T

NASA Contractor Report 175073

High Temperature Dispersion

Strengthening of NiAl

f {NASA-CB-175073} HIGH TEMPERATURE N86-22688

iDISPERSION STRENG1HENIKG OF NiAl Final.

'Report {Case Western Reserve Uciv.) 15 p

|HC &02/MF A01 CSCL 11F Dnclas

I G3/26 05918

Mark Sherman and Krishna Vedula

Case Western Reserve University

Cleveland, Ohio

March 1986

Prepared for

Lewis Research Center

Under Grant NAG 3-387

NASA

National Aeronautics and

*f.Space*A"dministration
HIGH TEMPERATURE DISPERSION STRENGTHENING OF NiAl

Mark Sherman

Sohio Research Center, Cleveland, Ohio

Krishna Vedula

Department of Metallurgy and Materials Science,

Case Western Reserve University, Cleveland, Ohio

SUMMARY

A potential high temperature strengthening mechanism for alloys based on

the interraetalllc compound NiAl has been investigated. This study forms part

of an overall program at NASA Lewis Research Center for exploring the

potential of alloys based on NiAl for high temperature applications. An alloy

containing 2.26 at% Nb and produced by hot extrusion of blended powders has

been examined in detail using optical and electron microscopy. Interdiffusion

between the blended Nb and NiAl powders results in the formation of

intermediate phases. A fine dispersion of precipitates of a hexagonal,

ordered NiAlNb phase in a matrix of NiAl can be produced and this results in

strengthening of the alloy by interfering with dislocation motion at high

temperature. These precipitates are, however, found to coarsen during the

high temperature (1300 K) deformation at slow strain rates and this may

impose some limitations on the use of this strengthening mechanism.

1. INTRODUCTION

Ordered interraetallics have long been recognized as candidates for high

temperature applications because thermally activated processes such as

diffusion are inhibited by the ordered crystal structure. Recently, interest

has been generated in developing certain intermetallic compounds for high

temperature applications [1-4]. The ordered intermetallic NiAl is especially

attractive since this compound offers oxidation and corrosion resistance,

remains ordered up to its high congruent melting point of 1913 K, and has a

density advantage over conventional superalloys because of its high aluminum

content.

The drawbacks of NiAl, however, are its room temperature brittleness and

its inadequate creep resistance at high temperatures, in spite of its ordered

crystal structure. In this paper, we are concerned only with its high

temperature mechanical properties. Previous work [5] has shown that the high

temperature creep resistance of NiAl can be improved to levels comparable

with those of some superalloys by inacroalloying with ternary additions, using

a powder metallurgy process. The powder metallurgy process involves blending

of elemental ternary additions with binary NiAl followed by hot extrusion in

a sealed steel container. Additions of Nb, Ta and Hf at the 2 at% level, in

particular, have shown the most promise. The reasons for this improvement in

high temperature creep resistance are, nevertheless, not clear.
The purpose of this study was to understand the interaction between the

Nb powder and the NiAl powder during the powder metallurgy processing of an

alloy containing 2.26 at% Nb and to identify the mechanism of high

temperature strengthening in this alloy. This alloy has been microscopically

examined before and after high temperature deformation in order to obtain

this understanding.

2. MATERIALS AND PROCEDURE

The alloy was prepared at NASA Lewis Research Center by blending

approximately 2 at% of -325 mesh commercially pure Nb powder, with -80 mesh

prealloyed NiAl powder. The fine Nb powder was used in order to enhance

interdiffusion between the matrix NiAl and the ternary addition. The powder

blend was filled in a mild steel can (5 cm. dia. and 12.5 cm. long) which was

evacuated and sealed. The sealed can was preheated to a temperature of 1366 K

and extruded through a vertical die at an area reduction ratio of 16:1. The

extruded can was air cooled and the steel can was then removed by grinding.

The alloy specimens were subsequently homogenized at 1523 K in argon for 175

hours to ensure maximum interdiffusion between the ternary addition and the

matrix NiAl. Spectrographic analyis of the as-homogenized alloy showed the

alloy to have a nominal chemical composition of 50.19 at% Ni, 47.58 at% Al

and 2.26 at% Nb.

Cross-sections of the as-extruded and as-homogenized specimens were

examined' using optical and scanning electron microscopy. Cylindrical

specimens (5 ram. dia. and 10 mm. long) of the as-homogenized material were

prepared by centreless grinding. Specimens were tested (at NASA Lewis

Research Center) in compression on an Instron Testing machine at a

temperature' of 1300 K or 0.68 Tm and at slow strain rates. The specimen

tested at an initial strain rate of 3 x 10 /sec to a final strain of about

7% was examined in considerable detail using optical and scanning electron

microscopy .

