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RAPID SOLIDIFICATION PROCESSING

PRINCIPLES AND TECHNOLOGIES, 1 1

1

%ThirdiConference[on

Rapid Solidification

Processing

held at the National Bureau

.of. Standards

Gaithersburg, Maryland

December, 1982

Editor—Robert Mehrabian

This document contains information which is subject to special export controls. It should not be transferred to foreign nationals in the U.S. or abroad without a validated export license. (Reference Export Administration

Regulations, Section 387.1, 1 October 1980 and Federal Register, 1 October 1980. Volume 45, No. 192,

Page 65014.)

0

RAPID SOLIDIFICATION PROCESSING

PRINCIPLES AND TECHNOLOGIES. Ill

Proceedings of the

Third Conference on

Rapid Solidification

Processing

held at the National Bureau

of Standards

December 6-8, 1 982

Gaithersburg, Maryland, US.A.

Edited by: Robert Mehrabian

PREFACE

The general aim of this Third Conference on Rapid Solidification

Processing was to review the advances in our understanding of principles as well as technological progress achieved during the past two and one-half

years. As in the previous two Conferences, emphasis was placed on rapidly

solidified crystal 1 ine materials, rather than on amorphous solids which have

been adequately treated at other meetings. However, attention was given to the amorphous state and its controlled thermal decomposition as a possible

"pathway" toward attaining unique crystalline structures.

The Conference comprised five plenary sessions, five poster sessions of

relevance to the plenary sessions, and a panel /workshop session. The topics

covered in the plenary and the corresponding poster sessions were: fundamentals of rapid solidification, microstructures and properties, surface layers

and thin filaments, consolidation and forming, and applications. A panel/

workshop session was also convened to address critical questions in rapid

solidification processing (RSP) such as: roadblock to RSP, alternatives to

RSP, and new horizons for RSP. In these proceedings no distinction has been

made between the invited lectures given at the plenary sessions and the

shorter poster presentations. Instead, the eighty-three papers have been

arranged sequentially, according to the subject areas covered. A summary of

the presentations and discussions from the panel/workshop is given at the

end of this volume.

The members of the Steering Committee are listed in front of this book.

The Committee wishes to thank Ms. Joyce Harris and Mrs. Kathy C. Stang for

organization, arrangements, and operations of the Conference, and for efficient

handling of the finances involved, Ms. CaRole Lamb for her invaluable help

in the final typing of some of the manuscripts and the overall supervision

over reproduction of this book. The Committee also acknowledges the financial

assistance provided by the following sponsors: the Air Force Wright

Aeronautics Laboratory, the Army Research Office, the Defense Advanced

Research Projects Agency, the Office of Naval Research, the National

Aeronautics and Space Administration, and the National Bureau of Standards.

Robert Mehrabian

Washington, DC
CHAIRMEN AND COMMITTEE MEMBERS

STEERING COMMITTEE

A. M. Adair

M. Cohen

L. A. Davis

M. A. Greenfield

Lt. Col. L. A. Jacobson

B. H. Kear

J. R. Manning

R. Mehrabian—Chairman

J. B. Moore

P. Parrish

D. E. Polk

B. B. Rath

R. A. Sprague

E. C. van Reuth

J. C. Williams

SESSION CHAIRMAN AND MODERATORS

I : M. Cohen and J. R. Manning

II : B. B. Rath and J. C. Williams

III: B. H. Kear and D. E. Polk

IV : A. M. Adair and R. A. Sprague

V : L. A. Jacobson and J. E. Moore

VI : Panel/Workshop (organizers and reporters)

Roadblocks to RSP: M. Cohen and J. B. Moore

M. A. Greenfield and J. R. Manning

Alternatives to RSP: J. C. Williams and R. A. Sprague

A. M. Adair and L. A. Jacobson

New Horizons for RSP: B. H. Kear and L. A. Davis

D. E. Polk and B. B. Rath

TV
AFFILIATIONS OF CHAIRMEN AND COMMITTEE MEMBERS

A. M. Adair, Air Force Aeronautical Laboratory, Wright-Patterson Air Force

Base, OH 45333

M. Cohen, Massachusetts Institute of Technology, Cambridge, MA 02139

L. A. Davis, Allied Corporation, Morristown, NJ 07960

M. A. Greenfield, National Aeronautics and Space Administration,

Washington, DC 20546

Lt. Col. L. A. Jacobson, Lawrence Livermore National Laboratory,

Livermore, CA 94550

B. H. Kear, Exxon Research and Engineering Company, Linden, NJ 07036

J. R. Manning, National Bureau of Standards, Washington, DC 20234

R. Mehrabian, National Bureau of Standards, Washington, DC 20234

J. B. Moore, Pratt and Whitney Aircraft, West Palm Beach, FL 33402

P. Parrish, U.S. Army Research Office, Research Triangle Park, NC 27709

D. E. Polk, Office of Naval Research, Arlington, VA 22217

B. B. Rath, U.S. Naval Research Laboratory, Washington, DC 20375

R. A. Sprague, General Electric Company, Cincinnati, OH 45215

E. C. van Reuth, Defense Advanced Research Projects Agency,

Arlington, VA 22209

J. C. Williams, Carnegie-Mel Ion University, Pittsburgh, PA 15213

ACKNOWLEDGMENT

This Conference was supported by grants from the following U.S. Government

agencies:

Defense Advanced Research Projects Agency

Office of Naval Research

Army Research Office

Air Force Wright Aeronautical Laboratory

National Aeronautics and Space Administration

National Bureau of Standards

V
TABLE OF CONTENTS

Page

I. FUNDAMENTALS

SOME FUNDAMENTAL ASPECTS OF RAPID SOLIDIFICATION--M. Cohen and R. Mehrabian 1

SOLIDIFICATION OF UNDERCOOLED LIQUIDS--J. H. Perepezko, Y. Shiohara, J. S. Paik, and

M. C. Flemings 28

MECHANISMS OF SEGREGATION-FREE SOLIDIFICATION--W. J. Boettinger, S. R. Coriel 1, and

R. F. Sekerka 45

MICROSTRUCTURAL CHARACTERIZATION OF RAPIDLY SOLIDIFIED MATERIALS--H. L. Fraser and

J. B. Vander Sande 56

ASPECTS OF SOLIDIFICATION IN THE RSR POWDER PROCESS— R. J. Patterson, II and C. M. Adam 70

OSTWALD RIPENING OF RAPIDLY SOLIDIFIED SOLID-LIQUID MIXTURES— P. W. Voorhees and

M. E. Glicksman 83

SOLIDIFICATION MICROSTRUCTURES OF UNDERCOOLED DROPLETS— J. J. Richmond, J. H. Perepezko,

S. E. LeBeau, and K. P. Cooper 90

MICROSTRUCTURAL STUDIES ON SUPERCOOLED SUB-MICROMETER POWDER PARTICLES— D. Schechtman,

S. D. Ridder, and R. Mehrabian 96

AN EXPERIMENTAL TECHNIQUE FOR DETERMINING THE DEGREE OF PARTICLE MELTING IN A PLASMA JET

OPERATING AT LOW PRESSURE— R. W. Smith, M. Paliwal, and D. Apelian 105

AN ATOMISTIC MODEL OF SOLUTE TRAPPING— M. J. Aziz 113

TEMPERATURE EFFECTS ON PLASMA DEPOSITION OF Ni-BASE MATERIALS—T. L. Cheeks,

M. E. Glicksman, M. R. Jackson, and E. L. Hall 118

SUPERSATURATED SOLID SOLUTIONS BY MASSIVE SOLIDIFICATION: OBSERVATIONS AND IMPLICATIONS--T. F. Kelly, M. Cohen, and J. B. Vander Sande 124

