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Influence of thermomechanical processing on microstructural evolution in near-α alloy IMI834
Materials Science and Engineering A 416 (2006) 300–311
Influence of thermomechanical processing on microstructural evolution in near-␣ alloy IMI834 P. Wanjara a,∗ , M. Jahazi a , H. Monajati b , S. Yue b a
National Research Council, Institute for Aerospace Research, Aerospace Manufacturing Technology Center, 5145 Decelles Avenue, Campus Universit´e de Montr´eal, P.O. Box 40, Montr´eal, Que., Canada H3S 2S4 b McGill University, Mining, Metals and Materials Engineering, 3610 University Street, Montr´ eal, Que., Canada H3A 2B2 Accepted 27 October 2005
Abstract The effect of processing parameters on the evolution of the microstructural characteristics of IMI834 were examined over the temperature range 950–1125 ◦ C at constant strain rates of 0.001–1 s−1 and true strains up to 1.2. During isothermal deformation, differences in the flow softening behavior were observed for processing in the single-phase beta region, as compared to the two-phase alpha–beta region, which has been related to differences in the structural changes occurring in IMI834. In particular, deformation processing of the alpha–beta structure was observed to work the lamellar alpha structure of the transformed beta grains without recrystallization of either phase, whilst microstructural development in a predominately beta structure resulted in the formation of dynamically recrystallized grains that were necklaced around the deformed and elongated beta phase. For the later case, the variation in the dynamically recrystallized grain size was determined to follow a Zener–Hollomon relationship between a strain rate range 0.001 and 1 s−1 . © 2005 Elsevier B.V. All rights reserved. Keywords: Titanium alloys; Dynamic recrystallization; Isothermal deformation; Shear banding; Flow softening
1. Introduction Over the last 50 years, the evolution of titanium alloys has been driven by the ever-increasing need to improve the efficiency of gas turbines by elevating operating temperatures and improving component durability. Titanium alloys have met the stringent requirements of the aerospace industry through alloy chemistry development, and more recently through microstructural design, both of which are necessary to remain responsive to the increasing needs for engine performance [1–4]. For manufacturing of specific component shapes from titanium alloys, it is recognized that the thermomechanical processing conditions can be optimized to control the microstructuremechanical property characteristics. However, considering the restrictive workability limit (or processing window) of titanium alloys in general, an increased understanding of the relationship between processing and microstructure is particularly critical for sustaining further improvements in performance and reliability. For instance, near-alpha titanium alloy, IMI834 has been
∗
Corresponding author. Tel.: +1 514 283 9380; fax: +1 514 283 9445. E-mail address: [email protected] (P. Wanjara).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.10.042
designed for compressor disks with improved tensile strength, fatigue resistance and creep performance at temperatures up to 600 ◦ C by means of a tailored bimodal microstructure [5–8]. Processing of IMI834 usually involves a series of primary and secondary deformation operations that incorporate a combination of non-homogeneous thermomechanical conditions involving complex metallurgical phenomena including phase transformations, both dynamic and static restoration mechanisms as well as grain growth [9–11]. The final component microstructure is dependent on the various structural features inherited from the primary stages (cogging or ingot breakdown) and the last forming and heat treatment operations [12]. A number of studies [13–16] have indicated that a bimodal structure consisting of primary alpha particles and fine transformed beta grains improves the mechanical properties of IMI834. In general, during processing of IMI834, the evolution in microstructure is dependent on dynamic recovery and recrystallization occurring under an applied load as well as static recovery, static recrystallization and grain growth that proceed upon load removal. Moreover, the occurrence of these dynamic and static phenomena during thermomechanical processing can influence the kinetics of restoration and grain growth in both the beta and alpha phases of IMI834 and produce considerable
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inhomogeneity in the final microstructure [17]. Presently, it has been reported that during high temperature deformation processing of IMI834 the flow stress exhibits a strong dependence on the temperature with the probability of considerable flow softening, both of which are related to the dependence of the alpha and beta phases on variations in the deformation parameters such as temperature, strain and strain rate [6,10,18]. It has also been suggested that the flow softening behavior is related to the restoration of the material during processing [19]. However, systematic evaluation of the role of various processing parameters (strain, strain rate, temperature) on the deformation characteristics (flow stress–strain behavior) and microstructural development in IMI834 remains a matter of an extended investigation. In this paper, the research work reported formed part of a larger program on developing the technology for processing titanium materials through optimizing the forging conditions to allow control of microstructure and properties in the finished product. Part of this program involved characterization of the flow behavior of IMI834 under hot forging conditions. The results of this research work are described elsewhere [20]. In this paper, the results of subsequent examination of the deformationinduced microstructures are reported to establish the influence of processing (ingot breakdown and closed die forging) conditions on structural evolution. 2. Experimental procedure 2.1. Starting materials Hot rolled IMI834 material was received in bar form with an outer diameter of 18 cm and chemical composition given in Table 1. The beta transformation temperature of this material was approximately 1045 ◦ C. The microstructure of the as-received material consisted of a fine, bimodal alpha–beta structure of approximately 25.6% equiaxed primary alpha (with a standard deviation of 3%) within a fine transformed beta matrix, as shown in Fig. 1. Cylindrical specimens of 7.5 mm in diameter and 11.4 mm in height were electro-discharged machined from the bar for the high temperature compression experiments. 2.2. Hot working equipment The hot working studies were conducted on a computerized 100 kN Materials Testing System (MTS 810) that was adapted for elevated temperature compression testing to 1200 ◦ C Table 1 Chemical composition (wt.%) of IMI834 material Al Sn Zr Nb Mo Si C Ti
5.5 4.0 3.5 0.7 0.5 0.35 0.06 Balance
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Fig. 1. Bimodal microstructure of as-received IMI834 material.
using a Research Incorporated radiant furnace interfaced with a Micristar digital programmer and controller. During high temperature forming, the compression specimen and the anvils were enclosed within a quartz tube, in which argon gas was passed to prevent oxidation of the work piece and tooling. To minimize friction during hot working, the ends of the specimens were coated with delta glaze and a thin layer of mica1 was placed between the face of the specimen and the anvils in order to maintain uniform deformation and avoid sticking problems during quenching. 2.3. Deformation schedule Hot working was performed using two distinct thermal cycles designed for: (1) beta forging corresponding to simulation of ingot breakdown of a cast microstructure to a wrought product and (2) alpha–beta forging corresponding to the simulation of closed-die forging to a final shape. In the first cycle, beta forging was conducted at constant temperature between 1050 and 1125 ◦ C after holding for 15 min at temperature to allow for thermal homogenization. For alpha–beta forging, the deformation was applied at constant temperature between 950 and 1030 ◦ C after solutionizing at 1050 ◦ C for 10 min and holding at the test temperature for 5 min. Solutionizing prior to deformation was applied to permit a consistent initial prior-beta microstructure for all test conditions performed below the beta-transus temperature of 1045 ◦ C. Deformation of the specimens was performed at a constant strain rate between 0.001 and 1 s−1 and a total true strain2 between 0.2 and 1.2 followed by water quenching. For each test temperature, one specimen was heat treated and water-quenched just before deforming to represent the initial microstructure of the undeformed material. 1
Lubrication selection was performed by examining the influence of various lubricant combinations on the flow behavior; the optimum results of which defined these lubrication conditions. 2 The true strain (ε ) was calculated as ln(L/L ), where L is the final specimen t 0 height and L0 is the initial specimen height.
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Table 2 Automated metallographic procedure applied for IMI834 Stage
Surface Lubricant Suspension (m) Time (min) Force (lb/sample) Platen speed (rpm)
Coarse grind
Fine grind
Coarse polish
Fine polish
280 grit SiC Water None Until planar 3 150
600 grit SiC Water None 2 3 150
Woven-silk cloth Alcohol-based 9 diamond 4 3 100
Porous pad None 0.02 silica 5 3 100
2.4. Metallographic evaluation The specimens were sectioned parallel to the compression axis using a low concentration diamond wafering blade on a slow speed (200 rpm) cutoff saw that was operated with an oil coolant. Sectioned IMI834 specimens were mounted in Bakelite and prepared for metallographic examination using automated techniques for grinding and polishing as outlined in Table 2. After ultrasonic cleaning, the specimens were etched with a solution of 60 ml H2 O2 , 10 ml HF and 40 ml H2 O. Microstructural examination was performed using optical light microscopy with polarizing capabilities. To measure the beta grain size for specimens processed under various conditions, the grain boundaries of the transformed beta structure were manually demarcated using image analysis software. The size of the prior-beta grains was then measured manually using the linear intercept method according to standard ASTM E112 and with the automated functions of the image analysis software. On average, 150–200 grains were sampled for each processing condition for statistically representative results.
in the lower two-phase alpha–beta region (950–975 ◦ C). For all deformation temperatures, flow softening was observed to occur after the peak stress. However, there is a significant difference in the flow softening behavior of the stress–strain curves observed for deformation in the two-phase alpha–beta region as compared to the single-phase beta-region. For deformation in the single-phase beta region, the stress rose sharply, however an early yield point was reached before continuing on to a true peak stress. Beyond this peak stress, the degree of flow softening observed for all forging conditions above the transus was lower than that for below the transus. Also, above 1030 ◦ C, the flow stress appeared to reach a steady state value at true strains greater than 0.6. The stress–strain curve for the material that was deformed below the beta-transus (less than 1030 ◦ C) showed a sharp increase in stress at the outset of forging, eventually culminating in a peak stress. Beyond the peak stress, the curve dropped continuously and steady state conditions were not observed even up to a true strain value of 1.0 applied in the present work. 3.2. Microstructural features prior to forging
3. Results and discussion 3.1. Forging characteristics An example of the true stress–true strain curves obtained at a strain rate of 1 s−1 are presented in Fig. 2 for deformation above the beta transus (1125 ◦ C), at the beta transus (1050 ◦ C), in the upper two-phase alpha–beta region (1000–1030 ◦ C) and
Fig. 2. Flow behavior of IMI834 at temperatures in the beta and alpha–beta phase regions at a strain rate of 1 s−1 .
