Investigation of strain aging in the ordered intermetallic compound B-NiAl


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Investigation of strain aging in the ordered intermetallic compound B-NiAl
Alternate title:
Investigation of strain aging in the ordered intermetallic compound Beta-NiAl
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ix, 195 leaves : ill. ; 29 cm.
Weaver, Mark Lovell, 1965-
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Materials Science and Engineering thesis, Ph. D
Dissertations, Academic -- Materials Science and Engineering -- UF
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Thesis (Ph. D.)--University of Florida, 1995.
Includes bibliographical references (leaves 184-194).
Statement of Responsibility:
by Mark Lovell Weaver.
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University of Florida
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Copyright 1995


Mark Lovell Weaver

To Mom and Dad.


The feat of obtaining a Doctor of Philosophy degree would be impossible without

the emotional, technical and financial support of others. Fortunately, I have been blessed
with a large number of individuals to support me during this period and am happy to

recognize them. First, I thank God for giving me guidance and insight when it seemed as

though nothing would work. I thank my parents, my brother, and my sister for their

unconditional love and support. Without them this never would have been possible. In
addition, I would like to thank my thesis advisor, Dr. M.J. Kaufman, whose assistance,
encouragement ,and warm personal friendship have made my graduate work both
rewarding and enjoyable. I would also like to thank the members of my supervisory

committee, Drs. F. Ebrahimi, R. Abbascian, R.E. Reed-Hill, and A.V. Kumar, for all of

the valuable advice and helpful comments. I especially would like to acknowledge the

friendship and technical guidance of Dr. Ron Noebe during my pursuit of this degree. I

would also like to acknowledge Drs. John J. Lewandowski, John Hack, Ben Oliver,

Vladmir Levit, and Ram Darolia for supplying some of the material used in this study and

for their helpful comments, enlightening discussions, and candid criticisms.

Finally, I would like to acknowledge the following co-workers and friends for the

technical assistance, emotional support, and for making my graduate school tenure most

enlightening, enjoyable, and bearable: Chris Moen (my mentor--thanks for teaching me

how to get things done); Andy Duncan and Tim Wattleworth (the best friends that anyone
could ask for); Jon Shults, Chris O'Gara, John Bockman, Dean Paxton, Thad Adams,
Andre Costa E Silva, Alex and Sunday Cozzi, Jim Cotton, Randy and Cheryl Bowman,

Patrick Wilson, Gail Anderson, Cindy Link, Don and Adrienne Jones, Steve Trail, Paul

Crofts, Andy Ibbotson, Jesse Mitrani, and Ryan Kaufman.


ACKNOWLEDGEMENTS ...................................................................... iv

ABSTRACT .................................................................................. viii


1 INTRODUCTION ....................................................................... 1

Background ...............................................................................1
Approach ..................................................................................4

2 LITERATURE REVIEW................................................................ 5

Introduction........................................................................ 5
Physical Metallurgy of NiAl and NiAl Alloys ...................................... 5
Physical/Thermnodynamic Properties........................................ 5
Lattice Parameter, Density and Defect Structures......................... 6
Elastic Properties............................................................ 7
Flow and Fracture Behavior ............................................................ 8
Slip Systems ................................................................... 8
Yield Strength.................................................................. 10
Ductility and Fracture......................................................... 10
Tensile behavior...................................................... 10
Fracture toughness................................................. 12
Influence of Microalloying Additions and Impurities..................... 12
Strain Aging ............................................................................. 13
Static Strain Aging (SSA).................................................... 13
Mechanisms of SSA ............................................... 14
Yield point return..................................................... 16
Dynamic Strain Aging (DSA).................................................. 16
Serrated flow curves................................................. 17
Types of serrations..................................................... 18
Theories of DSA ............................................................... 20
The Cottrell model.................................................... 20
The McCormick model ............................................ 21
The van den Beukel model.......................................... 23
The Reed-Hill model................................................... 26
Other models.......................................................... 27
Static and Dynamic Strain Aging in Ordered Alloys ............................... 28

SOFT-ORIENTED NiAl SINGLE CRYSTALS................................... 41

Experimental............................... ........................................... 42
Summary and Conclusions............................................................ 49

OF POLYCRYSTALLINE NiAl ..................................................... 57

Background............................................................................. 57
Experimental Details ..................................................................... 58
Material Characterization ..................................................... 58
Tensile Testing................................................................. 59
Experimental Results................................................................... 60
Composition and Microstructure ............................................ 60
Mechanical Properties......................................................... 62
Influence of Prestraining and Annealing on Baseline Properties ........ 62
TEM Observations of Deformed Specimens................................. 63
Discussion............................................................................... 65
Species Responsible for Strain Aging in NiAl...................................... 65
Influence of Prestraining .............................................................. 66
Conclusions............................................................................. 67

NiAl-BASED ALLOYS .................................... 97

Introduction............................................................................. 97
Experimental........................................................................... 98
Materials........................................... 98
Mechanical Testing.............................................................99
Results................................................................................... 100
Discussion............................................................................... 104
Summary and Conclusions ............................................................ 109

POLYCRYSTALLINE NiAl .......................................................... 119

Introduction............................................................................. 119
Materials and Methods................................................................. 119
Results................................................................................... 121
Microstructural Characterization.......................... 121
Tensile Properties ................................... 122
Discussion............................................................................... 124
Summary and Conclusions ............................................................ 126

ORIENTED NiAl SINGLE CRYSTALS ........................................... 136

Introduction............................................................................... 136
Experimental............................................ 137

Results..................................................................................... 139
Composition and Microstructure .............................................. 139
Mechanical Properties......................................................... 139
TEM observations of deformed samples ..................................... 142
Discussion............................................................................... 143
Conclusions............................................................................. 147

NiAl...................................................................................... 170

9 CONCLUSIONS....................................................................... 182

LIST OF REFERENCES ...................................................................... 184

BIOGRAPHICAL SKETCH.................................................................... 195

Abstract of Thesis Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy



May, 1995

Chairman: Dr. Michael J. Kaufman
Major Department: Materials Science and Engineering

The phenomenon of strain aging has been investigated in polycrystalline and single

crystal NiAl alloys at temperatures between 300 and 1200 K. Static strain aging studies

revealed that after annealing at 1100 K for 7200 s (i.e., 2 h) followed by furnace cooling,

high purity, nitrogen-doped and titanium-doped polycrystalline alloys exhibited continuous

yielding, while conventional-purity and carbon-doped alloys exhibited distinct yield points

and Lfiders strains. Prestraining by hydrostatic pressurization removed the yield points,

but they could be reintroduced by further annealing treatments. Yield points could be re-

introduced more rapidly if the specimens were prestrained uniaxially rather than

hydrostatically, owing to the arrangement of dislocations into cell structures during uniaxial
deformation. The time dependence of the strain aging events followed a t3 relationship

suggesting that the yield points observed in polycrystalline NiAl were the result of the
pinning of mobile dislocations by interstitials, specifically carbon.

Between 700 and 800 K, yield stress plateaus, yield stress transients upon a ten-
fold increase in strain rate, work hardening peaks, and dips in the strain rate sensitivity

(SRS) have been observed in conventional-purity and carbon-doped polycrystals. In single

crystals, similar behavior was observed; in conventional-purity single crystals, however,
the strain rate sensitivity became negative resulting in serrated yielding, whereas, the strain
rate sensitivity stayed positive in high purity and in molybdenum-doped NiAl. These

observations are indicative of dynamic strain aging (DSA) and are discussed in terms of
conventional strain aging theories. The impact of these phenomena on the composition-

structure-property relations are discerned. Finally, a good correlation has been
demonstrated between the properties of NiAl alloys and a recently developed model for
strain aging in metals and alloys developed by Reed-Hill et al. [1-3].



The development of more efficient gas turbine engines will depend upon the
advancement of new high temperature materials with improved mechanical properties.

These materials must have higher specific strengths and must maintain these strengths to
higher temperatures than the nickel-base superalloys currently in use. Ordered intermetallic

compounds are prime candidates to replace the superalloys. These materials form long

range ordered crystal structures which are associated with the formation of strong A-B type
bonds that typically result in high elastic moduli, high melting points, and high strengths.

Furthermore, aluminide-based intermetallics exhibit good oxidation resistance and greater

microstructural stability due to lower self-diffusion rates, another consequence of the

strong bonding.
Of the many intermetallic systems, alloys based on f3-NiAl are particularly attractive

for development. NiAl has a simple B2 (CsC1) crystal structure which is similar to the

body-centered-cubic (BCC) structure. It exists over a wide range of stoichiometries

(greater than 20 at.% at 1673 K) which allows for significant alloying to improve its

mechanical properties. In addition, it exhibits an appreciably higher melting point (-1950 K

versus 1573 K), lower density (5.9 g/cm3 versus 9.0 g/cm3), and a thermal conductivity
up to eight times greater than that of Ni-base superalloys. The potential benefits of using
NiAl in gas turbine applications would include: (1) decreased cooling requirements; (2)

decreased weight; and (3) higher operating temperatures resulting in increased operating

efficiencies and higher thrust-to-weight ratios. In addition, NiAl has been used for years as

an oxidation resistant coating on turbine blades [4]. Like most intermetallics, however,

NiAl is not currently a suitable replacement for superalloys due to its meager high
temperature strength (above 0.45Tmp) and its poor fracture resistance and ductility below

its brittle-to-ductile transition temperature (BDT).

NiAl undergoes a dramatic brittle-to-ductile transition at temperatures between 400
and 1000K with the actual transition temperature varying strongly with alloy composition,

processing history (i.e., thermal or mechanical), strain rate, bulk form (i.e., polycrystal or
single crystal), and orientation in single crystals. In polycrystals for example, Noebe et al.
[5] have reported a 200K increase in the BDTT coinciding with a three order of magnitude
increase in strain rate. Lahrman and coworkers [6] have reported similar increases in the
BDTT of "hard" [001] oriented single crystals. In "soft" non-[001] orientations however,
the BDTT was found to be less strain rate sensitive.

The BDTT in polycrystals and hard single crystals have been attributed to the onset
of localized dislocation climb processes driven by short circuit diffusion [5]. In soft single
crystal orientations however, the mechanism responsible for the BDTT is not as obvious.
In crystals with soft orientations the BDIT occurs as low as 0.25Tmp and, as mentioned

above, is much less sensitive to strain rate [6]. Possible explanations for the BDTT in soft
orientations include enhancement of cross-slip leading to slip homogenization the operation

of thermally activated deformation processes as in the other forms of NiAl or the unlocking

of dislocations from point defects or impurities [7]. These explanations are only

speculative and remain to be proven. Considering the peculiarly low BDTT in soft single
crystals, an understanding of the mechanisms responsible for the BDTT in soft single

crystals could reveal methods to lower the BDTT and lead to the development of more
ductile alloys.
At ambient temperatures, research has shown that NiAl deforms predominantly by
dislocation glide on the <100>(011) and <100>{001) slip systems of which only three are
independent [8]. According to the Von Mises criterion, at least five independent slip

systems are required for macroscopic plastic deformation in a polycrystalline body; this

may account for the lack of ductility in polycrystalline NiAl [9-13].
In single crystals, limited tensile ductilities (-2%) are also observed. However,
recent efforts have resulted in room temperature tensile elongations of up to 6% in

conventional-purity binary NiAl annealed above the BDTr and rapidly cooled to room
temperature [14,15] and in ternary NiAl crystals containing microalloying additions of Fe,
Mo and Ga [16]. The mechanism behind this increased ductility is currently unknown but
is suspected to be related to a gettering phenomenon since it is realized that some
substitutional elements and interstitial impurities can cause significant embrittlement in BCC
metals [17] and B2 ordered compounds such as NiAl [18]. Recent observations are
consistent with this viewpoint. For example, room temperature tensile elongations of up to

5% have been observed in low interstitial binary NiAl single crystals [19] and the ductility
and fracture strength in biaxial bending of high purity NiAl was found to be significantly
greater than for commercial purity crystals [20]. Still, the effects of particular substitutional
and interstitial elements on the mechanical behavior of NiAl is relatively unknown and the
mechanisms by which various elements may enhance or hinder tensile ductility still remain

a matter of conjecture.