SOLUTE REDISTRIBUTION DURING SOLIDIFICATION OF SUPERCOOLED ALLOYS— S. D. Peteves,

D. D. McDevitt, and G. J. Abbaschian 129

II. MICROSTRUCTURES AND PROPERTIES

MICROSTRUCTURE-PROPERTY RELATIONS AND THEIR ROLE IN RSP MATERIALS—J. C. Williams

and J. P. Hirth 135

THE EFFECT OF RAPID SOLIDIFICATION ON THE MICROSTRUCTURES AND PROPERTIES OF

ALUMINUM POWDER METALLURGY ALLOYS--J. R. Pickens and E. A. Starke, Jr. 150

STRUCTURE-PROPERTY RELATIONSHIPS IN RSP STEELS— G. B. Olson, M. Cohen, and

G. J. Yurek 171

STRUCTURE AND PROPERTIES OF RSP COPPER-BASED ALLOYS— I. E. Anderson, R. A. Masumura,

B. B. Rath, and C. L. Void 178

FUNDAMENTALS OF RAPID SOLIDIFICATION PROCESSING IN NICKEL-BASE SUPERALLOYS—

E. H. Aigeltinger, A. F. Giamei, and D. M. Dimiduk 186

THE MECHANICAL BEHAVIOR OF RSPD MATERIALS--M. R. Jackson, R. W. Smashey, and

L. G. Peterson 198

NOTCHED TENSILE PROPERTIES OF RAPIDLY SOLIDIFIED AND INGOT-BASED ALUMINUM ALLOYS—

S. Kang and N. J. Grant 209

RAPIDLY SOLIDIFIED O.D. IN-100: STRUCTURE AND PROPERTIES— C. H. Smith and N. J. Grant 215

vi
THE MICROSTRUCTURE OF TOOL STEELS PROCESSED FROM RAPIDLY SOLIDIFIED POWDERS— S. Kumar,

A. Lawley, and M. J. Koczak 223

STRUCTURE AND PROPERTIES OF A POWDER-PROCESSED HIGH-CHROMIUM ION— S. Mohan, A. Lawley,

L. R. Woodyatt, and K. S. Brosius 229

THE EFFECT OF POWDER COOLING RATE AND THERMAL EXPOSURE ON THE TENSILE PROPERTIES OF

A RAPIDLY SOLIDIFIED NICKEL BASE SUPERALLOY— D. B. Miracle, K. A. Williams, and

H. A. Lipsitt 234

TENSILE BEHAVIOR OF RSR IRON-ALUMINIDES— M. G. Mendiratta, S. K. Ehlers,

D. K. Chatterjee, and H. A. Lipsitt 240

MICROSTRUCTURE AND MECHANICAL PROPERTIES OF RAPIDLY-SOLIDIFIED WHITE-CAST IRON

P0WDERS--L. E. Eiselstein, 0. A. Ruano, 0. D. Sherby, and J. Wadsworth 246

MICROSTRUCTURE EVOLUTION IN RAPIDLY SOLIDIFIED Ni,(Mo,X) ALLOYS— P. L. Martin and

J. C. Williams 252

COMPARATIVE FINE STRUCTURE, MICROCHEMISTRY AND MORPHOLOGY OF RENE' 95 SUPERALLOY

POWDERS PRODUCED BY ARGON ATOMIZATION AND CENTRIFUGAL ATOMIZATION— J. A. Domingue,

R. D. Kissinger, W. J. Boesch, and J. K. Tien 259

RAPID SOLIDIFICATION STUDIES OF Ni^Al ALLOYS— C. C. Koch, J. A. Horton, C. T. Liu,

O. B. Cavin, and J. 0. Scarbrough ^ 264

MICROSTRUCTURAL CHARACTERIZATION OF RSPD STRUCTURES— A. M. Ritter and M. R. Jackson 270

PRECIPITATION SEQUENCES IN A RAPIDLY SOLIDIFIED Ni-Al-Mo-W ALLOY—D. M. Dimiduk 276

EFFECT OF THE RAPID-SOLIDIFICATION METHOD ON MICROSTRUCTURES AND PROPERTIES OF Al-Li

ALLOYS— P. J. Meschter, J. E. O'Neal, and R. J. Lederich 281

AGING BEHAVIOR OF DEVITRIFIED Ti RICH ALLOYS CONTAINING METALLOIDS— S. H. Whang and

Y. Z. Lu 286

III. SURFACE LAYERS AND THIN FILAMENTS

DIRECTED ENERGY SURFACE PROCESSING OF METALS— T. R. Tucker and A. H. Clauer 291

INTERFACE RESPONSE IN LASER-MELTED ALLOYS— S. M. Copley, D. Beck, and M. Bass 303

INFLUENCE OF ELECTRON BEAM SCANNING PARAMETERS ON SURFACE MELTING AND HEAT TREATMENT

CHARACTERISTICS—P. R. Strutt, M. Kurup, and A. Taquir 312

IMPROVED IRON BASE ALLOYS MADE USING A RAPID SOLIDIFICATION PROCESS— R. Ray,

V. Panchanathan, and S. Isserow 323

BORIDE AND CARBIDE DISPERSION IN MELT SPUN SUPERALLOYS— S. C. Huang, K. M. Chang,

and E. L. Hall ‘ 337

HIGH TEMPERATURE OXIDATION OF RAPIDLY SOLIDIFIED Ni-Al-Mo-W ALLOYS— G. H. Meier,

F. S. Pettit, and A. S. Khan 348

RAPIDLY SOLIDIFIED CARBIDE POWDERS PRODUCED BY SPARK MACHINING OF Ti AND Zr—

C. S. Pande and J. D. Ayers 360

PROCESSING, MICROSTRUCTURE, AND PROPERTIES OF MELT SPUN AND CONSOLIDATED SUPERALLOYS—

R. H. VanStone, J. R. Hughes, R. E. Maxwell, H. H. Liebermann, and R. G. Rowe 365

M

2

B

2 BORIDE IN MELT-SPUN SUPERALLOYS— S. C. Huang, E. L. Hall, and T. K. Glasgow 377

PREPARATION OF METALLIC GLASSES WITH SECOND PHASE PARTICLES— D. G. Ast and

P. G. Zielinski 384

THE EFFECT OF SOLIDIFICATION VELOCITY ON THE MICROSTRUCTURE OF Ag-Cu ALLOYS—

W. J. Boettinger, R. J. Schaefer, F. Biancaniel lo, and D. Shechtman 390
NONDESTRUCTIVE CHARACTERIZATION OF RAPIDLY SOLIDIFIED A1-Mn ALLOYS--R. J. Schaefer,