To effectively assess the influence of the deformation temperature, strain and strain rate on microstructural evolution, the structural features prior to deforming above and below the betatransus temperature were examined. Since for IMI834 deformation at temperatures above 1045 ◦ C occurs in the single-phase beta region, and below 1045 ◦ C in the two-phase alpha–beta region, the observation of transformed structures can be related to microstructure at the onset of deformation. 3.2.1. Single-phase beta region For the specimens heated above the beta transus, the quenched microstructure was observed to consist of a continuous matrix of transformed beta grains that had an acicular alpha structure with a martensitic or Widmanst¨atten appearance as shown in Fig. 3a. Hence at the deformation temperature, the microstructure before hot working consisted only of single-phase beta, and the transformation structure of the grains observed is a result of the fast cooling rate that causes the alpha phase to have an acicular appearance. The average prior-beta grain size (GS ) in the single-phase beta region was determined to be approximately 304 m after holding for 10 min at the processing temperatures between 1050 and 1125 ◦ C, as plotted in Fig. 4.
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Fig. 4. The prior-beta grain size of the IMI834 material after holding for 10 min at the processing temperature.
3.2.2. Two-phase alpha–beta region The transformation microstructures for the specimens held at 1030 and 1000 ◦ C were observed to consist of a continuous matrix of transformed beta grains with discontinuous alpha platelets at the prior-beta grain boundaries as shown in Fig. 3b. This discontinuous alpha structure appeared as a bright protruding phase in the microstructure, which has been attributed to the relatively high resistance of alpha to etching due to its hexagonal close packed crystal structure [21]. The transformed beta grains had a fine acicular alpha structure, that was martensitic or Widmanst¨atten in appearance, and the alpha phase at the prior-beta grain boundaries was lenticular in morphology, as illustrated in Fig. 3b. The presence of a martensitic or Widmanst¨atten alpha phase structure in the transformed beta grains suggests that, at temperatures of 1030 and 1000 ◦ C, the microstructure before forming mostly consisted of beta grains with some alpha phase at the grain boundaries. It is noteworthy that, at 1030 and 1000 ◦ C, the alpha phase fraction expected from the beta approach curve given in Fig. 5 is approximately 8% and 45%,
Fig. 3. The microstructure of IMI834 prior to forming at temperatures in the single-phase  region (>1050 ◦ C), the upper two-phase ␣– region (1000–1030 ◦ C) and the lower two-phase ␣– region (<1000 ◦ C): (a) 1050 ◦ C, (b) 1000 ◦ C and (c) 950 ◦ C. At processing temperature: (1)  grain boundaries, (2)  phase region, (3) ␣ at  grain boundaries and (4) transformed  phase region. Fig. 5. Extrapolated beta approach curve for IMI834 from 900 to 1050 ◦ C [21].
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respectively, and considerably greater than the evaluated values of 1.2% and 3.8%, respectively. Hence, it appears that the extent of the beta to alpha phase transformation after 5 min holding at 1030 and 1000 ◦ C has, most likely, not reached a steady state. Thus further increases in the alpha phase fraction are expected with increasing time during processing at higher strain values and/or decreasing strain rate. Moreover, it is anticipated that the greatest increase in the phase transformation will occur with decreasing strain rate considering that at the highest value of 1 s−1 deformation to a true strain of 1 is completed within 1 s while at the lowest value of 0.001 s−1 more than 15 min are spent hot working at temperature. For the specimens held at processing temperatures of 975 and 950 ◦ C, the transformation microstructures were observed to consist mostly of a continuous matrix of transformed beta grains having a lamellar alpha structure with a continuous layer of the alpha phase at the prior-beta grain boundaries, as illustrated in Fig. 3c. Within the lamellar alpha microstructure, some “islands” of a transformed beta phase having an alpha structure with a martensitic or Widmanst¨atten appearance were also observed, especially at 975 ◦ C. For these processing temperatures in the lower bound of the alpha–beta field, the beta approach curve, given in Fig. 5, indicates that the matrix of the IMI834 alloy should be dominated by the alpha phase and is consistent with the structural observations noted for the transformed beta phase. Specifically, the morphology of the transformed microstructure within the prior-beta grains is a result of the cooling conditions during processing, and the rate of cooling affects both the thickness and orientation of the alpha plates. A small driving force for transformation, at relatively slow cooling rates, achieves selective growth of the nucleated alpha plates that allows the lamellae to be arranged in packets or colonies; however, with increasing cooling rate, rapid growth of the nucleated alpha gives the multi-oriented appearance of the lamellae in the Widmanst¨atten structure. In this work, when cooling the IMI834 specimens from the solutionizing temperature (1050 ◦ C) to 950 or 975 ◦ C, the relatively slow transformation of beta to alpha, under furnace holding conditions of 5 min at temperature, led to the formation of a microstructure consisting mostly of lamellae having a colony appearance, as indicated in Fig. 3c. This alpha was present before deformation and represented roughly 78% and 90% of the microstructure at 975 and 950 ◦ C, respectively. In addition, longer holding times at the deformation temperature prior to quenching did not increase the lamellar fraction at either 950 or 975 ◦ C, indicating that the microstructure after holding for 5 min had reached steady state. Quenching of the specimen after holding in the furnace at the deformation temperature caused rapid transformation of the remaining beta phase. This led to the formation of Widmanst¨atten alpha that represented roughly 22% and 10% of the microstructure at 975 and 950 ◦ C, respectively. Hence, the beta fraction at the deformation temperature prior to hot working at 975 and 950 ◦ C was determined to be approximately 22% and 10%, respectively. Moreover, the beta fraction values obtained in this work for the lower alpha–beta temperature region were observed to alight on the extrapolated curve of the equilibrium phase data [22] given in Fig. 5 and further support the finding that holding for 5 min at
950 and 975 ◦ C is sufficient for equilibrating the microstructure. The faster equilibration for temperatures in the lower two-phase region may be attributed mainly to a greater undercooling effect at 950 and 975 ◦ C as compared to 1000 and 1030 ◦ C as well as a secondary effect of a decrease in the cooling rate from 140 to 43 ◦ C min−1 with decreasing temperature from 1030 to 950 ◦ C during furnace cooling. Hence, during cooling from the solutionizing stage in the single-phase beta region to the processing temperature, 975 or 950 ◦ C, the beta phase was observed to predominately transform to alpha. This phase transformation of beta to alpha was observed to first occur as a layer of alpha at grain boundary sites followed by nucleation and growth of alpha plates. On account of the slower furnace cooling conditions from the beta solutionizing stage to the processing temperature, the formation of lamellar colonies of thick alpha plates occurred, inevitably due to the reduction in the nucleation sites available as compared to air cooling or water quenching that usually develop martensitic or Widmanst¨atten structures [21]. Therefore, for processing temperatures of 975 and 950 ◦ C, the microstructure before quenching (i.e. at the onset of hot working) consists of a continuous layer of grain boundary alpha around predominately transformed beta grains having a lamellar alpha plate structure. The microstructural observation of martensitic or Widmanst¨atten alpha “islands” is then due to the quenching of the specimens from the processing temperature that gave a rapid phase transformation of the remnant beta regions, which represented 22% and 10% of the microstructure at 975 and 950 ◦ C, respectively. 3.3. Microstructural features associated with forging 3.3.1. Deformation in the single-phase beta region Knowing that the initial microstructure in the single-phase beta region (1050–1125 ◦ C) prior to hot working consisted of beta grains having an average grain size of 304 m (Fig. 4), the effect of deformation processing was examined by compressing IMI834 specimens to true strains of 0.2, 0.6, 0.8 and 1.2. Between 1050 and 1125 ◦ C, the morphology of the beta grains becomes elongated (pancake-shape) with increasing strain in the plane perpendicular to the forging direction, as typified for processing at 1050 ◦ C in Fig. 6a and b. The beta grain boundaries were observed to become increasingly jagged and serrated with increasing strain and, at a true strain of 0.2, the formation of small-recrystallized grains was apparent in the vicinity of the deformed grain boundaries (Fig. 6a). In particular, deformation at 1050 ◦ C with a strain rate of 1 s−1 to a true strain of 0.2 was observed to attain an average recrystallized fraction of 2.3%. With increasing strain, further recrystallization of beta grains was apparent at the prior-beta grain boundaries, which gave a necklace-appearance of fine equiaxed beta grains around the deformed and elongated beta phase (Fig. 6b). Since the nucleation of dynamically recrystallized grains is initiated when a critical dislocation density in the deforming material is reached during hot working, nuclei for dynamic recrystallization (DRX) typically appear at the grain boundaries or other internal crystallographic features such as shear bands or inclusions. Hence, the
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Fig. 6. Microstructure of IMI834 deformed at 1050 ◦ C (single-phase  region) with a strain rate and true strain of (a) 1 s−1 and 0.2, (b) 1 s−1 and 0.8 and (c) 0.01 s−1 and 0.8. At processing temperature: (1) original  grains and (2) recrystallized  grains.
