SUB-SOLVUS RECRYSTALLIZATION MECHANISMS IN UDIMETB ALLOY 720LI

B. Lindsley, X. Pierron

Special Metals Corporation Middle Settlement Road New Hartford, NY, 13413

Abstract

Sub-solvus recrystallization mechanisms were studied in a nickel base superalloy and the initial microstructure was found to influence the operating mechanism. Processing maps were developed to describe recrystallization for 3 different initial microstructures. In the cast-wrought condition, recrystallization takes place by nucleation and growth along grain boundaries between y particles in a similar fashion to single phase alloys. For large grain, small y structures, growth of recrystallized grains was found to occur primarily by boundary looping of they, leaving the y misoriented with the new matrix. Electron backscatter diffraction was used to measure the orientation of the y matrix and the y. After a heat treatment that resulted in static recrystallization, three y orientations were found in the recrystallized grain, includingy with a twin relationship to the matrix. A mechanism is proposed to describe the formation of the cast-wrought structure consisting of fine grains and y misoriented with both they matrix and neighboring f.

Introduction

as 50% to increase high temperature mechanical properties. The presence of the coherent ordered y precipitates in the microstructure at deformation temperatures results in complex recrystallization micro-mechanisms which need to be understood if one wants to optimize the thermo- mechanical processes necessary to manufacture forged components. Second phase particles act as barriers to grain boundary motion. This effect, known as Zener pinning, plays a major role in retarding recrystallization, and may also affect grain nucleation. The volume fraction second phase (FJ and particle size (r) are critical variables for grain boundary pinning, as they determine interparticle spacing. It is generally accepted that, if FJr > 0.2 pm-, recrystallization will be inhibited( 1).

UDIMET@* alloy 720LI (720) is a challenging alloy as the y solvus temperature is quite high for a cast and wrought superalloy. Above the solvus temperature, grain growth is rapid and the workability of the alloy decreases dramatically; therefore the material is worked sub-solvus. The volume fraction of y at working temperatures is approximately 20%, so according to the above rule, recrystallization will be retarded when the particle radius is less than lpm. Particle radii < 0.5pm are typically

Recrystallization in multiphase microstructures is an important mechanism for conversion of nickel base superalloy ingot to wrought product. Typical cast-wrought nickel base superalloys contain a y volume fraction as high

* UDIMET is a registered trademark of Special Metals Corporation.

Superalloys 2000 Edited by T.M. Pollock, R.D. Kissinger, R.R. Bowman, K.A. Green, M. McLean, S. Olson, aad J.J. Schirra

ThfS (The Minerals, Metals &Materials Society), 2ooO

59

found in air cooled and furnace cooled specimens after supersolvus solutioning. Since grain boundary motion is suppressed by particle pinning, additional mechanisms must exist to enable boundary motion through a sub-solvus material.

The interaction of the grain boundary and the second phase during recrystallization in superalloys has been extensively studied(2-4). Several processes have been reported for Ni based superalloys: 1. Dissolution of the y precipitates after contact with the

boundary and coherent re-precipitation of they either at or behind the boundary;

2. Grain boundary pinning by the coherentywhich have coarsened at the grain boundary, and subsequent nucleation of new grains to form a necklace structure about the original grain boundary;

3. Grain boundary cutting through a y precipitate and reorienting the precipitate to the orientation of the recrystallized grain (rarely seen).

The most common mechanism for Ni based superalloys was found to be dissolution and re-precipitation of the y.

Little work has been done on dynamic or static subsolvus recrystallization with a y diameter 2 0.5pm. The effect of y size and interparticle spacing has not been addressed within a given alloy system, nor has the associated mechanistic changes in recrystallization. The interparticle spacing also affects the competition between recovery and recrystallization. Recovery can be quite rapid since the working temperature is approximately 0.9T,, and they pm grain boundaries and inhibit recrystallization. The current work attempts to expand the current understanding of subsolvus recrystallization in nickel base superalloys by examining the different recrystallization mechanisms within a single alloy system containing varying microstructures deformed under a range of conditions.