M. Rosen, J. J. Smith, D. Shechtman, and R. Mehrabian 397

EXPERIMENTAL STUDIES OF THE DYNAMICS OF RAPID SOLIDIFICATION USING ACOUSTIC EMISSION--

R. B. Clough, H. Wadley, and R. Mehrabian 403

CRYSTALLIZATION KINETICS OF AMORPHOUS ALLOYS BY ULTRASONIC MEASUREMENTS--M. Rosen,

E. Horowitz, H. N. G. Wadley, and R. Mehrabian 416

OXIDATION RESISTANCE OF A FINE-GRAINED RAPIDLY SOLIDIFIED 18-8 STAINLESS STEEL--

G. J. Yurek 424

CHARACTERIZATION OF METASTABLE PHASE POWDERS PREPARED BY ELECTROLESS DEPOSITION--

L. V. Meisel, P. J. Cote, and D. M. Quinn 429

PREPARATION OF GLASSY METALS AS POWDERS VIA HYDROGEN EMBRITTLEMENT--B. C. Giessen,

S. H. Whang, and F. Dabkowski 435

AN ARC FURNACE MELT SPINNER FOR THE RSR PROCESSING OF REFRACTORY AND REACTIVE ALL0YS--

S. H. Whang and B. C. Giessen 439

A NEW HIGH STRENGTH NICKEL-BASE MICROCRYSTALLINE ALLOY VIA RAPID SOLIDIFICATION

TECHNOLOGY--V. Panchanathan, S. Sastri , and S. Hahn 443

DYNAMIC WETTING/SOLIDIFICATION PHENOMENA— R. P. I. Adler and S. C. Hsu 448

RAPID-SOLIDIFICATION PROCESSING OF TITANIUM ALLOYS— T. C. Peng, S. M. L. Sastry, and

J. E. O'Neal 452

THE MICRO-PARTICLE PROCESSOR TO PRODUCE FINE POWDERS AND COATINGS—J. Perel

J. F. Mahoney, S. Taylor, Z. Shanfield, and C. Levi 458

RAPID SOLIDIFICATION SURFACE TREATMENT OF DISCS AND RINGS BY ELECTRON BEAM

PROCESSING—M. Tuli and P. R. Strutt 464

PLASMA DEPOSITION EFFECTS ON STRUCTURES AND PROPERTIES OF RSPD SUPERALLOYS— R. W. Smith,

D. V. Rigney, J. S. Smith, and J. R. Rairden 468

THE MICROSTRUCTURE OF RAPIDLY SOLIDIFIED NiAl-Cr QUASIBINARY EUTECTIC— D. Shechtman,

T. Z. Kattamis, F. S. Biancaniel lo, and W. J. Boettinger 476

WEAR PROPERTIES OF RAPIDLY SOLIDIFIED SURFACES OF A PSEUDO-BINARY Fe-TiC ALLOY—

P. R. Strutt, B. G. Lewis, S. F. Wayne, S. L. Rice, and B. H. Kear 483

IV. CONSOLIDATION AND FORMING

FABRICATION OF ORIENTED SINGLE CRYSTAL WAFER STOCK FROM Ni-Al-Mo-X ALLOY POWDERS—

R. E. Anderson, H. A. Chin, S. E. Hughes, and B. B. Rath 488

DEFECT ANALYSES OF P/M SUPERALLOYS—W. H. Chang, H. M. Green, and R. A. Sprague 500

INTERPRETIVE REVIEW OF COMPARATIVE MICROSTRUCTURES OF POWDERS AND CONSOLIDATES OF P/M

SUPERALLOYS PRODUCED BY MODERN ATOMIZATION PROCESSES— R. D. Kissinger and J. K. Tien 510

CONSOLIDATION AND FORMING OF P/M POROUS BILLETS— H. L. Gegel , J. S. Gunasekera,

S. M. Doraivelu, J. C. Malas, J. T. Morgan, and L. E. Matson 519

PROCESSING OF P/M ALUMINUM ALLOYS X7090 AND X7091— F. R. Billman, J. C. Kuli, Jr.,

G. J. Hi 1 deman, J. I. Petit, and J. A. Walker 532

POST RSPD PROCESSING EFFECTS ON GRAIN GR0WTH--R. A. Amato, M. R. Jackson, and

P. A. Siemers 547

Ni-Mo-B ALLOYS: METALLIC GLASS TO DUCTILE CRYSTALLINE SOLID— S. K. Das, L. A. Davis,

J. R. Y. Wang, and D. Kapoor 559
MICROSTRUCTURAL MODIFICATION OF ALUMINUM-SILICON ALLOY POWDERS DURING DYNAMIC

CONSOLIDATION—W. H. Gourdin and J. E. Smugeresky 565

MICROSTRUCTURE OF RAPIDLY SOLIDIFIED STAINLESS STEEL AND SUPERALLOY POWDERS: A COMPARISON OF POWDER-MAKING TECHNIQUES—J. E. Smugeresky 572

CONSOLIDATION, THERMOMECHANICAL PROCESSING, AND MECHANICAL PROPERTIES OF RAPIDLY

SOLIDIFIED, DISPERSION-STRENGTHENED TITANIUM ALLOYS— S. M. L. Sastry, T. C. Peng,

J. E. O'Neal, and L. P. Beckerman 579

EFFECT OF AGING ON THE MICROSTRUCTURE OF RAPIDLY COOLED Ti-5Al-2.5Sn WITH Si ADDITIONS—

A. G. Jackson, T. F. Broderick, F. H. Froes, and J. Moteff 585

ON THE TEXTURE DEVELOPMENT— DIRECTIONAL RECRYSTALLIZATION INTERACTION OF RAPIDLY

SOLIDIFIED Ni-Mo-Al-W AND Ni-Cr-Al-W ALLOYS— H. A. Chin, S. E. Hughes, J. S. Andrus,

and J. L. Larson 590

THERMAL STABILITY OF HIGH TEMPERATURE P/M ALUMINUM ALLOYS— Y.-W. Kim, M. M. Cook, and

W. M. Griffith 597

EVALUATION OF VOIDS IN POWDER CONSOLIDATES USING LONG WAVELENGTH NEUTRONS—

M. Fatemi and B. B. Rath 603

V. APPLICATIONS

RAPID SOLIDIFICATION— PAST PROMISES AND CURRENT REALITY— A. M. Adair, L. A. Jacobson,

E. C. van Reuth, and J. Dickson 609

AEROSPACE STRUCTURAL APPLICATIONS OF RAPIDLY SOLIDIFIED ALUMINUM-LITHIUM ALLOYS—

R. E. Lewis, I. G. Palmer, J. C. Ekvall, I. F. Sakata, and W. E. Quist 615

APPLICATIONS OF RAPIDLY SOLIDIFIED ALUMINUM ALLOYS AT ELEVATED TEMPERATURES—

C. M. Adam, J. W. Simon, and S. Langenbeck 629

GAS TURBINE APPLICATIONS OF RST ALLOYS— D. B. George, R. G. Bourdeau, and J. A. Miller 640

AIRCRAFT ENGINE GAS TURBINE COMPONENT FABRICATION CONCEPTS USING RSPD—A. M. Johnson,

J. S. Kelm, R. W. Smashey, D. V. Rigney, and T. G. Wakeman 650

RECENT APPLICATIONS OF ULTRA RAPID SOLIDIFICATION TECHNOLOGY (URST) MICROCRYSTALLINE

AND AMORPHOUS POWDER MATERIALS—J. H. Kushnick and D. L. Sawhney 662

SHOCK COMPACTION OF FERROUS ALLOY POWDERS— T. J. Ahrens, D. Koska, P. Kasiraj,

T. Vreeland, Jr., A. W. Hare, F. D. Lemkey, and E. R. Thompson 672

ULTRASONIC GAS ATOMIZED POWDERS OF MAGNETIC ALLOYS— R. C. O' Handley, P. Domalavage,

