Author's Accepted Manuscript

Hot deformation characteristics of a polycrys-talline γ-γ′-δ ternary eutectic Ni-base super-alloy

Martin Detrois, Randolph C. Helmink, SammyTin

PII: S0921-5093(13)00891-5DOI: http://dx.doi.org/10.1016/j.msea.2013.07.089Reference: MSA30189

To appear in: Materials Science & Engineering A

Received date: 3 June 2013Revised date: 16 July 2013Accepted date: 19 July 2013

Cite this article as: Martin Detrois, Randolph C. Helmink, Sammy Tin, Hotdeformation characteristics of a polycrystalline γ-γ′-δ ternary eutectic Ni-basesuperalloy, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2013.07.089

This is a PDF file of an unedited manuscript that has been accepted forpublication. As a service to our customers we are providing this early version ofthe manuscript. The manuscript will undergo copyediting, typesetting, andreview of the resulting galley proof before it is published in its final citable form.Please note that during the production process errors may be discovered whichcould affect the content, and all legal disclaimers that apply to the journalpertain.

www.elsevier.com/locate/msea

Hot deformation characteristics of a polycrystalline �-�’-� ternary eutectic Ni-base superalloy

Martin Detrois a,*, Randolph C. Helmink b, Sammy Tin a

a Illinois Institute of Technology, 10 W. 32nd Street, Chicago, IL 60616, USA b Rolls Royce Corporation, PO Box 420, Indianapolis, IN 46241, USA

* Corresponding author. Tel.: +1 312 576 3789; E-mail address: mdetrois@iit.edu

Keywords Deformation; anisotropy; superalloy; delta; behavior; texturing

Abstract

Nickel-base superalloys possess an unparalleled combination of mechanical and physical properties for

elevated temperature applications and are likely to remain the first choice for structural components in turbine

engine applications in the future. In order to keep pace with increasingly demanding design requirements associated

with advanced gas turbine technologies, the properties of Ni-base superalloys can be greatly enhanced via thermo-

mechanical and/or compositional changes. Recent studies have revealed that ternary eutectic �-�’-� Ni-base

superalloys exhibit promising high temperature mechanical properties that may potentially be superior to state-of-

the-art polycrystalline Ni-base superalloys that are currently used in advanced gas turbines. As the properties of this

novel class of alloys are largely dependent on microstructural strengthening mechanisms, both the composition and

thermo-mechanical processing parameters need to be optimized concurrently. Deformation at temperatures and

strain rates between 1040�C to 1140�C and 0.1/s to 0.001/s resulted in a transition of dominant deformation

mechanisms from dislocation based plasticity to grain boundary rotation and sliding. The governing deformation

mechanism was found to influence the alignment of � phase precipitates and the degree of anisotropy that occurs

during hot deformation of �-�’-� Ni-base superalloys. Coupling high deformation temperatures with low strain rates

promotes grain boundary sliding and rotation leading to less internal cavitation damage and a more isotropic

microstructure.

Introduction

Nickel-base superalloys exhibit an excellent combination of properties that make them ideal for service at

elevated temperatures and corrosive environments [1-4]. For hot section structures that do not operate in the direct

path of the combustion gases and are maintained at temperatures below 700�C where creep deformation is not life

limiting [5-7], polycrystalline Ni-base superalloys provide the best balance of high strength and resistance to cyclic

loading, or fatigue. Polycrystalline Ni-base superalloys typically consist of a two phase microstructure composed of

a disordered FCC matrix � (Ni) and ordered L12 precipitates �’ (Ni3Al). This microstructure provides an excellent

combination of strength and toughness at elevated temperatures. In addition, the properties of these alloys can be

modified by the addition of other elements, such as Mo and W for solid solution strengthening of the � matrix, Ti,

Ta, Al and Nb for strengthening and control of the volume fraction of the �’ precipitates, Cr for environmental

resistance and minor alloying additions of B, C, Zr and Hf to control the grain size and grain boundary properties

[1-4].

As the operating temperatures of advanced gas turbine engines are continuously increasing to enhance

performance and improve efficiency, the development of polycrystalline Ni-base superalloys with conventional �-�’

microstructures capable of handling the increased temperatures has not kept pace. In recent years, only incremental

improvements in properties of polycrystalline Ni-base superalloys have been achieved via minor changes in

composition or variations in thermo-mechanical processing as the intrinsic limitations of these materials are

approached. In addition, these improvements have tended to further impair the formability of these high

temperature structural materials [8]. Thus, new alloying approaches or concepts are needed that provide

transformational improvements in properties and enable the design of advanced gas turbine engine concepts.