Transmission electron microscopy of the specimens before and after

compression testing was used to identify the second phases and to examine the

"substructure of the specimens in order to understand the strengthening

mechanism. Thin foils for TEM were prepared by electropolishing in a 2:1

methanol / nitric acid solution cooled to 253 K.

3. RESULTS AND DISCUSSION

(A) AS-EXTRUDED AND AS-HOMOGENIZED SPECIMENS

The as-extruded micrestructure (illustrated in Figure la) contains

'recrystallized NiAl grains with very little porosity. The second phase

particles observed are the undissolved Nb particles which have a wide range

of sizes, the average size being roughly 4 micrometers.

The as-homogenized microstructure (illustrated in Figure Ib) shows a

well developed grain structure with an average grain size of about 25

micrometers which is somewhat larger than the as-extruded grain size. A

slight increase in the amount of porosity is observed and is suspected to be
due to pull out of particles during specimen preparation. The average size of

the second phase particles has increased by about 50% to roughly 5.5

micrometers. The increase in size of the second phase particles, which were

originally Nb powder particles, is believed to be a result of transformation

of the Nb to a ternary intermetallic of Ni, Al and Nb by interdiffusion.

"The distribution of the second phase particles is clearly observed in

the scanning electron micrograph in Figure 2a. Some of the larger particles

reveal the presence of two phases, one surrounding the other, as illustrated

in Figure 2b. Energy Dispersive Analysis of X-Rays from the two regions shows

that the inner phase contains more Nb than the outer phase and the

compositions of the phases correspond in 'general to the phases NiAlNb« and

NiAlNb respectively (Figure 3). A more definite confirmation of their

compositions could not be made since standards were not used for the

analyses. These two phases are in agreement with an earlier study by Kornilov

[6] on the phase equilibrium between NiAl and Nb. This kind of a layered

structure results from incomplete interdiffusion between the matrix of NiAl

and the elemental addition of Nb.

The pseudo-binary phase diagram for NiAl-Nb is reproduced in Figure 4

from Kornilov [6], With increasing Nb, the phases present in this pseudo

binary section are NiAl (with a very limited solubility of Nb), NiAlNb and

NiAlNb-. NiAl and NiAlNb exhibit a eutectic at about 1710 K. The solubility

of Nb In NiAl, according to this diagram is less than 1 at% at low

temperatures, but increases significantly above about 1450 K and can approach

5 at% near the eutectic temperature. Hence, during the homogenization

treatment (1523 K), the matrix of NiAl and the Nb powder particles act as a

finite spherical diffusion couple and the phases corresponding to the two

intermediate ternary interraetallics form between the NiAl and the Nb.

Since the solubility of Nb in NiAl is much higher at the homogenization

temperature than at room temperature, the slow cooling after the

homogenization treatment results in reprecipitation of a second phase of

NiAlNb in the matrix of NiAl. A distribution of very fine precipitates is, in

fact, observed when a thin foil is examined on the Transmission Electron

Microscope at a high magnification (Figure 5a). These precipitates are,

however, only 10 to 50 nm. in size and too small to be detected at the low

magnifications used in the optical and scanning electron micrographs. A

diffraction pattern showing the spots for the matrix as well as the

precipitates is shown in Figure 5b. The diffraction spot used to obtain the

dark field image of the precipitates in Figure 5a is indicated by the arrow

in Figure 5b.

(B) DEFORMED SPECIMEN

Cylindrical specimens were tested at slow strain rates in compression

and the behavior of a specimen with an initial strain rate of 3 x 10 / sec

is plotted in the form of a true stress- true plastic strain curve in Figure

6 and compared with similar curves for binary NiAl » Plastic flofr dn the

ternary alloy is initiated at a stress of about 35.0 MPa which is

significantly higher than that for the binary alloy, indicating a substantial

strengthening due to the ternary addition. The flow stress, however, drops

*Unpublished results from J.D. Whittenberger of NASA Lewis Research Center.
during the deformation to a lower value of about 29 MPa after a plastic

strain of about 4%. A possible explanation for the strengthening can be

obtained from an understanding of the microstructure of the tested specimen.