Experimental Approach

Three initial microstructures of UDIMET alloy 720LI were compression tested and the recrystallization response was evaluated using light optical microscopy, scanning electron microscopy and electron backscatter diffraction. The composition of the alloy used in this study is given in Table I. The alloy was heat treated to three different starting microstructures to evaluate the effect ofy size and distribution on subsolvus recrystallization. The first microstructure was typical cast-wrought material rolled to 28.6mm (1 I/S) bar, containing large y and small grains (material A). For the other two microstructures, the as- rolled material was solutioned and furnace cooled to form a large grain, fine y microstructure (material B). This material was then aged below the solvus to coarsen the y to an intermediate size (material C). The y sizes given in the paper were evaluated using image analysis after a heat

treatment at 1110C. Cylindrical compression samples were made from each microstructure and the samples were deformed 30% at several temperatures (1066C (1950F), lllOC (2030F) and 1132C (2070F)) and three strain rates (1, 0.1 and 0.01 .sec-). The compression test resulted in a barrel shaped specimen with non-uniform deformation from the bottom to the center of the sample. A wedge shaped dead zone was found in each end of all samples where die lock occurred. Samples were then air cooled or oil quenched to room temperature and evaluated both after deformation and after subsequent heat treatments to evaluate dynamic and static mechanisms. The heat treatment temperature used for all samples was 1132V. Samples were etched with either modified Kallings (1 OOml methanol, lOOm1 HCl, 50g CuC&) or Chromate (460ml HsP04, 25ml H$Od, 50g Cr0.J. The resulting microstructures were analyzed using light optical and scanning electron microscopy.

Table I. Al10 corn osition in wei ht ercent.

71

The orientations of the y precipitates and the surrounding matrix were investigated using the electron backscattered diffraction technique (EBSD) in the SEM. This technique allows the orientation of the observed crystals to be described by indexing the diffraction patterns generated by the backscattered electrons on a detector. By mapping the orientations as a function of position on the sample, an orientation map or OM can be generated and analyzed. The data was collected by TSL Inc. on specimens which were first polished down to 0.05 micron using a silica solution. A backscattered electron image (BEI) of the scanned area was then taken in order to associate each phase with its orientation. The orientation data were analyzed using TSL software and various representations of the orientation data were constructed. The misorientation between two adjacent areas can be visualized by associating different colors to the two misoriented regions. In addition, regions of same orientation were highlighted in the orientation maps. Those orientation images were then superimposed with the backscattered electron images of the same area.

Results

Microstructures

Material A (as-rolled material) contained large y (2.8um) and small grains (15um). EBSD of this structure showed that the y are misoriented with the surrounding grains, Figure 1. An SEM image and an EBSD image of the same area are shown in Figure 1, and the two are overlayed in

60

l

Figux 1. SEM image and EBSD orientation map of material A. (a) and (c) arc SEM and orientation maps of the same region rcspcctivcly. (b) is a superposition of image (a) and (c). The arrows indicate some y precipitates misoriented with the surround- ing grains. (A color version of rhis figure appears on page 841.)

Fig. I b. Each color represents a misorientation of greater than 15 with its neighbors, and the different colors of the y and surrounding y grains reveal their different orientations. The solution and furnace cooled samples (materials B and C) had a large grain size after the supersolvus solution treatment (3 mm). Material B had an average y size of 0.7pm and material C had an average y size of 0.95pm. EBSD of the as-heat treated structure in material C showed that the y was coherent with the matrix, as expected in a non-deformed, as heat treated microstructure.

Processing Maps

Processing maps of strain rate vs. temperature were generated for the three microstructures, Figure 2. The process maps were created to evaluate the recrystallization response of the three microstructures to the testing variables that control recrystallization, namely strain rate and temperature (constant strain used for all tests). The result of the thermomechanical treatments to the material can generally be described by three responses: no recrystallization; a necklace of new grains at the original grain boundaries; or as some percentage of dynamic recrystallization. [Values for percent recrystallization are average values across the entire sample. The effect of changing strain and strain rates along the radius of the sample was therefore not addressed.] For material A, the processing map is separated into these three responses, although necklacing is the beginning of dynamic recrystallization for this microstructure. No dynamic recrystallization (DRX) was found at 1066C or at low strain rates at 1 I 10C. Some recrystallization was found in the form of a necklace as the temperature and strain rate was increased, Figure 3a. The percent recrystallization is also given in Figure 2a. 100% recrystallization was found

at I 132C and strain rates 2 0.1 set, Figure 3b. It should be noted that the dead zone of the compression sample was not included in the percent recrystallization values. The recrystallization of cast-wrought 720 was found to occur by new grains forming at the prior grain boundaries followed by grain growth until the deformed grains are consumed. This is a classic recrystallization mechanism in single phase alloys. Since the grain size of material A is on the order of the y spacing, little to no y exists within each grain, so on this scale the material behaves similarly to a single phase alloy.