Initially investigated as candidate materials for turbine airfoil applications during the 1970’s and 1980’s,

directionally solidified �-��-� Ni-Al-Nb eutectic alloys were shown to possess excellent high temperature strength

and creep resistance [9-14]. High thermal gradients combined with slow withdrawal rates during directional

solidification processing of these alloys promoted the formation of an in-situ composite microstructure comprised

of aligned � lamellae contained within a �-�� matrix. Despite possessing excellent mechanical properties at elevated

temperature, the high degree of anisotropy and lack of transverse strength exhibited by these alloys made them

unsuitable for the intended applications. Coupled with the mismatch in thermal expansion coefficients, the low

interfacial strength between the � lamellae and surrounding �-�� matrix in directionally solidified structures resulted

in high notch sensitivities and severely degraded fatigue resistance. As such, these intrinsic attributes exhibited by

directionally solidified �-��-� eutectic alloys made them unsuitable for high temperature turbine airfoil applications.

Although the properties of these alloys render them unsuitable for use at temperatures above 1000�C, many of their

unique intrinsic attributes may potentially be harnessed for high performance structural applications where the

temperatures are limited to below 800�C and creep resistant columnar or single crystal grain structures are typically

not required.

Compared to state-of-the-art high strength polycrystalline Ni-base superalloys such as Udimet 720 and

Rene 95, this novel class of �-��-� alloys may potentially offer a number of benefits. Firstly, a substantially higher

level of dispersion/precipitate strengthening is utilized in this class of materials. With up to ~40% � and ~30% �� by

volume, the microstructure may contain up to 70% intermetallic phases to provide high temperature strength. The

high combined volume fraction of intermetallic precipitates rivals that of creep resistant single crystal Ni-base

superalloys produced via investment casting and provides potent levels of Orowan and composite strengthening

[1,3,15]. In addition to improved strength, microstructural stability after prolonged exposures to elevated

temperatures is becoming increasingly important in a number of structural applications. Highly alloyed

polycrystalline Ni-base superalloys are susceptible to the precipitation of topologically-close-packed (TCP) phases

when subjected to high temperature service environments for extended durations [16-18]. The precipitation of TCP

phases is primarily due to supersaturation of the � matrix with refractory alloying additions that have been added to

provide solid solution strengthening. Elevated levels of Cr, Mo and W combined with high equilibrium �� volume

fractions tend to promote the precipitation of TCP phases in �-�� Ni-base superalloys following long term thermal

exposures. The �-��-� alloys, however, are near eutectic alloys that possess high levels of microstructural stability

even at temperatures close to the melting point. The preferred solidification sequence results in the formation of a

pseudo-binary �-� eutectic structure from which the �� precipitates out in the solid state from the � phase upon

cooling. As the intermetallic � is a primary phase that forms upon solidification, the microstructure exhibits

excellent stability and additional TCP phases are unlikely to form within this class of alloys as confirmed both

experimentally and with ThermoCalc predictions using the TCNI5 database. Finally, the �-��-� alloys being

proposed as part of this study possess lower densities and raw materials costs when compared to commercial high

strength Ni-base superalloys. The majority of advanced powder processed Ni-base superalloys rely on the additions

of W, Mo and Ta to provide high levels of strengthening in both the � and �� phases [19-23]. Not only do these

refractory alloying additions increase the overall density of the alloy, they also increase the raw materials costs

associated with the alloy. The pseudo-eutectic �-��-� alloys are based on ternary Ni-Al-Nb or quaternary Ni-Al-Nb-

Cr systems and contain limited amounts of expensive refractory alloying additions. With up to a ~10% lower alloy

density, the comparatively higher density normalized specific strength may be ideal for producing low cost, high

strength structures for high performance applications. While the mechanical properties of this novel ternary eutectic

Ni base superalloy are being quantified, the capacity to forge these materials using isothermal or conventional

forging techniques has yet to be determined. This present investigation was performed in order to assess the

feasibility of forming near net shape structures with these �-��-� alloys by identifying a range of deformation

temperatures and strain rates amenable for inducing plastic or superplastic flow. Moreover, microstructural changes

such as the texturing and alignment of the � phase precipitates during hot deformation and forming was also

quantified as a function of the processing parameters.

Table 1: Composition of V204H (wt. %).

Ni Cr Co Al Ti Nb Ta Mo W B C Hf Zr Si Bal. 5.2 2 2.6 0.2 16.5 2 1.5 2 0.01 0.04 0.15 0.03-0.05 0.5

Procedure

The material considered in this investigation was a powder processed, Ni-base superalloy containing 16.5

wt. % Niobium to induce the formation of a �-�’-� microstructure. The composition of the alloy is listed in Table 1.

A 20kg alloy billet was produced via hot isostatic pressing (HIP) of spray atomized powder consolidated within a

mild steel can. Cylindrical rods measuring 10mm in diameter and 100mm in length were extracted from the billet

using wire electro-discharge machining (EDM).