Hack [14,15] attributes the profound effects of heat treatment on mechanical
properties to the creation mobile dislocations and suggests that the low ductilities reported

in other studies are due to strain age embrittlement in which interstitial atoms segregate to
mobile dislocations at moderate temperatures. The interstitials then pin the dislocations

resulting in a lower density of mobile dislocations. Evidence in support of a strain aging
effect are provided by the observation of serrated yielding [15,21-26], plateaus in the
temperature dependence of yield strength, and low strain rate sensitivities near the BDTT
and in the temperature regimes where serrated flow is observed [27-29]. In addition, strain
aging has also been reported in Ni-rich NiAl [22], stoichiometric polycrystalline NiAl [30],
mechanically alloyed NiAl [30], NiAl deformed under hydrostatic pressure [23,31], and in

other B2 intermetallics [32-34]. These phenomena are similar to those commonly observed

in mild steels and BCC refractory metals [21,35-43]. However, the species responsible for
this behavior in NiAl or other ordered alloys (i.e., interstitials, substitutional impurities,

precipitates, vacancies, etc.) are currently unknown. A comprehensive understanding of

these factors is required to fully discern the flow behavior of NiAl.

Thus, the objectives of this investigation are to examine, in a systematic fashion,
the phenomenon of strain aging and it's influence on the mechanical behavior of near

stoichiometric NiAl.


In order to examine the effects of strain aging, several nominally stoichiometric
polycrystalline and single crystal alloys containing varying interstitial contents were
subjected to static strain aging (SSA) experiments and strain rate change experiments to
reveal the temperature dependence of the strain rate sensitivity (SRS). The resulting data

were then used to determine activation energies for deformation processes and to determine
the species responsible for strain aging in NiAl. The resulting data was also analyzed using

a recently developed theory by Reed-Hill et al. [1,3].



In the sections that follow, the physical metallurgy of NiAl alloys and the general

aspects of static and dynamic strain aging phenomena are described. As the physical
metallurgy of NiAl alloys has been described in detail in recent review articles [7,18,44],

only a summary of the properties pertinent to the deformation of NiAl is provided.

Physical Metallurgy of NiAl and NiAl Alloys

PhysicalThermodnamic Properties

The ordered intermetallic NiAl is a Hume-Rothery P-phase electron compound with

a valence electron-to-atom ratio of 3/2. As a result, NiAl crystallizes with a primitive cubic

CsCl (cP2, B2) crystal structure which may be described as two interpenetrating primitive

cubic unit cells where Al atoms occupy one sublattice and Ni atoms the second. This is

illustrated in Figure 1. NiAl exists as a single phase ordered intermetallic over the

composition range of 45 to 60 at.% Ni at 1000 K and has the highest melting temperature

of any compound in the Ni-Al binary system although the melting point of the

stoichiometric compound (i.e., Ni-50 at.%Al) is in dispute. For example, the phase
diagram of Singleton et al. [45] (Figure 2) indicates that stoichiometric NiAl melts

congruently at 1911 K. More recent evaluations, however, place the melting temperature

of stoichiometric NiAl near 1955 K [46]. It has been suggested that the lower value might

be attributed to the steep drop-off in melting temperature with deviation from stoichiometry

or to unintentional contamination by ternary elements [7]. In addition, NiAl exhibits a high

degree of thermodynamic stability, as indicated by it's large negative heat of formation

(approximately -72 kJ/mol) [47].

Lattice Parameter. Density and Defect Structures

Lattice parameter and density have been investigated thoroughly and have been used
to deduce the types of defect structures occurring in the NiAl lattice [48-50]. Nickel, being
smaller and heavier than Al, should cause a reduction in the lattice parameter and an
increase in the density when it is substituted for Al. This is consistent with observations
for alloys containing more than 50 at% Ni (Figure 3). In Al-rich alloys, however, both
the lattice parameter and the density decrease with increasing Al content and the decrease is
more rapid than would be expected by replacement of Ni atoms by Al. This behavior is
rationalized by the creation of vacancies on the Ni-sublattice rather than by substitutional

defects as observed in Ni-rich alloys. In addition to influencing the lattice parameter and
density, the defect structures induced due to deviations from stoichiometry also
dramatically influence the mechanical behavior. For example, Vedula and Khadkikar [50]
have shown that the yield strength shows a minimum at the stoichiometric composition.
Furthermore, Hahn and Vedula [9] have shown that deviations of less than 1% from

stoichiometry result in brittle behavior at room temperature as well as an increase in the
BDTT. The influence of stoichiometry on the yield stress of NiAl is described in more

detail in the next section. In all cases, the yield stress decreases with increasing
temperature. The behavior of near-stoichiometric polycrystals resembles that of soft-
oriented single crystals while the strengths of off-stoichiometric alloys approach that of
hard-oriented single crystals. Interestingly, the increases in 0.2% offset yield stress, a0o2,

with deviation from stoichiometry are not equivalent on both sides of the stoichiometric
composition [50]. For Ni-rich alloys, for example, the hardening rate was shown to be
120 MPa/at.% while in Al-rich alloys the hardening rate was approximately 350 MPa/at%

[18]. The greater hardening rate on the Al-rich side of stoichiometry suggests that Ni

vacancies provide more resistance to dislocation motion than antisite atoms. These strength

increases, however, become irrelevant above 1000 K where stoichiometric NiAl becomes

stronger than nonstoichiometric compositions [51,52] primarily due to diffusional


In addition to the constitutional vacancies described above, another type of vacancy

defect can exist in NiAl. These are thermal vacancies which can be introduced by rapid
quenching from elevated temperatures. Bowman et al. [53] have shown that a 50-fold

increase in cooling rate from temperatures above 1000 K can increase the compressive yield
stress by almost 30 percent for near-stoichiometric binary NiAl. However, when the

material is microalloyed with Zr, the dependence of strength on cooling rate disappears.
Similarly, Nagpal and Baker [54] have shown that deviations from stoichiometry reduce

the sensitivity of the binary alloy to cooling rate. Cooling rate has been shown to similarly
influence the yield stress of NiAl single crystals [55].

Elastic Properties

The elastic behavior of NiAl has also been studied in some detail and has been

shown to vary with processing technique and temperature. For example, Rusovic and

Warlimont [56] have summarized the single-crystal elastic constants for NiAl as a function

of temperature, cooling rate and stoichiometry showing the overall elastic properties of
NiAl to be anisotropic with an anisotropy factor, E0o/El110, close to 3.3 [57] and showing

a mild temperature dependence but a strong stoichiometry dependence. This is illustrated in

Figure 4a along with the recent results of Walston and Darolia [46] which shows the single

crystal dynamic Young's moduli, E, for near-stoichiometric NiAl for a variety of
orientations. It has also been shown that minor alloying additions have relatively little
influence on the dynamic Young's modulus of <001> single crystals [46]. In polycrystals,

Young's modulus is relatively insensitive to stoichiometry, but is very dependent on

processing technique [58-60] (Figure 4b); extruded materials exhibit higher moduli and

different temperature dependencies than conventional cast and homogenized ingots or hot-

pressed prealloyed powders. This can be rationalized, in comparison to single crystals, in

terms of the crystallographic texture that develops during processing. Extruded NiAl-based

materials commonly exhibit a preferred <111> orientation [61,62], whereas cast or hot-

pressed materials are not expected to exhibit a strong preferred orientation. As a result, cast
or hot-pressed materials have lower moduli in comparison to the higher moduli observed in
<11 >-textured material.

The remaining physical properties have not been characterized to the same degree as
the lattice parameter, density or the elastic modulus. However, it has been reported that

alloying NiAl with Ti and Re significantly reduces the thermal conductivity, whereas

additions of 2.5 at.% Hf had minor effects, decreasing the thermal conductivity of NiAl
single-crystals by only 15% [46]. These and other properties have been reviewed recently


Flow and Fracture Behavior

Slip Systems

The operative slip systems in NiAl single crystals and polycrystals are described in

detail in recent review articles by Miracle [44] and Noebe et al. [7,18]. NiAl single crystals

exhibit two different types of slip behavior depending upon crystal orientation. In NiAl,

the shortest translation vector that will maintain the B2 structure is that along the cube edge
(i.e., in the <100> crystallographic direction). For single crystals in "soft" orientations and

in polycrystals, the dominant slip vector is <001>. However, if the loading direction is
along [001], the "hard" orientation, then the operative slip vector at low and intermediate

temperatures is <111>, and at elevated temperatures is a combination of <110> and <100>.

Soft orientations include all non-<001> loading directions where <100> slip dominates.
Orientations near [001] are considered hard because <001> Burgers vectors have a zero or

near-zero resolved shear stress resulting in the operation of alternative slip systems at very
large yield stresses.
Early slip system determinations between 300 and 1273 K were completed by Ball
and Smallman [51,63] who identified a <001>( 110) slip system in all soft oriented single
crystals. In addition, they also observed cross slip or pencil glide on orthogonal ( 110)
planes. Shortly afterward, Wasilewski et al. [64] reported duplex cube slip, <001>( 100),
in [110] single crystals. This observation was later confirmed by Field et al. [65]. Cube
slip was also reported by Loretto and Wasilewski [66] in [112] crystals deformed between
77 and 1053 K. Only <001> slip is observed in soft oriented single crystals due to the
nondissociated, compact structure of the <001> dislocation core [67] making <001>
dislocations much more mobile than those with other slip vectors. In addition, it has been
observed that NiAl in soft orientations deforms by <001> slip on either (100) or (110)
slip planes in accordance with Schmid's law whereas for hard orientations, Schmid's law
fails [68]. In hard oriented single crystals, <100> slip does not occur because the resolved
shear stress for <100> slip approaches zero. As a result, deformation occurs by non-
<001> dislocations giving rise to elevated yield stresses at low temperatures [64] and
enhanced creep strengths at elevated temperatures [69].
Deformation in polycrystals occurs in accordance with that in soft-oriented single
crystals. Investigators have reported the operation of <001> 110) and <001>{ 100) slip
systems [53,70,71]. Isolated dislocation segments with non-<001> Burgers vectors have
been identified in as-extruded NiAl [72,73]; their presence is attributed to interactions
between <001> dislocations due to the extensive deformation that occurs during the
extrusion process [74]. The operation of <100> slip vectors on planes other than (001)
and (011) has been reported [75,76] under conditions of constrained flow, but are not a
common aspect of the deformation of NiAl.

Yield Strength

Similar to BCC transition metals, the yield and flow behavior of NiAl is extremely

sensitive to temperature and composition. At low temperatures, the yield stress exhibits a

strong temperature dependence. This temperature dependence is attributed to a high Peierls

stress. At intermediate temperatures, a yield stress plateau is observed where the yield

stress is mildly temperature dependent Finally at high temperatures, the yield stress again

drops with temperature. Some results for single crystal and polycrystalline NiAl are shown
in Figure 5. As shown in Figure 5a, the yield stress of single crystals exhibits a strong

dependence on orientation. The yield stress versus temperature curves for soft oriented

single crystals is similar to those observed for low yield strength polycrystalline NiAl.

At low temperatures, hard oriented single crystals, on the other hand, exhibit yield

stresses several times larger than those for other orientations and exhibit a lower
dependence on temperature. When the temperature is increased above approximately 600

K, however, the yield stresses of these crystals become extremely temperature dependent

exhibiting a sharp decrease in yield stress over a narrow range of temperatures. In this

regime, the slip vector has been shown to change from <111> to <001> and <011> [77].

Finally, at temperatures exceeding 1000 K, bulk diffusional processes dominate resulting

in yield strengths similar to those observed in soft-oriented single crystals and in

polycrystalline NiAl.

Ductility and Fracture

Tensile behavior

Single crystals of nominally stoichiometric NiAl exhibit different behaviors

depending on crystallographic orientation. Hard-oriented single crystals exhibit essentially

zero plastic strain to failure in tension at room temperature but undergo a sharp BDTT at

temperatures near 600 K [78,79]. Similarly, soft-oriented single crystals also exhibit a

sharp BDTT ranging from 475 to 525 K [6,80]. In hard-oriented crystals, the BDTT
corresponds to the temperature where the steep decrease in yield stress begins with

increasing temperature. This decrease has been attributed to a change from <111> slip to
climb of <100> and <110> dislocations. Initially, it was reported that soft-oriented single

crystals also exhibited extremely low plastic strains to failure (on the order of one percent)
at room temperature [68,80]. Slightly above the BDTT, anomalously large tensile
elongations (greater than 100 percent) have been reported for soft-oriented single crystals
while at even higher temperatures the ductility decreased to approximately 45 percent

[6,78,80,81]. Takasugi et al. [80] attributed this increase to a balance between work
hardening caused by glide and relaxation processes due to climb resulting in a large
resistance to necking.