•

The matrix NiAl grains appear to have grown somewhat during the testing

as evidenced by comparing the optical microstructure of the tested specimen

in Figure 7a with that of the untested specimen In Figure Ib. The average

grain size measured shows an Increase from about 25 micrometers to about 33

micrometers. A considerable amount of etch pitting within the grains is

suspected to be associated with dislocations within these grains. Voids

appear to have coalesced along the triple points and grain boundaries during

the testing, presumably due to grain boundary sliding.

A striking feature of the deformed specimen when compared with the

undeformed specimen is the presence of many precipitates within the grains of

the deformed specimen when compared on the scanning electron microscope in

the backscatter mode (Figure 7b compared with Figure 2a). These precipitates

are between 0.5 to 1 micrometer In size and much smaller than the original

Nb powder particles. These precipitates have been shown to be NiAlNb

precipitates and could be a result of coarsening of much finer precipitates

which were not resolved on the SEM in the homogenized specimen.

An interesting aspect of the microstructure of the deformed specimen is

the presence of regions around the large particles which are devoid of the

fine precipitates. Such a precipitate-free zone must be a consequence of

dissolution of smaller precipates in the vicinity of the larger particles and

redeposltion on the surface of the larger particles due to Ostwald ripening.

Transmission electron microscopy of the deformed specimen reveals

precipitates of the size of about 500 run (Figure 8), thus, providing direct

evidence that the precipitates have coarsened significantly during the high

temperature testing (compare Figure 5 with Figure 8). The size of the

precipitates observed in the transmission electron micrograph is in agreement

with the size observed in the scanning electron micrograph of the tested

specimen. A possible reason for the rapid coarsening during deformation may

be Increased diffusion through dislocation pipes.

(C) STRENGTHENING MECHANISM

The mechanism of strengthening In the ternary alloy is believed to be

due to the Interaction between the fine precipitates and dislocations.

Dislocation bowing as well as development of dislocation networks between the

precipitates are observed throughout the thin foil of the deformed specimen.

Typical examples are illustrated in Figures 9a and 9b. This kind of

Interaction is generally observed when the precipitates act as obstacles to

dislocation motion and hence strengthen the material [7,8,9].

Dislocation networks with segments parallel to <101> and <110>g vectors

developed between precipitate particles. Curved dislocation lines are

dominant at and near precipitates , whereas straight parallel dislocations

were observed in groups in the matrix and piled up at elongated particles.

Subgrain boundaries were observed at grain boundaries and at the
precipitates. The burgers vectors of dislocations in the deformed specimen

were not uniquely determined. However, in general, the dislocations were

strongly visible with (110) reflections and invisible with (112) reflections.

The identification of the precipitates which are believed to cause the

strengthening was carried out using selected area and microdiffraction on the

transmission electron microscope using the deformed specimen. The

precipitates in the undeformed specimen were too small to be properly

analyzed and it is assumed that they have the same composition and crystal

structure as those identified in the tested specimen. The precipitates in the

deformed specimen are found to be NiAlNb with a hexagonal crystal structure.

A TEM bright field image of one of the precipitates and the corresponding SAD

and microdiffraction patterns of a zone axis are shown in Figure 10.

Measurements from several diffraction patterns confirm the hexagonal phase

NiAlNb. The d-spacings from the Inorganic Powder Diffraction File for NiAlNb

and those measured from electron diffraction patterns are presented in

Table 1. X-Ray diffraction of the deformed specimen in a powder form, also

agreed with the conclusion that the second phase present in the matrix of

NiAl is the hexagonal phase NiAlNb. These data are also presented in the same

table for comparison. The possibility of the second phase being Ni^AlNb (a

Heusler alloy) was very carefully checked and ruled out.

In terms of the strengthening mechanism in the ternary alloy, since the

ternary intermetallic phase NiAlNb is harder than the binary NiAl [6],

dislocation motion at the testing temperature appears to be hindered by the

presence of these fine precipitates. Such a mechanism is very sensitive to

the average spacing between and hence the size of the precipitates for a

given volume fraction. Coarsening of these precipitates could, therefore,

have dramatic consequences for strengthening. In fact, the drop of flow

stress during the test is suspected to have been caused by coarsening of the

precipitates from about 50 nm observed in the undeformed as-homogenized

specimen to about 500 nm in the deformed specimen. At the same time, the

dislocations present between the precipitates may act as rapid paths for

diffusion in the matrix, thereby enhancing the kinetics of the coarsening of

the precipitates.