In order to quantitatively assess the characteristic microstructural changes as a function of hot deformation,

metallographic samples were systematically prepared from the material at various processing stages, starting with

the as-HIP condition. The as-HIP rod was sectioned using a high speed diamond saw and metallographically

prepared finishing with a polish using 0.05 micron colloidal silica. Following the polish, surface relief revealed the

presence of the three phases: � (Ni), �’ (Ni3Al) and � (Ni3Nb). Scanning Electron Microscopy using a JEOL 5900

was then used to characterize the microstructures and measure the composition of the constituent phases using

energy dispersive spectroscopy (EDS). As the � phase precipitates form as elongated rods or platelets, the aspect

ratio (length over width) and relative orientations of these phases within the �-�’ matrix was assessed as a function

of the deformation processing parameters. To ensure that the results were statistically significant, the orientations

and aspect ratios of respectively 115 and 50 discrete � precipitates was measured over multiple fields and averaged

to obtain the reported results.

Cylindrical samples measuring 15mm height and 10mm diameter were sectioned from the rods for a series

of elevated temperature compression tests designed to simulate isothermal forming conditions using a Gleeble 3500

system. Twenty four compression samples were prepared to explore the effects associated with a broad range of

deformation temperatures and strain rates. Three strain rates (0.1, 0.01 and 0.001/s) and eight deformation

temperatures ranging from 1000°C to 1140°C at 20°C intervals were explored. All of the deformation temperatures

used were below the �’ solvus temperature of 1207°C for V204H. For all of the compression tests, the final strain

under compression was set to a constant value of 0.5. To minimize microstructural changes occurring prior to

deformation, the samples were rapidly heated to the set temperature in one minute and allowed to thermally soak

for three minutes prior to deformation. Three thermocouples were spot welded onto each compression sample and

used to monitor the temperature of the sample during deformation. Barreling of the samples during deformation

was minimized as variations of less than ~20�C were measured along the length of the specimen. During

deformation, graphite was used to minimize friction occurring between the compression anvil and sample.

Following compression tests, the samples were rapidly cooled to retain microstructures representative of those at

temperature and sectioned parallel to the deformation axis. These samples were metallographically prepared and

microstructural characterization was performed along the geometric centers of the samples where strains were

uniform and constant. Fiducial reference markers were placed on the sample to permit alignment of the samples

parallel to the deformation axis during characterization in the SEM. This provided a reference orientation for the

determination of the relative orientation distributions of the � phase precipitates.

The corresponding SEM images were then analyzed using digital image analysis software, GIMP. Lines

were manually drawn over the � phase precipitates to quantify their relative orientations. After subtraction of the

background, the angle between every � phase precipitates and the horizontal was measured and recorded.

Considering the accuracy of this method, results were binned into ranges of 10°. The number of precipitates in

these ranges was determined and using the total number of precipitates over the sample the percentage of

precipitates in every range was calculated. For each sample, a minimum of three different fields were characterized

at a magnification of 5000X and approximately 300 distinct � phase precipitates were considered. Figure 1

illustrates the process of quantifying the relative orientation distribution of � phase precipitates in a single field.

Finally, the deformation texture present in the deformed samples was quantified using an Oxford Instruments

Nanoscan Electron Back Scatter Diffractometer (EBSD).

Results

The constituent �-�’-� phases present within the as received Ni-base superalloy microstructure are shown in

Figure 2. The � phase precipitates appear as blocky plates or rods that exhibit a range of morphologies. In the as-

received or as-HIP condition, these � phase precipitates were distributed randomly along both the grain boundaries

and intergranularly within the underlying �-�’ matrix. To understand the plastic flow behavior of the alloy during

hot deformation, ternary eutectic Ni-base superalloys such as the one investigated in this study, can be viewed as

exhibiting a composite structure comprised of hard intermetallic � phases contained within a relatively ductile �-�’

matrix. Compositions of the constituent phases measured via energy dispersive spectroscopy (EDS) are reported in

Table 2. Compared to the �-�’ matrix, the � phase contains a significantly higher concentration of niobium (27.1%)

and a noticeably lower level of aluminum (0.5%). Interestingly, Cr also appears to exhibit limited solubility in the �

phase and partitions preferentially to the �-�’ matrix. To quantify the morphologies of the � phase precipitates, the

aspect ratios were systematically measured and assessed for each processing condition. In the as-received condition,

the � lattices were measured and the average aspect ratio of the � phase was determined to be approximately 4.7

with a maximum of 8.6 and a minimum of 1.9. Similar values were obtained after deformation, for example, the

average aspect ratio of the � phase after deformation at 1040°C for a strain rate of 0.01/s was found to be

approximately 4.6 with a maximum of 10.3 and a minimum of 1.2, and 4.8 with a maximum of 8.9 and a minimum

of 1.5 after deformation at 1140°C for a strain rate of 0.1/s.