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recrystallization of the beta phase that occurs in the single-phase region (1050–1125 ◦ C) can be predominately related to dynamic processes, as equiaxed small and slightly elongated intermediate grains are apparent at the original beta grain boundaries. This finding is consistent with previous results on hot working of IMI834 that associated the occurrence of recrystallization under isothermal forging conditions to deformation induced dynamic and/or metadynamic recrystallization [13,19,23–25]. Such a structural evolution has also been reported to occur regularly during beta processing of titanium alloys such as Ti–6Al–4V [21,25,26]. Moreover, the incidence of serrated grain boundaries during beta phase deformation has been considered as a special case of DRX, referred to as necklace recrystallization, since this softening process concentrates at the grain boundaries during hot working [26,27]. In addition, it was observed that beta phase working in IMI834 using high strain rates (0.1 and 1 s−1 ) and true strains of up to 1.2 leads to the formation of a mixed grain microstructure consisting of small equiaxed recrystallized grains as well as large deformed prior-beta grains as revealed in Fig. 6b. This microstructural non-uniformity has been related to the incidence of selective recrystallization in the high strain regions along the beta grain boundaries combined with a lower driving force for recrystallization in the interior regions of the beta grains where recovery is most likely dominant [26]. The occurrence of selective recrystallization during beta phase hot working in titanium-based materials has been noted to be particularly problematic for grain size and structure uniformity because the mixed grain microstructure of both large and small grains that is produced can not be eliminated by any heat treatment process [28–30]. During beta phase hot working of IMI834, control of the uniformity in the grain structural characteristics such as size and morphology, which is critical for optimizing secondary processing and final mechanical properties, necessitates prudent selection of the thermomechanical processing parameters. Hence a high degree of dynamic recrystallization during processing is crucial for producing uniformity in both microstructural and mechanical property characteristics. However, as noted in this work, processing conditions at 1050 ◦ C must surpass a strain rate of 1 s−1 and a true strain of 1.2 to achieve just over 50% recrystallization, as illustrated in Fig. 7. Thus for the forging conditions in this work (i.e. strain rates up to 1 s−1 , true strains up to 1.2 and processing temperatures of 950–1125 ◦ C), the elongated beta grains dominated the high temperature deformation microstructure of IMI834 and the amount of recrystallization was, overall, quite limited. These observations are similar to the findings of less than 30% recrystallization in Ti–6Al–4V reported in previous work for processing conditions of 1000–1050 ◦ C and strain rates of 0.05–1 s−1 [31]. In general, this limitation for grain refinement by dynamic recrystallization in titanium alloys has been attributed to the difficulty in establishing a sufficiently high dislocation density. Specifically, the high self-diffusivity of titanium in the beta phase readily enables dislocation mobility and decreases the stored energy through dynamic recovery, which in turn homogenizes the dislocation substructure and reduces the tendency (or driving force) for recrystallization [32].
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Fig. 7. Microstructure of IMI834 deformed at 1050 ◦ C at a strain rate of 1 s−1 to a true strain of 1.2. At processing temperature: (1) original  grains and (2) recrystallized  grains.
In the single-phase beta region (1050–1125 ◦ C), a decrease in the strain rate was observed to result in the growth of the recrystallized beta grains, such that at the lowest strain rate, 0.001 s−1 , the once fully recrystallized (dynamic) fine beta grains (25–40 m depending on the temperature) have grown (120–150 m depending on the temperature) and elongated during the continuing deformation, as depicted in the Fig. 6c. Also, with increasing strain rate the degree of dynamic recrystallization was observed to increase and is consistent with previous findings for thermomechanical processing of Ti–6Al–4V in the single-phase beta region [31]. 3.3.2. Deformation in the upper two-phase alpha–beta region For processing temperatures of 1000 and 1030 ◦ C, the microstructure before forming was similar to that in the singlephase beta region, with the exception that discontinuous alpha platelets were present at the beta grain boundaries during deformation in the upper alpha–beta region. Notwithstanding this structural difference, the effect of deformation was similar to that observed for hot working in the single-phase beta region, as demonstrated in Fig. 8a and b. Specifically, with increasing strain, the beta grains elongated and the discontinuous alpha platelets at the grain boundaries became aligned in the plane perpendicular to the forging direction. Additionally the formation of small recrystallized beta grains were observed, but, as their concentration at prior-beta grain boundaries was lower than that observed during deformation in the single-phase beta region, only limited necklace like structure of the fine equiaxed beta grains was apparent around the deformed beta grains. From the beta phase fractions given in Fig. 5, it was remarked that the microstructure after holding for 5 min at 1000 and 1030 ◦ C had not equilibrated and perhaps further transformation
Fig. 8. Microstructure of IMI834 deformed at 1000 ◦ C (upper two-phase ␣– region) with a strain rate and true strain of (a) 1 s−1 and 0.2, (b) 1 s−1 and 0.8 and (c) 0.001 s−1 and 0.8. At processing temperature: (1) original  grains, (2) recrystallized  grains and (3) ␣ at  grain boundaries.
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of the beta phase to alpha could occur with longer holding time at higher applied strain or lower strain rate conditions. In this work, the extent of the beta to alpha phase transformation with increasing strain at deformation temperatures of 1000 and 1030 ◦ C was observed to be most marked at strain rates of 0.001 and 0.01 s−1 . For example, comparison of IMI834 deformed at 1000 and 0.001 s−1 to a true strain of 0.8, as shown in Fig. 8c indicates that additional holding during hot working resulted in the formation of nearly 40% alpha, either at the beta grain boundaries or within the beta grains, because the structure before forming was not under equilibrium conditions as predicted by the beta approach curve given in Fig. 5. This additional transformation of beta to alpha at a strain rate of 0.01 or 0.001 s−1 produced a change in the microstructural characteristics during deformation at 1000 and 1030 ◦ C. Specifically, at the start of hot working, the predominantly beta grain structure was observed to deform concomitantly with the formation of alpha at the beta grain boundaries and within the prior-beta grain structure, thereby giving a continuous alpha layer and alpha platelets in the grain interior, respectively. Hence, at a strain rate of 0.1–1 s−1 , deformation processing at 1030 and 1000 ◦ C was similar to that in the single-phase beta region since the microstructure consisted of a continuous matrix of beta grains with discontinuous alpha platelets present at the beta grain boundaries as revealed in Fig. 8b. So processing with a strain rate of 0.1–1 s−1 at either 1000 or 1030 ◦ C was observed to result in the elongation of the beta grains and the discontinuous alpha platelets with the formation of recrystallized beta grains. At lower strain rates of 0.01–0.001 s−1 , processing at either 1000 or 1030 ◦ C resulted in the deformation of the beta grains and the alpha platelets, the latter of which formed additionally during deformation at the grain boundaries and within the beta grains, without evidence of beta recrystallization as shown in Fig. 8c. This kinetic effect on phase transformation during deformation is consistent with the previous laboratory work that showed that the fraction of alpha during extrusion of IMI834 was a function
Fig. 10. Microstructure of IMI834 deformed at 950 ◦ C (lower two-phase ␣– region) with a strain rate and true strain of (a) 1 s−1 and 0.2, (b) 1 s−1 and 0.8 and (c) 0.001 s−1 and 0.8. At processing temperature: (1) transformed  grains, (2) remnant  phase and (3) lenticular ␣ phase. Fig. 9. Microstructure of IMI834 deformed at 975 ◦ C with a strain rate 1 s−1 and true strain of 0.8. At processing temperature: (1) transformed  grains, (2) remnant  phase and (3) lenticular ␣ phase.
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of the strain rate and strain [22]. Hence for the upper alpha–beta temperature region, the processing window for dynamic recrystallization of the beta grains in IMI834 is at a strain rate greater than 0.01 s−1 at 1000 ◦ C and 0.001 s−1 at 1030 ◦ C. 3.3.3. Deformation in the lower two-phase alpha–beta region At processing temperatures of 950 and 975 ◦ C, the microstructure equilibrated within 5 min of holding and, before forming, there were remnant beta phase regions in a continuous matrix of transformed beta grains having a lamellar alpha structure with continuous alpha at the prior-beta grain boundaries. With increasing strain, for the remnant beta phase regions at both 975 and 950 ◦ C, no recrystallization was observed as shown in Fig. 9. Nonetheless, the larger “islands” of the beta phase present before forming were observed to elongate in the plane perpendicular to the forging direction with increasing strain. In contrast, for all strain rate conditions, deformation processing at 950 and 975 ◦ C was observed to work the alpha structure within the transformed beta grains, i.e. the elongated lamellar alpha platelets distorted (wavy appearance) increasingly with increasing strain as depicted in the Fig. 10a and b. These prior-beta phase regions having a martensitic or Widmanst¨atten structure were easily differentiable from the continuous lamellar alpha microstructure. Particularly, bending, distortion, shearing and break up of the lamellar structure was observed to occur to an increasing extent as the strain increased and the processing temperature decreased. These deformation characteristics of the lamellar structure are consistent with previous work on Ti–6Al–4V [31,33]. With increasing strain rate, deformation processing at 950 and 975 ◦ C was observed to work, to a greater degree, the alpha structure within the transformed beta grains. Similar to the effect of increasing strain, with increasing strain rate, the elongated lamellar alpha platelets distorted increasingly and break up of the lamellar structure was apparent. Weiss et al. have attributed lamellar alpha phase separation into shorter segments to two possible mechanisms that can occur during thermomechanical processing [33]. In their work it was shown that hot working led to the development of both low and high angle sub-boundary formation across the alpha plates. In the present work, separation of the alpha lamellae was observed to occur for the specimens processed at 950 and/or 975 ◦ C at strain rate conditions of 0.001 s−1 or greater, as illustrated in Fig. 10c. In contrast to this evolution in the lamellar alpha microstructure, the beta phase regions simply elongated with increasing strain and strain rate for deformation at either 975 or 950 ◦ C and no evidence of dynamic recrystallization was observed (Fig. 9) using field emission gun scanning electron microscopy and electron backscattered diffraction techniques [17].