B. W. Corb, Y. Hara, and N. J. Grant 678

PRODUCTION APPLICATIONS OF RAPIDLY SOLIDIFIED TOOL STEELS, SUPERALLOYS, TITANIUM

ALLOYS, AND CORROSION-RESISTANT ALLOYS— J. H. Moll and F. J. Rizzo 686

COMMERCIAL RAPID SOLIDIFICATION (CRS) PROCESSING FOR ADVANCED TURBINE DISK APPLICATIONS—

R. Beck, W. Kent, and C. S. Kortovich 692

VI. PANEL/WORKSHOP ON CRITICAL QUESTIONS IN RAPID SOLIDIFICATION PROCESSING

ROADBLOCKS TO RAPID SOLIDIFICATION PROCESSING—T. Tietz, M. DeCrescente, S. Copley,

and T. Kearns 697

ALTERNATIVES TO RAPID SOLIDIFICATION PROCESSING— C. Kortovich, C. Blankenship,

L. Johnson, and W. Boesch 699

NEW HORIZONS FOR RAPID SOLIDIFICATION PROCESSING--R. Schaefer, P. Strutt, R. Maringer,

B. Giessen, H. Fraser, and M. Flemings 701

IX
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SOME FUNDAMENTAL ASPECTS OF RAPID SOLIDIFICATION

M. Cohen

Massachusetts Institute of Technology

Cambridge, MA 02139

and

R. Mehrabian

National Bureau of Standards

Washington, DC 20234

ABSTRACT

Detailed consideration is given to the interplay of supercooling and rapid heatextraction rate in controlling the mode and kinetics of solidification processes. Achievable external heat extraction rates impose strict limitations on solid/liquid interfacial

velocities in the absence of bulk supercooling. On the other hand, for the case of supercooling, the recalescence rate at the interface almost always predominates the external

heat-extraction rate. Particular attention is directed to the resulting degrees of solute

partitioning at the moving solid/liquid interface, and also to the special conditions which

achieve partitionless (or massive) solidification. The required conditions for the latter

L /(y are an interfacial temperature below the T^ line, and an interfacial velocity which

exceeds the diffusive velocity of solute in the liquid. The complete trapping of solute

elements in the solidifying phase is potentially of major significance from the standpoint

of microstructural control and alloy design because extremely high supersaturations of

very stable (low solid solubility) phases can thereby be obtained. Subsequent precipitaO tion leads to ultrafine (~ 100 A) dispersions of second-phase particles which resist

coalescence and re-solution in the solid matrix at elevated temperatures. Such dispersions

are unusually effective in pinning grain boundaries and inhibiting grain growth.

A further question now arises concerning an appropriate definition of rapid solidifi - cation , in view of the fact that the conventionally specified "cooling rate" does not

always properly express a "solidification rate." Accordingly, an attempt will be made to describe various modes of solidification in terms of interfacial velocities.

Introduction

At the Second International Conference on Rapid Solidification Processing: Principles and Technologies [1] about three years ago, considerable emphasis was directed to the microstructural refinement achieved in cellular and dendritic morphologies by rapid cooling from the liquid state. Even

though almost equilibrium partitioning of solute and impurity atoms may still occur at the advancing

solid/liquid interfaces, there is substantial refinement in the scale of the resulting microsegregations and second-phase dispersions due to the fact that rapid cooling allows less time for the

coarsening of morphological features during the solidification growth process. The conditions for

such rapid cooling are typically attained by setting up positive temperature gradients in the melt

ahead of the moving interfaces, and efficiently extracting heat through the solid phase to the external surroundings as the solidification proceeds.

In the present paper, more attention is given to solidification under conditions of bulk melt

supercooling such that the temperature gradients ahead of the solidifying front become negative, and

then the heat of solidification can flow into the remaining supercooled liquid even more rapidly than

to the surroundings. This mode of solidification leads to high interfacial velocities and, in turn,

to solute trapping in the advancing solid, i.e., nonequilibrium partitioning at the interfaces. We

shall be especially interested in the limiting case in which the partition ratio reaches unity, thus

attaining partitionless (or massive) solidification. Under these circumstances, the solid phase not

only inherits the same compositional uniformity as prevails in the melt, but it may also be highly

supersaturated with respect to very stable second phases which have been previously dissolved in the

liquid state. This supersaturation in the solid state then permits the precipitation of extremely

fine dispersions which are otherwise impossible to dissolve in the solidified alloy.

Some microstructural consequences of massive solidification will then be discussed, particularly

from the standpoint of high-temperature grain-growth resistance arising from the grain-boundary

pinning effect of stable second-phase dispersions. These characteristics have important implications

in structure/property relationships.

1
Modes of Solidification

In dealing with rapid solidification phenomena, it is well to recognize two conditions of

extreme heat flow. In the first, the solid/liquid interface moves into a superheated melt, and the

latent heat of fusion is withdrawn through the solid. Typical cases are atomization, melt spinning,

and surface melting when there is no bulk supercooling. The resulting achievable interfacial velocities are then limited by the rate of external heat extraction and the thickness of the material being

processed, assuming that no restraints are imposed on the interfacial motion by local atomic rear- rangements [2]. The second situation involves solid growth into a supercooled melt, where most of

the latent heat of fusion can be absorbed by the liquid itself and so the external rate of heat

extraction is less important. In such instances, the temperature of the liquid rises (recal escence)

and must be taken into account in the accompanying heat-flow kinetics. Nevertheless, sufficiently I high interfacial velocities can be attained in this type of solidification to reach complete solute I trapping. '

1 In this paper, we shall describe modes of solidification in two ways: (1) according to the

morphology of growth (i.e., plane-front, cellular, or dendritic), and (2) according to compositional

changes at the solid/liquid interface (i.e., with equilibrium, partial, or no partitioning of solute

atoms between the liquid and solid phases). Clearly, these are related but different kinds of classification. The variations in solidification behavior will depend on the transport of heat and solute

from the interface, and in some cases on atomic rearrangement processes at the interface. It is

necessary to take these interdependent factors into account in order to deal with the microstructural

consequences of rapid solidification and, hence, with processing/structure/property relationships.

Special emphasis will be placed on the solid/liquid interfacial velocity which is generated under the

operative conditions of nucleation (usually heterogeneous) and heat flow.

We shall first review some fundamentals of sol idification when compositional equilibrium is maintained locally at the moving growth front (no supercooling). Figure 1 shows that the interfacial

morphology is a function of the thermal gradient (Gj^) in the melt at the interface and the interfacial velocity or growth rate (R). The slope of the line AB is given by [3]:

AVERAGE COOLING RATE f = GlR

T2 >

INTERFACIAL VELOCITY, R

PLANAR

INTERFACE

CELLULAR

INTERFACE

DENDRITIC

INTERFACE

Figure 1. Mode of solidification as a function of temperature gradient in the liquid (G|^) and interfacial velocity or growth rate (R). Microstructural refinement due to increasing the

cooling rate (T=G|^R) is indicated by the arrow [1]. Line AB is given by eq. (1).

2
( 1 )

il

ir

[I

II _ mLC^(k^-l)

il where = solute concentration in the bulk liquid

j D|^ = diffusity of solute in liquid,

i k^ = equilibrium partition ratio (< 1 as in figure 12), and

m|^ = slope of liquidus line (negative as in figure 12).

Ij When G^^/R is greater than the right-hand side of eq. (1), the interface advances on a planar (or

I smooth) front as heat is withdrawn through the solid phase. Under these conditions, the interface

I tends to be stabilized in that chance protuberances on the solid front are "melted back" by the steep

I temperature gradient in the liquid relative to the interfacial velocity. Moreover, at steady state,

' there will be no microsegregation of solute atoms behind the advancing front, except possibly at the

! very end of solidification.