Table 2: Composition of different phases using EDS (wt. %).

Al Cr Co Ni Nb Ta W Mo �-�’ 4.1 6.8 1.6 70.1 11.1 2.9 2.2 1.3

� 0.5 2.5 1.5 62.4 27.1 2.8 2.4 0.8

From the high temperature compression tests in the Gleeble, the maximum stress was found to vary as a

function of both temperature and strain rate, Table 3. Decreases in the maximum flow stress were observed as

deformation temperatures increased or as the strain rate decreased, Figure 3. The flow stress curves also show that

at temperatures below 1100°C, significant softening occurs once the maximum stress is achieved. For samples

deformed at temperatures above 1100°C, neither softening nor hardening was observed and the flow stresses were

not found to vary significantly as a function of strain. At a nominally constant temperature of 1000°C, the

maximum stress ranged from 592MPa at a strain rate of 0.1/s to 169MPa at a strain rate of 0.001/s, Table 3. For the

flow stress variations at a constant strain rate of 0.1/s, a difference of 486MPa was measured between deformation

conditions at 1000°C and 1140°C. The magnitude of the flow stresses was found to vary strongly as a function of

both temperature and strain rate. Deformation at temperatures above 1100°C or at strain rates of 0.001/s both serve

to lower the maximum stress.

All of the high temperature compression testing was conducted such that the samples were subject to a

nominally constant strain level of 0.5. Following deformation, the samples were sectioned to determine the effect of

the hot deformation parameters on the underlying microstructure. Due to powder processing of the alloy, no

damage was observed in the as-HIP condition, however, most of the samples were found to exhibit damage in the

form of porosity or cavitation damage following deformation. However, the extent of damage was found to

correlate to temperature and strain rate as higher temperatures and slower the strain rates resulted in a reduced

overall level of damage. Two distinct forms of damage were observed: cavitation damage and interfacial debonding

of the � phase precipitates. Quantification of the magnitude of overall damage associated with the various

processing parameters revealed that the extent of damage was most pronounced in samples deformed at relatively

low temperatures with a strain rate of 0.1/s. Damage within the samples was not uniform as near surface regions

where plastic flow induced tensile stresses were present exhibited significantly more damage than regions in which

primarily compressive stresses were located. For this reason, only the magnitude of damage along the geometric

centers of the compressed samples was assessed in this study. Figure 4 shows representative microstructures of the

�-�’-� alloy exhibiting different degrees of damage following hot deformation at various temperatures and strain

rates. The most damage corresponding to an area fraction of 1%, was observed in the sample deformed at 1000°C at

a strain rate of 0.1/s, while no measureable amount of damage was observed in the sample deformed at 1140°C at a

strain rate of 0.001/s.

Table 3: Maximum stress from the samples tested at various temperatures and strain rates.

Temperature Max Stress (MPa) 0.1/s 0.01/s 0.001/s

1000°C 592 384 169 1020°C 511 304 119 1040°C 399 229 91 1060°C 360 172 53 1080°C 278 129 30 1100°C 207 89 19 1120°C 158 61 11 1140°C 106 34 5

Following deformation, the initially random distribution and relative orientations of the � phase precipitates

became anisotropic as the precipitates became aligned in the direction of plastic flow, Figure 5. The degree of

anisotropy was found to be dependent on the processing parameters as increasing strain rates and lowering

deformation temperatures increased the extent of alignment of the � phase precipitates. For example, following

deformation at 1000°C and a strain rate of 0.001/s, only a modest degree of texturing was observed within the

resulting microstructure, while a significant degree of anisotropy was observed in the sample deformed at 1000°C

and a strain rate of 0.1/s, Figure 5. The degree of anisotropy in each of the samples was quantified by measuring the

overall distribution of relative orientations associated with the � phase precipitates. Comparison of these results

allows for correlation of the degree of anisotropy as a function of deformation temperature and strain rate, Figure

6a-c. For a constant strain rate, the relative orientation distribution of the � phase precipitates was nominally similar

irrespective of the range of temperatures investigated (1000°C to 1120°C). However, the degree of alignment of the

� phase precipitates following deformation was found to vary as a function of strain rate as the distribution of

relative precipitate orientations was less pronounced as strain rates decreased. Since deformation over the range of

temperatures investigated did not significantly affect the distribution of precipitate orientations, results were

compiled for each strain rate and the average � phase precipitate orientation distribution was compiled in Figure 6d.