and strain rate on structural evolution in IMI834 during hot working in the single-phase beta field and the two-phase alpha–beta region, it is apparent that the processing window for beta phase recrystallization is restricted to thermomechanical conditions giving limited formation of discontinuous alpha lamellae at the beta grain boundaries. Specifically, it was explicitly demonstrated that even when the alpha fraction was constrained to a discontinuous thin film at the beta grain boundaries the extent of dynamic recrystallization was lower than that observed for processing in the single-phase beta field. In view of the mechanisms for deformation-induced recrystallization, the steady state flow stress conditions can be considered to evaluate the influence of the processing conditions on the grain size [34–36]. For the beta processing conditions in this work, full recrystallization was not attained and the flow curves continued to exhibit a decrease in the stress with increasing strain. Nonetheless, it was noticeably established that the grain size of the recrystallized beta phase depends on both the processing temperature and strain rate, with higher temperatures and lower strain rates giving an increase in the average grain size of the recrystallized beta phase, as shown in Figs. 11 and 12, respectively. Also, the overall effect of strain rate appears to be greater than that of temperature and corroborates with the results of the dependence of prior-beta grain size on deformation temperature given in Fig. 4. Previous work has indicated that for temperatures that are close to, or exceeding the beta transformation temperature, where the volume fraction of alpha is small, the constraint on grain growth during recovery or recrystallization is minor and a strong dependence of grain size on strain rate exists [13,35]. However, the effect of temperature appears to be increasingly marked with decreasing strain rate from 1 to 0.001 s−1 , which inevitably is related to a kinetic effect during the longer holding period at temperature. Hence, high strain rates (greater than 1 s−1 ) and temperatures just above the beta transus offer an optimal potential for recrystallization of the beta phase from an initial grain size of over 300 to less than 30 m.
4. Discussion In this work, the thermomechanical processing of IMI834 over a wide range of testing conditions enabled evaluation of the effect of thermomechanical parameters on the microstructural characteristics. From the dependence of the alpha and beta phase fractions on processing temperature, as well as the role of strain
Fig. 11. Effect of processing temperature on the recrystallized beta grain size for various strain rate conditions and a true strain of 0.8.
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Fig. 14. Shear band formation in IMI834 through the aligning of adjacent alpha lamellae during deformation processing at a temperature of 950 ◦ C and strain rate of 1 s−1 . Fig. 12. Effect of strain rate on the recrystallized beta grain size for various processing temperatures and a true strain of 0.8.
The observed increase in the recrystallized beta grain size with increasing temperature and decreasing strain rate (Figs. 11 and 12), i.e. with decreasing Zener–Hollomon parameter (Z), was correlated (Fig. 13) and the linear dependence that was found could be fitted to a single equation of the form: Drex = kZ−0.19
(1)
where Drex is the recrystallized grain size of the beta phase, Z the Zener–Hollomon parameter (˙ε exp(Q/RT )), and k is a constant. From Fig. 13, the refinement in the beta grain size with increasing Z stems from the greater driving force for recrystallization (a higher dislocation density). The R2 -value of 0.98 indicates that overall there is a good approximation of the model with the experimental data in Fig. 13. It is also noteworthy that although the fraction recrystallized, for certain Z values, was in the early stages of dynamic recrystallization, a reasonably good
Fig. 13. Variation in the recrystallized grain size of the beta phase in IMI834 with the Zener–Hollomon parameter for true strains of 0.6 (filled markers) and 0.8 (open markers).
correlation between the Z and the average beta grain size was obtained. This may be related to the dynamically recrystallized beta grain size having reached steady state conditions beyond a true strain value of 0.6, as illustrated by the similarity in the beta grain size data for the two strain conditions plotted in Fig. 13. The observed linearity between the dynamically recrystallized beta grain size and the Zener–Hollomon parameter has also been reported recently for other titanium alloys [37,38]. Moreover, the exponent in the power law relationship of Eq. (1), i.e. 0.19, is in good agreement with value of 0.17 obtained for beta phase hot working of Ti–6Al–4V by Tamirisakandala et al. [38]. As compared to deformation processing above 1000 ◦ C, for which recrystallization of the beta phase was observed, the considerable flow softening apparent at 950 and 975 ◦ C may be attributed to dislocation processes that were observed in this work to result primarily in platelet distortion as well as crystallographic changes [17]. In particular, slip along alpha lamellae during deformation was physically manifested in the micro- and macro-structural observations of adjacent lamellae aligning in the shear bands and shear band formation as shown in Fig. 14. These findings are supported by the recent work of Semiatin and Biehler [39] that related dislocation processes, platelet distortion and crystallographic texture as contributing factors to flow softening mechanisms. In particular, Semiatin and Biehler have indicated that during deformation of alpha–beta titanium alloys in the two-phase field, dislocation pileups/slip transfer processes may be instigated at alpha–beta interfaces due to the dissimilarity in the sub-grain formation characteristics of the two-phases [40]. In their work, the alpha–beta interfaces have been hypothesized as being analogous to grain boundaries in single-phase materials and, on account of the limited number of independent slip systems in the hexagonal crystal structure of alpha titanium, a strong grain boundary or Hall–Petch strengthening dependence on the alpha lath/platelet thickness was determined. Regarding the relevance of slip transmission across alpha–beta interface in relation to Hall–Petch strengthening in titanium alloys, previous work has suggested that during uniaxial compression of Ti–6Al–4V, with a colony microstructure, aligning of shear bands in adjacent lamellae occurs with a resulting
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localization in the flow characteristics [39]. This development of microscopic shear bands within the alpha titanium lamellae indicates that sub-grain formation in the lamellar microstructure is difficult and it has been suggested that the principal mechanism responsible for the considerable flow softening behavior in Ti–6Al–4V can be attributed to slip transmission coupled with a loss of alpha–beta interface strengthening [41]. For IMI834, the deformation processing conditions in the two-phase alpha–beta region that exhibited considerable flow softening behavior (i.e. at 950 and 975 ◦ C) may then be related to similar phenomena as supported by the platelet distortion and the presence of shear band formation (Fig. 14).