The processing map for material B is given in Figure 2b. It was found that strain rate was a more important variable for dynamic recrystallization than temperature for this microstructure (as compared to material A). Again, the three responses were found in the material. No recrystallization was found at low temperature and low strain rate. At higher temperatures and strain rates I 0. I set -I, a necklace of fine grains 5 - 10 pm wide formed at the original grain boundaries. At a strain rate of 1 set-, dynamic recrystallization was found for all three temperatures. The recrystallization was primarily within the grains and no longer associated with the boundaries. Preferential recrystallization was also found at nitrides within the matrix.

The processing map for material C is given in Figure 2c. Again, no recrystallization was found at low temperatures and strain rates. A necklace of recrystallized grains formed at the original grain boundaries at intermediate strain rates and temperatures. The band of recrystallized grains at the old boundary expanded in width to approximately 25 to 50pm at 1 l32*C and 0.01 set-. Relative to material B, the region of the map containing dynamic recrystallization shifted toward higher temperatures. No recrystallization

Microstructure A

None i Necklace ~ 5yo DRX ; 100% DRX

I I I I I I b

1066 1110 1132

Temperature (C)

4 Microstructure B

h 0

1 JO-15%DRX 70%DRX 60% DRX

% (I OO%Rx) (lOO%Rx) (7O%RX)

----___ ----- .___.__ a, -.._ ----.____ --._ ---___ % 0.1 -- None ...L._ Necklacd-.---~~~~ d (lOO%Rx) c

(4.pO%RX) l\ (lo-15%RX) .A

s l\ L.

m 0.01 -- None -*... Necklace (5%Rx) (Necklace)

I I I I b

1066 1110 1132

Temperature (C)

Microstructure C

1

5 t

None 40% DRX 0 $4

( 1 OO%Rx) ( I OO%Rx) %. -..

1066 1110 1132

Temperature (C) Figure 2. Processing maps for materials A, B, and C. The amount of dynamic recrystallization (DRX) is given for each deformation condition. Total recrystallization after an 1132C heat treatment for 5 hours is given in parenthesis for B and C.

Figure 3. Recrvstallized microstructures in material A: (a) necklace struc;ure (b) fully recrystallized. modified Kallings.

Etched in

was found at 1066C and a strain rate of 1 set-I, as compared to 1 O-1 5% DRX for material B. Increased recrystallization at 1132C and strain rates I 0.1 set was also found. Examples of dynamic recrystallization are given in Figure 4, where Figure 4a shows a finer grain size along the original grain boundary and larger grains towards the center of the grains, and Figure 4b shows the recrystallization front, with the newly recrystallized grain on the right. No change in they size was found due to the dynamic recrystallization, although some evidence of the y inhibiting boundary motion was seen.

62

(b) Figure 4. Dynamically recrystallized grains in material C. (a) Fine grains near the original grain boundary with coarser grains near the center of the original grain. (b) SEM image of a recrystallized grain boundary. Arrows indicate grain growth direction prior to quenching. Letters mark various stages of boundary looping around y precipitates. Etched in modified Kallings.

Effect of Heat Treatment

Samples from material B & C were heat treated at I 132C for 5 hours and it was found that select samples statically recrystallized. The amount of recrystallization after the heat treatment is given in parentheses on the process maps (Figures 2b & c) for each sample. For both materials,

samples that were tested at temperatures I 11 10C and strain rates 2 0.1 see- were found to be 100% recrystallized after the heat treatment. Those samples tested at lower strain rates and higher temperatures tended to have less static recrystallization. Material B showed very little static recrystallization under these test parameters, whereas increased static recrystallization was observed for material C. The recrystallized grain size changed from the center to the edge of the prior grains. Near the prior grain boundary, the grain size ranged from 20 to 50 pm while in the center of the prior grains, grain sizes up to 5OOpm were found. Similar results were found during dynamic recrystallization for material C.

Figure 5 shows a boundary between a recrystallized grain (lower right) and an original grain (upper left). Pinning of the grain boundary by y particles is evident by the looping

Figure 5. Statically recrystallized grain in material B after heat treatment. Arrows indicate grain growth direction orior to auenchina. Etched in modified Kallings.

Figure 5. Statically recrystallized grain in material B after heat treatment. Arrows indicate grain growth direction prior to quenching. Etched in modified Kallings.

Figure 6. Necklace grain structure after heat treatment in material B. Larger y (white) lie primarily on the original grain boundary. Electrolytically etched in Chromate.

63

grain. (A color versisrr of this figure uppears an page 84 I.)