Comparing these results show that a uniform distribution of precipitates exists in the as-HIP condition and the

deformation induces a certain degree of alignment of the � phase precipitates. Interestingly, the degree of alignment

and the distribution functions were found to vary as a function of strain rate. Following deformation at a strain rate

of 0.001/s, most of the � precipitates were oriented at angles between 40° and 90° with the highest distribution of

17% corresponding to 65°, Figure 6a. As strain rates increased to 0.01/s, the peak of the orientation distribution

function shifted to higher misorientation angles between 50° and 90° which correspond to higher degrees of

anisotropy, Figure 6b. For strain rates of 0.1/s, most � phase precipitates were oriented at angles between 80° and

90°, Figure 6c.

Discussion

Ternary eutectic �-�’-� Ni-base superalloys have been shown to possess excellent mechanical properties

and exhibit potential for use in structures for advanced turbine engines. Due to their in situ composite

microstructure comprised of a high volume fraction of intermetallic reinforcement, these alloys are capable of

maintaining good strength and resist both creep and environmental degradation at temperatures up to 800°C [24-26].

However, the bulk formability characteristics and effect of thermal-mechanical processing on this class of materials

is not well understood. For this reason, we have conducted this investigation to quantify the effects of hot

deformation on the microstructure of these eutectic �-�’-� Ni-base superalloys. In the as-HIP condition, the

microstructure resembles a particulate reinforced metal matrix composite (MMC) in which the � phase precipitates

are distributed randomly within the �-�’ matrix. The � phase precipitates appear as elongated or blocky rods that

contain a significantly higher concentration of niobium and lower concentration of aluminum compared to the �-�’

matrix. Therefore, in order to understand the plastic flow behavior of the alloy during the hot compression, the �

phase precipitates were treated as rigid, non-deforming particles that are contained within the plastically flowing �-

�’ matrix. The assumption that the � phase precipitates are rigid and non-deforming is consistent with the nominally

constant aspect ratios of the precipitates measured before and after deformation.

From the hot deformation tests, relatively high flow stresses and moderate levels of accumulated cavitation

damage within the microstructure were observed in samples deformed at temperatures between 1000�C and 1040�C

at a strain rate of 0.1/s. However, when the strain rates were reduced to 0.001/s, the measured flow stresses were

significantly lower and the level of observed damage became negligible at temperatures above 1040�C. This

indicates that isothermal forging operations are required to physically shape the structure. With an average grain

size of less than 10microns in diameter, the combination of the high deformation temperatures and low strain rates

enables diffusion based mechanisms to play a dominant role during deformation as the macroscopic strains can be

accommodated via superplastic flow mechanisms such as grain rotation and sliding. Under these conditions, less

texturing or alignment of the � phase precipitates occurs as the intragranular precipitates rotate in conjunction with

the grains. These findings are also consistent with the observed relative orientation distributions of � phase

precipitates in the samples deformed at 0.001/s. There was noticeably less texturing and alignment of the

precipitates under hot deformation conditions that promote superplastic flow. Moreover, with less plastic strain

being accumulated within the �-�’ grain and at the � precipitate interfaces, the occurrence of cavitation damage also

decreased. Furthermore, deformation at temperatures approaching the �’ and/or � solvus temperatures require lower

flow stresses as the overall volume fraction of intermetallic phases within the microstructure decreases thereby

making it more conducive for grain rotation and sliding. Conversely, as strain rates increase, dislocation based

plasticity mechanisms become dominant and are responsible for accommodating strain since there is insufficient

time for the grains to rotate and slide under these conditions.

The flow stress behavior of the �-�’-� Ni-base superalloy was studied in detail to better understand the

characteristic behavior of the alloy during hot deformation [27]. The Zener-Hollomon parameter was used to model

the flow stress behavior of the �-�’-� Ni-base superalloy as a function of deformation temperature and strain rate.

Characteristic flow stresses corresponding to strains of 0.4 were chosen to avoid regions of significant softening

induced by the segmentation of prior grains during deformation following the peak stress. When coupled with low

deformation temperatures and relatively high strain rates, � phase precipitates have been shown to increase the

generation of dislocations during deformation. This tends to result in higher magnitudes of peak stresses followed

by significant softening [28-30]. Since the mechanisms governing deformation were found to vary within the range

of deformation temperatures and strain rates investigated, two set of model parameters were used to delineate the

transition in mechanisms from dislocation based plasticity to grain boundary rotation and sliding. One set of model

parameters denoted with a subscript p for “plastic” corresponds to characteristic flow stresses >50MPa at a strain of

0.4 which are representative of the stresses that are required for dislocation based plasticity. The other set of model

parameters denoted with a subscript c for “creep” corresponds to flow stresses <50MPa since the magnitude of

stresses required for grain boundary sliding are comparatively lower. The value of this transition stress was

determined by optimization of the model to match with the experimental data and is equivalent to a maximum

stress of approximately 100MPa. The functional form of the Zener-Hollomon parameter is shown in equations (1)

and (2).