In contrast, when Gj^/R is less than the right-hand side of eq. (1), the interface becomes morphoi logically unstable in as much as protuberances tend to grow ahead preferentially into combinations

A of liquid temperature and composition where the melt would otherwise be supercooled (or supersatuI] rated) with respect to the solid phase. Another way of putting it is that, when R is increased rela-

:j tive to G|^, the solute concentration gradient ahead of the interface can become steep enough to

' induce morphological instability. This situation leads to cellular or dendritic growth morphologies,

j| with solute-enriched microsegregations building up in the remaining liquid between the cells or

Ij dendritic arms*. These microstructural regions are also locations where other solid phases usually

I' precipitate near the end of the solidification process.

il For the case of cellular or dendritic solidification, the temperature gradient and interfacial

velocity in figure 1 could represent averages^ for the liquid-plus-solid "mushy" region. The product

G|^R then expresses the average cooling rate (f) and, as the sloping arrow indicates, the cellular or dendritic arm spacings decrease in scale with increasing t. This trend comes about because interfaI cial energy associated with these morphological features causes them to coarsen as a function of the

available time for solidification. Accordingly, such coarsening can be reduced by increasing the

average cooling rate during cellular or dendritic solidification. Thus, even with equilibrium partitioning being substantially maintained at the moving interfaces, accelerated cooling can provide

significantly smaller distances between microsegregated regions and precipitated second phases in

the final microstructure.

Figure 1 must be modified in important ways at very high interfacial velocities (very rapid

solidification rates), and then the regime of morphological instability is capped off, as illustrated

in figure 2. When R is sufficiently large, beyond line EF, morphological stability is reestablished

and plane-front growth ensues, still under the assumption of equilibrium partitioning or no supercooling. The reason for this interesting behavior is that, when the limiting interfacial velocity

at EF is approached, potential perturbations in the growth front become so fine in scale that surface

tension imposes a dominant stabilizing influence toward establishing a smooth interface. Coriell

and Sekerka [5] have solved this complex solidification problem, taking into account these capillary

forces as well as the solute and thermal diffusion fields in front of the interface (including the

effect of latent heat evolution). The solute field is always destabilizing, while the capillary

forces always oppose the growth of perturbations. The thermal field will be stabilizing if the net

' flow of heat flow is toward the solid. Calculated results for an aluminum-silicon alloy are reprei sented in figure 2 [6].

i

j The line AB in figure 2 has the same significance as in figure 1, but the regime of cellular

and dendritic solidification is now confined to the dome-shaped region in figure 2. Therefore, planej front growth is obtained not only to the left of line AB, as discussed in connection with eq. (1),

' but also above and to the right of line BFE. The condition represented by line EF is often referred

I to as absolute morphological stability. The corresponding velocity is given by [5,7]:

i

' *The morphological instability attending cellular or dendritic solidification is often termed constiI tutional supercooling [4]. However, local compositional equilibrium at the advancing growth front is

I still assumed, and so this is not supercooling in the usual sense.

I

3
INTERFACIAL VELOCITY, R (m/s)

Figure 2 Regimes of liquid temperature gradient (Gj^) and interfacial velocity (R) for plane-front

growth (morphologically stable interface) and for cellular/dendritic growth. The boundary

ABFE between these two regimes is calculated for an Al-0.1 wt.% Si alloy, assuming equilibrium partitioning of Si at the moving interface [6]. Cross-hatched band indicates estimated

interfacial velocity which is fast enough to achieve partitionless solidification.

R(at EF) =

ni|_DL(kE-l)Co

( 2 )

where = melting temperature of the interface without solute, and

r = is the ratio of solid/liquid surface energy to the latent heat of fusion per unit volume. !

According to eq. (2), the boundary EF is shifted to higher velocities with increase and with

|

decreasing k^. Some typical calculations in table I for aluminum alloys, with solute contents in

|

the range of their solid solubility limits, indicate that interfacial velocities of the order of 3 to ^ 13 m s~^ are required to achieve absolute morphological stability, under the assumption of equilib- J rium partitioning. However, such velocities lie beyond the range that can be attained by maximum ' external heat-extraction rates or feasible cooling rates, as shown by the heat-flow calculations in if table II for 10 pm melts of A1 , Fe, and Ni when cooled by processes of atomization, melt spinning,

and substrate self-quenching. ,

4
Table I. Interfacial Velocity for Absolute Morphological

Stability in Binary Aluminum Alloys; Equilibrium

at Interface

Al-Cu Al-Fe Al-Si

Maximum Solid

Solubility (at.%) 2.5 0.025 1.6

(kg - 1)

2 "’l

•^E

“V 265 -V 7000 -V 388

R (m s ^) ^ 13.2 3.5 “V 12.4

kg Equilibrium Partition Coefficient

mg Slope of Liquidus Line

Table II. Cooling-Rate and Interfacial-Velocity Limitations in RSP

of 10 pm Size A1 , Fe, Ni Specimens; No Supercooling

Heat Transfer

Coefficient

Avg. Max. Cooling Rate

in Liquid

Max. Interfacial

Velocity

Fraction

Sol id

Atomization 4 -2 -1

-V 5x10 W m K -V 7x10^ K s"^ 0.1 -0.27 m s"^* 0.1 to 0.95

Melt Spinning ^ 10^ -V- 2.5 X 10^ 0.6 -0.75* 0 to 1.0

Self Quenching Tends to infinity

-V 10® 1.5 -2.3 Al**

0.35-0.65 Fe,Ni

0.1 to 1.0

* Independent of Size for Newtonian Cooling

** Inversely Proportional to Melt Thickness

The equilibrium-partitioning assumption is not realistic at the high growth rates in question.

Various relationships have been proposed for the dependence of nonequilibrium partitioning on interfacial velocity [8]. For present purposes, a convenient form is:

RA/D. +kp ^ RA/Dg+1 ’

where A = effective diffusion distance in the liquid concentration

gradient at the interface, -v- 10 x interatomic distance (0.

k(R) is plotted as a function of R in figure 3 for Al-Si alloys, taking Dg = 2.3 x 10“^ m^ s"^, and

C ~ 2.3 X 10“^° m. It will be seen that k(R) increases progressively (signifying more solute trapping) with increasing interfacial velocity. If the form of eq. (2) is now adopted for general values

of k (instead of only for kg), the resulting connection between alloy composition and the velocity

for absolute morphological stability can also be plotted in figure 3. Because of the enhanced values

of k, absolute stability may be attained at growth rates of less than 1 m s“^ even with concentrated

alloys (up to 5 wt.% Si), but the solidification is then far (k s 0.5) from being partitionless.

5
INTERFACIAL VELOCITY, R (m/s)

Figure 3 Calculated changes in partition coefficient (k(R)) and composition (C^) for partitionless

solidification as a function of interfacial velocity (R) in Al-Si alloys. k(R) vs R is

given by eq. (3) and vs R is given by eq. (2) with k(R) substituted for k^.

Shortcomings of the above approach include the fact that eq. (3) has been derived for dilute

alloys. Furthermore, substitution of k(R) for k^ in the absolute morphological stability eq. (2)

should only be made when the change of partition coefficient with velocity (3k(R)/8R) is small and/or

the interfacial kinetics are sluggish [9]. The latter condition implies a small interfacial kinetic

coefficient = 3R/8T^. ^ 1.0 cm s“^K“^, where T^. is the interfacial temperature.