single linear relationship between a strain rate range 0.001 and 1 s−1 for experimental conditions where the microstructure was predominately beta. (6) Processing at temperatures in the lower two-phase field, 975 and 950 ◦ C was observed to work the lamellar alpha structure within the transformed beta grains and bending, distortion, shearing and break up of the lamellar structure was observed to occur to an increasing extent as the strain and strain rate increased. At both 975 and 950 ◦ C no recrystallization was observed in the remnant (interlamellar and islands) beta phase regions. Acknowledgments
5. Conclusions Manufacturing of titanium alloys is a multi-step process that involves consecutive heating and deformation at beta and alpha–beta phase temperatures. The evolution of component characteristics (shape, structure and properties) that develop during deformation processing require a methodical understanding of the role of various parameters on the hot working behavior. In this work isothermal hot working of IMI834 was studied in the temperature range 950–1125 ◦ C at strain rates of 0.001–1 s−1 and true strains up to 1.2 with the follow observations: (1) Deformation processing of IMI834 indicated that at temperatures in the single-phase beta region, the flow stress behavior exhibited strain hardening and minor flow softening, while in the two-phase alpha–beta region considerable flow softening occurred. (2) Processing in the single-phase beta region (1050–1125 ◦ C, 0.001–1 s−1 ) resulted in the serration of beta grain boundaries and the formation of small dynamically recrystallized grains (by a true strain of 0.2) which concentrated at the prior-beta grain boundaries and gave a necklace-appearance of fine equiaxed beta grains around the deformed and elongated beta phase (true strains greater than 0.6). (3) Processing at 1000 and 1030 ◦ C involved working of a predominately beta microstructure with alpha at the beta grain boundaries, especially for strain rates greater than 0.01 s−1 , since the microstructure was not equilibrated. Under such conditions, hot working was observed to recrystallize the beta phase albeit in lower concentration and the processing window for dynamic recrystallization of the beta grains (beta phase working) was determined to be a strain rate greater than 0.01 s−1 at 1000 ◦ C and 0.001 s−1 at 1030 ◦ C. (4) For conditions giving a dominant beta phase microstructure, decreasing the strain rate from 1 to 0.001 s−1 was observed to increase the size of the recrystallized grains from 25–40 to 120–150 m, respectively. Also, the average size of the dynamically recrystallized beta grains was observed to increase from 25–120 to 40–150 m with increasing processing temperature from 1000 to 1125 ◦ C, respectively, depending on the strain rate condition. (5) Variation in the dynamically recrystallized beta grain size with the Zener–Hollomon parameter was observed to give a
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Influence of thermomechanical processing on microstructural evolution in near-␣ alloy IMI834 P. Wanjara a,∗ , M. Jahazi a , H. Monajati b , S. Yue b a
National Research Council, Institute for Aerospace Research, Aerospace Manufacturing Technology Center, 5145 Decelles Avenue, Campus Universit´e de Montr´eal, P.O. Box 40, Montr´eal, Que., Canada H3S 2S4 b McGill University, Mining, Metals and Materials Engineering, 3610 University Street, Montr´ eal, Que., Canada H3A 2B2 Accepted 27 October 2005
Abstract The effect of processing parameters on the evolution of the microstructural characteristics of IMI834 were examined over the temperature range 950–1125 ◦ C at constant strain rates of 0.001–1 s−1 and true strains up to 1.2. During isothermal deformation, differences in the flow softening behavior were observed for processing in the single-phase beta region, as compared to the two-phase alpha–beta region, which has been related to differences in the structural changes occurring in IMI834. In particular, deformation processing of the alpha–beta structure was observed to work the lamellar alpha structure of the transformed beta grains without recrystallization of either phase, whilst microstructural development in a predominately beta structure resulted in the formation of dynamically recrystallized grains that were necklaced around the deformed and elongated beta phase. For the later case, the variation in the dynamically recrystallized grain size was determined to follow a Zener–Hollomon relationship between a strain rate range 0.001 and 1 s−1 . © 2005 Elsevier B.V. All rights reserved. Keywords: Titanium alloys; Dynamic recrystallization; Isothermal deformation; Shear banding; Flow softening
1. Introduction Over the last 50 years, the evolution of titanium alloys has been driven by the ever-increasing need to improve the efficiency of gas turbines by elevating operating temperatures and improving component durability. Titanium alloys have met the stringent requirements of the aerospace industry through alloy chemistry development, and more recently through microstructural design, both of which are necessary to remain responsive to the increasing needs for engine performance [1–4]. For manufacturing of specific component shapes from titanium alloys, it is recognized that the thermomechanical processing conditions can be optimized to control the microstructuremechanical property characteristics. However, considering the restrictive workability limit (or processing window) of titanium alloys in general, an increased understanding of the relationship between processing and microstructure is particularly critical for sustaining further improvements in performance and reliability. For instance, near-alpha titanium alloy, IMI834 has been
∗
Corresponding author. Tel.: +1 514 283 9380; fax: +1 514 283 9445. E-mail address: [email protected] (P. Wanjara).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.10.042
designed for compressor disks with improved tensile strength, fatigue resistance and creep performance at temperatures up to 600 ◦ C by means of a tailored bimodal microstructure [5–8]. Processing of IMI834 usually involves a series of primary and secondary deformation operations that incorporate a combination of non-homogeneous thermomechanical conditions involving complex metallurgical phenomena including phase transformations, both dynamic and static restoration mechanisms as well as grain growth [9–11]. The final component microstructure is dependent on the various structural features inherited from the primary stages (cogging or ingot breakdown) and the last forming and heat treatment operations [12]. A number of studies [13–16] have indicated that a bimodal structure consisting of primary alpha particles and fine transformed beta grains improves the mechanical properties of IMI834. In general, during processing of IMI834, the evolution in microstructure is dependent on dynamic recovery and recrystallization occurring under an applied load as well as static recovery, static recrystallization and grain growth that proceed upon load removal. Moreover, the occurrence of these dynamic and static phenomena during thermomechanical processing can influence the kinetics of restoration and grain growth in both the beta and alpha phases of IMI834 and produce considerable
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inhomogeneity in the final microstructure [17]. Presently, it has been reported that during high temperature deformation processing of IMI834 the flow stress exhibits a strong dependence on the temperature with the probability of considerable flow softening, both of which are related to the dependence of the alpha and beta phases on variations in the deformation parameters such as temperature, strain and strain rate [6,10,18]. It has also been suggested that the flow softening behavior is related to the restoration of the material during processing [19]. However, systematic evaluation of the role of various processing parameters (strain, strain rate, temperature) on the deformation characteristics (flow stress–strain behavior) and microstructural development in IMI834 remains a matter of an extended investigation. In this paper, the research work reported formed part of a larger program on developing the technology for processing titanium materials through optimizing the forging conditions to allow control of microstructure and properties in the finished product. Part of this program involved characterization of the flow behavior of IMI834 under hot forging conditions. The results of this research work are described elsewhere [20]. In this paper, the results of subsequent examination of the deformationinduced microstructures are reported to establish the influence of processing (ingot breakdown and closed die forging) conditions on structural evolution. 2. Experimental procedure 2.1. Starting materials Hot rolled IMI834 material was received in bar form with an outer diameter of 18 cm and chemical composition given in Table 1. The beta transformation temperature of this material was approximately 1045 ◦ C. The microstructure of the as-received material consisted of a fine, bimodal alpha–beta structure of approximately 25.6% equiaxed primary alpha (with a standard deviation of 3%) within a fine transformed beta matrix, as shown in Fig. 1. Cylindrical specimens of 7.5 mm in diameter and 11.4 mm in height were electro-discharged machined from the bar for the high temperature compression experiments. 2.2. Hot working equipment The hot working studies were conducted on a computerized 100 kN Materials Testing System (MTS 810) that was adapted for elevated temperature compression testing to 1200 ◦ C Table 1 Chemical composition (wt.%) of IMI834 material Al Sn Zr Nb Mo Si C Ti
5.5 4.0 3.5 0.7 0.5 0.35 0.06 Balance
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Fig. 1. Bimodal microstructure of as-received IMI834 material.
using a Research Incorporated radiant furnace interfaced with a Micristar digital programmer and controller. During high temperature forming, the compression specimen and the anvils were enclosed within a quartz tube, in which argon gas was passed to prevent oxidation of the work piece and tooling. To minimize friction during hot working, the ends of the specimens were coated with delta glaze and a thin layer of mica1 was placed between the face of the specimen and the anvils in order to maintain uniform deformation and avoid sticking problems during quenching. 2.3. Deformation schedule Hot working was performed using two distinct thermal cycles designed for: (1) beta forging corresponding to simulation of ingot breakdown of a cast microstructure to a wrought product and (2) alpha–beta forging corresponding to the simulation of closed-die forging to a final shape. In the first cycle, beta forging was conducted at constant temperature between 1050 and 1125 ◦ C after holding for 15 min at temperature to allow for thermal homogenization. For alpha–beta forging, the deformation was applied at constant temperature between 950 and 1030 ◦ C after solutionizing at 1050 ◦ C for 10 min and holding at the test temperature for 5 min. Solutionizing prior to deformation was applied to permit a consistent initial prior-beta microstructure for all test conditions performed below the beta-transus temperature of 1045 ◦ C. Deformation of the specimens was performed at a constant strain rate between 0.001 and 1 s−1 and a total true strain2 between 0.2 and 1.2 followed by water quenching. For each test temperature, one specimen was heat treated and water-quenched just before deforming to represent the initial microstructure of the undeformed material. 1
Lubrication selection was performed by examining the influence of various lubricant combinations on the flow behavior; the optimum results of which defined these lubrication conditions. 2 The true strain (ε ) was calculated as ln(L/L ), where L is the final specimen t 0 height and L0 is the initial specimen height.
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Table 2 Automated metallographic procedure applied for IMI834 Stage
Surface Lubricant Suspension (m) Time (min) Force (lb/sample) Platen speed (rpm)
Coarse grind
Fine grind
Coarse polish
Fine polish
280 grit SiC Water None Until planar 3 150
600 grit SiC Water None 2 3 150
Woven-silk cloth Alcohol-based 9 diamond 4 3 100
Porous pad None 0.02 silica 5 3 100
2.4. Metallographic evaluation The specimens were sectioned parallel to the compression axis using a low concentration diamond wafering blade on a slow speed (200 rpm) cutoff saw that was operated with an oil coolant. Sectioned IMI834 specimens were mounted in Bakelite and prepared for metallographic examination using automated techniques for grinding and polishing as outlined in Table 2. After ultrasonic cleaning, the specimens were etched with a solution of 60 ml H2 O2 , 10 ml HF and 40 ml H2 O. Microstructural examination was performed using optical light microscopy with polarizing capabilities. To measure the beta grain size for specimens processed under various conditions, the grain boundaries of the transformed beta structure were manually demarcated using image analysis software. The size of the prior-beta grains was then measured manually using the linear intercept method according to standard ASTM E112 and with the automated functions of the image analysis software. On average, 150–200 grains were sampled for each processing condition for statistically representative results.
in the lower two-phase alpha–beta region (950–975 ◦ C). For all deformation temperatures, flow softening was observed to occur after the peak stress. However, there is a significant difference in the flow softening behavior of the stress–strain curves observed for deformation in the two-phase alpha–beta region as compared to the single-phase beta-region. For deformation in the single-phase beta region, the stress rose sharply, however an early yield point was reached before continuing on to a true peak stress. Beyond this peak stress, the degree of flow softening observed for all forging conditions above the transus was lower than that for below the transus. Also, above 1030 ◦ C, the flow stress appeared to reach a steady state value at true strains greater than 0.6. The stress–strain curve for the material that was deformed below the beta-transus (less than 1030 ◦ C) showed a sharp increase in stress at the outset of forging, eventually culminating in a peak stress. Beyond the peak stress, the curve dropped continuously and steady state conditions were not observed even up to a true strain value of 1.0 applied in the present work. 3.2. Microstructural features prior to forging
3. Results and discussion 3.1. Forging characteristics An example of the true stress–true strain curves obtained at a strain rate of 1 s−1 are presented in Fig. 2 for deformation above the beta transus (1125 ◦ C), at the beta transus (1050 ◦ C), in the upper two-phase alpha–beta region (1000–1030 ◦ C) and
Fig. 2. Flow behavior of IMI834 at temperatures in the beta and alpha–beta phase regions at a strain rate of 1 s−1 .