W Figure 8. SEM image and EBSD orientation map of statically recrystallized material C. (a) and (c) are SEM and orientation maps olthc same region respectively. (bj is a superposition of image (a) and (c). The red region is the original, unrccrystallizcd glIi11. (A c-olor ~Y~IY~IIII of this figure q?pears OYL page 841.)

of the boundary about the particles. Significant coarsening of the y and an associated decrease in the number of y particles in recrystallized grains was observed relative to the y in both the unrecrystallized regions and in dynamically recrystallized grains. The y size increased from less than lprn to 2-3pm. A necklace structure that formed after heat treatment of a low strain rate / high temperature sample is shown in Figure 6. Large y reside along the original grain boundary and the new necklace grains are primarily devoid of y.

Electron Backscatter Diffraction

EBSD was used to analyze the orientation relationship between the various phases and grains in order to provide insight on the dynamic and static recrystallization mechanisms that were observed in materials B & C.

An orientation map of a dynamically recrystallized region in material C is presented in Figure 7. Recrystallization

occurred after compression testing at 1 132OC, with a strain rate of O.lHz. The region in Figure 7 contains two recrystallized grains that were growing toward the top of the image when the sample was quenched. One can note that the y size and distribution in the recrystallized grains and in the unrecrystallized grain at the top of the Figure are similar. The orientation map (OM) of the unrecrystallized grain shows that all the y precipitates have the same orientation as the matrix, as no change in color can be seen in the OM. In this Figure the misorientation treshold for color change was set at a 10 misorientation angle. In the recrystallized grains, the y is evidently misoriented with the new matrix. By highlighting all the regions with the same orientation as the unrecrystallized grain in red, it was found that all of the misoriented y in the OM had the same orientation as the original top grain (Fig 7~). A comparison of the microstructure and the EBSD results revealed that less than 5% of they did not show up on the misorientation map. These y either have the same

64

orientation as the matrix in the new grains or were too small to be detected during the EBSD scan.

In the same material, material C, the orientation relationship between the precipitates and the matrix was analyzed in the recrystallized region after deformation at 1 I 10C with a lower strain rate of O.OlHz. After deformation, no recrystallized regions could be observed in this sample. The deformed material was then heat treated for 5 hours at 1 132C and recrystallized statically. The resulting microstructure can be seen in Figure 8. An unrecrystallized grain can be seen in the upper left part of the figure, and the orientation image shows that the y phase in this grain is oriented cube on cube with the surrounding matrix. By comparing the BEI image and the orientation image, it can be seen in the recrystallized grain that the majority of they precipitates are misoriented with the surrounding matrix, as only a few precipitates are visible on the BEI image but not on the OM. By highlighting the phases with the same orientation, it was found that in the misoriented y in the recrystallized grain are of two kinds. The first kind (red in the OM) have the same orientation as their parent grain, the unrecrystallized top grain. This is similar to the precipitates found in the dynamically recrystallized grain. The second kind (blue in the OM) are not oriented with the original grain, but rather have a special orientation relationship with the recrystallized grain. A region containing the blue precipitates and the recrystallized matrix was selected and pole figures were generated with only these orientations, Figure 9. It follows from close examination of the pole figures that the blue? precipitates have a twin relationship with the matrix, and are rotated 60 around a [l I l] matrix

TD

KU

\ll 3 Common [llO]

ommon [l I I]

RD

Figure 9. Pole figures for the recrystallized grain and the blue y in Figure k. (Tlzisfiguw uppears in color CM page 842.)

direction, The blue y and the matrix share one common (I I 1) plane and three common [ 1 IO] directions. Annealing twins were also present in the recrystallized grain and are visible through misorientation boundaries in the white region of Figure 8c. The four annealing twins share a common (Ill) twin plane with the matrix that is different to the (111) twin plane between the matrix than the blue y.

Discussion

It was found that the y matrix and the y precipitates retained no residual orientation relationship in the as-rolled (cast-wrought) condition. Further, they were misoriented with neighboring y. [It should be noted that they referred to here only include those that precipitated at the working temperatures. y that precipitate during the final strengthening heat treatment will be a fine, cube on cube oriented precipitate.] However, when 720 is processed from ingot, the material is first homogenized, and after cooling, a large grain, fine y structure is produced, similar to material B, where all of they have the same orientation within a large grain. It has been well established that such precipitates effectively inhibit grain boundary motion. Previous recrystallization mechanisms that have been reported in subsolvus superalloys include solutioning of the fine y ahead of the boundary and coherent re- precipitation in or behind the boundary, the boundary moving through the y precipitate, or the pinning of the original boundary by y and the formation of a necklace structure(2-4). All of these mechanisms may eventually lead to a random grain orientation, however, they tend to maintain a common orientation within each grain. The questions of what recrystallization mechanism(s) occurs in the subsolvus condition and how they eventually change from a cube on cube orientation with the matrix to being misoriented with both the matrix and neighboring y have not yet been answered.

