workability of superalloys through hot deformation was assessed using dynamic material model (DMM) taking advantage of processing maps for identifying deformation mechanism in safe regions and damaging ones in unstable zones that use processing maps to identify the safe regions and the unstable zones of deformation mecha- nisms [11]. Based on primary works of Frost and Ashby on entropy production through deformation process and defects occurred during the forming at different strain, strain rate, and temperature, Prasad has developed a workability evalu- ation paradigm following the principle of DMM [12]. For the identification of desired hot working parameters, it has been broadly believed that a hot deformation processing map is an applied and beneficial workflow [13]. Processing map was first introduced by Prasad et al. [12,13] has been used for modelling hot workability of different superalloys like Co, Fe and Ni- based superalloys. In spite of its wide application of these processing maps, some researchers claimed that modifica- tions of its their development approach should be considered by performing comparisons between instability criteria introduced by different researchers in the literature [11,14]. To the best of author's knowledge, rare investigations has been done on the recrystallization and hot forming characteristics of Inconel X-750 alloy. Therefore, the key purpose of this research was to examine the DRX kinetics and microstructure evolution caused by hot deformation behaviour of Inconel X- 750 compressed at dissimilar temperatures and strain rates.
temperature properties are shall meet the required specifica- tions to ensure about their performance at severe thermal situations. To attain good high-temperature properties and hot forging technology of superalloys, it is essential to meet the required specifications that ensure their performance in severe thermal situations [5]. Correlating flow behaviour with microstructure through hot deformation is meaningfully affected by temperature and strain rate variation of the pro- cess route [6]. Consequently, it is necessary to study the hot deformation behaviour of superalloys by fully understanding the damaging/restoration phenomena and their influence of on the flow behaviour and microstructure evolution of Inconel X-750. In this study, the Arrhenius [7], Zerilli-Armstrong [8] and Johnson—Cook [9] constitutive equations have been suc- cessfully utilized for prediction of hot flow curves through hot deformation. The Arrhenius-type equation, also known as stress compensated temperature-dependent model, can be stated in terms of Zener—Hollomon (Z) parameter. In this equation, superimposed recompense of working temperature, strain rate at any strain level is considered for flow stress description at elevated temperatures, accurately. The micro- structure is analyzed by different restoration mechanisms like dynamic recovery (DRV) and dynamic recrystallization (DRX) [10]. The ability of the other phases rather than matrix is related to the pinning effect and particle-stimulated nucle- ation depending upon the size and volume fraction of such phases [10]. Lately, hot workability of superalloys through hot deformation was assessed using dynamic material model (DMM) taking advantage of processing maps for identifying deformation mechanism in safe regions and damaging ones in unstable zones that use processing maps to identify the safe regions and the unstable zones of deformation mecha- nisms [11]. Based on primary works of Frost and Ashby on entropy production through deformation process and defects occurred during the forming at different strain, strain rate, and temperature, Prasad has developed a workability evalu- ation paradigm following the principle of DMM [12]. For the identification of desired hot working parameters, it has been broadly believed that a hot deformation processing map is an applied and beneficial workflow [13]. Processing map was first introduced by Prasad et al. [12,13] has been used for modelling hot workability of different superalloys like Co, Fe and Ni- based superalloys. In spite of its wide application of these processing maps, some researchers claimed that modifica- tions of its their development approach should be considered by performing comparisons between instability criteria introduced by different researchers in the literature [11,14]. To the best of author's knowledge, rare investigations has been done on the recrystallization and hot forming characteristics of Inconel X-750 alloy. Therefore, the key purpose of this research was to examine the DRX kinetics and microstructure evolution caused by hot deformation behaviour of Inconel X- 750 compressed at dissimilar temperatures and strain rates.
Fig. 1 — The overall schematic image of the experimental procedure and the thermomechanical process diagram
Fig. 1 — The overall schematic image of the experimental procedure and the thermomechanical process diagram
Fig. 2 — Modified variables of isothermal compression test at strain of 0.65 for (a) and (b) the degree of temperature deviation deviated temperature of in the specimens at different strain rates and temperatures at strain of 0.65, and (c, d) the increased stress after adiabatic heat modification at strain of 0.65 for studied conditions.
Fig. 2 — Modified variables of isothermal compression test at strain of 0.65 for (a) and (b) the degree of temperature deviation deviated temperature of in the specimens at different strain rates and temperatures at strain of 0.65, and (c, d) the increased stress after adiabatic heat modification at strain of 0.65 for studied conditions.