To effectively assess the influence of the deformation temperature, strain and strain rate on microstructural evolution, the structural features prior to deforming above and below the betatransus temperature were examined. Since for IMI834 deformation at temperatures above 1045 ◦ C occurs in the single-phase beta region, and below 1045 ◦ C in the two-phase alpha–beta region, the observation of transformed structures can be related to microstructure at the onset of deformation. 3.2.1. Single-phase beta region For the specimens heated above the beta transus, the quenched microstructure was observed to consist of a continuous matrix of transformed beta grains that had an acicular alpha structure with a martensitic or Widmanst¨atten appearance as shown in Fig. 3a. Hence at the deformation temperature, the microstructure before hot working consisted only of single-phase beta, and the transformation structure of the grains observed is a result of the fast cooling rate that causes the alpha phase to have an acicular appearance. The average prior-beta grain size (GS ) in the single-phase beta region was determined to be approximately 304 m after holding for 10 min at the processing temperatures between 1050 and 1125 ◦ C, as plotted in Fig. 4.
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Fig. 4. The prior-beta grain size of the IMI834 material after holding for 10 min at the processing temperature.
3.2.2. Two-phase alpha–beta region The transformation microstructures for the specimens held at 1030 and 1000 ◦ C were observed to consist of a continuous matrix of transformed beta grains with discontinuous alpha platelets at the prior-beta grain boundaries as shown in Fig. 3b. This discontinuous alpha structure appeared as a bright protruding phase in the microstructure, which has been attributed to the relatively high resistance of alpha to etching due to its hexagonal close packed crystal structure [21]. The transformed beta grains had a fine acicular alpha structure, that was martensitic or Widmanst¨atten in appearance, and the alpha phase at the prior-beta grain boundaries was lenticular in morphology, as illustrated in Fig. 3b. The presence of a martensitic or Widmanst¨atten alpha phase structure in the transformed beta grains suggests that, at temperatures of 1030 and 1000 ◦ C, the microstructure before forming mostly consisted of beta grains with some alpha phase at the grain boundaries. It is noteworthy that, at 1030 and 1000 ◦ C, the alpha phase fraction expected from the beta approach curve given in Fig. 5 is approximately 8% and 45%,
Fig. 3. The microstructure of IMI834 prior to forming at temperatures in the single-phase  region (>1050 ◦ C), the upper two-phase ␣– region (1000–1030 ◦ C) and the lower two-phase ␣– region (<1000 ◦ C): (a) 1050 ◦ C, (b) 1000 ◦ C and (c) 950 ◦ C. At processing temperature: (1)  grain boundaries, (2)  phase region, (3) ␣ at  grain boundaries and (4) transformed  phase region. Fig. 5. Extrapolated beta approach curve for IMI834 from 900 to 1050 ◦ C [21].
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respectively, and considerably greater than the evaluated values of 1.2% and 3.8%, respectively. Hence, it appears that the extent of the beta to alpha phase transformation after 5 min holding at 1030 and 1000 ◦ C has, most likely, not reached a steady state. Thus further increases in the alpha phase fraction are expected with increasing time during processing at higher strain values and/or decreasing strain rate. Moreover, it is anticipated that the greatest increase in the phase transformation will occur with decreasing strain rate considering that at the highest value of 1 s−1 deformation to a true strain of 1 is completed within 1 s while at the lowest value of 0.001 s−1 more than 15 min are spent hot working at temperature. For the specimens held at processing temperatures of 975 and 950 ◦ C, the transformation microstructures were observed to consist mostly of a continuous matrix of transformed beta grains having a lamellar alpha structure with a continuous layer of the alpha phase at the prior-beta grain boundaries, as illustrated in Fig. 3c. Within the lamellar alpha microstructure, some “islands” of a transformed beta phase having an alpha structure with a martensitic or Widmanst¨atten appearance were also observed, especially at 975 ◦ C. For these processing temperatures in the lower bound of the alpha–beta field, the beta approach curve, given in Fig. 5, indicates that the matrix of the IMI834 alloy should be dominated by the alpha phase and is consistent with the structural observations noted for the transformed beta phase. Specifically, the morphology of the transformed microstructure within the prior-beta grains is a result of the cooling conditions during processing, and the rate of cooling affects both the thickness and orientation of the alpha plates. A small driving force for transformation, at relatively slow cooling rates, achieves selective growth of the nucleated alpha plates that allows the lamellae to be arranged in packets or colonies; however, with increasing cooling rate, rapid growth of the nucleated alpha gives the multi-oriented appearance of the lamellae in the Widmanst¨atten structure. In this work, when cooling the IMI834 specimens from the solutionizing temperature (1050 ◦ C) to 950 or 975 ◦ C, the relatively slow transformation of beta to alpha, under furnace holding conditions of 5 min at temperature, led to the formation of a microstructure consisting mostly of lamellae having a colony appearance, as indicated in Fig. 3c. This alpha was present before deformation and represented roughly 78% and 90% of the microstructure at 975 and 950 ◦ C, respectively. In addition, longer holding times at the deformation temperature prior to quenching did not increase the lamellar fraction at either 950 or 975 ◦ C, indicating that the microstructure after holding for 5 min had reached steady state. Quenching of the specimen after holding in the furnace at the deformation temperature caused rapid transformation of the remaining beta phase. This led to the formation of Widmanst¨atten alpha that represented roughly 22% and 10% of the microstructure at 975 and 950 ◦ C, respectively. Hence, the beta fraction at the deformation temperature prior to hot working at 975 and 950 ◦ C was determined to be approximately 22% and 10%, respectively. Moreover, the beta fraction values obtained in this work for the lower alpha–beta temperature region were observed to alight on the extrapolated curve of the equilibrium phase data [22] given in Fig. 5 and further support the finding that holding for 5 min at
950 and 975 ◦ C is sufficient for equilibrating the microstructure. The faster equilibration for temperatures in the lower two-phase region may be attributed mainly to a greater undercooling effect at 950 and 975 ◦ C as compared to 1000 and 1030 ◦ C as well as a secondary effect of a decrease in the cooling rate from 140 to 43 ◦ C min−1 with decreasing temperature from 1030 to 950 ◦ C during furnace cooling. Hence, during cooling from the solutionizing stage in the single-phase beta region to the processing temperature, 975 or 950 ◦ C, the beta phase was observed to predominately transform to alpha. This phase transformation of beta to alpha was observed to first occur as a layer of alpha at grain boundary sites followed by nucleation and growth of alpha plates. On account of the slower furnace cooling conditions from the beta solutionizing stage to the processing temperature, the formation of lamellar colonies of thick alpha plates occurred, inevitably due to the reduction in the nucleation sites available as compared to air cooling or water quenching that usually develop martensitic or Widmanst¨atten structures [21]. Therefore, for processing temperatures of 975 and 950 ◦ C, the microstructure before quenching (i.e. at the onset of hot working) consists of a continuous layer of grain boundary alpha around predominately transformed beta grains having a lamellar alpha plate structure. The microstructural observation of martensitic or Widmanst¨atten alpha “islands” is then due to the quenching of the specimens from the processing temperature that gave a rapid phase transformation of the remnant beta regions, which represented 22% and 10% of the microstructure at 975 and 950 ◦ C, respectively. 3.3. Microstructural features associated with forging 3.3.1. Deformation in the single-phase beta region Knowing that the initial microstructure in the single-phase beta region (1050–1125 ◦ C) prior to hot working consisted of beta grains having an average grain size of 304 m (Fig. 4), the effect of deformation processing was examined by compressing IMI834 specimens to true strains of 0.2, 0.6, 0.8 and 1.2. Between 1050 and 1125 ◦ C, the morphology of the beta grains becomes elongated (pancake-shape) with increasing strain in the plane perpendicular to the forging direction, as typified for processing at 1050 ◦ C in Fig. 6a and b. The beta grain boundaries were observed to become increasingly jagged and serrated with increasing strain and, at a true strain of 0.2, the formation of small-recrystallized grains was apparent in the vicinity of the deformed grain boundaries (Fig. 6a). In particular, deformation at 1050 ◦ C with a strain rate of 1 s−1 to a true strain of 0.2 was observed to attain an average recrystallized fraction of 2.3%. With increasing strain, further recrystallization of beta grains was apparent at the prior-beta grain boundaries, which gave a necklace-appearance of fine equiaxed beta grains around the deformed and elongated beta phase (Fig. 6b). Since the nucleation of dynamically recrystallized grains is initiated when a critical dislocation density in the deforming material is reached during hot working, nuclei for dynamic recrystallization (DRX) typically appear at the grain boundaries or other internal crystallographic features such as shear bands or inclusions. Hence, the
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Fig. 6. Microstructure of IMI834 deformed at 1050 ◦ C (single-phase  region) with a strain rate and true strain of (a) 1 s−1 and 0.2, (b) 1 s−1 and 0.8 and (c) 0.01 s−1 and 0.8. At processing temperature: (1) original  grains and (2) recrystallized  grains.