The first step in answering these questions was to develop processing maps to understand recrystallization conditions in 720 with different microstructures. In the as-rolled condition, material A, the grain size is on the order of the y spacing. The y tend to reside at the grain boundary triple points and inhibit further grain growth, but few y reside within the grains or at the grain boundaries between triple points. Recrystallization via necklacing at the boundaries was not significantly inhibited by they. This was found to be especially true at higher temperatures where the y spacing increased due to solutioning of the y. Further grain growth after necklacing leads to full recrystallization, which was also accelerated at high temperatures. Hence, once the material has been worked enough so that the grain size is on the order of they size, recrystallization occurs in a similar manner to a single phase material.

Material B has a tine y size, so for all test temperatures below solvus, the y spacing remains small. For this reason, temperature is a less significant recrystallization variable in this material, and strain rate appears to be the dominant recrystallization variable. It is interesting to note that dynamic recrystallization was found at a temperature of 1066C and a strain rate of 1 set- for this condition, while none was found in the other materials. This is likely related to how the material accommodates the deformation, but is outside the scope of this paper and will not be addressed further.

The processing map for material C appears to be a hybrid between materials A and B. A large shift in the dymanic recrystallization regime was noted between materials B and C, with increased recrystallization occurring at hotter temperatures and lower strain rates for material C. The coarser y in material C equate to a larger interparticle spacing, which apparently encourages dynamic recrystallization at higher temperatures. However, the size of this shift was a bit surprising given that the starting microstructures are not vastly different for the two materials.

The heat treatment given to materials B and C resulted in 100% recrystallization for those samples tested at lower temperatures (I1 110C) and higher strain rates (2 0.1 set- ). These test conditions lead to an elevated amount of stored work in the material, which upon heat treatment, resulted in full static recrystallization. Samples tested at higher temperatures and especially those tested at lower strain rates exhibited lower amounts of static recrystallization. This finding was attributed to recovery in the material. Materials that have high stacking fault energies, such as nickel, undergo recovery when deformed at high temperatures(5). It has also been found that in low stacking fault materials, such as stainless steel, if recrystallization is inhibited by a second phase, then at sufficiently high temperatures, recovery can occur( 1). When the materials are deformed at lower strain rates, more time is available for dynamic recovery and cell wall formation. Reduced levels of static recrystallization in material B as compared to C under these conditions may be a result of increased boundary pinning by the smaller inter-particle spacing.

For recrystallization to occur in materials B and C, the grain boundary must be able to move through an area containing y. As described previously, the y precipitates inhibit boundary motion. The mechanism by which dynamic recrystallization occurred in these materials was grain boundary looping of the y. The stages of looping can be seen in Figure 4b. The boundary initially contacts a y (A) and is halted at the ylrecrystallized grain interface. As the rest of the boundary continues to move, the boundary begins to bow around the particle (B). Finally,

the boundary loops around the particle (C) and continues on leaving y with their original orientation in its wake. The orientation map in Figure 7 clearly shows they retain their original orientation after recrystallization. Both the microstructural and EBSD analysis support the boundary looping mechanism. It should be noted that grain boundary looping is not a new mechanism and often occurs for incoherent particles( 1). However, looping of coherent particles has rarely been reported. Looping of the coherent NbC phase in y-Fe was found by Jones and Ralph(6) but has not been widely reported for nickel base superalloys.

One possible reason that looping is rarely seen for coherent, cube on cube oriented precipitates is that the coherent particles are more effective at inhibiting grain boundary motion than incoherent particles. It has estimated(2) that coherent precipitates are up to 4 times more effective in resisting grain boundary motion than incoherent particles, due to the additional interfacial energy required to transform the particle/matrix interface from coherent to incoherent. Inter-particle spacing is also an important variable for this mechanism. The stress of the bowing boundary is inversely related to the radius of curvature of the boundary. If a very fine, uniform distribution of particles exists, the radius of curvature decreases between particles on the boundary and the resistance to boundary movement increases. As the particle spacing increases, the energy required for boundary movement decreases. It should be noted that once the first recrystallization front has passed through a region containing y and the particles become misoriented with the matrix, further recrystallization should be easier in this region. Any future boundary would encounter incoherent particles and would no longer need to supply the additional energy to transform the coherent interface to an incoherent interface.