In Equations (1) and (2), Z is the Zener-Hollomon parameter, is the strain rate, Q is the hot deformation

activation energy, R is the gas constant equal to 0.008314 kJ/mol, T is the deformation temperature, � is the flow

stress corresponding to a strain of 0.4, n is the strain rate sensitivity exponent relative to the hardening or softening

behavior of the material and finally, A and � are material parameters. Using data from the experimental flow stress

curves, a unified and global set of parameters capable of describing the behavior of the material over the two

deformation regimes was back calculated. The hot deformation activation energy Q was found to be 400 kJ/mol for

both conditions, which is close to the expected value for hot deformation relative to common nickel based

superalloys. Values for the other parameters are as follows: �p = 0.35/MPa, �c = 4.30/MPa, np = 2.7, nc = 1.3, Ap =

3.24x1010/s and finally, Ac = 2.65x1010/s. Figure 7 shows a comparison of the model and experimental data. This

plot reveals good agreement between the model and the experimental data. Plastic deformation was observed above

the transition stress while creep deformation was predominant during deformation involving stresses below the

transition stress. Results from the flow stress modeling were coupled with observations of the levels of texturing

within the samples after deformation at 1140°C for strain rates of 0.001/s, 0.01/s, and 0.1/s, Figure 8 and Figure 9.

The orientation maps, Figure 8, were obtained at magnification 500X on the SEM with an EBSD step size of 0.3.

To better observe the grain orientations within the �-�’ matrix, only the FCC � and L12 �’ phase were indexed and

the orthorhombic � phase precipitate corresponds to the unindexed regions of the orientation map. Quantification of

the degree of texturing was extracted from the EBSD orientation maps in the form of inverse pole figures using the

software HKL Channel 5, Figure 9. Consistent with the statements above, a higher overall level of texturing was

observed with increasing strain rate. At 1140�C, a nearly random grain orientation distribution was found following

deformation at 0.001/s on the inverse pole figure Figure 9a. As the strain rate increased to 0.01/s, texturing along

the {001} plane was observed with a maximum density of 1.34, Figure 9b. Finally, an even higher degree of

texturing was measured following deformation at 0.1/s, Figure 9c, with a maximum density of 1.50. Common grain

orientations were observed for all of the samples, Figure 8. Moreover, intragranular misorientations that correspond

to stored strain energy or structures comprised of high dislocation densities were observed within the microstructure

following deformation at 0.1/s at a temperature of 1140°C, Figure 8a. These observations are consistent with the

delineation of dominant deformation mechanisms as predicted by the flow stress model, Figure 7. Conversely, at a

strain rate of 0.001/s, Figure 8b, a moderate extent of grain growth was observed to occur within regions of the

microstructure that did not contain significant fractions of grain boundary �. Based on the results from the flow

stress modeling, Figure 7, superplastic flow dominates this deformation regime. Observation of orientation maps

for the sample deformed at 1140°C and 0.01/s showed a nominally identical image as for the sample deformed at

0.001/s and is consistent with the transition boundary between dislocation based plasticity and superplastic flow

defined using the flow stress model in Figure 7. The same observation was made when comparing samples

deformed at 1080°C and 0.01/s to those deformed at 1140°C and 0.1/s since both samples are located in the

dislocation based plasticity regime.

To summarize, ternary eutectic �-�’-� Nickel-base superalloys exhibit a unique set of physical and

mechanical properties that may be potentially harnessed for high temperature structural applications. In this

investigation, the effect of hot deformation processing parameters on the flow stresses, flow behavior,

microstructure and � phase precipitate orientation distributions were investigated. During deformation, a transition

flow stress was identified that corresponded to a change in the dominant deformation mechanism. Deformation

temperatures below 1030�C at strain rates of 0.001/s, below 1110�C at 0.01/s, and for all of the temperatures

investigated at strain rates of 0.1/s, dislocation based plasticity was responsible for high flow stresses, texturing or

alignment of the � phase precipitates with the plastic flow. Under these conditions, cavitation damage was also

accumulated at the � phase interfaces. Following deformation involving flow stresses below the transition stress i.e.

temperatures above 1030�C at a strain rate of 0.001/s and above 1110�C at 0.01/s, flow stresses were significantly

reduced as grain boundary sliding and rotation were dominant. As a result, this led to less texturing and alignment

of the � phase precipitates and cavitation damage. These results clearly indicate that these novel ternary eutectic �-

�’-� Nickel-base superalloys may be practically formed into near net shape products using conventional isothermal

forging processes.