More recently, however, tensile elongations approaching seven percent have been
measured in soft-oriented binary single crystals of low interstitial high purity NiAl [19],
conventional purity NiAl [14,15,28] and in nearly stoichiometric crystals doped with

approximately 1000 appm of Fe, Mo or Ga [16,26]. In the conventional purity material,

the dramatic increase in ductility occurred after rapid cooling from elevated temperatures
(1573 K) whereas the increased ductility in the low interstitial, high purity material was not
dependent on heat treatment and cooling rate. The influence of the ternary dopants is
illustrated in Figure 6. Interestingly, the ductility passes through a maximum at small

alloying additions and the benefits of doping vanish as the dopant level exceeds 0.5 at%.

It has been suggested that the ductilizing effect is due to the gettering of interstitials,

although, the real reasons for this behavior remain unknown. Interestingly, similar
alloying schemes in polycrystalline alloys have been unsuccessful [82,83]. This is not
really surprising since deformation occurs by <001> slip with only three independent slip
systems available for deformation [63], rather than the five required for extensive uniform
deformation of a polycrystal, leaving little room for significant room temperature ductility,

independent of the other factors.

Fracture toughness
Room temperature fracture toughness of notched NiAl bend samples has been
reported to be in the range 7 to 12 MPa-Im when the notch is cut normal to the <100>
direction and in the range 4 to 6 MPa4-vm when the notch is cut normal to the <110>

direction [84,85]. In notched polycrystalline bend samples of nominally stoichiometric
single phase NiAl, fracture toughness values have been observed to be independent of

grain size, stoichiometry for Ni-rich alloys, and processing technique [86-89] with
measured values in the range 4 to 7 MPa .
As mentioned previously however, Hack and co-workers [14,15,28] recently

observed that the fracture toughness of commercial purity single crystals is extremely
sensitive to heat treatment and cooling rate. In double cantilever beam specimens with the
notch plane perpendicular to the <110> direction, single crystals rapidly cooled to room
temperature from 1573 K exhibited fracture toughness values of nearly 16 MPa- 'm.
When rapidly cooled specimens were subsequently re-annealed at 473 K and slowly cooled
to room temperature, however, the fracture toughness dropped to 3 MPa'm. Comparable
heat treatments had no influence on low-interstitial, high purity single crystals tested in
four-point bending with the crack plane normal to <100>. There, fracture toughness

values in the range 10 to 12 MPa-mi were observed independent of heat treatment [19].

More experiments using a miniaturized disk bend method [20,90] indicate that low-
interstitial NiAl has an intrinsically larger room temperature ductility and fracture toughness

than commercial purity material.

Influence of Microalloving Additions and Impurities

Finally, substitutional and interstitial elements appear to significantly influence the
yield and flow behavior of NiAl. For polycrystals, abundant solid-solution alloying data
exist. Some of the ternary and quaternary additions have included Be, B, C, Cr, Cu, Fe,
Ga, La, Mo, N, Nb, Mo+Ti, V, Nb and Y to name a few (reviewed in references

[6,18,44,91,92]). In all cases, the flow stress was enhanced by the presence of solute and
the hardening rate was generally shown to be dependent on solute size. This is illustrated
in Figure 7.
In single crystals, a study [16] of the influence of Fe, Ga and Mo on the yield

strength of <110> oriented single crystals showed Mo to be a potent solid solution
strengthener with very limited solubility, while Ga exhibited a mild strengthening effect and
Fe, at levels of less than 1 at.%, reduced the yield stress. In addition, recent studies
indicated that the critical resolved shear stress in low-interstitial high purity single crystals
is significantly lower than that for commercial purity material [55].
As discussed above, microalloying additions of Fe, Ga and Mo in the 0.1 to 0.2

at.% range consistently increase the room--temperature tensile ductility of soft-oriented
single crystal NiAl. The mechanisms) for this increase in ductility have not been
determined, although it is speculated that it is the result of gettering of interstitials [16].


Static Strain Aging (SSA)

The term strain aging characterizes a time-dependent strengthening or hardening
process resulting from elastic interactions of solute atoms with strain fields of dislocations
in plastically deformed metals and alloys [93]. Strain aging is most common in alloys
containing interstitial or substitutional solute atoms capable of segregating to and pinning
dislocations. The aging reactions can occur in either static or dynamic modes depending

upon whether they occur prior to or during plastic deformation. Static strain aging (SSA)
typically occurs in metals and alloys following restraining, unloading (either partially or
fully), aging for a prescribed time and then reloading at the same strain rate as the prestrain.

A schematic illustration of SSA is provided in Figure 8. SSA is typically manifested by an

increase in yield stress or flow stress following aging and the return of a sharp yield point
in the deformed alloy [94].

Mechanisms of SSA
In BCC metals, SSA can typically be separated into four processes: the Snoek

effect, Cottrell locking, Suzuki locking and precipitate formation. Each mechanism is
described below.
The Snoek effect The Snoek effect is a strain-induced ordering of interstitial solute

atoms around dislocations [95-97]. In the BCC lattice, interstitial atoms typically occupy
octahedral sites at the center of cube edges and cube faces (Figure 9). However, the
interstitial atoms are larger than the space available for them in an octahedral site. For
example, an atom at position a (Figure 9a) will cause the substitutional atoms A and B to be
displaced in the z direction. If four of the octahedral positions lying parallel to the z axis

were to become occupied by interstitial atoms as illustrated in Figure 9b, then the unit cell
would become elongated in the z direction and would assume a tetragonal shape. In the
absence of an applied stress, a statistically equal number of interstitials will occupy sites
parallel to each of the x, y and z axes. Thus, the unit cell remains cubic. The application of
an external stress in the z direction, for example, causes the interstitial sites parallel to the z

axis to enlarge while the openings perpendicular to the z axis decrease in size making it

energetically more favorable for atoms in position a of Figure 9c to jump to position b.
Schoeck and Seeger [97] have critically evaluated this mechanism and have concluded that,

since no long range diffusion is required, this process occurs very rapidly and is normally
completed within the time interval of one atomic jump of the species responsible for
pinning. In addition, Nakada and Keh [98] have indicated that the apparent intercept of
yield point return data plotted as Aou versus t2/3 is positive when Snoek ordering occurs

prior to Cottrell atmosphere formation. Rosinger [42] has found that the activation energy

for this process in ferritic steels is approximately 60 kJ mol-1.

Cottrell locking. Cottrell locking (also known as Cottrell atmosphere formation),
which has been treated in detail by Cottrell and Bilby [43], involves the time- and

temperature-dependent growth of solute atmospheres near dislocations. During atmosphere
formation, elastic interactions between solute atoms and the strain fields of dislocations
create a driving force for the diffusion of interstitial solutes toward the strain field. This
mechanism results in a lowering of the total energy of the system and effectively locks the
dislocations in the sites that they occupy during the process. The end result is an increase
in the stress required to move a dislocation. Important characteristics of Cottrell locking are
an aging time dependence for yield point return which follows a 2/3 power law and an
activation energy for yield point return which is equivalent to the migration energy for the

solute causing Cottrell locking. Though a 2/3 aging time dependence has been reported for
several metals and alloys (for example, see reference [94]), the 2/3 power law often fails
[99]. In ferritic steel, the activation energy for atmosphere formation is approximately 90
kJ mol-1 [42], which is equivalent to the activation energy for volume diffusion of
interstitial solutes.

Suzuki locking. Suzuki locking has its origin in the chemical interaction between
solute atoms and stacking faults [100]. This mechanism has been commonly observed in

superalloys and is only expected to be significant in metals exhibiting low stacking fault

energies in which stacking fault widths are large, i.e., FCC and HCP metals and alloys.
As NiAl exhibits a high stacking fault energy and no evidence of stacking faults before or
after deformation, this mechanism inapplicable.
Precipitate formation. Precipitate formation only occurs when the metal is

supersaturated with solute atoms. Strain aging occurs when interstitial or substitutional
atoms, or compounds composed of those solutes (e.g., carbides, nitrides, oxides, borides,
etc.) precipitate on dislocations during aging, effectively pinning them.

Yield point return
Yield points that repeatedly return after aging are associated with the formation of
solute atom atmospheres around dislocations. Mobile dislocations that were once active
during deformation prior to unloading are pinned as a result of aging. For this pinning to

occur, solute atoms must diffuse through the lattice to accumulate around dislocations. As
a result, the reappearance of the yield point is a function of time which depends on the
temperature since diffusion is a temperature-dependent function. Rosinger et al. [101] have
shown that two plateaus are typically observed when the increment of yield stress, Ao,
observed after SSA is plotted versus the aging time ta (Figure 10). The first plateau, which
occurs at shorter ta's, is associated with Snoek strain aging while the second plateau,
which occurs at longer ta's, is associated with Cottrell strain aging. The time

corresponding to the first inflection point between the two stages has been shown to

correspond to the time required for a single interstitial solute atom to undergo a single jump
[101]. Easy identification of each regime can be made by plotting Ao versus ta2/3. The
decrease in Ao observed after the second plateau is associated with saturation which occurs

when the interstitial atoms have migrated to dislocations in sufficient numbers to either
relieve the strain energy induced by dislocations in the lattice or to set up concentration

gradients restricting the further migration of solute [43]. Saturation also occurs when the

migration of interstitial atoms to dislocations results in the depletion of interstitial atoms in

the surrounding lattice. At even greater times, the concentration of solute atoms about the

dislocations may exceed the solid solubility limit of the solute in the solvent metal at that
temperature. As a result, precipitates may form resulting in a decrease in Ao, or softening


Dynamic Strain Aging (DSA)

Dynamic strain aging is a phenomenon exhibited by many metals and alloys [103].
It is the result of interactions between diffusing solute atoms and mobile dislocations during

plastic deformation. This process tends to occur over a wide temperature range which is

dependent upon strain rate 6.
Dynamic strain aging is manifested by: the appearance of serrations, load drops,

jerkiness or other discontinuities in the stress-strain curves obtained in constant-extension-
rate tensile or compression tests; peaks or plateaus in the variation of flow stress, work-
hardening rate, O=Ao/Ae, and Hall-Petch slope, kg, with temperature, T, and minima in the

variation of ductility and strain-rate sensitivity, s=Aa/Aln or n=Alno/Aln with T; and

low or negative values of s and n in the temperature region of DSA [104]. Some of these
manifestations are illustrated in Figure 11. These phenomena are associated with the
dynamic formation and migration of solute atmospheres around dislocations during
Serrated flow curves

Serrations, load drops or jerkiness in the stress-strain curves obtained in constant-
extension-rate tensile or compression tests are perhaps the best known manifestation of
DSA. In creep tests under constant load or stress, or in constant-loading rate tests, DSA is
manifested as staircase creep whereby sudden bursts of plastic strain periodically occur

resulting in staircase-like creep curves [104].
Johnston [105] and Hahn [21] have proposed that load drops or yield point
phenomena are related to an initially low mobile dislocation density and a low dislocation-

velocity stress sensitivity. In reference to the latter, Gilman and Johnston [106,107] and
Stein and Low [108] have demonstrated that the dislocation velocity V is related to the

resolved shear stress according to the equation:

V =, Dm (1)

where r is the applied resolved shear stress and D and m are material properties. In tensile
or compression tests, specimens are deformed at constant nominal strain rates E such that

S= te + ip (2)

where ip is the plastic strain rate in the specimen and ie is the elastic strain rate of the

specimen and the machine. For a load drop to occur during deformation, the plastic strain

rate must exceed the imposed strain rate. Assuming that the plastic strain rate under

dislocation glide obeys the Orowan equation

ep = PmbV (3)

where Pm is the mobile dislocation density, V is the average dislocation velocity and b is

the Burgers vector, load drops or serrations can occur when there is an instantaneous
increase in either Pm V, or both. It is commonly accepted that the serrations observed

during DSA are the result of repeated locking and unlocking of dislocations from solute
atmospheres which results in a sudden increase in Pm. The solute atoms, however,

repeatedly diffuse to form atmospheres resulting in repeated yielding (i.e., serrations).