In the case of static recrystallization, some looping ofy is still observed. The red y in Figure 8 indicate some of the y retain their original orientation. In addition, boundary looping can be seen in Figure 5. However, additional processes occur during the heat treatment. Coarsening of the y occurs during the heat treatment, and is at least partially a result of accelerated growth on the grain boundary. The boundary is a fast diffusion path for Al and Ti hence larger y in the boundary will coarsen rapidly at the expense of smaller y that come in contact with the boundary. Some matrix diffusion ahead of the boundary can also occur. The coarsening of y in the boundary can be seen in Figure 5. In fact, in cases of low driving force, grain boundary movement may not occur without this coarsening mechanism. Coarsening of the y at the boundary reduces the number of precipitates and barriers to boundary motion. The y continue to coarsen at high temperature after the boundary has swept past (Oswald ripening) and the coarsening kinetics may be accelerated

66

due to the higher interfacial energy of the misoriented y with the matrix. Regardless of whether the accelerated coarsening is due to the boundary interaction, the misoriented y, or both, the coarsening is more rapid in the recrystallized region. The initial 48 hour heat treatment at 1132C used to coarsen the structure in material B to material C resulted in a change in y size from 0.7pm to 0.95pm. After the 5 hours at 1132C for the static recrystallization, y 2 to 3 pm in diameter were observed in the recrystallized regions.

Finding y with a twin relationship to the y in the heat treated samples was surprising. To the authors knowledge, y with a twin relationship to the y have not been previously found in nickel based superalloys. The mechanism for their formation is unknown. One potential mechanism for their formation may be similar to the formation of annealing twins. As the boundary moves and solutions some of the smallery, new y may precipitate in the boundary, which has been documented by other investigators(2-4). It may be energetically favorable for the y to grow with a twin relationship to the new grain in order to minimize the total free energy of the resultant boundaries between the y and both the recrystallized grain and the original grain, In other words, if by forming the twin, they lowers the interfacial energy between it and the old grain enough to offset the additional energy associated with the twinned interface, then this mechanism may be possible. For this mechanism to be viable, a special (low energy) orientation relationship should exist between the twinned y and the old matrix which would minimize the interfacial energy. However, no special relationship was found between these y and the original matrix.

Another possible mechanism for the twinned y could involve the dissolution of the misoriented y left behind from the looping process and the re-precipitation ofy with some orientation to the matrix. The y with the twinned relationship represent a lower interfacial energy precipitate, and while not being the lowest energy precipitate (cube on cube orientation), may be metastable. Nucleation of such a particle may require a defect in the matrix, such as a pre-existing annealing twin. Both of the above mechanisms also allow for the formation of y oriented cube on cube with the matrix. This is an important point since y oriented with the matrix were also found.

The original question of how a structure withy that are misoriented to both the surrounding matrix and neighboring y can now be addressed. It was found that, for dynamic recrystallization, one passage of a grain boundary through the material lead toy misoriented with the matrix, but oriented with the surrounding y. Additional dynamic recrystallization will not change the

orientation of the y. However, after heat treatment and static recrystallization, 3 y orientations were found. If another recrystallization front was to pass through this region via the looping mechanism, three populations of misoriented y would exist, and with additional holds at temperature, two more orientations ofy (cube on cube and twinned with the new matrix) could exist. It can be seen that repeated deformation and heat treatments could lead to condition where the y are now misoriented with neighboring y. In addition, as the material undergoes many repetitions of deformation and heat treatment, they will coarsen to a point at which recrystallization can occur between the y. The process described above represents a possible mechanism to go from a large grain, cube on cube coherent y structure to a tine grain, misoriented y structure.

Conclusions

Analysis of the recrystallization mechanisms has shown how the starting microstructure influences the operating mechanism for a given set of processing conditions. In the cast-wrought condition, recrystallization occurs by nucleation and growth along grain boundaries between y particles in a similar fashion to single phase alloys. For a large grain, fine y structure, growth of recrystallized grains was found to occur primarily by boundary looping of the y, leaving the y misoriented with the new matrix. After a heat treatment that resulted in static recrystallization, three y orientations were found in the recrystallized grain, including y with a twin relationship to the matrix. A mechanism is proposed to describe the transformation of a large grain, fine y structure containing y with a cube on cube orientation to the cast-wrought structure consisting of fine grains and y misoriented with both the y matrix and neighboring y.

Acknowledgements

The authors would like to thank B. Antolovich for his insight during technical discussions and K. Sellitti who conducted SEM analysis and documentation.

References

1. F. J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena (Pergamon, Elsevier Science Inc., Tarrytown, NY, 1996).

2. R. D. Doherty, Metal Science, 16 (1982) l-13.

3. V.. Randle, B. Ralph, Acta Metallurgica, 34 (1986) 891- 898.

4. A. Porter, B. Ralph, Journal of Materials Science, 16 (1981) 707-713.

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5. R. E. Reed-Hill, Physical Metallurgy Principles (PWS- Kent, Wadsworth Inc., Boston, MA, 1973).