In general, solidification under conditions of absolute morphological stability is not neces- sarily partitionless. From its very nature, the plane-front growth characteristic of absolute

stability means that there will then be no lateral partitioning, but partitioning may still take

place between the solid and the liquid in the direction of growth. In fact, in order to avoid the

latter type of partitioning, the interface will have to move at a rate in excess of the diffusive

velocity (Dj^/^), or at least with an effective diffusive velocity such as D|^/A. Dj^/lO This range . of velocities is shown as a cross-hatched bar in figure 2 and, for the case at hand, happens to lie J beyond the velocity where absolute stability should set in. Experimental evidence of such a situa-

| tion has been observed [10] and will be discussed later.

| It is worth noting that in the modes of solidification which have been analyzed in this section,

| the net heat flux is always from the liquid to the solid. In contrast, when bulk melt supercooling \ is involved, as will be covered in the following section, the heat flow may be reversed, and then

the latent heat of fusion can be discharged into the supercooled liquid. This circumstance contrib- , utes substantially to the interfacial velocities that can be attained during, the solidification

process. )

Solidification of Supercooled Droplets

j

i In this section, the thermal history during solidification of supercooled droplets is described

using enthalpy/temperature diagrams. Newtonian heat-flow conditions, i.e., negligible temperature j- gradients inside the droplets, are first assumed in order to portray the thermal history of pure

aluminum droplets during solidification. Conditions for hypercooling are described and related to i

6
i the overall recalescence of the droplet. It is shown that the Newtonian model is reasonable under

I either of two conditions: first, when the atomic rearrangement process at the interface is slow

j compared to thermal diffusion (motion is then "interface limited"), and second, when the number of

' nucleation events is large. Next, the general case of solidification of a supercooled droplet is

! described when the temperature gradients are not negligible. Sample calculations are presented from

reference [11] to describe temperature profiles and interfacial velocities as a function of initial

supercooling, droplet size, and interfacial kinetics. The necessary conditions for parti tionl ess

solidification of an alloy are given in terms of curves and interfacial velocities. Finally,

microstructural observations on submicron and commercial-sized powders are correlated with the predictions of the heat-flow solidification model.

Three general "paths" portraying the thermal history for Newtonian solidification of pure aluminum droplets are designated in the enthalpy/temperature diagram of figure 4 [10]. These 'paths

express the relative enthalpy per mole of the system (H - as a function of temperature and

j weight fraction solid (g) based on the following equation:

TEMPERATURE (K)

X 600 700 800 900

DIMENSIONLESS TEMPERATURE

9 = K\_ (T—T 1^ ) /A H

(O

X

>-Q.

<

X

K

LU

LU

>

<

-I

liJ CC

Figure 4 Enthalpy- temperature diagram for aluminum, showing three solidification paths. In path

C, 0,. is the dimensionless nucleation temperature on supercooling. Other symbols are

N defined in text [10].
H (4) - = [AHm + K|_(T - T^)][l - g] + K^(T - T^)g

where: H = enthalpy per mole of the system

H$m = enthalpy per mole of the solid phase at the melting

temperature,

AHu = latent heat of fusion,

M Ks,Kl = specific heats of solid and liquid (assumed constant

unless otherwise indicated).

The two terms on the right-hand side of eq. (4) represent the enthalpies of the liquid and solid

fractions in the droplet. The relationship in eq. (4) can also be written in terms of dimensionless

enthalpy and temperature 6, where:

point (T ^ T|^^, g = 0), 6^0 and i\> ^ 1.0. On the other hand, for a volume of solid at or below its

melting point (T ^ T^^, g = 1.0), 0^0 and i|j 5 0. A volume of partially solid material at Tj^ (6 = 0)

will have a dimensionless enthalpy equal to the fraction of liquid (1 - g). This is because lines of

constant fraction solid intersect the melting temperature axis at i|< = 1 " 91 the lines for g = 0,

g = 0.5, and g = 1.0 are drawn in figure 4.

Isothermal solidification in the diagram of figure 4 follows a verticle "path". For example,

in the simple case of a liquid droplet solidifying without supercooling (T = Tj^, 0 = 0), "path" A is

followed from ijj = 1.0, g = 0to4i = 0, g = 1.0 where the droplet is completely solid. A horizontal

"path" represents adiabatic solidification without any external heat extraction. If a liquid droplet

is supercooled to the temperature Tu (0 = -1.0), an amount of enthalpy equivalent to the heat of

fusion is withdrawn prior to nucleation, and the droplet can then solidify adiabatical ly along "path"

B. In general, the critical hypercooling temperature, T^, of a pure metal is defined by:

A pure metal droplet is then said to be hypercool ed when its temperature at nucleation, T^, is

equal to or less than T^ (0|^ ^ -1.0).

When nucleation is suppressed below the melting temperature Tj^, but not into the hypercooled

region, the system can follow an intermediate "path" such as C in figure 4, starting at the nucleation temperature -1 <

0^^ < 0 on the liquid (g = 0) line and ending at some point on the solid (g = 1)

line. Two heat-flow regimes are identified when the thermal history follows "path" C. The first is

a "rapid solidification" stage accompanied by recal escence , where the supercooled droplet absorbs

most of the liberated heat of fusion, and external heat extraction from the droplet surface plays a relatively minor role. For the case at hand (Newtonian model), no distinction is made between the

interfacial temperature and the remainder of the solidifying droplet. Near-adiabatic solidification

exhausts the available supercooling and the partially solid droplet continues to solidify at a much

reduced rate dg/dt on a near isothermal (vertical i]; - 0) "path." This slower solidification stage

is predominantly controlled by the rate of heat extraction through the droplet surface.

For the Newtonian model, the thermal history and the smooth-front interfacial velocity of a single nucleus formed at the droplet surface are only functions of the nucleation temperature 0j^ and

the dimensional ratio K/Bi [11]. K is a dimensionless interfacial kinetic coefficient defined as:

= (H - Hj|^)/AH|^ , and

e = K|_(T - T^)/AH^ .

(5)

( 6 )

These parameters are also shown in figure 4. For a volume of liquid at or above its melting

T,M

(7)

T,H

K = (r^of^)(AH^/K^) ( 8 ) where r^ = droplet radius, = thermal diffusivity of liquid;

8
Bi = (r^/k^)/(l/h), (9)

where kj^ = thermal conductivity of liquid, and

h = heat transfer coefficient at droplet/environment interface.

Bi , the Biot number, is the ratio of resistance to heat flow inside the droplet to that at the

droplet/environment interface. Thus, the K/Bi parameter is a measure of the degree of predominance

of the recalescence rate over the rate of heat loss to the environment.

Figure 5 shows the effect of initial supercooling, on the ij) - 9 "path" for K/Bi = 100 [11].

This value corresponds to a relatively low interfacial kinetic coefficient (K|^ = 2 cm s~^K~^) and a very high rate of external heat extraction (h = 2 x 10^ W from aluminum droplets [11]. The

TEMPERATURE (K)

DIMENSIONLESS TEMPERATURE

6 = AC

|_

(T—T

^ ) /A H

Figure 5 Calculated enthalpy-temperature curves for the solidification of aluminum droplets after

different degrees of supercooling over the range of 0|^ = -0.01 to -1.0. Newtonian heat

flow is assumed; the ratio K/Bi = 100 corresponds to extremely rapid cooling conditions

for atomized aluminum droplets [11].