DSA is, therefore, expected to be more pronounced when V is equal to the drift velocity of

solutes in the stress field of a dislocation. Furthermore, DSA can only occur in a range of

intermediate strain rates and temperatures. At low temperatures and high strain rates, the

solute velocity will be too small (compared to the dislocation velocity) to cause strain aging.
At high temperatures and low strain rates, any solute atmospheres that form will be able to

keep up with the dislocation velocity and the serrated flow will again disappear.
Types of serrations

Serrated flow can exhibit itself in a number of forms. Five types of serrations

resulting from DSA have been identified (Figure 12) and were summarized by Rodriguez

[104] as follows:

1. Type A serrations arise from periodic Liiders bands initiating at one end of the

specimen and propagating along its gauge length. They are considered to be

"locking" serrations and are characterized by an abrupt rise followed by a drop

below the general level of the stress-strain curve. They occur in the low T, high i

portion of the DSA regime.

2. Type B serrations are oscillations about the general level of the stress strain curve.

They also arise from the formation of Liiders bands, however, these bands do not

propagate and are normally observed at higher temperatures and lower strain rates

than type A bands; in other words they occur when there is an increased diffusion

rate of solute atoms.

3. Type C serrations, which occur at higher temperatures and lower strain rates than

types A and B, are yield drops below the general level of the flow curve. These

serrations are considered to be the result of dislocation unlocking.

4. Type D serrations are plateaus in stress-strain curves due to the propagation of

deformation bands with no work hardening or strain gradient ahead of the moving

front. These serrations can occur alone or with type B serrations.

5. Type E serrations develop from type A serrations at high strains. They resemble

type A serrations but exhibit little or no work-hardening during band propagation.

As noted by Reed-Hill [109], regardless of the type of serrations observed, serrated flow is

discontinuous and involves the immobilization of dislocations.
Often associated with the occurrence of serrations is a critical strain, ec, for the

onset of serrated yielding. This strain is associated with the buildup of a sufficient

dislocation density for serrated flow to occur and its value is dependent on both T and e.
At higher strain rates and lower temperatures, ec typically increases with increasing strain

rate and decreasing temperature [104]. However, in regions of higher temperature and
lower strain rate, ec exhibits inverse behavior and increases with increasing temperature and

decreasing strain rate [104]. This phenomenon is known as the inverse Portevin-Le

Chatelier effect and is normally associated with type C serrations. Even though the
occurrence of this "inverse" phenomenon has been substantiated, the reasons for its

occurrence are unknown.

Theories of DSA

Several theories have been advanced to explain DSA phenomena and they typically

fall into three main categories: solute drag models based on the model of Cottrell, "static"
aging models based on the work of McCormick [110,111] and van den Beukel [112], and
dislocation interaction models based on the work of Kocks [113,114]. These theories will
be discussed below along with a theory recently advanced by Reed-Hill et al. [1-

The Cottrell model

The Cottrell model [43,117-119] considers DSA in terms of its most visible
manifestation, the Portevin-Le Chatelier effect. This model assumes that serrated flow

begins when the velocity of a dislocation exceeds the critical drag stress exerted by a solute
atmosphere. The critical velocity, Vc, is given by:

Vc=4D/I (4)
where D is the solute diffusion coefficient and I the effective radius of the solute
atmosphere. Above this velocity, the stress decreases with an increase in dislocation
velocity making it logical to assume that Vc represents a critical condition for the appearance

of serrations on a stress-strain curve.
If the Orowan equation is assumed valid, then Vc can be expressed as:

Vc = e/4bpm (5)

or the equation may be rewritten as:

t = Vc4bpm = 4D1bpm/t (6)
where i is the applied strain rate, 0 is a Schmid orientation factor, b is the Burgers vector,

pm is the mobile dislocation density, D the solute diffusion coefficient, and I the radius of

the dislocation atmosphere. The dislocation density is normally considered to be a function
of the strain, e, such that:

Pm = NE (7)
where N and P are constants. In the case of substitutional alloys, diffusion occurs by the

vacancy mechanism such that the diffusion coefficient, D, is given by the relation:

D= Doexp[(Qm+f = DC exp[ ] (8)


+ Jk- (9)

and represents the thermal equilibrium concentration of vacancies, Do is the frequency
factor, Qm represents the activation energy for the movement of vacancies and Qf is the
work to form a vacancy. During plastic deformation, it is generally agreed that the total
vacancy concentration increases with strain according the relation:

Cv=CKem (10)

where Cv is the vacancy concentration and K and m are constants. Based upon this, the
diffusion equation can be rewritten as:

D= DoKmexp( -) (11)

Substituting this relation back into the original Orowan equation yields:

p+m &expQm /kT
4Cb ND(12)

This equation suggests several experiments for determining the parameters P3+m and Qm,

but fails when it comes to predicting the critical strain itself [111].
The McCormick model
The inability of the Cottrell model to predict ec prompted McCormick to propose an
alternative model [110,111]. The basis for this model is the assumption that dislocation
movement is discontinuous as assumed in the dislocation arrest theory of Sleeswyk [120]
who proposed that during the time when dislocations wait at obstacles, mobile solute atoms

can be drawn to them resulting in strain aging. According to this model the dislocation
velocity, v, can be expressed as:

v = (13)
tw + tf

where L is the average spacing between obstacles, tf the mean flight time between

obstacles, and tw the mean waiting time at an obstacle. In most cases, the average
dislocation velocity is determined by the arrest time such that:

v= L (14)

McCormick defined the aging time, ta, as the time required to lock moving dislocations and
proposed that serrations on stress-strain curves occur when ta =s tw. When ta > tw, upon the

onset of plastic deformation dislocations arrested at obstacles will not be locked and flow

will be continuous. During straining, however, ta can decrease due to vacancy production
while tw increases due to dislocation multiplication such that a critical strain will exist where

ta= tw.
To evaluate ta, McCormick assumes a Cottrell-Bilby [43] 2/3 power law where:

2/3 kT2
ta aCo) 3UmD (15)

and CI is the solute concentration at the dislocation required to lock it, Co is the solute
concentration in the alloy, Um is the binding energy between the solute and the dislocation,
D is the solute diffusion coefficient and a is a constant equal to about 3.

Assuming that both dislocation density and vacancy concentration are functions of
strain as expressed above, McCormick arrives at an expression for the critical strain in
substitutional alloys of the form:

SC 3/2 kTbexpQm/kt (16)
c CaCo ) 3ONKUmDoL

This expression can be simplified for BCC interstitial alloy systems where diffusion occurs
independently of vacancy concentration. In this case, the equation becomes:

P C1I %11 kTbexp(Q/kT)
C -aC0oj NLUmDo

where Do is the interstitial diffusion frequency factor and Q is the activation energy for the
diffusion of interstitial solute atoms.
This theory has advantages in that it is able to accurately predict the critical strain for
the appearance of serrated yielding as well as the temperature and strain rate dependence of
e~. This theory, however, does not take into account other aspects of DSA such as the

yield stress plateau, abnormal and rate dependent work hardening, flow stress transients
that occur on changing the strain rate and the development of negative strain rate
The van den Beukel model
In an attempt to create a more universal theory of DSA, van den Beukel [112,121],
starting with Sleeswyk's hypothesis [120], developed a theory based on the idea that a
moving dislocation can be subject to strain-aging during its waiting time at an obstacle.
The major advance in this model is the inclusion of the activation enthalpy, H, in the
thermally activated strain rate equation:

e= oexp-H) (18)

where i is the strain rate, io is a constant and k and T have their usual meanings. It is

assumed in this model that the activation enthalpy is a function of the effective stress, a ,
and the local solute concentration at the dislocation, C. The value of C is a function of the
time that the dislocation waits at an obstacle, tw, and the rate of solute drift to the
dislocation which is, in turn, a function of the diffusion coefficient or.
C=-f(Dtw) (19)

He further shows the quantity Dtw to be a function of the strain, strain rate and temperature
and makes use of a t2/3 relation between the concentration and the time formulated by
Friedel [122] assuming that Dtw is small. Thus we have:

C = Co + (KDtw)2/3 (20)
where Co is the nominal solute concentration of the alloy and K is given by:

K = 3Um o)2 (21)
In this equation Um represents the binding energy between the solute and the dislocation.
By further considering H to be a function of the local solute concentration at the
dislocation and the effective stress, van den Beukel obtained an activation enthalpy equation
of the form:

H=-T*V +T d (22)
a aC dT
where V* is the activation volume. The first term on the right hand side of this equation
assumes that a single thermally activated mechanism controls the flow stress [123] and is
considered to give the activation enthalpy in the presence of DSA. The second term on the
right hand side of the equation is considered to represent the DSA component of the
activation enthalpy.
He also showed that the strain rate dependence of the flow stress could be
expressed as:

u kT DH dC
= -- Dt ) (23)
ai tV 5E d(Dtw)

where the first term on the right side is the normal strain rate dependence in the absence of
diffusion and the second is due to DSA.
Finally, van den Beukel obtained a relationship for the strain rate sensitivity by
assuming that, upon an increase in strain rate from a low to a high rate, the change in
activation enthalpy is given by:

AH = -kT. In (24)

which allowed him to write:

Ao kT 1 aH dC
A(Ini) V V M7aCdln )

In this equation, the first term on the right side corresponds to the strain rate sensitivity in
the absence of DSA while the second term is the component due to DSA. It can be seen
from these equations that any solute mobility makes a negative contribution to the total

strain rate sensitivity and that this contribution increases with increasing strain. When the
total strain rate sensitivity becomes negative, plastic flow becomes unstable and serrated

flow is observed.
Van den Beukel further extends his investigation to show that the temperature
dependence of the flow stress as well as the work hardening can be expressed by two
additive terms; one friction term and one forest strength term. The friction term is assumed
to be independent of strain but is said to be affected by aging while the forest strength term

is assumed to be independent of solute concentration except through bulk material
properties such as stacking fault energy [121] (i.e., the friction term is affected by DSA

while the forest term is not).

Although this model represents an improvement over the earlier models of Cottrell
and McCormick, it does have shortcomings:
1. Quantitative predictions depend upon a detailed knowledge of the variation of

activation enthalpy with solute concentration near a dislocation. This data is

generally not available.
2. The mathematical approach of this model makes visualization of the physics of DSA
3. This theory deals only with long range diffusion of solute atoms to dislocations.

The Reed-Hill model
The shortcomings of the Cottrell, McCormick and van den Beukel theories led
Reed-Hill [1,3] to propose a new theory of DSA. In this model it is assumed that the total
flow stress at can be written as follows:

ct = O + ODSA (26)
where a represents the stress in the absence of DSA and ODSA represents the stress
associated with DSA. It is further shown that o also has two parts, the internal stress, GE,
and the thermally activated or effective stress, o*. Thus the total stress can be written as:

Ot = GE + + DSA (27)

In BCC interstitial systems, however, it has often been demonstrated that ODSA may also

consist of two additive parts due to Snoek and Cottrell aging such that:

GDSA=Gsn+Ocot (28)
where asn represents the Snoek component and acot represents the Cottrell component. He

also developed a new method to evaluate the internal stress [115,124] and demonstrated
that the effective stress can be approximated by a power law of the form:

S kT/H
a =aF o -O (29)

where Co is the effective stress at 0 K, i is the nominal strain rate, io and HO are material

constants with units of energy, and k is Boltzmann's constant.
The component of the flow stress attributed to DSA is given by the equation:

ODSA= sm 1-exp [+.m} 1-exp 2/ (30)

where smasx and cmx are the isothermal maximum obtainable magnitudes of the Snoek

and Cottrell contributions to the flow stress, respectively, and cs and tc are the relaxation

times for Snoek and Cottrell aging, respectively. The decrease in max and max due

to dynamic recovery was achieved by multiplying the DSA component by the value:

exp(Bx(T-To)) (31)

where To represents a temperature below the DSA regime and B is defined by:

B = x In (32)

It has been shown that this theory can model the temperature dependence of the flow stress
and, when properly differentiated, the temperature dependence of the strain rate sensitivity
in a wide variety of alloy systems [1-3,116].
Other models

Other models have been developed by Kocks and co-workers
[113,114,121,125,126], and by Estrin and co-workers [127-130]. Kocks' model
suggested that mobile dislocations are temporarily arrested at forest dislocations; solute
atmospheres form on the forest dislocations and then drain by pipe diffusion from the
forest dislocations to the mobile dislocations during their waiting times. Since this model
relies on pipe diffusion rather than bulk diffusion, it allows for rapid atmosphere formation
without the need for enhancement by vacancies. As a result, solute atmospheres need only
pin portions of the dislocation line at the forest dislocation junctions rather than entire
dislocation segments and the obstacle strength increases with waiting time. Like the van
den Beukel model [112,121] described above, this model also assumes that the flow stress
is composed of two additive terms but that the forest term rather than the friction term is
influenced by aging. This assumption implies that the strength of the dislocation-
dislocation interactions is altered. The model of Estrin and co-workers [127-130], is a
further refinement of the Penning model [131] which is based on an N-shaped curve

representing the strain rate dependence of the flow stress. Although both of these models
have their merits and have been successfully used to model aspects of DSA
[29,114,129,132], they will not be addressed here.