Fig. 3 — Modified stress (lowered) of isothermal compression test at strain of 0.65 for (a) different strain rates at studiec femperatures and (b) different temperatures at studied strain rates. The negative values stand for stress reduction afte ‘riction effect correction, which stems from barreling and imperfect lubrication.
Fig. 3 — Modified stress (lowered) of isothermal compression test at strain of 0.65 for (a) different strain rates at studiec femperatures and (b) different temperatures at studied strain rates. The negative values stand for stress reduction afte ‘riction effect correction, which stems from barreling and imperfect lubrication.
Fig. 5 — Critical strain variation for the onset of DRX with respect to temperature (°C) at different strain rates (s—*).
Fig. 5 — Critical strain variation for the onset of DRX with respect to temperature (°C) at different strain rates (s—*).
Fig. 6 — Flow softening percentage with respect to temperature (°C) at different strain rates (s~*). where a, @,n’,n and A are material constants, a is flow stress at a given strain. The hyperbolic sine law, exponential law, and the power law constitutive equations are employed to calcu- late the material constants and flow stress according to the steps presented in Fig. 7. Following In accordance with this computational flowchart, the calculated constitutive equa- tions of the Inconel X-750 superalloy at peak strain can be reported as expressed in Eq. (4) [20]:
Fig. 6 — Flow softening percentage with respect to temperature (°C) at different strain rates (s~*). where a, @,n’,n and A are material constants, a is flow stress at a given strain. The hyperbolic sine law, exponential law, and the power law constitutive equations are employed to calcu- late the material constants and flow stress according to the steps presented in Fig. 7. Following In accordance with this computational flowchart, the calculated constitutive equa- tions of the Inconel X-750 superalloy at peak strain can be reported as expressed in Eq. (4) [20]:
Fig. 7 — The flow chart for calculation progress to obtain constant parameters of constitutive equations. The reliability of the reported constitutive equations was evaluated by the standard statistical parameters [27] (average absolute error-AARE) and also correlation coefficient (R) [28]. This statistical information was calculated for the result of all three constitutive equations at various temperatures and strain rates and presented in Fig. 8. Apparently, when ag is greater than 1.2, the flow stress is best predicted by the exponential law at various deformation conditions (Fig. 8b). Similarly, when ac The activations energy for all equations are computed from 461 to 478 kJ.mol*, which is considered a rather large value compared to the energy for the diffusion of Cr, Ti, Ni, and Al atoms in the Ni matrix. i.e., the activation energy for these elements in Ni was reported 278 [22], 256 [15], 275 [23] and 269 [15] kj.mol~’, respectively. Hence, the activation energy
Fig. 7 — The flow chart for calculation progress to obtain constant parameters of constitutive equations. The reliability of the reported constitutive equations was evaluated by the standard statistical parameters [27] (average absolute error-AARE) and also correlation coefficient (R) [28]. This statistical information was calculated for the result of all three constitutive equations at various temperatures and strain rates and presented in Fig. 8. Apparently, when ag is greater than 1.2, the flow stress is best predicted by the exponential law at various deformation conditions (Fig. 8b). Similarly, when ac The activations energy for all equations are computed from 461 to 478 kJ.mol*, which is considered a rather large value compared to the energy for the diffusion of Cr, Ti, Ni, and Al atoms in the Ni matrix. i.e., the activation energy for these elements in Ni was reported 278 [22], 256 [15], 275 [23] and 269 [15] kj.mol~’, respectively. Hence, the activation energy
Table 2 — Table of calculated constants for constitutive equations: hyperbolic sine, exponential law, and power law equation for peak strain.
Table 2 — Table of calculated constants for constitutive equations: hyperbolic sine, exponential law, and power law equation for peak strain.
Table 3 — The reported activation energy for other selected Ni-based superalloys.
Table 3 — The reported activation energy for other selected Ni-based superalloys.
Fig. 8 — The validation of predicted peak flow stress with experimental results for (a) hyperbolic sine equation, (b) exponential equation, and (c) power law equation.
Fig. 8 — The validation of predicted peak flow stress with experimental results for (a) hyperbolic sine equation, (b) exponential equation, and (c) power law equation.
Fig. 9 — The 2D processing map of Inconel X-750 superalloy at various strains. The orange highlights depict the powe instability zones.
Fig. 9 — The 2D processing map of Inconel X-750 superalloy at various strains. The orange highlights depict the powe instability zones.