It should be pointed out here that these results are in agreement with

the proposed strengthening mechanism in an alloy containing 2 at% Ta [10],

investigated as part of the same overall program for investigating the effect

of ternary additions on NiAl at NASA Lewis Research Center.

4. CONCLUSIONS

The results of this investigation suggest the following conclusions:

1) Homogenization of ternary alloys of NiAl and Nb, prepared by hot extrusion

of powder blends of NiAl and Nb, results in the formation of the

intermediate phases NiAlNb and NiAlNb-.

2) The high temperature strengthening of this alloy appears to be due to a

dispersion of fine precipitates ("50 nm in size) of NiAlNb which have a

hexagonal crystal structure. These precipitates interfere with dislocation
motion during high temperature deformation.

3) The strengthening precipitates coarsen significantly during the testing

(from about 50 nm to 500 nm) and this coarsening may be responsible for

the drop in flow stress during the deformation.

ACKNOWLEDGEMENTS

The authors would like to acknowledge the support of Sohio Research

Center for the use of their microscopy facilities and NASA Lewis Research

Center for supplying the materials.
REFERENCES

1. "Structural Uses for Ductile Ordered Alloys", NMAB Report

No. 419, August 31, 1984, National Research Council,

Washington, D.C.

2. J.R. Stephens, "The B2 Alurainides as Alternative Materials",

NASA Technical Memorandum 86937, prepared for the 1984 Fall

Meeting of the Materials Research Society, Boston, Nov.,

1984.

'3. N.S. Stoloff, International Metals Reviews, 1984, Vol. 29.

No. 3, p. 123.

4. C.T. Liu, 'Design of Ordered Intermetallic Alloys for High

Temperature Structural Uses'. Paper presented at the AIME

Symposium on High Temperature Alloys, Theory and Design,

April 1984, Bethesda, Maryland.

5. K.M. Vedula, V. Pathare, I. Aslanidls and R.H. Titran, to be

published in the Proceedings of a Symposium on High

Temperature Ordered Interraetallic Alloys at the Materials

Research Society Meeting, Boston, November, 1984.

6. I.I. Kornilov, R.S. Mints, L.N. Guseva, Yu.S. Malkov, Russian

Metallurgy, No.6, 1965, p.83.

7. O.D. Sherby, R.H. Klundt and A.K.. Miller, Met. Trans. 1977,

Vol. 8A, p.843.

8. G.S. Ansell and J. Weertman, Trans. AIME, 1959, Vol. 215,

p.838.

9. J.J. Petrovlc and L.J. Ebert, Met. Trans. Vol. 4, 1973,

p.1309.

10. V.M. Pathare, M.S. Thesis, Department of Metallurgy and

Materials Science, Case Western Reserve University,

Cleveland, Ohio, 1984.
TABLE 1: COMPARISON OF d-VALUES FROM POWDER DIFFRACTION (PD) FILE

FOR NiAlNb WITH THOSE DETERMINED BY ELECTRON DIFFRACTION

(ED) FROM PRECIPITATES AND X-RAY DIFFRACTION (XRD) OF

PULVERIZED HOMOGENIZED ALLOY IN THIS STUDY.

Hexagonal 12H a = 5.00 c = 8.093

ANGSTROMS

hkl PD FILE MEASURED, ED MEASURED, XRD

100

002

101

102

201

110

103

213

220

4.33

4.05

3.82

2.96

2.09

2.50

2.29

1.39

1.25

4.35

4.06

3.83

2.96

2.11

2.25

2.10

2.47

2.26

1.38

1.23
ORIGINAL PAGE IS

OF POOR QUALITY

I «>

(a) ( b)

Figure 1. Comparison of Optical Micrographs of Cross-sections of

(a) As-Extruded and (b) As-Homogenized specimens of the

NiAl-2.26 at% Nb alloy (Marbles Reagen Etch)

(a) (b)

Figure 2. SEM Backscattered Electron Images of the Homogenized

Alloy Showing (a) A Typical Distribution of Nb-rich

Particles and (b) The Existence of a Layered Two Phase

Structure within some of the Larger Particles.

(Cross-section).

9
I I

wt. % a t . %

A1 10.48 24.97

Mb 56.93 39-37

Ni 32.59 35.66

(a) (b)

Figure 3. EDS Analysis of the Two Regions of Particle in Figure 2b

(a) The Inner Phase Corresponding Roughly to NiAlNb2 and

(b) The Outer Phase Corresponding Roughly to NiAlNb.