Static and Dynamic Strain Aging in Ordered Alloys

As mentioned above, most metals and alloys are subject to some sort of strain aging
phenomenon [103]. And although several intermetallics have been shown to exhibit
manifestations of strain aging (e.g., FeAl [133,134], AgMg [33], CuZn [32], Ni3Fe
[135], NiAl [14,15,22,24-29,136-139], Al3Ti-X [140-143], and TiAl [144,145]), the

influence of strain aging on the mechanical behavior of intermetallic alloys and the species
and mechanisms responsible for this behavior (i.e. Snoek effect, Cottrell aging, etc.) have,
until recently, been essentially ignored.
A perfect example involves the L12 intermetallic alloys based on A13Ti. Yield point

plateaus and serrated flow stress curves have often been reported for these compounds
[25,140-142]. Lerf and Morris [142], however, initially elected to attribute the occurrence
of serrations at intermediate temperatures to the repeated dissociation of <110>
superdislocations into pairs of mobile superpartials with higher mobilities. As Potez et al.
[141] pointed out, however, such an explanation would imply strain softening rather than
strain rate softening via dynamic strain aging and, consequently, is insufficient to explain
the related strain anomalies observed in these alloys. Potez et al. [141] indicated that
oxygen might be the species responsible for DSA in A13Ti-Cu based on analysis of two
Al3Ti-Cu alloys with the same bulk compositions but different oxygen concentrations.

However, it was later shown [146-149] that DSA in these alloys was actually the result of
the precipitation of complex Al2Ti-based precipitates on dislocations.
For the other intermetallic alloys mentioned above, the authors have often alluded to
the occurrence of strain aging and speculated on the nature of the species responsible (e.g.,
N and 0 in AgMg [33]); however, a detailed study of the species responsible and the
resulting strain-aging mechanism in ordered alloys is lacking.

Ni-- ,

*-I I
... .... I .


Figure 1. The CsCl (cP2, B2) crystal structure of NiAl illustrating how this structure
is composed of interpenetrating simple cubic sublattices of Ni and Al atoms.

Weight Percent Nickel
0 10 20 30 40 SO 80 70 80 0 100

1400 AIN! UM
SAtomic Percent Nickel (Ni)

1000 .

0 10 20 30 40 60 10 70 so s0 100
Al Atomic Percent Nickel Ni

Figure 2.

The Ni-Al phase diagram [47].

0.290 ,. 7

E -6
c 0.289

m 0.288

4 0

-j lattice parameter 3

0.286 .**. .... .. 2
40 45 50 55 60 65
Composition (at.% Ni)

Figure 3. Variation in room-temperature lattice parameter and density of NiAl as a
function of stoichiometry (Data compilation from [18]).








350 '
-----Walston an
300 < > Wasilewski,


200 <110>


100 <100>
100 ----

50 I *
200 400 600 800 1000
Temperature (K)

260 .
Moose, 1991
240 "-----... Ni-50.6AI
"-.'" "----..._ (powder extrud

220 -----.
Harmouche and Wolfenden, 1987
200 (powder extruded)

1 80 ... Rusovic and Warl
180- .... Ni-50,
Moose, 1991 cas
160 Ni-50.6AI
(hot pressed prealloyed powder)

140 1 *, .
200 400 600 800 1000
Temperature (K)


imont, 1979

1200 1400

Figure 4. Elastic properties of single crystal and polycrystalline NiAl alloys: (a)
single crystal NiAl as a function of orientation and temperature; (b) polycrystalline NiAl as
a function of processing route and temperature (Data compilation from [7]).

d Darolia, 1993

1200 1400



1600 4 1uuJ a [100o o.maa. Noobe. nd 0.rolla 1tl8
SIa (to0) Pasoe eand N.. ey 1968b
SAI (123) Kim 1000
1400 a 1 [1,10 Lahrm..n. n.eld d ea o ll. 1061
V &+ (1] Lharoman. Field and Deralla 1081
1200 v 0 *[O(110) al and Smallman 1S66a

1000 (123]

800 -

600 -

400 [110

200 0' 0 0

0 200 400 600 800 1000 1200 1400






" 1200

O 1000


, 600


400 -

200 -[ a

0 200 400 600 800 1000 1200 1400
Temperature (K)

Figure 5. The variation in yield stress as a function of temperature for polycrystalline
NiAl and for several different single crystal orientations: (a) single crystals and (b)
polycrystals (Data compilation from [18]).

1 rM


0.5 1.0

-r 7








Atomic Percent of Ternary Element

Figure 6. Room-temperature tensile ductility of NiAl single crystals doped with iron,
gallium or molybdenum [16].

6000 1 I I I '
Interstitial o Sul
5000 -

4000 1 Zr

03000 -

1000 Fe

SB. ..... \B.e.Ga

0.00 0.05 0.10 0.15 0.20
Goldschmidt Radii (nm)

0.25 0.30

Figure 7. Relationship between hardening rate and element size for various alloying
additions in NiAl [7,18]


Figure 8. Schematic illustration of static strain aging.

(a) z




^-- -- -- ---T

'/ 0 interstitial

Figure 9. Tetragonal distortion of the BCC unit cell due to the presence of interstitial
atoms: (a) interstitial solute atoms occupying octahedral positions at the cube edges and
faces; (b) four octahedral positions lying parallel to the z axis are occupied by interstitial
atoms; (c) external stress applied parallel to the z axis causes atoms in position a to shift to
more energetically favorable position b.



Figure 9 -- continued.



Aging Time

Figure 10. Schematic illustration of the variation in the yield stress increment, Aa, with
aging time. The first plateau is due to Snoek strain aging while the second is due to Cottrell
strain aging.



........ flow stress, O

---------- -- --------------Ao/A

I t

m i _I'

,%E ,%E


: A D.E :C


Figure 11. Manifestations of dynamic strain aging (after [106]).

Type E


Se(C) YType C

S Type D

,ce(C) > E (A) or E(B)


Figure 12. Examples of the different types of flow stress serrations observed during
dynamic strain aging (after [106]).



Alloys based on the intermetallic compound NiAl are considered good candidates
for high temperature structural applications due to their excellent oxidation resistance, low
density and high thermal conductivity compared to nickel-base superalloys. In spite of
these advantages, their development into a viable engineering material has been limited by
low tensile ductility (2%) at ambient temperatures. However, recent efforts have resulted

in room temperature tensile elongations of up to 6% for doped NiAl single crystals

containing microalloying additions of Fe, Mo or Ga [16,68]. The mechanism behind this
increased ductility is unknown but may be related to a gettering phenomenon since it is
realized that some substitutional elements and interstitial impurities can cause significant

embrittlement in BCC metals [17,114] and B2 ordered compounds such as NiAl
[7,18,150-152]. Current observations are consistent with this viewpoint. For example,

room temperature tensile elongations of up to 5% have been observed in low interstitial
binary NiAl single crystals [19] and the ductility and fracture strength in biaxial bending of

high purity NiAl was found to be significantly greater than for commercial purity crystals

[20]. Still, the effects of particular substitutional and interstitial elements on the mechanical
behavior of NiAl are relatively unknown and the mechanisms by which various elements
may enhance or hinder tensile ductility still remain a matter of conjecture. Consequently,
the purpose of this chapter is to provide preliminary details concerning the effects of alloy
purity on the mechanical properties of NiAl single crystals. To accomplish this goal, two

stoichiometric NiAl single crystals with differing impurity contents were studied. In

addition, a NiAl alloy intentionally doped with Mo was investigated and a comparison of

the mechanical properties of all three alloys is discussed in terms of relative compositions.


Conventional purity (CP-NiAl) and Mo-doped (NiAl-0.2Mo) NiAl single crystal
slabs, 25 mm x 32 mm x 100 mm, were grown in argon by a modified Bridgman
procedure. The slabs were then homogenized at 1589 K for 48 h in argon and cooled at
-0.213 K/s by back filling the furnace chamber with gaseous argon. A high purity NiAl
(HP-NiAl) single crystal, 25 mm diameter x 120 mm length, was produced by

containerless float zone refining of a vacuum induction melted polycrystalline starter ingot
of NiAl. A description of the equipment and process used to grow the low interstitial NiAl

ingot has been presented in detail elsewhere [19,153].
Chemical analyses were conducted at the NASA-Lewis Research Center by the
following techniques deemed to be the most accurate for the particular elements: Ni and Al

were determined using analytical wet chemistry/titration techniques, Si was determined by
ultraviolet/visible spectrophotometry, and Mo was determined by flame atomic absorption
spectrophotometry. Oxygen, nitrogen, carbon and sulfur contents were determined by

combustion techniques using LECO oxygen/nitrogen and carbon/sulfur determinators.

The results of these analyses on the alloys studied are summarized in Table 1.

The crystals were oriented using the back reflection Laue technique and EDM wire
cut into cylindrical compression specimens parallel to the <123> axis. Specimen

dimensions were 3.0 mm and 6.4 mm for the diameter and height, respectively. Prior to
testing, several CP-NiAI and HP-NiAl specimens were wrapped in Ta foil and annealed in
argon at 1000 K for 2 h followed by either furnace cooling at -0.083 K/s (FC) or water

quenching (WQ) to room temperature. All compression tests were performed on an Instron
Model 1125 load frame at a constant crosshead velocity corresponding to an initial strain
rate of 1.4x104 s'I. Tests were run in air between 300 and 700 K by heating the samples

in a clamshell type resistance furnace. True stress-strain data were calculated from the

load-time plots and yield stresses were determined by the 0.2% offset method.

Table 1. Chemical Compositions of Single Crystals Examined

at.% Impurities (appm)
Alloy Ni Ala Mo I Sib Os Nc Cd Sd
CP-NiAl 50.40.2 49.50.2 -- 0.15 92 <31 169 <7
HP-NiAl 50.20.2 49.80.2 -- 0.06 <27 <31 <36 <13
NiAl-0.2Mo 50.00.2 49.60.2 0.20 0.20 105 21 143 <4
b Analysis performed using analytical wet chemistry/tilration techniques, relative accuracy 1%
Analysis performed on an Ultraviolet/Visible Spectrophotometer, Shimadzu, Model UV-160,
relative accuracy 10%
Analysis performed on a Simultaneous Nitrogen/Oxygen Determinator, LECO Corp., Model TC-
136 or Model TC-436, relative accuracy 10%
Analysis performed on a Simultaneous Carbon/Sulfur Determinator, LECO Corp., Model CS-244,
relative accuracy 10%
Analysis performed on a Flame Atomic Absorption/Emission Spectrophotometer, Perkin Elmer
Model 5000, relative accuracy 5%

Samples for transmission electron microscopy (TEM) were cut from the tested

compression specimens with a low-speed diamond saw and twin jet-electropolished in a

solution of 70% ethanol, 14% distilled water, 10% butylcellusolve, and 6% perchloric acid

at 273 K, 25V, and 0.15 mA. TEM examinations were conducted in JEOL 100C and

JEOL 200CX microscopes operating at accelerating voltages of 120 kV and 200 kV,



Within experimental accuracy, the Ni and the AI+Mo contents of the three alloys are

not significantly different from each other. The major differences between the materials are

the oxygen and carbon contents between the HP-NiAl and the other two materials and the

addition of 0.2 Mo to one of the alloys. The Si contamination in the CP-NiAl and NiAl-

0.2Mo alloys is the result of reaction with the ceramic mold walls during Bridgman growth

while the lower Si level in HP-NiAl is residual contamination from the original induction

melting of the NiAl feed rod used for directional solidification.

Pronounced effects of alloy purity and thermal history (i.e., prior heat treatment and
subsequent cooling rate) on the compressive properties of NiAl were observed and are
summarized in Table 2. First, the as-grown HP-NiAl samples exhibited yield and flow

stresses similar to the as-received CP-NiAl samples but both binary NiAI alloys were much

lower in strength than the as-received NiAl-0.2Mo crystal. Typical room temperature

stress-strain curves for the three alloys are shown in Figure 13. Following identical
treatments (i.e., 1000 K(2h)+FC), both HP-NiAl and CP-NiAl experienced large decreases

in yield strength (i.e., >25%). It is evident from the data in Table 2 that the WHR in as-
received HP-NiAl (13.80.7 MPa/%%p) and CP-NiAl (24.59.6 MPa/%ep) decreased to
nearly equal values (~10.4 and 10.6 MPa/%ep, respectively) after this particular heat

treatment. However, the yield stresses of the HP-NiAl specimens were nearly 30% lower

than those of CP-NiAl. The yield stress and WHR of the Mo-doped NiAl were generally
higher than those of the other alloys and were relatively insensitive to the thermal treatments

employed in this study.