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recrystallization of the beta phase that occurs in the single-phase region (1050–1125 ◦ C) can be predominately related to dynamic processes, as equiaxed small and slightly elongated intermediate grains are apparent at the original beta grain boundaries. This finding is consistent with previous results on hot working of IMI834 that associated the occurrence of recrystallization under isothermal forging conditions to deformation induced dynamic and/or metadynamic recrystallization [13,19,23–25]. Such a structural evolution has also been reported to occur regularly during beta processing of titanium alloys such as Ti–6Al–4V [21,25,26]. Moreover, the incidence of serrated grain boundaries during beta phase deformation has been considered as a special case of DRX, referred to as necklace recrystallization, since this softening process concentrates at the grain boundaries during hot working [26,27]. In addition, it was observed that beta phase working in IMI834 using high strain rates (0.1 and 1 s−1 ) and true strains of up to 1.2 leads to the formation of a mixed grain microstructure consisting of small equiaxed recrystallized grains as well as large deformed prior-beta grains as revealed in Fig. 6b. This microstructural non-uniformity has been related to the incidence of selective recrystallization in the high strain regions along the beta grain boundaries combined with a lower driving force for recrystallization in the interior regions of the beta grains where recovery is most likely dominant [26]. The occurrence of selective recrystallization during beta phase hot working in titanium-based materials has been noted to be particularly problematic for grain size and structure uniformity because the mixed grain microstructure of both large and small grains that is produced can not be eliminated by any heat treatment process [28–30]. During beta phase hot working of IMI834, control of the uniformity in the grain structural characteristics such as size and morphology, which is critical for optimizing secondary processing and final mechanical properties, necessitates prudent selection of the thermomechanical processing parameters. Hence a high degree of dynamic recrystallization during processing is crucial for producing uniformity in both microstructural and mechanical property characteristics. However, as noted in this work, processing conditions at 1050 ◦ C must surpass a strain rate of 1 s−1 and a true strain of 1.2 to achieve just over 50% recrystallization, as illustrated in Fig. 7. Thus for the forging conditions in this work (i.e. strain rates up to 1 s−1 , true strains up to 1.2 and processing temperatures of 950–1125 ◦ C), the elongated beta grains dominated the high temperature deformation microstructure of IMI834 and the amount of recrystallization was, overall, quite limited. These observations are similar to the findings of less than 30% recrystallization in Ti–6Al–4V reported in previous work for processing conditions of 1000–1050 ◦ C and strain rates of 0.05–1 s−1 [31]. In general, this limitation for grain refinement by dynamic recrystallization in titanium alloys has been attributed to the difficulty in establishing a sufficiently high dislocation density. Specifically, the high self-diffusivity of titanium in the beta phase readily enables dislocation mobility and decreases the stored energy through dynamic recovery, which in turn homogenizes the dislocation substructure and reduces the tendency (or driving force) for recrystallization [32].
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Fig. 7. Microstructure of IMI834 deformed at 1050 ◦ C at a strain rate of 1 s−1 to a true strain of 1.2. At processing temperature: (1) original  grains and (2) recrystallized  grains.
In the single-phase beta region (1050–1125 ◦ C), a decrease in the strain rate was observed to result in the growth of the recrystallized beta grains, such that at the lowest strain rate, 0.001 s−1 , the once fully recrystallized (dynamic) fine beta grains (25–40 m depending on the temperature) have grown (120–150 m depending on the temperature) and elongated during the continuing deformation, as depicted in the Fig. 6c. Also, with increasing strain rate the degree of dynamic recrystallization was observed to increase and is consistent with previous findings for thermomechanical processing of Ti–6Al–4V in the single-phase beta region [31]. 3.3.2. Deformation in the upper two-phase alpha–beta region For processing temperatures of 1000 and 1030 ◦ C, the microstructure before forming was similar to that in the singlephase beta region, with the exception that discontinuous alpha platelets were present at the beta grain boundaries during deformation in the upper alpha–beta region. Notwithstanding this structural difference, the effect of deformation was similar to that observed for hot working in the single-phase beta region, as demonstrated in Fig. 8a and b. Specifically, with increasing strain, the beta grains elongated and the discontinuous alpha platelets at the grain boundaries became aligned in the plane perpendicular to the forging direction. Additionally the formation of small recrystallized beta grains were observed, but, as their concentration at prior-beta grain boundaries was lower than that observed during deformation in the single-phase beta region, only limited necklace like structure of the fine equiaxed beta grains was apparent around the deformed beta grains. From the beta phase fractions given in Fig. 5, it was remarked that the microstructure after holding for 5 min at 1000 and 1030 ◦ C had not equilibrated and perhaps further transformation
Fig. 8. Microstructure of IMI834 deformed at 1000 ◦ C (upper two-phase ␣– region) with a strain rate and true strain of (a) 1 s−1 and 0.2, (b) 1 s−1 and 0.8 and (c) 0.001 s−1 and 0.8. At processing temperature: (1) original  grains, (2) recrystallized  grains and (3) ␣ at  grain boundaries.
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of the beta phase to alpha could occur with longer holding time at higher applied strain or lower strain rate conditions. In this work, the extent of the beta to alpha phase transformation with increasing strain at deformation temperatures of 1000 and 1030 ◦ C was observed to be most marked at strain rates of 0.001 and 0.01 s−1 . For example, comparison of IMI834 deformed at 1000 and 0.001 s−1 to a true strain of 0.8, as shown in Fig. 8c indicates that additional holding during hot working resulted in the formation of nearly 40% alpha, either at the beta grain boundaries or within the beta grains, because the structure before forming was not under equilibrium conditions as predicted by the beta approach curve given in Fig. 5. This additional transformation of beta to alpha at a strain rate of 0.01 or 0.001 s−1 produced a change in the microstructural characteristics during deformation at 1000 and 1030 ◦ C. Specifically, at the start of hot working, the predominantly beta grain structure was observed to deform concomitantly with the formation of alpha at the beta grain boundaries and within the prior-beta grain structure, thereby giving a continuous alpha layer and alpha platelets in the grain interior, respectively. Hence, at a strain rate of 0.1–1 s−1 , deformation processing at 1030 and 1000 ◦ C was similar to that in the single-phase beta region since the microstructure consisted of a continuous matrix of beta grains with discontinuous alpha platelets present at the beta grain boundaries as revealed in Fig. 8b. So processing with a strain rate of 0.1–1 s−1 at either 1000 or 1030 ◦ C was observed to result in the elongation of the beta grains and the discontinuous alpha platelets with the formation of recrystallized beta grains. At lower strain rates of 0.01–0.001 s−1 , processing at either 1000 or 1030 ◦ C resulted in the deformation of the beta grains and the alpha platelets, the latter of which formed additionally during deformation at the grain boundaries and within the beta grains, without evidence of beta recrystallization as shown in Fig. 8c. This kinetic effect on phase transformation during deformation is consistent with the previous laboratory work that showed that the fraction of alpha during extrusion of IMI834 was a function
Fig. 10. Microstructure of IMI834 deformed at 950 ◦ C (lower two-phase ␣– region) with a strain rate and true strain of (a) 1 s−1 and 0.2, (b) 1 s−1 and 0.8 and (c) 0.001 s−1 and 0.8. At processing temperature: (1) transformed  grains, (2) remnant  phase and (3) lenticular ␣ phase. Fig. 9. Microstructure of IMI834 deformed at 975 ◦ C with a strain rate 1 s−1 and true strain of 0.8. At processing temperature: (1) transformed  grains, (2) remnant  phase and (3) lenticular ␣ phase.
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of the strain rate and strain [22]. Hence for the upper alpha–beta temperature region, the processing window for dynamic recrystallization of the beta grains in IMI834 is at a strain rate greater than 0.01 s−1 at 1000 ◦ C and 0.001 s−1 at 1030 ◦ C. 3.3.3. Deformation in the lower two-phase alpha–beta region At processing temperatures of 950 and 975 ◦ C, the microstructure equilibrated within 5 min of holding and, before forming, there were remnant beta phase regions in a continuous matrix of transformed beta grains having a lamellar alpha structure with continuous alpha at the prior-beta grain boundaries. With increasing strain, for the remnant beta phase regions at both 975 and 950 ◦ C, no recrystallization was observed as shown in Fig. 9. Nonetheless, the larger “islands” of the beta phase present before forming were observed to elongate in the plane perpendicular to the forging direction with increasing strain. In contrast, for all strain rate conditions, deformation processing at 950 and 975 ◦ C was observed to work the alpha structure within the transformed beta grains, i.e. the elongated lamellar alpha platelets distorted (wavy appearance) increasingly with increasing strain as depicted in the Fig. 10a and b. These prior-beta phase regions having a martensitic or Widmanst¨atten structure were easily differentiable from the continuous lamellar alpha microstructure. Particularly, bending, distortion, shearing and break up of the lamellar structure was observed to occur to an increasing extent as the strain increased and the processing temperature decreased. These deformation characteristics of the lamellar structure are consistent with previous work on Ti–6Al–4V [31,33]. With increasing strain rate, deformation processing at 950 and 975 ◦ C was observed to work, to a greater degree, the alpha structure within the transformed beta grains. Similar to the effect of increasing strain, with increasing strain rate, the elongated lamellar alpha platelets distorted increasingly and break up of the lamellar structure was apparent. Weiss et al. have attributed lamellar alpha phase separation into shorter segments to two possible mechanisms that can occur during thermomechanical processing [33]. In their work it was shown that hot working led to the development of both low and high angle sub-boundary formation across the alpha plates. In the present work, separation of the alpha lamellae was observed to occur for the specimens processed at 950 and/or 975 ◦ C at strain rate conditions of 0.001 s−1 or greater, as illustrated in Fig. 10c. In contrast to this evolution in the lamellar alpha microstructure, the beta phase regions simply elongated with increasing strain and strain rate for deformation at either 975 or 950 ◦ C and no evidence of dynamic recrystallization was observed (Fig. 9) using field emission gun scanning electron microscopy and electron backscattered diffraction techniques [17].