9
curves demonstrate that as the supercooling prior to nucleation increases, so does the extent of adiabatic solidification; this manifests itself in flatter ijj - 6 curves and larger fractions of solid

at the end of recalescence. The same trend is observed when the K/Bi ratio is raised for a given

initial supercooling An increase in K/Bi signifies faster interfacial kinetics or increased

solid/liquid interfacial area (i.e., increased number of nucleation events or dendritic vs smoothfront solidification), and a corresponding decrease in the contribution made by the external heat

extraction (i.e., more adiabatic solidification during recalescence). Figure 6 illustrates this

trend; here a factor 2r^/d|^j^ is introduced to represent multiple nucleation events is the mean distance between adjacent nuclei). Starting with a K/Bi ratio of 50 and a single nucleus, the successive curves indicate the thermal histories of droplets with 1, 8, 64, and 1000 nucleation events,

respectively [11]. It is important to note that with increasing number of nuclei, the distances over which temperature gradients occur decrease from r^ to thus improving the validity of the

Newtonian approximation.

TEMPERATURE (K)

700 800 900

DIMENSIONLESS TEMPERATURE ^ = l(T-T,^)/AH|^

Figure 6 Effect of the number of nucleation events on the thermal history of supercooled droplets

solidifying under Newtonian conditions [11].

For the case of one or a few nucleation events, the Newtonian assumption becomes invalid when

substantial supercooling prior to nucleation is achieved-- the local recalescence of the interface is

likely to be faster than the overall recalescence of the droplet. In the extreme, when the rates of

interfacial rearrangement processes are high, the crystal growth into the bulk supercooled liquid

approaches the "heat transport limited" regime. In a recent paper. Coriel 1 and Turnbull [2] estimated interfacial kinetic coefficients as high as K^^ = 200 cm s”^K"^ for growth of pure nickel

dendrites into bulk supercooled melts. This kinetic coefficient was based on the speed of sound in

liquid nickel which establishes an upper bound for the collision frequency of atoms from the liquid

to the i nterface.

Figure 7 shows calculated [11] temperature _profi 1 es , at discrete fractions of solid, along the

growth axis of a 10 pm droplet supercooled to

0^^ = -0.5. Substantial temperature gradients in the

liquid, in front of the interface, are evidenced even for as low an interfacial kinetic coefficient

as K|^ = 2 cm s~^K~^. The corresponding calculated interfacial profiles and velocities are illustrated in figure 8 along with two other droplet sizes of 1 pm and 100 pm in diameter [11]. In all

10
0

-0.3

-0.1

II t -0.2

II 05

UJ

QC3

H

<

QCUJ0. S

UJH

(/>

(/> UJz

o

0)

z

UJ

-0.4

-0.5

- 0.6

- 900

- 850 ^

UJ

QC3

I-<

QC

UJ

0.

S

800 UJ

I-

- 750

-1.0 -0.8 -0.6 -0.4 -0.2 0.0 0.2 0.4 0.6 0.8 1.0

DIMENSIONLESS DISTANCE ALONG GROWTH AXIS

Figure 7 Thermal profiles developed in a supercooled aluminum droplet during solidification from

a single nucleation event at the droplet surface [11].

three sizes, the interfacial velocity (R) decreases with increasing fraction solid (g) due to reca- lescence as expected, but remains substantially faster than the highest values of 0.1 to 0.27 m s

^ reported for the case of no supercooling in table II.

Figure 9 presents calculated [11] interfacial velocities as a function of fraction solid for

different initial supercoolings. R is larger than 1.5 m s~^ throughout the solidification (recalescence) process when the droplet is hypercooled (6|^ S -1.0) prior to nucleation.

For the non-Newtonian cases noted above, the enthalpy/temperature curves are evaluated in terms

of the weighed averages of and 6 over the entire droplet. The thermal history can no longer be

expressed just in terms of 6|^ and K/Bi , but depends on the particular values of droplet diameter,

interfacial kinetic coefficient, and heat-transfer coeffi_cient at the droplet surface. Figure 10

indicates how increasing the droplet size shifts the iji - 6 curve away from the Newtonian case [11].

Interfacial recalescence introduces large temperature gradients, especially in the coarser particles.

11
b. DIAMETER = 10/im

6 ="0.5 K|y^ = 2 cm s"^ K"^

K/Bi = 400 h = 5 X 10^ Wm”^ K

Figure 8 Interfacial positions and velocities as a function of fraction solidified (g) for three

sizes of atomized aluminum droplets, supercooled to 6|^ = -0.5. Calculations based on non- Newtonian heat-flow model [11].

Locations far removed from the interface are not influenced by the heat released at the interface,

and may even cool slightly below their original supercooling temperature due to heat extraction by

the surroundings. It can be postulated that the larger particles maintain these supercooled regions

for longer times, and are therefore better candidates for multiple nucleation events.

Increasing the interfacial kinetic coefficient (K^^) increases the growth velocity, interfacial

recalescence rate, and temperature gradients. Figure 11 shows calculated [11] temperature profiles

12
7

CO ALUMINUM POWDER

Figure 9 Calculated interfacial velocities (R) as a function of volume fraction solidified (g) for

different initial average supercoolings (6|yj)-

in 1 and 10 pm diameter droplets at g = 0.1. Increasing the kinetic coefficient from 2 to

50 cm s~^K~^ increases the maximum temperature difference inside a 1 pm droplet from 7 to_82 K, and

in a 10 pm droplet from 52 to 154 K, both for an initial average supercooling of 182 K

(0^^ = -0.5)

[11]. The larger interfacial kinetic coefficient is in line with current estimates for "heat transport limited" growth into bulk supercooled melts [2]. Again, note the near-absence of temperature

gradients in the solid compared to the large negative gradients in the liquid. The influence of

on interfacial velocities is given in table III. For high interfacial rearrangement processes (Kj^ =

50 cm s“^K~^), the interfacial velocity (R) remains above 1 m s~^ for over half of the solidification

process. Minimum R values of 0.14 m s~^ are calculated after the loss of bulk supercooling at which time interfacial velocities are controlled by the rate of heat extraction from the droplet surface. The data in figure 9 and table III illustrate that in atomized droplets the supercooling prior

to nucleation is perhaps the only realistic way of achieving interfacial velocities necessary for

parti ti onl ess sol i di f i cati on

Extension of the above kinetic considerations to the solidification of alloy (in contrast to pure metal) droplets requires a brief description of thermodynamic restrictions on partitionless

solidification. Consider the binary phase diagram in figure 12. The diagram includes metastable

L/cy extensions of the solidus a/L and liquidus L/a curves (dashed lines), and the T curve for

• L/01 parti ti onl ess solidification. The T^ curve is the locus of temperatures and compositions where

the a-solid and liquid have the same free energy. Partitionless solidification to a crystalline

solid can only occur at temperatures below this curve. A liquid solution, such as C^' ,

' , and

C ''', can solidify to the a-phase of the same composition if the interfacial temperature lies below

L/a the T^ curve. Kinetic considerations also require interfacial velocities in excess of the diffusive velocity of the solute in front of the interface; otherwise partitioning will occur in the

13
}

!

750

TEMPERATURE (K)

800 850 900

Figure 10 Departure from Newtonian thermal history as droplet size increases. Note that dimensionless enthalpy and temperature are now averaged over the entire droplet [11].

Table III. Effect of Kinetic Coefficient on Interfacial Velocity, R(m s"^);

from Reference (11).