Table 2. Average Room Temperature Compressive Properties for <123> Oriented NiAl
Single Crystals.

Alloy 0.2% Yield Critical Resolved Shear Samples Work Hardening
Stress (MPa) Stress (MPa) Tested Rate* (# Tested)
HP-NiA: 212.6t14.4 96.56.6 13.80.7 (
HP-NiAl: 125.95.7 57.22.6 11 10.41.1 (4)
CP-NiAl: 233.121.0 105.99.5 5 24.59.6 (3)
CP-NiAl: 174.78.9 79.44.0 8 10.63.4 (4)
NiAl-0.2Mo: 349.214.6 158.66.6 10 20.11.2 (4)
NiAl-0.2Mo: 369.211.0 167.75.0 5 19.12 (2)
1000K(2h)+FC _______
FC=urnace cooled at 0.083 K/s
"MPa/%ep at 4% plastic strain.

The temperature dependence of the yield stress for all three alloys in the as-received
condition and the room temperature yield strengths for the annealed materials are presented
in Figure 14 along with the data reported in previous studies of [123] oriented crystals
[154,155]. As expected, the yield strength decreased gradually with increasing
temperature. Figure 14 indicates that the yield stresses observed in other studies were
higher than those observed for annealed HP-NiAl which suggests an influence of both
purity and prior thermal history. It was shown above that low temperature heat treatments
prior to mechanical testing at room temperature resulted in large yield strength drops for the

binary alloys. This observation is confirmed and extended to tests run at elevated
temperatures by the data of Takasugi et al. [155] who homogenized their single crystals at
1323 K for two days followed by furnace cooling at 0.368 K/s. The three NiAl alloys in
this study exhibit slight plateaus in the yield strength vs. temperature data between 500 and
650 K. Similar behavior is frequently observed in BCC metals [17,114] and has been
reported by Kim [154] for [123] oriented Ni-48AI and Takasugi et al. [80,155] for [123]
oriented Ni-50Al single crystals. Considering the dependence of yield stress on

temperature, it is convenient to represent the data in an Arrhenius form (Figure 15), i.e.,
oy,=AeQT, where A is a constant, R is the universal gas constant, Q is the activation

energy, and T is the absolute temperature. In Figure 15, the changes in slope near 500 K
for CP-NiAl, HP-NiAl, and NiA1-0.2Mo imply a change in deformation behavior and may

be associated with the location of the brittle-to-ductile transition temperature (BDTT) [53].
Transmission electron microscopy of the dislocation microstructures revealed no

differences from what has been previously reported. Prior to deformation, the overall
dislocation density was low, consisting of randomly distributed <001>( 110) dislocations
(Figure 16). The NiAl-0.2Mo alloy contained small a-Mo precipitates in the 10 to 30 nm

size range. Following approximately 1% deformation at 300 K, the deformation

microstructure consisted of elongated dislocation cells and randomly distributed dislocation
tangles (Figure 17). Once again only { 110)<001> dislocations were observed.

Faceted voids were observed in all three alloys annealed at 1000 K(2h)+FC (Figure
18) but not in the as-received materials. Similar voids have been reported previously in

rapidly solidified NiAl and in NiAl alloys annealed at lower temperatures after quenching
from high temperatures [156-158]. These voids are attributed to the coalescence of thermal
vacancies during heat treatment. All alloys exhibited <001> slip, consequently, alloy
composition and purity had no effect on slip mode.


Past studies have clearly shown that the yield stress and its temperature dependence
for BCC metals can be significantly altered by either reducing or increasing the impurity
contents or by causing a change in the point defect structures via irradiation and/or heat
treatment [17,114,159,160]. Not surprisingly, this is also the case for NiAl. In agreement
with prior studies of BCC metals, the yield stress of NiAl tends to increase as a function of
interstitial and substitutional impurity contents [7,18], though great care to accurately
measure composition is very rarely taken. Typically in BCC metals, interstitial solutes
form Cottrell and Snoek atmospheres that pin dislocations resulting in increased strengths
while substitutional additions such as Mo and Si result in solid solution hardening by
increasing the frictional resistance of the lattice to dislocation motion [17,43,96,114].

Microalloying additions of Mo, Ga and Fe have been reported to decrease the
BDTT of <110> oriented NiAl single crystals [16,68]. In polycrystals however,
microalloying additions of Ga and Fe have been shown to increase the BDTT by as much
as 250 K [82,83] while Zr additions have been shown to nearly double the BDTT to >1050
K [53,83]. In the case of Zr, the increase in the BDTT has been attributed to the pinning of
extrinsic grain boundary dislocations due to grain boundary segregation of the
microalloying addition [53,83].

Based on the Arrhenius plots, the BDTr for HP-NiA1 and NiAl-0.2Mo appears to
be the same as that for CP-NiAl at approximately 500 K which is consistent with a reported

BDTT of between 473 and 573 K [16,155]. It is important to note that the assessment of
the BDTT in this study is based solely on compressive data; tensile testing should be

performed in order to validate this speculation. The mechanism for the BDTT in "soft"
oriented single crystals is unknown although it has been proposed that the BDTT probably
arises from the thermally activated climb of [001] dislocations [155], similar to what has

been previously proposed for polycrystalline NiAl [53].

Heat treatment temperature and cooling rate are other important variables effecting
the mechanical properties of NiAl. Previous studies have shown that high concentrations
of thermal vacancies can be quenched into NiAl [161] and these vacancies result in
increased yield stresses by providing resistance to the motion of dislocations

[7,18,53,54,157]. Low temperature anneals, however, result in decreased yield strengths
and may be attributed to reduced thermal vacancy content. This hypothesis is supported by
the presence of faceted voids and significantly reduced yield stresses and WHR's in CP-
NiAl and HP-NiAl following low temperature annealing. These voids form when alloys

are annealed at high temperatures, rapidly cooled to produce a supersaturation of thermal

vacancies, and subsequently re-annealed at intermediate temperatures resulting in the
coalescence of the vacancies into faceted voids [18,156]. Low temperature heat treatments
designed to produce similar defect contents in CP-NiAl and HP-NiAl result in a decrease in

yield stress for both alloys and decreases in WHR to nearly equal levels. This suggests
that WHR is dependent on the presence of thermal defects but independent of impurity
content for the composition range encompassed in this study.

The effects of interstitial elements (e.g., B, C, 0, N, S, etc.) cannot be ignored
when considering the mechanical properties of NiAL Previously, it was mentioned that the
CP-NiAl alloy exhibited higher yield stresses than the HP-NiAl alloy. The higher yield
stresses of the CP-NiAl crystal are believed to be due to a solid solution strengthening

effect due to the presence of interstitial C and 0 impurities. The solid solution

strengthening due to C and 0 in CP-NiAl relative to HP-NiAl is -2500MPa/at% C+O; this
is in general agreement with the observations of George et al. [151] who reported the solid

solution strengthening due to C in polycrystalline NiAl to be greater than 1700 MPa/at.%


As already mentioned, the NiAl-0.2Mo alloy exhibited higher yield stresses and
WHR's than the binary NiAl alloys and was unaffected by the low temperature anneal.
The higher yield stresses of the Mo-containing crystals are believed to be due to a solid
solution strengthening effect as opposed to hardening due to the presence of the a-Mo

precipitates. Mo was substituted in the place of Al (Table 1) and since all published phase

diagrams [162] suggest that Mo preferentially substitutes on the Al sublattice, it is

suspected that the positioning of Mo atoms on these sites leads to lattice strain and overall
hardening. Furthermore, the formation of a-Mo precipitates may result in a shift in

stoichiometry. Similar observations were reported recently by Cotton et al. [163] in the

similar NiAl+Cr system. More convincing evidence in support of a solid solution
strengthening mechanism is provided by the recent observations of Darolia et al. [16] who

report the critical resolved shear stress for <110> oriented NiAl doped with 0.1% Mo to be

154 MPa which is virtually identical to the value (-159 MPa) reported for the <123>

oriented NiAl-0.2Mo in this study. Specifically, if precipitation hardening were the cause

of the increased strengths in Mo-doped NiAl, it would be expected that the alloy doped

with 0.1%Mo would exhibit lower critical resolved shear stresses because it lies closer to

the solubility limit for NiAI.

The reasons for insensitivity of the Mo-doped alloy to the lower temperature heat
treatment is currently unclear. One possible explanation is that the NiAl/Mo interfaces
could act as sinks for the thermal vacancies. For example, Locci et al. [156] observed

crystallographic voids in stoichiometric NiAl melt spun ribbons subjected to intermediate

temperature anneals. However, no such voids were observed in W-doped NiAl melt spun

ribbons subjected to the same heat treatments. The W-doped alloy contained fine W

precipitates which were presumed to form effective sinks for the thermal vacancies thus
allowing them to anneal out rapidly at the resulting W/NiAl interfaces. Thus, vacancy

hardening of the as-received materials might be similar to that for the material given the low
temperature anneal. An alternative explanation for the observed insensitivity to the low
temperature anneal is the possible formation of Mo-vacancy pairs or clusters. Such
interactions might preclude the formation of faceted voids as well as the annealing out of

the thermal vacancies at 1000 K. However, neither mechanism is consistent with the
observation of faceted voids in the annealed NiAl-0.2Mo alloy in similar densities as
observed in the binary alloy.

Summary and Conclusions

Reductions in the interstitial and substitutional levels result in reduced yield
strengths in NiAl. Heat treatment also results in reduced yield and flow stresses in both
CP-NiAl and HP-NiAl due to a reduction in the concentration of thermal vacancies

resulting from vacancy coalescence during heat treatment
Conventional purity and HP-NiAl exhibit similar work hardening rates after similar
heat treatments (i.e., 1000 K(2h)+FC) which implies that the WHR might be dependent on

the concentration of thermal vacancies and independent of purity level over the range
encompassed by these alloys.

The yield stress and WHR of NiAl single crystals doped with 0.2%Mo exhibited no
dependence on the pre-test heat treatment studied but displayed higher work hardening rates

and yield stresses than both CP and HP-NiAl. These increases have been attributed to
solid solution hardening effects due to the addition of Mo.







0 0.02 0.04 0.06 0.08 0.1
True Compressive Strain

Figure 13. Typical room-temperature compressive stress-strain response for HP-NiAl,
CP-NiAl and NiAI-0.2Mo single crystals.


"A + Takasugi et al., 1992

_-^ ,, .
s ^A -

X 0


" 300









Temperature (K)

Figure 14. The temperature dependence of the 0.2% offset yield stress for HP-NiAl,
CP-NiAl, and NiAl-0.2Mo single crystals and additional data from the literature.


0.5 1 1.5 2 2.5 3 3.5
Temperature (K)

Figure 15. Arrhenius representation of the compressive yield stress as a function of
1000/T for HP-NiAl, CP-NiAl and NiAl-0.2Mo single crystals.

-&- HP-NiAI
-- --NiAl-0.2Mo
HP-NiAI: 1000 K(2h)+FC
CP-NiAI: 1000 K(2h)+FC
V, V.; n191




Figure 16. Bright Field TEM (BFTEM) micrographs of the dislocation morphology
observed in NiAl alloys prior to deformation. (a) HP-NiAl, (b) CP-NiAl, and (c) NiAl-
0.2Mo. Note the presence of a non=uniform distribution of 10 to 30 nm a-Mo precipitates
(indicated by the arrows) in NiAl-0.2Mo.

Figure 16 -- continued

Figure 17. BFTEM micrographs showing the dislocation morphology observed in
NiAl alloys after approximately 1% plastic strain in compression at 300 K. (a) HP-NiAl,
(b) CP-NiAl, and (c) NiAl-0.2Mo. Note the reduced tendency for cell formation in

Figure 17 -- continued

Figure 18. Faceted voids observed in HP-NiAl. Voids result from point defect
agglomeration during the 1000 K/2 h/FC anneal.