Fig. 10 — (a) The 2D processing map of Inconel X-750 superalloy at strain of 0.65; (b) dependency of power efficiency value t temperature at different strain rates strain rate at different temperatures; and (c) dependency of power efficiency value t strain rate at different temperatures. temperature at different strain rates. The orange highlights depict the power instability zones.
Fig. 10 — (a) The 2D processing map of Inconel X-750 superalloy at strain of 0.65; (b) dependency of power efficiency value t temperature at different strain rates strain rate at different temperatures; and (c) dependency of power efficiency value t strain rate at different temperatures. temperature at different strain rates. The orange highlights depict the power instability zones.
Fig. 11 — (a) Optical microstructure and (b, c, and d) SEM micrograph of the Inconel X-750 as-received bar in different magnitudes: (a) 100 pm (b) 25 pm, (c) 1 pm, and (d) 400 nm. The marked areas by P, S, and C are primary y’, secondary y’ anc carbides, respectively.
Fig. 11 — (a) Optical microstructure and (b, c, and d) SEM micrograph of the Inconel X-750 as-received bar in different magnitudes: (a) 100 pm (b) 25 pm, (c) 1 pm, and (d) 400 nm. The marked areas by P, S, and C are primary y’, secondary y’ anc carbides, respectively.
Table 5 — The size and volume fraction of recrystallized grains at studied thermomechanical conditions.
Table 5 — The size and volume fraction of recrystallized grains at studied thermomechanical conditions.
Fig. 13 — The TG graph of solution-treated Inconel X-750 superalloy indicating the y’ dissolution temperature range. curve in Fig. 10c (which represent the change in n for 1000 and 1050 °C at various strain rates), shows an uprising trend with strain rate. While the overall power efficiency portends a higher DRX at strain rates of 0.001 and 0.01s~*, no higher DRXed grains are present in their related microstructures (Fig. 12g and h). Likewise, despite the lower estimated n for strain rates of 0.1 and 1 s~’, their final structure at strain of 0.65, depicts a great advance in DRX at all former grain boundaries (Fig. 12e and f). Thus, the precipitation of y’ not only prohibits the DRX and softening mechanisms, but it also misleads the power effi- ciency nature (compared to 950 °C) in this temperature zone. Elevating the temperature to 1050 °C has presented relatively Z- dependent progress of DRX with increasing strain rate (Fig. 16a—d). Within this temperature zone, although the excessive softening due to adiabatic heat is fading (Fig. 16a and b), two other phenomena control the final morphology at strain of 0.65. While at higher strain rates (1 and 0.1 s-%), the commencement of dynamic precipitation inhibits the DRX process (Fig. 16a and b), the precipitation of 7’ still lingers at lower strain rates (1 and 0.1 s~’), and it slightly weakens the
Fig. 13 — The TG graph of solution-treated Inconel X-750 superalloy indicating the y’ dissolution temperature range. curve in Fig. 10c (which represent the change in n for 1000 and 1050 °C at various strain rates), shows an uprising trend with strain rate. While the overall power efficiency portends a higher DRX at strain rates of 0.001 and 0.01s~*, no higher DRXed grains are present in their related microstructures (Fig. 12g and h). Likewise, despite the lower estimated n for strain rates of 0.1 and 1 s~’, their final structure at strain of 0.65, depicts a great advance in DRX at all former grain boundaries (Fig. 12e and f). Thus, the precipitation of y’ not only prohibits the DRX and softening mechanisms, but it also misleads the power effi- ciency nature (compared to 950 °C) in this temperature zone. Elevating the temperature to 1050 °C has presented relatively Z- dependent progress of DRX with increasing strain rate (Fig. 16a—d). Within this temperature zone, although the excessive softening due to adiabatic heat is fading (Fig. 16a and b), two other phenomena control the final morphology at strain of 0.65. While at higher strain rates (1 and 0.1 s-%), the commencement of dynamic precipitation inhibits the DRX process (Fig. 16a and b), the precipitation of 7’ still lingers at lower strain rates (1 and 0.1 s~’), and it slightly weakens the
Fig. 15 — The schematic phase diagram of a typical Ni-base superalloy at the vicinity of dissolution limit for y’ in a y matrix.
Fig. 15 — The schematic phase diagram of a typical Ni-base superalloy at the vicinity of dissolution limit for y’ in a y matrix.
Fig. 16 — Scanning Electron Microscopy (SEM) of hot deformed Inconel X-750 alloy at 1050 and 1100 °C to strain of 0.65 at various strain rates.