and strain rate on structural evolution in IMI834 during hot working in the single-phase beta field and the two-phase alpha–beta region, it is apparent that the processing window for beta phase recrystallization is restricted to thermomechanical conditions giving limited formation of discontinuous alpha lamellae at the beta grain boundaries. Specifically, it was explicitly demonstrated that even when the alpha fraction was constrained to a discontinuous thin film at the beta grain boundaries the extent of dynamic recrystallization was lower than that observed for processing in the single-phase beta field. In view of the mechanisms for deformation-induced recrystallization, the steady state flow stress conditions can be considered to evaluate the influence of the processing conditions on the grain size [34–36]. For the beta processing conditions in this work, full recrystallization was not attained and the flow curves continued to exhibit a decrease in the stress with increasing strain. Nonetheless, it was noticeably established that the grain size of the recrystallized beta phase depends on both the processing temperature and strain rate, with higher temperatures and lower strain rates giving an increase in the average grain size of the recrystallized beta phase, as shown in Figs. 11 and 12, respectively. Also, the overall effect of strain rate appears to be greater than that of temperature and corroborates with the results of the dependence of prior-beta grain size on deformation temperature given in Fig. 4. Previous work has indicated that for temperatures that are close to, or exceeding the beta transformation temperature, where the volume fraction of alpha is small, the constraint on grain growth during recovery or recrystallization is minor and a strong dependence of grain size on strain rate exists [13,35]. However, the effect of temperature appears to be increasingly marked with decreasing strain rate from 1 to 0.001 s−1 , which inevitably is related to a kinetic effect during the longer holding period at temperature. Hence, high strain rates (greater than 1 s−1 ) and temperatures just above the beta transus offer an optimal potential for recrystallization of the beta phase from an initial grain size of over 300 to less than 30 m.
4. Discussion In this work, the thermomechanical processing of IMI834 over a wide range of testing conditions enabled evaluation of the effect of thermomechanical parameters on the microstructural characteristics. From the dependence of the alpha and beta phase fractions on processing temperature, as well as the role of strain
Fig. 11. Effect of processing temperature on the recrystallized beta grain size for various strain rate conditions and a true strain of 0.8.
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Fig. 14. Shear band formation in IMI834 through the aligning of adjacent alpha lamellae during deformation processing at a temperature of 950 ◦ C and strain rate of 1 s−1 . Fig. 12. Effect of strain rate on the recrystallized beta grain size for various processing temperatures and a true strain of 0.8.
The observed increase in the recrystallized beta grain size with increasing temperature and decreasing strain rate (Figs. 11 and 12), i.e. with decreasing Zener–Hollomon parameter (Z), was correlated (Fig. 13) and the linear dependence that was found could be fitted to a single equation of the form: Drex = kZ−0.19
(1)
where Drex is the recrystallized grain size of the beta phase, Z the Zener–Hollomon parameter (˙ε exp(Q/RT )), and k is a constant. From Fig. 13, the refinement in the beta grain size with increasing Z stems from the greater driving force for recrystallization (a higher dislocation density). The R2 -value of 0.98 indicates that overall there is a good approximation of the model with the experimental data in Fig. 13. It is also noteworthy that although the fraction recrystallized, for certain Z values, was in the early stages of dynamic recrystallization, a reasonably good
Fig. 13. Variation in the recrystallized grain size of the beta phase in IMI834 with the Zener–Hollomon parameter for true strains of 0.6 (filled markers) and 0.8 (open markers).
correlation between the Z and the average beta grain size was obtained. This may be related to the dynamically recrystallized beta grain size having reached steady state conditions beyond a true strain value of 0.6, as illustrated by the similarity in the beta grain size data for the two strain conditions plotted in Fig. 13. The observed linearity between the dynamically recrystallized beta grain size and the Zener–Hollomon parameter has also been reported recently for other titanium alloys [37,38]. Moreover, the exponent in the power law relationship of Eq. (1), i.e. 0.19, is in good agreement with value of 0.17 obtained for beta phase hot working of Ti–6Al–4V by Tamirisakandala et al. [38]. As compared to deformation processing above 1000 ◦ C, for which recrystallization of the beta phase was observed, the considerable flow softening apparent at 950 and 975 ◦ C may be attributed to dislocation processes that were observed in this work to result primarily in platelet distortion as well as crystallographic changes [17]. In particular, slip along alpha lamellae during deformation was physically manifested in the micro- and macro-structural observations of adjacent lamellae aligning in the shear bands and shear band formation as shown in Fig. 14. These findings are supported by the recent work of Semiatin and Biehler [39] that related dislocation processes, platelet distortion and crystallographic texture as contributing factors to flow softening mechanisms. In particular, Semiatin and Biehler have indicated that during deformation of alpha–beta titanium alloys in the two-phase field, dislocation pileups/slip transfer processes may be instigated at alpha–beta interfaces due to the dissimilarity in the sub-grain formation characteristics of the two-phases [40]. In their work, the alpha–beta interfaces have been hypothesized as being analogous to grain boundaries in single-phase materials and, on account of the limited number of independent slip systems in the hexagonal crystal structure of alpha titanium, a strong grain boundary or Hall–Petch strengthening dependence on the alpha lath/platelet thickness was determined. Regarding the relevance of slip transmission across alpha–beta interface in relation to Hall–Petch strengthening in titanium alloys, previous work has suggested that during uniaxial compression of Ti–6Al–4V, with a colony microstructure, aligning of shear bands in adjacent lamellae occurs with a resulting
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localization in the flow characteristics [39]. This development of microscopic shear bands within the alpha titanium lamellae indicates that sub-grain formation in the lamellar microstructure is difficult and it has been suggested that the principal mechanism responsible for the considerable flow softening behavior in Ti–6Al–4V can be attributed to slip transmission coupled with a loss of alpha–beta interface strengthening [41]. For IMI834, the deformation processing conditions in the two-phase alpha–beta region that exhibited considerable flow softening behavior (i.e. at 950 and 975 ◦ C) may then be related to similar phenomena as supported by the platelet distortion and the presence of shear band formation (Fig. 14).
single linear relationship between a strain rate range 0.001 and 1 s−1 for experimental conditions where the microstructure was predominately beta. (6) Processing at temperatures in the lower two-phase field, 975 and 950 ◦ C was observed to work the lamellar alpha structure within the transformed beta grains and bending, distortion, shearing and break up of the lamellar structure was observed to occur to an increasing extent as the strain and strain rate increased. At both 975 and 950 ◦ C no recrystallization was observed in the remnant (interlamellar and islands) beta phase regions. Acknowledgments
5. Conclusions Manufacturing of titanium alloys is a multi-step process that involves consecutive heating and deformation at beta and alpha–beta phase temperatures. The evolution of component characteristics (shape, structure and properties) that develop during deformation processing require a methodical understanding of the role of various parameters on the hot working behavior. In this work isothermal hot working of IMI834 was studied in the temperature range 950–1125 ◦ C at strain rates of 0.001–1 s−1 and true strains up to 1.2 with the follow observations: (1) Deformation processing of IMI834 indicated that at temperatures in the single-phase beta region, the flow stress behavior exhibited strain hardening and minor flow softening, while in the two-phase alpha–beta region considerable flow softening occurred. (2) Processing in the single-phase beta region (1050–1125 ◦ C, 0.001–1 s−1 ) resulted in the serration of beta grain boundaries and the formation of small dynamically recrystallized grains (by a true strain of 0.2) which concentrated at the prior-beta grain boundaries and gave a necklace-appearance of fine equiaxed beta grains around the deformed and elongated beta phase (true strains greater than 0.6). (3) Processing at 1000 and 1030 ◦ C involved working of a predominately beta microstructure with alpha at the beta grain boundaries, especially for strain rates greater than 0.01 s−1 , since the microstructure was not equilibrated. Under such conditions, hot working was observed to recrystallize the beta phase albeit in lower concentration and the processing window for dynamic recrystallization of the beta grains (beta phase working) was determined to be a strain rate greater than 0.01 s−1 at 1000 ◦ C and 0.001 s−1 at 1030 ◦ C. (4) For conditions giving a dominant beta phase microstructure, decreasing the strain rate from 1 to 0.001 s−1 was observed to increase the size of the recrystallized grains from 25–40 to 120–150 m, respectively. Also, the average size of the dynamically recrystallized beta grains was observed to increase from 25–120 to 40–150 m with increasing processing temperature from 1000 to 1125 ◦ C, respectively, depending on the strain rate condition. (5) Variation in the dynamically recrystallized beta grain size with the Zener–Hollomon parameter was observed to give a
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