Body-centered cubic (BCC) metals and alloys exhibit an extreme sensitivity to (1)
point defects introduced during processing and/or heat treatment, (2) the level and type of
prestrain in the material, and (3) minute additions of interstitials which can lead to strain

aging phenomena. Not surprisingly, several aspects of strain aging have been identified as
playing a role in the deformation of polycrystalline and single crystal NiAl. They are the

occurrence of yield points and serrated stress-strain curves [15,22,23,26,28,164,165],
strain rate sensitivity minima [12,166], yield stress plateaus as a function of temperature

[18] and flow stress transients on changes in strain rate [139,166]. In addition, extensive
work by Margevicius et al. [23,167-169] has shown that a sharp yield point can be formed
in binary NiAl following annealing at 1100 K and furnace cooling (FC). This yield point

can be removed by subsequent restraining of the material by hydrostatic pressurization
prior to testing and recovered by aging the prestrained material for 7200 s (i.e., 2 hours) at

673 K. Similarly, Pascoe and Newey [12] observed the formation of room temperature

yield points in near stoichiometric NiAl annealed for 3600 s (1 hour) at 350 K following a
uniaxial prestrain. In addition, preliminary investigations by Weaver et al. [138] have
shown that these yield points can be removed by water quenching (WQ) from high
temperature as opposed to FC. Despite these observations, no complete investigation of
the interrelated effects of interstitial content, annealing and prestrain on mechanical behavior

has been conducted on NiAl. The purpose of this chapter is to describe the interrelated

effects of interstitial content, annealing and prestrain on the tensile flow and fracture
behavior of polycrystalline NiAl.

Experimental Details

NiAl alloys in the form of: (1) one titanium doped ingot (NiAl-Ti); (2) two
conventional purity (CPNiAl-1, CPNiAl-2) ingots; (3) two carbon doped induction melted
castings with varying carbon and oxygen concentrations (NiAl-100C and NiAl-300C); (4)
two low interstitial high-purity zone refined ingots, one of which was subsequently zone
leveled with carbon (HP-NiAl and HPNiAl-C respectively); and (5) a nitrogen doped
powder (NiAl-N) were the basic starting materials used in this investigation. All starting
materials were extruded at 1200 K at either a 12:1 or 16:1 reduction ratio. Descriptions of
the equipment and processes used to fabricate the high purity, zone leveled and nitrogen
doped alloys are presented elsewhere [19,92,153].

Material Characterization

Chemical analyses of the eight extrusions were conducted at the NASA-Lewis
Research Center by the following techniques deemed to be the most accurate for the

particular elements. Ni and Al were determined using analytical wet chemistry/titration
techniques and Si was determined by inductively coupled plasma atomic emission
spectroscopy. Oxygen, nitrogen, carbon and sulfur contents were determined by
combustion techniques using LECO oxygen/nitrogen and carbon/sulfur determinators.
Optical, scanning electron microscopy (SEM) and transmission electron microscopy
(TEM) were used to assess the microstructure of the materials. Polished optical
microscopy specimens were etched by swabbing with a mixture of 0.10 kg MoO3, 50 ml
HF and 150 ml H20.
Samples for transmission electron microscopy (TEM) were cut from the gage of
tested tensile specimens with a low-speed diamond saw and twin jet-electropolished in a

solution of 70% ethanol, 14% distilled water, 10% butylcellusolve and 6% perchloric acid
at 273 K, 20-25 V and 0.15 mA. TEM examinations were conducted in either a JEOL

100C or a Philips EM420 microscope operating at accelerating voltages of 120 kV.

Fracture surfaces of selected tensile samples were examined using a Cambridge 200
scanning electron microscope. Quantitative fractography was performed to determine the

percentage of intergranular fracture for most alloys. This was accomplished by taking at
least five random micrographs of appropriate magnification relative to the grain size from
each test specimen and using a point counting technique.

Tensile Testing

Round button-head tensile specimens were ground from the extruded rods so that
the gage lengths of the samples were parallel to the extrusion direction. Sample dimensions
were 3.1 mm for the tensile gage diameters and 30.0 mm for the tensile gage lengths. Prior
to testing, all samples were electropolished in a 10% perchloric acid-90% methanol solution
that was cooled to 208 K. Tensile tests were performed on an Instron Model 1125 load
frame at a constant crosshead velocity corresponding to an initial strain rate, i, of
1.4 x 10-4 s-1. All tests were performed in air at 300 K. True stress-strain data were

calculated from the load-time plots and yield stresses were determined by the 0.2% offset


The tensile testing was accomplished in four steps:
First, baseline mechanical properties were determined for all eight alloys by testing
them as follows: (1) as-extruded and (2) as-extruded + 1100 K/7200 s/FC. (3) Four
alloys (CPNiAl-1, HP-NiAl, HPNiAl-C, and NiAl-N), having received treatment (2), were
prestrained via pressurization to 1.4 GPa. The prestrain pressurization treatment was
selected based on the observations of Margevicius et al., [23,167-169].
Second, the temperature regime resulting in the maximum recovery of the yield
point was determined for a series of CP-NiAl specimens, having received treatment (3).

Specimens were annealed at temperatures ranging from 500 K to 1100 K for times ranging
between 60 s and 604,800 s (168 h) followed by FC, air cooling (AC) or WQ.

Furthermore, some specimens were uniaxially prestrained in tension approximately 0.2%
prior to annealing to determine the influence of heat treatment and type of prestrain on the

baseline flow and fracture behavior of the eight alloys.

Third, some specimens were statically strain aged as follows: specimens were
restrained approximately 0.2% at room temperature, unloaded, aged in situ on the load

frame for aging times varying between 50 s and 113,000 s (30 h), and then restrained at
room temperature approximately 0.2%. Aging temperatures were selected based on the
results from the test sequences described above. On several occasions, specimens were
subjected to recovery anneals of 1100 K/1800 s (30 min.)/AC following an aging cycle.
These procedures were repeated several times until fracture occurred in an effort to
elucidate the influence of strain aging on flow and fracture behavior of NiAl. The results of

one such experiment are presented in Figure 19 which shows the stress strain curves for an
alloy following multiple strain aging cycles. A more detailed accounting of the test method

used to study strain aging in NiAl is provided in references [137] and [138].

Experimental Results

Composition and Microstructure

The results of the chemical analyses of the eight alloys used in this study are shown
in Table 3. Within experimental accuracy (0.2 at.% for Ni and Al), the Ni and Al

contents of the eight alloys are not significantly different from each other. The major
differences between the materials are the residual silicon, carbon, oxygen and nitrogen
contents and the presence of Ti as an alloying addition to NiAl-Ti.

The microstructures of all the NiAl alloys were similar as observed by optical and
transmission electron microscopy, and consisted of fully dense, recrystallized and equiaxed

grains as demonstrated in Figure 20. The only differences were the observation of semi-

continuous stringers of nanometer-size precipitates in the NiAl-N (see reference [92]) and

NiAl-Ti alloys (Figure 21). In the case of NiAl-N, previous studies have revealed the

precipitates to be AIN [92]. Energy dispersive spectroscopic analysis in the TEM indicated
that the stringers in NiAl-Ti were rich in Ti. The individual particles composing the
stringers were, in general, too fine for analysis. On occasion, however, larger precipitates
were observed within individual grains or along grain boundaries. These precipitates were
typically elongated in shape as indicated in Figure 21. Analysis of microdiffraction patterns
taken from these particles indicated that they were TiC precipitates which suggests that the
Ti gettered C from the matrix in the NiAl-Ti alloy.

Table 3. Compositions and Grain Sizes of Polycrystalline NiAl-Alloys Investigated in This

Grinat.% appm
Alloy Size
(Heat) (m) Ni Al Ti Si C O N
CPNiAI-i 18.71.5 50.10.2 49.70.2 ---- 0.15 147 550 <9
CPNiAl-2 18.0M2.0 50.1T0.2 49.8UT 0.- 186 94 <9
HP-NiAI 51.42.3 49.90.2 501+0.2 0T W3 52 9
HPNiAI- 44.64.0 50.20.2 949.80-2 -- 05-92~- 30 9
NiA'-N 4.00.3 10.2 49.70.2 .02 57 347 904
NiAl-100( 20.0.0 49.9:0.2 50.00.2 0-0- 490 T ---
NIA_300( 20.02.0 49.9:k0.2 50002 00~1 1153 131 <9
NiAl-Ti 20.0.0 49.90.2 50.00.2 0.03 000 214 113 T15

The baseline mechanical properties are summarized in Table 4. Typical room-
temperature stress-strain curves for each alloy are shown in Figures 22 through 26. From
this data it is observed that the yield stress of each alloy generally decreased following the
1100 K/7200 s/FC anneal. In addition, a tendency for discontinuous yielding was apparent
in the CPNiAl-1, CPNiAl-2, HPNiAl-C, NiAl-100C and NiAI-300C alloys but not in HP-

NiAl, NiAl-N or in NiAl-Ti. Further decreases in yield stress and elimination of the
tendency for discontinuous yielding were achieved in the CPNiAl-1 and HPNiAl-C alloys
if the specimens were subsequently hydrostatically pressurized at 1.4 GPa whereas no
additional decreases in yield stress or other apparent changes in flow behavior were
observed in the powder processed NiAl-N alloy or in the HP-NiAl material (Figures 22-
25). Pressurization treatments were not conducted on the remaining alloys. Recovery of
the discontinuous yield behavior following pressurization could be accomplished by re-
annealing pressurized specimens at 1100 K/2 h followed by furnace cooling (Figure 23).
Interestingly, even though the yield stress could be reduced by annealing and in some cases
by hydrostatic pressurization, these treatments had no obvious influence on the tensile

ductility or fracture behavior of the various alloys. Similar observations have been recently

reported by Margevicius and Lewandowski for single crystal and polycrystalline NiAl
alloys [170]. Scanning electron micrographs of the fracture surfaces are exhibited in
Figures 27-34. In all specimens, failure was always by a combination of intergranular
separation and transgranular cleavage. Surprisingly, the HP-NiAl and NiAl-Ti exhibited a
greater tendency for intergranular failure and lower tensile ductility than the other alloys.

Influence of Prestraining and Annealing on Baseline Properties

To determine whether the observed yield points resulted from the hold at
temperature or during cooling from the annealing temperature, specimens of CPNiAl-1

previously prestrained hydrostatically, were reannealed at 1100 K/7200 s followed by AC
or WQ. The resulting properties are summarized in Figure 35 and in Table 4. After WQ,

only continuous yielding was observed while after AC, there was some evidence of a yield

plateau which initially suggests that the yield points observed following FC are the result of
the pinning of dislocations by mobile solute atoms during cooling through lower

temperatures. As a result, annealing experiments were initiated at lower temperatures to
determine the critical temperature for the migration of solute atoms to dislocations. The
results of these experiments are summarized in Figure 36. It is observed that yield plateaus
formed in CP-NiAl after hydrostatic restraining followed by annealing treatments of 700
K/7200 s/FC but not following anneals of 500 K/7200 s/FC. Conversely, if the specimens

were prestrained uniaxially, notable yield points formed readily after as little as 900 s (15
min.) at 522 K and in as little as 60 s (1 min.) at 700 K. Additionally, no yield points were

observed when annealing temperatures exceeded 900 K.

TEM Observations of Deformed Specimens

Figure 37 shows a series of TEM bright field images of CPNiAl-1 that was
deformed at room-temperature following anneals of 1100 K/2 h/FC. Tensile tests were

interrupted at plastic strains of 0.05%, 0.31% and 2.04% corresponding to the yield stress

peak, the Liiders region and after fracture respectively. At 0.05 and 0.31% strain (Figures
37a and b), the dislocation structure consisted of a low density of inhomogeneously
distributed dislocations arranged into poorly defined cells and dense tangles. Some
intercellular/inter-tangle dislocation debris was observed although the cell interiors

remained largely dislocation free. As the strain was increased to 2.04% (Figure 37c), the
dislocation density increased and the cells became more well defined. The dislocations
observed were predominantly <100> dislocations of mixed character. The predominant
slip plane was (011) although some dislocation debris lying on (001) slip planes was

occasionally detected.