Fig. 16 — Scanning Electron Microscopy (SEM) of hot deformed Inconel X-750 alloy at 1050 and 1100 °C to strain of 0.65 at various strain rates.
Fig. 18 — The DTA graph of solution-treated Inconel X-750 indicating the MC carbides precipitation temperature range. In the course of During the thermo-mechanical process, the pinning effect of the y’ particles and to a lesser degree ,MC carbides can restrict the grain mobility and reduce the work- ability of the alloy. To address this issue, prior to hot defor- mation, the samples were annealed to the super-solvus temperature of 1150 °C and kept for 90 min. Owing to this solution-annealing process, the y’ particles were dissolved, and the workability of the alloy would enhance (com pere Fig. 21a to b). Although, all y’ particles were dissolved during solution annealing treatment, some carbides have remai ined intact. Some of these carbides could precipitate throug. hout the solution annealing treatment or within the subsequent water quenching. The metal component of these newly engendered MC carbides stems from the dissolution of y’' particles and the carbon component can either be provided by the dissolution of low-solvus (1150 °C) Mz3C, carbides or the existing carbon within the alloy matrix [31]. just
Fig. 18 — The DTA graph of solution-treated Inconel X-750 indicating the MC carbides precipitation temperature range. In the course of During the thermo-mechanical process, the pinning effect of the y’ particles and to a lesser degree ,MC carbides can restrict the grain mobility and reduce the work- ability of the alloy. To address this issue, prior to hot defor- mation, the samples were annealed to the super-solvus temperature of 1150 °C and kept for 90 min. Owing to this solution-annealing process, the y’ particles were dissolved, and the workability of the alloy would enhance (com pere Fig. 21a to b). Although, all y’ particles were dissolved during solution annealing treatment, some carbides have remai ined intact. Some of these carbides could precipitate throug. hout the solution annealing treatment or within the subsequent water quenching. The metal component of these newly engendered MC carbides stems from the dissolution of y’' particles and the carbon component can either be provided by the dissolution of low-solvus (1150 °C) Mz3C, carbides or the existing carbon within the alloy matrix [31]. just
Fig. 17 — The SEM image of dynamically precipitated MC carbides when the alloy deformed at strain rate of 1 s~* for temperature of (a) 1050 °C, and (b) 1100 °C. The red markings indicate the y matrix (y) and dynamically precipitated Nbc carbide (c) during deformation.
Fig. 17 — The SEM image of dynamically precipitated MC carbides when the alloy deformed at strain rate of 1 s~* for temperature of (a) 1050 °C, and (b) 1100 °C. The red markings indicate the y matrix (y) and dynamically precipitated Nbc carbide (c) during deformation.
Fig. 19 — MC carbides precipitation of Inconel X-750 superalloy at temperature of 1100 °C (a) after straining at strain rate of 1s“ strained to strain of 0.65 (b) just before straining.
Fig. 19 — MC carbides precipitation of Inconel X-750 superalloy at temperature of 1100 °C (a) after straining at strain rate of 1s“ strained to strain of 0.65 (b) just before straining.
Fig. 20 — Scanning Electron Microscopy (SEM) of hot deformed Inconel X-750 alloy at 1150 °C to strain of 0.65 at various strain rates.
Fig. 20 — Scanning Electron Microscopy (SEM) of hot deformed Inconel X-750 alloy at 1150 °C to strain of 0.65 at various strain rates.
Fig. 21 — SEM image of (a) gamma prime precipitants of Inconel X-750 before solution annealing and (b) after solution heat treatment at temperature of 1150 °C for 90 min.
Fig. 21 — SEM image of (a) gamma prime precipitants of Inconel X-750 before solution annealing and (b) after solution heat treatment at temperature of 1150 °C for 90 min.
Fig. 22 — The peak strain variation (a) at different strain rate temperatures and (b) at different temperatures strain rates. Table 6 — General flow behaviour and predominant softening mechanisms of Inconel X-750 at studied temperatures and strain rates. (PDRX: partials DRX — AH: Adiabatic heating — GG: Grain growth).
Fig. 22 — The peak strain variation (a) at different strain rate temperatures and (b) at different temperatures strain rates. Table 6 — General flow behaviour and predominant softening mechanisms of Inconel X-750 at studied temperatures and strain rates. (PDRX: partials DRX — AH: Adiabatic heating — GG: Grain growth).
Related topics:
Ni Based SuperalloysFlow Instability and Transition
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