6. A. R. Jones, B. Ralph, Acta Metallurgica, 23 (1975) 355-363.

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Table of Contents-------------------------Next PagePrevious Page-------------------------Next HitPrevious HitSearch ResultsNew Search-------------------------Keynote AddressSuperalloys: The Utility Gas Turbine Perspective

Ingot, Powder and Deformation Processing Characterization of Freckles in a High Strength Wrought Nickel SuperalloySimulation of Intrinsic Inclusion Motion and Dissolution during the Vacuum Arc Remelting of Nickel Based SuperalloysPredicting Grain Size Evolution of UDIMET(r) Alloy 718 during the "Cogging" Process through Use of Numerical AnalysisControl of Grain Size Via Forging Strain Rate Limits for R'88DTSub-Solvus Recrystallization Mechanisms in UDIMET(r) Alloy 720LIThe Mechanical Property Response of Turbine Disks Produced Using Advanced PM Processing TechniquesSegregation and Solid Evolution during the Solidification of Niobium-Containing SuperalloysMicrostructural Evolution of Nickel-Base Superalloy Forgings during Ingot-to-Billet Conversion: Process Modeling and ValidationRemoval of Ceramic Defects from a Superalloy Powder Using Triboelectric ProcessingProduction Evaluation of 718ER(r) AlloyQuench Cracking Characterization of Superalloys Using Fracture Mechanics ApproachDevelopment and Characterization of a Damage Tolerant Microstructure for a Nickel Base Turbine Disc AlloyThe Microstructure Prediction of Alloy 720LI for Turbine Disk ApplicationsCharacteristics and Properties of As-HIP P/M Alloy 720Enhanced Powder Metallurgy (P/M) Processing of UDIMET(r)Alloy 720 Turbine Disks - Modeling StudiesCharacterization and Thermomechanical Processing of Sprayformed Allvac(r) 720Alloy

Solidification and Casting ProcessingProperties of RS5 and Other Superalloys Cast Using Thermally Controlled SolidificationAdvanced Superalloys and Tailored Microstructures for Integrally Cast Turbine WheelsImproved Quality and Economics of Investment Castings by Liquid Metal Cooling - The Selection of Cooling MediaA Novel Casting Process for Single Crystal Gas Turbine ComponentsCarbon Additions and Grain Defect Formation in High Refractory Nickel-Base Single Crystal SuperalloysNew Aspects of Freckle Formation during Single Crystal Solidification of CMSX-4Competitive Grain Growth and Texture Evolution during Directional Solidification of SuperalloysRecrystallization in Single Crystals of Nickel Base SuperalloysStructure of the Ni-Base Superalloy IN713C after Continuous CastingThe Thermal Analysis of the Mushy Zone and Grain Structure Changes during Directional Solidification of SuperalloysFreckle Formation in SuperalloysModelling of the Microsegregation in CMSX-4 Superalloy and its Homogenisation during Heat TreatmentEnhancement of the High Temperature Tensile Creep Strength of Monocrystalline Nickel-Base Superalloys by Pre-rafting in Compression

Blade AlloysAlloying Effects on Surface Stability and Creep Strength of Nickel Based Single Crystal Superalloys Containing 12 Mass% CrEvaluation of PWA 1483 for Large Single Crystal IGT Blade ApplicationsEffect of Ru Addition on Cast Nickel Base Superalloy with Low Content of Cr and High Content of WPrediction and Measurement of Microsegregation and Microstructural Evolution in Directionally Solidified SuperalloysDevelopment of a Third Generation DS SuperalloyThe Development and Long-Time Structural Stability of a Low Segregation Hf-free Superalloy - DZ125LThe Growth of Small Cracks in the Single Crystal Superalloy CMSX-4 at 750 and 1000 CThe Influence of Load Ratio, Temperature, Orientation and Hold Time on Fatigue Crack Growth of CMSX-4Modelling the Anisotropic and Biaxial Creep Behaviour of Ni-Base Single Crystal Superalloys CMSX-4 and SRR99 at 1223KCBED Measurement of Residual Internal Strains in the Neighbourhood of TCP Phases in Ni-Base SuperalloysThe Influence of Dislocation Substructure on Creep Rate During Accelerating Creep Stage of Single Crystal Nickel-based Superalloy CMSX-4Oxidation Improvements of Low Sulfur Processed Superalloys