Table 4. Baseline Tensile Properties of NiAl Alloys
0.2% Yield Fracture Intergranular
Material Condition* Stress Stress Ductility Facture Observations
(MPa) (MPa) (%) (%) #
CPNiAl-1 as-extruded 269 379 2.11 -na- DY, YP
as-extruded 275 368 1.83 36.0 DY, YP
annealed/FC 184 301 2.08 36.1 DY, sharp YP
annealed/FC 197 228 1.04 -na- DY, sharp YP
annealed/AC 154 309 2.26 -na- DY, plateau
annealed/WQ 143 228 1.13 46.2 CY
pressurized 154 288 1.86 -na- CY
pressurized 159 317 1.81 37.5 CY
pressurized 154 322 2.16 37.4 CY
CPNiA1-2 as-extruded 164 241 1.05 30.8 CY
as-extruded 172 306 1.70 37.9 CY
annealed 117 257 2.05 35.2 DY, sharp YP
annealed 116 336 3.34 40.2 DY, sharp YP
HP-NiAl as-extruded 166 214 0.79 57.5 CY
as-extruded 157 235 1.16 60.4 CY
annealed/FC 98 174 1.17 63.2 CY
pressurized 118 176 0.92 58.2 CY
HPNiAl-C as-extruded 170 201 0.59 43.4 CY
annealed 113 157 0.68 -na- YP
pressurized 96 171 0.98 39.4 CY
NiAl- as-extruded 155 255 1.35 37.5 DY, plateau
as-extruded 151 319 2.25 33.0 DY, plateau
annealed 113 342 3.29 41.9 DY, sharp YP
annealed 115 285 2.35 40.1 DY, sharp YP
NiAl- as-extruded 1...80 322 2...2 32.4 CY
as-extruded 162 297 1.84 28.6 DY, plateau
annealed 109 343 3.29 42.8 DY, plateau
annealed 108 352 3.38 48.4 DY, plateau
NiA1-N as-extruded 298 409 1.32 -na- C .
as-extruded 297 476 2.20 34.1 CY
annealed/FC 265 468 2.45 31.9 CY
pressurized 266 434 2.16 -na- CY
pressurized 274 352 1.03 33.5 CY
NiAl-Ti as-extruded 170 235 0.86 57.5 CY
as-extruded 176 176 0.20 56.2 CY
annealed 123 256 1.76 64.7 CY
annealed 124 265 1.81 64.3 CY
*annealed/FC or annealed = as-extruded + 1100 K/7200 s/FC
*annealed/AC = as-extruded + 1100 K/7200 s/AC
*annealed/WQ = as-extruded + 1100 K/7200 s/WQ
*pressurized = as-extruded + 1100 K/7200 s/FC + pressurize 1.4 GPa
#DY = discontinuous yielding, upper yield point, sharp yield drop
#YP = yield point; #CY = continuous yielding; -na- = not available
Intergranular fracture: accuracy = +10%

Figure 37e shows the dislocation substructure observed in NiAl-N following

hydrostatic restraining to 1.4 GPa. In agreement with the reports of Margevicius and co-

workers [168,169] for conventional purity cast and extruded NiAl, the dislocation

substructures consisted of a more uniform distribution of long, straight dislocations
generated from grain boundaries. Diffraction contrast analysis revealed that all of the

dislocations were of <001> type. Similar microstructures were observed in CPNiAl-1 after
hydrostatic restraining.


In agreement with the many prior studies of extruded near-stoichiometric NiAl in

bulk form [71,171,172] and during in-situ TEM observations [173], the deformation

substructure following straining at room-temperature consisted of a network of poorly

defined dislocation cells, dense tangles and intercellular debris. Well defined deformation

bands, were not observed by TEM nor was there evidence of coarse slip bands intersecting
the specimen surfaces.

Species Responsible for Strain Aging in NiAl

A determination of the species most likely responsible for the strain aging effects in

NiAl can be made by examination of the aging behavior of all eight alloys. Discontinuous

yielding, in the form of yield points and yield plateaus, was observed in CPNiAl-1,

CPNiAl-2, NiAl-100C, NiAl-300C and HPNiAl-C, while continuous yielding was

observed in HP-NiAl, NiAl-Ti or NiAl-N following heat treatments known to produce

yield points in conventional cast and extruded NiAl [167]. In NiAl-N, the oxygen and
nitrogen contents were much higher than those observed in CPNiAl-1 or CPNiAl-2 while

the C-content was much lower, which suggests that nitrogen and oxygen are not the

species responsible for the observed yield points. When the excess interstitials were

reduced sufficiently, as in the case of HP-NiAl, no yield points were observed; however,

doubling the carbon concentration (HPNiAl-C) resulted in a well-defined yield point.

These observations are supported by the prior investigations of Noebe and Garg [92] who

observed sharp yield points in powder processed conventional purity NiAl (C=143 appm,

0=227 appm, and N=6 appm) but no yield points in powder processed nitrogen-doped

NiAl (C=57 appm, 0=347 appm, and N=904 appm). Also, in HPNiAl-C, it was observed

[137,138] that longer aging times are required to achieve the same yield increment as

observed in CPNiAl-1 and CPNiAl-2. It is believed that this behavior is a result of the

significant reduction in the concentrations of interstitials, particularly C. Since there is less

carbon to pin dislocations in NiAl-C, the carbon present must, presumably, diffuse longer
distances to cause pinning. Finally, in NiAl-Ti, the bulk interstitial levels were equivalent

to those observed in CPNiAl-2 which exhibited yield point behavior. However, no yield

points were observed in NiAl-Ti which was shown to contain a TiC stringers; this suggests

that the lack of a yield point in NiAl-Ti is due to the gettering of sufficient carbon from the

NiAl matrix and that carbon is the species responsible for the observed yield points.

Influence of Prestraining

As noted in the Results, the return of a sharp yield point is much more rapid when

the specimen has been prestrained uniaxially as opposed to hydrostatically. During uniaxial

deformation of NiAl, dislocations cross-slip easily forming cell structures [53,92] that

result in high work hardening rates at room temperature. As a result, the dislocations are in

essence pinned. In contrast, samples pressurized hydrostatically show a more even

distribution of dislocations which are not bound in cells [169]. In uniaxially prestrained

samples, since some of the dislocations are already locked up in cell structures, fewer

mobile dislocations are available. Thus, less solute is required to pin the available mobile

dislocations. Since more mobile dislocations are available in hydrostatically prestrained

samples, more carbon is required to cause pinning. As a result, longer aging times are

required to achieve the same yield point increments observed after uniaxial restraining. A

similar explanation has been applied to strain-aged steels prestrained in directions non-

parallel to the original tensile direction [174].

The results of this study also revealed a further reduction in the yield stress of

CPNiAl-1, HP-NiAl, and HPNiAl-C following pressurization to 1.4 GPa. However, no

such effect was observed in NiAl-N. Prior investigations by Margevicius et al. [169] on

pressurized NiAl alloys have shown that dislocation generation is enhanced by

compositional differences between neighboring grains. This suggests that higher

dislocation densities will be enduced in CPNiAl-1, HP-NiAl, and HPNiAl-C over NiAl-N

due to the larger compositional variations between the neighboring grains.

Despite it's influence on the tensile flow behavior of NiAl alloys, restraining,

whether uniaxial or hydrostatic, had no observable influence on the tensile ductility and

fracture behavior of NiAl. In CP-NiAl, for example, the fracture surfaces were

approximately 37% intergranular regardless of how the specimens were prestrained or

strain aged prior to testing. In HP-NiAl, however, there was an apparently higher

propensity for intergranular fracture over transgranular failure. Interestingly, this alloy,

although apparently closer to stoichiometry and free from high interstitial levels, exhibited

lower tensile ductilities than the CP-NiAl and the NiAl-N alloys. This observation

contradicts the observations in single crystals where soft-oriented specimens having low

interstitial levels and high purity were shown to exhibit nearly 5% tensile ductility at room

temperature, regardless of pretest treatments [19]. Although some of these observations

can be related to the differences in grain size between the alloys, more thorough

investigations of these issues are required.


The yield points observed in conventional purity and carbon-doped NiAl are the

result of strong dislocation pinning by interstitial carbon. Oxygen and nitrogen levels

below 0.035 and 0.09 at.%, respectively, do not appear to pin dislocations in NiAl and,

therefore, do not produce yield point phenomena.

Hydrostatic restraining as opposed to uniaxial restraining delays the kinetics of

the yield point return by forming random networks of free unpinned dislocations which

require more diffusion time for strong locking to occur.

Yield point phenomena can be removed by microalloying with sufficient levels of

reactive elements such as Ti to getter carbon from the matrix by forming semi-continuous

precipitates of TiC.

Despite their influence on the flow behavior of NiAl, restraining and/or strain

aging have no observable impact on the room-temperature fracture characteristics of NiAl

alloys. In fact, fracture always occurs by a mixture of intergranular failure and

transgranular cleavage.

Percent Strain

Series of true stress-strain curves for CP-NiAI after multiple strain aging

Figure 19.

Figure 20. Transverse sections of extruded material: (a) CPNiAl-1; (b) CPNiAl-2; (c)
HPNiAl-C; (d) HP-NiAl; (e) NiAl-100C; (f) NiA1-300C; and (g) NiAl-N.

Figure 20 -- continued

w'". '
*'.. ,* r" -*

,6 I -'

fg. 7' lO m .

Figure 20-- continued

Figure 20 -- continued


d Ti

Ni Al Ti iNi Ni

4- 0.000 Range- 10.230 keV 10.110 -9
integral 0 20476

Figure 21. Stringers observed in NiAl-Ti: (a) Backscattered scanning electron
micrograph of TiC stringers in NiAl-Ti; (b) BFTEM micrograph of a large TiC precipitate;
(c) TEM microdiffraction pattern from the TiC precipitate ([01 1] TiC zone axis); and (d)
EDS spectra for the particle in (c) indicating the presence of Ti.







1100 K/2 h/FC
1100 K/2 h/FC+
P (1.4 GPa)

1100 K/2 h/FC+
P(1.4 GPa)
1100 K/2 h/FC


J ____________________ _____________________ I.

Percent Strain

Figure 22. Room-temperature tensile stress-strain curves for CPNiAl-1 in the as-
extruded, annealed, pressurized and pressurized+annealed conditions.












50 -


Percent Strain

Figure 23. Room-temperature tensile stress-strain curves for HPNiAl-C in the as-
extruded, annealed, and pressurized conditions.



200 -
( As-extruded

S- 100

r As-extruded+
I 1100 K/2 h/FC+
50 1
P (1.4 GPa)


Percent Strain

Figure 24. Room-temperature tensile stress-strain curves for HP-NiAl in the as-
extruded, annealed, and pressurized conditions.





Percent Strain

Figure 25. Room-temperature tensile stress-strain curves for NiAl-N in the as-
extruded, annealed, and pressurized conditions







10 A = as-extrudei
B = annealed
Nil-OO,0 K.____ _____---LJ------
Percent Strain

Figure 26. Representative room-temperature tensile stress-strain curves for CPNiAl-2,
NiAl-100C, NiAI-300C and NiAl-Ti

Figure 27. Fracture surfaces of CPNiAl-1 samples tensile tested at room temperature:
(a) as-extruded; (b) after annealing at 1100 K for 2 h followed by furnace cooling; (c) after
annealing followed by subsequent pressurization.

Figure 27 -- continued

Figure 28. Fracture surfaces of HP-NiAl samples tensile tested at room temperature: (a)
as-extruded; (b) after annealing at 1100 K for 2 h followed by furnace cooling; (c) after
annealing followed by subsequent pressurization.

Figure 28 -- continued

Figure 29. Fracture surfaces of NiAl-C samples tensile tested at room temperature: (a)
as-extruded; (b) after annealing at 1100 K for 2 h followed by furnace cooling; (c) after
annealing followed by subsequent pressurization.

Figure 29 -- continued

Figure 30. Fracture surfaces of NiAl-N samples tensile tested at room temperature: (a)
as-extruded; (b) after annealing at 1100 K for 2 h followed by furnace cooling; (c) after
annealing followed by subsequent pressurization.

Figure 30 -- continued

Figure 31. Fracture surfaces of CPNiAl-2 following tensile testing at room-
temperature: (a) as-extruded and (b) after annealing at 1100 K for 2 h followed by furnace

Figure 32. Fracture surfaces of NiAl-100C following tensile testing at room-
temperature: (a) as-extruded and (b) after annealing at 1100 K for 2 h followed by furnace

Figure 33. Fracture surfaces of NiAl-300C following tensile testing at room-
temperature: (a) as-extruded and (b) after annealing at 1100 K for 2 h followed by furnace

Figure 34. Fracture surfaces of NiAl-Ti following tensile testing at room-temperature:
(a) as-extruded and (b) after annealing at 1100 K for 2 h followed by furnace cooling.

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