Conclusions

Based on the results and observations presented in this study, the following conclusions can be reported:

1. Deformation producing maximum flow stresses >100MPa due to high strain rates and/or low deformation

temperatures involve dislocation based plastic deformation resulting in higher flow stresses, higher levels

of microstructural damage and a more textured alignment of � phase precipitates and grain orientation.

2. Deformation involving maximum flow stresses <100MPa produced by low strain rates and/or high

deformation temperatures induce grain rotation and sliding resulting in lower flow stresses, lower levels of

microstructural damage and a less textured alignment of the � phase precipitates and grain orientation.

3. The � phase precipitates tend to orientate themselves in the direction of plastic flow and are more

influenced by variations in the strain rate than in deformation temperature.

4. The Zener-Hollomon parameter was used relate the flow stress behavior of the material as a function of

deformation temperature and strain rate.

Acknowledgements

Financial support for this work was provided by Rolls-Royce North American Technologies, Rolls-Royce

Corporation, Rolls-Royce plc., and NSF-DMR-1006953.

References [1] T.M. Pollock, S. Tin, Journal of Propulsion and Power 22.2 (2006) 361-374. [2] D. Furrer, H. Fecht, JOM 51.1 (1999) 14-17. [3] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, Wiley, New York, 1987. [4] R.F. Decker, C.T. Sims, The Metallurgy of Nickel-base Superalloys, Paul D. Merica Research Laboratory, New York, 1972. [5] R.R. Unocic, G.B. Viswanathan, P.M. Sarosi, S. Karthikeyan, J. Li, M.J. Mills, Materials Science and Engineering A 483-

484 (2008) 25-32. [6] R.R. Unocic, L. Kovarik, C. Shen, P.M. Sarosi, Y. Wang, J. Li, S. Ghosh, M.J. Mills, Superalloys, TMS, Warrendale, PA,

2008. [7] D. Locq, P. Caron, S. Raujol, F. Pettinari-Sturmel, A. Coujou, N. Clément, Superalloys, TMS, 2004. [8] R.M. Forbes Jones, L.A. Jackman, JOM (1999). [9] A.D. Cetel, M. Gell, J.W. Glatz, Conference on In Situ Composites III, Boston, MA, (1978) 292-302. [10] R.W. Farley, The Superalloys, Wiley, New York, 1972. [11] H.R. Gray, Material Show and Conference, NASA TM 73714, (1977). [12] R.L. Ashbrook, Specialists Meeting on Directionally Solidified In Situ Composites, NASA TM X-71514, (1974). [13] J. Stringer, D.M. Johnson, D.P. Whittle, Oxidation of Metals 13.3 (1978). [14] D.M. Johnson, D.P. Whittle, J. Stringer, Oxidation of Metals 12.3 (1978). [15] R.C. Reed, The Superalloys: Fundamentals and Applications, Cambridge University Press, New York, 2006. [16] J.X. Yang, Q. Zheng, X.F. Sun, H.R. Guan, Z.Q. Hu, Materials Science and Engineering: A 465 (2007) 100-108. [17] S. Zhao, X. Xie, G.D. Smith, S.J. Patel, Materials Science and Engineering A 355 (2003) 96-105. [18] S.E. Kim, M.P. Jackson, R.C. Reed, C. Small, A. James, N.K. Park, Materials Science and Engineering A 245 (1998) 225-

232. [19] J. Tiley, G.B. Viswanathan, R. Srinivasan, R. Banerjee, D.M. Dimiduk, H.L. Fraser, Acta Materialia 57 (2009) 2538-2549. [20] Y. Gao, J.S. Stölken, M. Kumar, R.O. Ritchie, Acta Materialia 55 (2007) 3155-3167. [21] C. Stöcker, M. Zimmermann, H.-J. Christ, Z.-L. Zhan, C. Cornet, L.G. Zhao, M.C. Hardy, J. Tong, Materials Science and

Engineering: A 518 (2009) 27-34. [22] Y.F. Gu, C. Cui, D. Ping, H. Harada, T. Fukuda, J. Fujioka, Materials Science and Engineering: A 510-511 (2009) 250-255. [23] C. Cui, Y. Gu, H. Harada, A. Sato, Metallurgical and Materials Transactions A 36 (2005) 2921-2927. [24] M. Xie, R.C. Helmink, S. Tin, Metallurgical and Materials Transactions A 43 (2011) 1259-1267. [25] M. Xie, R.C. Helmink, S. Tin, Superalloys 2012, Wiley, Hoboken, 2012, pp. 633-642. [26] S. Tin, A. Rodriguez, A. DiScuillo-Jones, R.C. Helmink, R. Hardy, Superalloys 2012, Wiley, Hoboken, 2012, pp. 833-841. [27] D. Snyder, E.Y. Chen, C.C. Chen, S. Tin, Metallurgical and Materials Transactions A 44 (2013) 479-493. [28] S. Azadian, L.-Y. Wei, R. Warren, Materials Characterization 53 (2004) 7-16. [29] H.M. Lalvani, M.A. Rist, J.W. Brooks, Advanced Materials Research 89-91 (2010) 313-318. [30] Y. Wang, W.Z. Shao, L. Zhen, B.Y. Zhang, Materials Science and Engineering A 528 (2011) 3218-3227.