Disk AlloysOptimisation of the Mechanical Properties of a New PM Superalloy for Disk Applicationsg' Formation in a Nickel-Base Disk SuperalloyMicrostructure and Mechanical Property Development in Superalloy U720LISub-Solidus HIP Process for P/M Superalloy Conventional Billet ConversionEffect of Oxidation on High Temperature Fatigue Crack Initiation and Short Crack Growth in Inconel 718The Effects of Processing on Stability of Alloy 718Long Term Thermal Stability of Inconel Alloys 718, 706, 909 and Waspaloy at 593 C and 704 CEffects of Microstructure and Loading Parameters on Fatigue Crack Propagation Rates in AF2-1DA-6The Common Strengthening Effect of Phosphorus, Sulfur and Silicon in Lower Contents and the Problem of a Net SuperalloySimulation of Microstructure of Nickel-Base Alloy 706 in Production of Power Generation Turbine Disks

Mechanical BehaviorInfluence of Long Term Exposure in Air on Microstructure, Surface Stability and Mechanical Properties of UDIMET 720LIEffects of Grain and Precipitate Size Variation on Creep-Fatigue Behaviour of UDIMET 720LI in Both Air and VacuumEffects of Local Cellular Transformation on Fatigue Small Crack Growth in CMSX-4 and CMSX-2 at High TemperatureMultiaxial Creep Deformation of Single Crystal Superalloys: Modelling and ValidationInvestigations of the Origin and Effect of Anomalous RaftingStress Rupture Behavior of Waspaloy and IN738LC at 600 C in Low Oxygen Gaseous Environments Containing SulfurIsothermal and Thermomechanical Fatigue of Superalloy C263Structure/Property Interactions in a Long Range Order Strengthened SuperalloyMicrostructural Changes in MA 760 during High Temperature Low Cycle FatigueHigh Temperature Low-Cycle Fatigue Behavior of Haynes 230 SuperalloyHigh Cycle Fatigue of ULTIMET AlloyThe Effect of Strain Rate and Temperature on the LCF Behavior of the ODS Nickel-Base Superalloy PM 1000Effect of Thermomechanical Processing on Fatigue Crack Propagation in INCONEL Alloy 783The Ductility of Haynes(r) 242 Alloy as a Function of Temperature, Strain Rate and Environment

Coatings, Welding and RepairProcessing Effects on the Failure of EBPVD TBCs on MCrAlY and Platinum Aluminide Bond CoatsCompositional Effects on Aluminide Oxidation Performance: Objectives for Improved Bond CoatsModelling and Neutron Diffraction Measurement of Stresses in Sprayed TBCsInterdiffusion Behavior in NiCoCrAlYRe-Coated IN-738 at 940 C and 1050 CEffect of Coating on the TMF Lives of Single Crystal and Columnar Grained CM186 Blade AlloyProcess Modelling of Electron Beam Welding of Aeroengine ComponentsNovel Techniques for Investigating the High Temperature Degradation of Protective Coatings on Nickel Base SuperalloysSintering of the Top Coat in Thermal Spray TBC Systems Under Service ConditionsOveraluminising of NiCoCrAlY Coatings by Arc PVD on Ni-Base SuperalloysThe Influence of B, P and C on Heat Affected Zone Micro-Fissuring in INCONEL type SuperalloyImproving Repair Quality of Turbine Nozzles Using SA650 Braze AlloyImproving Properties of Single Crystal to Polycrystalline Cast Alloy Welds through Heat Treatment

Alloy DevelopmentDevelopment of a New Single Crystal Superalloy for Industrial Gas TurbinesHigh g' Solvus New Generation Nickel-Based Superalloys for Single Crystal Turbine Blade ApplicationsDistribution of Platinum Group Metals in Ni-Base Single Crystal SuperalloysDevelopment of A Low Angle Grain Boundary Resistant Single Crystal Superalloy YH61Topologically Close Packed Phases in an Experimental Rhenium Containing Single Crystal SuperalloyA Low-Cost Second Generation Single Crystal Superalloy DD6The Development of Improved Performance PM UDIMET(r) 720 Turbine DisksMicrostructural Stability and Crack Growth Behaviour of a Polycrystalline Nickel-Base SuperalloyThe Application of CALPHAD Calculations to Ni-Based SuperalloysFormation of a Pt2Mo Type Phase in Long-Term Aged INCONEL Alloy 686Development of New Nitrided Nickel-Base Alloys for High Temperature ApplicationsMC-NG: A 4th Generation Single-Crystal Superalloy for Future Aeronautical Turbine Blades and Vanes