T ° K

2OOO

I5OO

IOOO

5OO

NIAI

I7IO

13 + NbNiAl

NbNiAl

/

Nb2NiAI

c

2OOO

I5OO

IOOO

50O

IO 20

Nb at%

30 40

Figur e 4. Pseudo-Binary Phase Diagram of the NiAl-Nb system.

(Kornilov, Ref. 6).

1

 10
ORIGINAL PAGE IS

OF POOR QUALfTY

(a) (b)

Figure 5. (a) TEM Dark Field Image of Precipitates in the

Homogenized Alloy and (b) The Corresponding Diffraction

Pattern with Arrow Indicating the Reflection used for

Imaging.

40

30

I

UJ

a

in

Ul

a

I10

STRAIN RATE 3.0»l6r

HOT COMPRESSION

I3OO-K AIR

Ni: 5019*

Al : 47 58«

Nb: 2 36*

"(THIS STUDY)

Ni: 4991 %

Al: 5O09*

(WHITTENBERGER)

REF.7

3 4 5

TRUE STRAIN ( % )

Figure 6. Comparison of the True Stress- True Plastic Strain

Behavior of the Nb Containing Alloy with a Binary NiAl

Alloy (Ref. 7). Specimens Tested in Compression at

1300 K and Slow strain Rates.

11
(a) (b)

Figure 7. (a) Optical Microstructure of Specimen after Testing

Showing Grain Boundary Void Formation (Etched),

(b) SEM Backscattered Electron Image of Same Specimen

Showing Coarsened NiAlNb Precipitates and Precipitate

Depleted Regions Around Original Nb-Rich Particles.

Figure 8. TEM Bright Field Image of Deformed Specimen Showing

Coarsened Precipitates of NiAlNb.

12
ORIGINAL PAGE IS

OF POOR QUALfTY

(a) (b)

Figure 9. Transmission Electron Micrographs of Deformed Specimen

Illustrating (a) Dislocations Pinned at and bowing

Between Precipitates and (b) Dislocation Networks Formed

Between Precipitates.

250nm

Figure 10. TEM Bright Field Image of a Large Precipitate in the

Deformed Specimen and Corresponding SAD and Micro;

diffraction Patterns of a Zone Axis Indexed as (2110)

13
1. Report No.

NASA CR-175073

2. Government Accession No. 3. Recipient's Catalog No.

4. Title and Subtitle

High Temperature Dispersion Strengthening of N1A1

5. Report Date

March 1986

6. Performing Organization Code

7. Author(s)

Mark Sherman and Krishna Vedula

8. Performing Organization Report No

None

10. Work Unit No

9. Performing Organization Name and Address

Case Western Reserve University

Dept. of Metallurgy and Materials Science

Cleveland, Ohio

11. Contract or Grant No.

NAG 3-387

12. Sponsoring Agency Name and Address

National Aeronautics and Space Administration

Washington, D.C. 20546

13. Type of Report and Period Covered

Contractor Report

14. Sponsoring Agency Code

505-63-01

15. Supplementary Notes

Final report. Project Manager, Robert H. THran, Materials Division, NASA Lewis

Research Center, Cleveland, Ohio 44135. Mark Sherman 1s presently at Sohio

Research Center, Cleveland, Ohio.

16. Abstract

A potential high temperature strengthening mechanism for alloys based on the

IntermetalUc compound N1A1 has been Investigated. This study forms part of an

overall program at NASA Lewis Research Center for exploring the potential of

alloys based on N1A1 for high temperature applications. An alloy containing

2.26 at% Nb and produced by hot extrusion of blended powders has been examined

1n detail using optical and electron microscopy. Interdlffusion between the

blended Nb and N1A1 powders results 1n the formation of Intermediate phases. A

fine dispersion of precipitates of a hexagonal, ordered NIAlNb phase 1n a matrix

of N1A1 can be produced and this results 1n strengthening of the alloy by Interfering with dislocation motion at high temperature. These precipitates are, however, found to coarsen during the high temperature (1300 K) deformation at slow

strain rates and this may Impose some limitations on the use of this strengthening mechanism.

17. Key Words (Suggested by Author(s))

Alum1n1de; Nickel; IntermetalUc;

Dispersion strengthening

19. Security Classif. (of this report)

Unclassified

18. Distribution Statement

Unclassified

STAR Category

20. Security Classif. (of this page)

Unclassified

- unlimited

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21. No. of pages

14

22. Price'

A02

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