Figure 1: Procedure to quantify the relative orientations of the � precipitates following deformation at 1120°C and 0.001/s. (a) Lines were manually drawn on the � phase precipitates and the (b) relative orientation of the precipitates was measured. Figure 2: Microstructure of V204H in the as-received or as HIP condition. Figure 3: Flow stress curves from Gleeble tests at various temperatures and strain rates of (a) 0.1/s, (b) 0.01/s and (c) 0.001/s. Figure 4: Microstructures following hot deformation at (a) 0.1/s – 1000°C, (b) 0.1/s – 1140°C, (c) 0.001/s – 1000°C, and (d) 0.001/s – 1140°C. Figure 5: Microstructures after deformation at 1000°C and a strain rate of (a) 0.001/s and (b) 0.1/s. Figure 6: Orientation of the � phase precipitates after deformation at various temperatures at a strain rate of (a) 0.001/s, (b) 0.01/s, (c) 0.1/s along with (d) a comparison for the different strain rates. Figure 7: Modeling of the flow stresses. Figure 8: EBSD Orientation maps following deformation at 1140°C and (a) 0.1/s and (b) 0.001/s. (c) Inverse pole figure legend corresponding to the grain orientations. Figure 9: Inverse pole figures for the observation of the texturing involved after deformations at 1140°C and strain rates of (a) 0.001/s, (b) 0.01/s, and (c) 0.1/s.

FIG

1A

FIG

1B

FIG

2

0

100

200

300

400

500

600

0 0.

1 0.

2 0.

3 0.

4 0.

5

Stress (MPa)

Stra

in

1000

°C

1020

°C

1040

°C

1060

°C

1080

°C

1100

°C

1120

°C

1140

°C

FIG

3A

0 50

100

150

200

250

300

350

400

0 0.

1 0.

2 0.

3 0.

4 0.

5

Stress (MPa)

Stra

in

1000

°C

1020

°C

1040

°C

1060

°C

1080

°C

1100

°C

1120

°C

1140

°C

FIG

3B

0 20

40

60

80

100

120

140

160

180

0 0.

1 0.

2 0.

3 0.

4 0.

5

Stress (MPa)

Stra

in

1000

°C

1020

°C

1040

°C

1060

°C

1080

°C

1100

°C

1120

°C

1140

°C

FIG

3C

FIG

4A

FIG

4B

FIG

4C

FIG

4D

FIG

5A

FIG

5B

0.0

5.0

10.0

15.0

20.0

25.0

Percentage of δ precipitates

Angl

e (°)

1000

°C

1040

°C

1080

°C

1120

°C

Mea

n

FIG

6A

0.0

5.0

10.0

15.0

20.0

25.0

Percentage of δ precipitates

Angl

e (°)

1000

°C

1040

°C

1080

°C

1120

°C

Mea

n

FIG

6B

0.0

5.0

10.0

15.0

20.0

25.0

30.0

Percentage of δ precipitates

Angl

e (°)

1000

°C

1040

°C

1080

°C

1120

°C

Mea

n

FIG

6C

0.0

5.0

10.0

15.0

20.0

25.0

Percentage of δ precipitates

Angl

e (°)

as-H

IP

0.00

1 0.

01

0.1

FIG

6D

1.0

2.0

3.0

4.0

5.0

6.0

-8

-7

-6

-5

-4

-3

-2

-1

ln(σ) (Mpa)

ln(S

trai

n Ra

te) (

1/s)

1000

°C M

odel

10

00°C

Exp

. 10

20°C

Mod

el

1020

°C E

xp.

1040

°C M

odel

10

40°C

Exp

. 10

60°C

Mod

el

1060

°C E

xp.

1080

°C M

odel

10

80°C

Exp

. 11

00°C

Mod

el

1100

°C E

xp.

1120

°C M

odel

11

20°C

Exp

. 11

40°C

Mod

el

1140

°C E

xp.

Tran

sitio

n

Dis

loca

tion

base

d pl

astic

ity

Supe

rpla

stic

flow

FIG

7

FIG

8A W

EB

FIG

8A P

RIN

T

FIG

8B W

EB

FIG

8B P

RIN

T

FIG

8C W

EB

FIG

8C P

RIN

T

FIG9A

FIG9B

FIG9C