Fraction

Solid, g

2

K|^ (cm s

8 50

0.1 2.4 6.4 14.6

0.3 1.5 4.0 10.1

0.5 0.41 0.64 1.3

Mi nimum

Velocity

0.14 0.14 0.12

Particle Diameter 10 pm, 0^^ = -0.5, h = 5 X 10^ W

-2./-1 m K

14
DIMENSIONLESS DISTANCE ALONG GROWTH AXIS

HI

oc3

I-<

QC

LU

Q.

UJ

hFigure 11 Temperature profiles during solidification of supercooled droplets, at a constant fraction

solid g = 0.1, for two droplet sizes and two values of the interfacial kinetic coefficient

stable or metastable two-phase (a + liquid) region. Calculations presented earlier show that these

two conditions can be met in significantly supercooled droplets. The alloy C '''' cannot be

\ /a • * • supercooled below the T curve and so partitionless solidification to the a-phase is thermodynamically forbidden in this instance.

The degree of supercooling required to achieve complete partitionless solidification of an alloy

system must be such that the growth velocity will remain sufficiently fast to avoid solute partitioning despite the rising interfacial temperature due to the evolving heat of fusion. Once solidification sets in after supercooling, both the degree of supercooling and resulting interfacial velocity

will decrease progressively because. of the attendent recalescence. Accordingly, during this process,

the growth rate may fall below that needed for complete solute trapping at some temperature, T^,

where T^ < In other words, the entire solidification will be partitionless if the interfacial

temperature remains below T^. Hence, complete solute trapping (massive solidification) requires an initial supercooling temperature, Tj^:

15
K^dT ^ AH^ ( 10 ) /

T

N Experimental observations on fine (less than 1 pm diameter) electron-transparent powders of

Al-4.5 wt.% Cu and Al-6 wt.% Si alloys confirm the above concepts [10], The segregated cellular

structures in figure 13 are evidence that this Al-4.5 wt.% Cu alloy particle soli_dified in a more-orless conventional manner. Solidification was probably initiated at (0|^ ~ 0), with interfacial breakdown soon thereafter due to constitutional supercooling. The thermodynamic conditions

at the solid/liquid interface, i.e., temperature and composition, would seem to have corresponded

fairly closely to those dictated by the equilibrium phase diagram. It appears that solidification

occurred over the temperature range T|^ ^ T ^ T^ with the solid and liquid compositions and temperatures at the interface following the liquidus and solidus lines, respectively.

Figure 14 shows the Al-rich corner of the Al-Si binary phase diagram. The diagram includes

\-/

metastable extensions of the solidus and liquidus lines as well as the T^ curve. Figures 15 and 16

show two particles of Al-6 wt.% Si alloy that solidified with different degrees of initial bulk

supercooling: Tj^ 664 K

(0^^ -0.5), and Tj^ < 482 K(0j^ < -1.0).

The micrograph of figure 15 suggests that region A_underwent partitionless solidification while

both the interface and the average droplet temperature 0 recalesced up to T^ along the vertical line

shown in this figure. During this stage, the solidification evidently proceeded with a smooth-front

morphology, and the entrapped solute precipitated out subsequent to solidification on cooling to room temperature. The reported [10] crystallographic relationship between the precipitates and the matrix

supports this observation. Region B corresponds to a solute-poor precipitation-free zone, which

delineates the smooth interface at a stage when the interfacial velocity had slowed down to the point

A % B

L /Q Figure 12 A schematic binary phase diagram, showing the T^ curve where liquid and a-solid solutions

of the same composition have the same free energy. The metastable extensions of the

a/L

solidus and liquidus lines are also shown. Alloy C

' '

' does not have a T temperature.

16
TEMPERATURE

Figure 13 Schematic phase diagram of Al-Cu system and electron-transparent micrograph of a rapidly

solidified Al-4.5 wt.% Cu submicron droplet. The cellular (segregated) structure in the

particle resulted from nucleation at T = T|^ (no supercooling) with interfacial breakdown

due to constitutional supercooling. Liquid and solid compositions at the interface followed the "paths" shown on the phase diagram. The micrograph is from reference [10].

that solute partitioning set in, although the velocity was high enough to maintain morphological

stability. Solute was rejected ahead of the solidification front as the solid and liquid composition

at the interface moved up along the dashed "path" in the phase diagram to a maximum recalescence

temperature Tj^. In region C the further reduction in growth rate rendered the interface unstable,

I and cellular solidification then took place with segregation of silicon to the cell walls. The postuI lated solidification "path" at the interface(s) is noted by arrows in figure 15. The calculated

thermal profiles in figures 7 and 11 show little or no thermal gradients in the solid during recalescence. It is conceivable that some of the supersaturated solid behind the interface could have been

I

j heated above its T^ temperature. However, even if this were the case, no remelting seems to have

I occurred due to the short times available. Remelting after recalescence has been observed in larger

' dendritical ly solidified particles where heat-flow conditions were more sluggish [12]. Finally, the

j smaller droplet in figure 16 was probably supercooled sufficiently so that complete partitionless

I solidification prevailed during the recalescence.

! Figure 17 presents an SEM view of a large (~ 500 pm in size) powder particle of a maraging 300

steel alloy. The corn-on-the-cob dendritic structure emanating from a nucleation point on the top

surface at the top right-hand corner is fairly characteristic of structures observed in large complex

alloy powders [13]. There is also evidence that a substantial number of droplets like the one in

I figure 17 were supercooled prior to solidification. The light etching in the center of the dendritic

I arms (figure 18) is evidence of high solute cores, which were verified by microprobe analysis [14].

' This indicates some solute trapping early in the solidification process when the dendrites were

j growing into a bulk supercooled melt.
oo

LU

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H

<

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LU

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L/ci ^ Figure 14 Al-rich corner of the Al-Si phase diagram, indicating the curve and metastable extensions of the solidus and liquidus lines. I

Figure 19 illustrates some calculated [11] growth fronts in large (1000 pm) diameter droplets

of pure aluminum with and without initial supercooling. The overall interfacial shapes are concaveij

toward the liquid in contrast to the case of the smaller particles shown in figures 8 and 19a where

the interfaces are convex. In large-sized droplets, positive temperature gradients exist in the melt

prior to nucleation, ^ 90 K in the 1000 pm particles of figure 19 [11]. Therefore, the solid growing

along the symmetry axis faces a superheated liquid, whereas dendrites initially growing along the

j surface encounter a supercooled liquid as well as some external heat extraction. In the limit, the

|

droplet can become rapidly covered with a solid "shell" after several nucleation events along itsj

surface, and further solidification then proceeds in a fairly concentric pattern. Increasing from

2 to 50 cm s“’K~* magnifies the concavity of the interface, as in figure 19c. j

Importance of Second-Phase Dispersions

The mode of solidification which we have referred to as partitionless (or massive) provides a pathway for obtaining especially fine dispersions of very stable second-phase precipitates. Such ’ solidification, from its very nature, is synonymous with complete solute trapping of all the atomic |i species previously dissolved in the liquid state. In this manner, second phases which are highly

insoluble in the solid matrix of the alloy system can be retained in solid solution, thus attaining

exceptional degrees of supersaturation as a stepping-stone to subsequent precipitation in the solid ij state. These circumstances result in finer and more uniformly distributed dispersions than can be

achieved when the second-phase precipitation occurs directly from the melt, i.e., as a part of the ;| solute-partitioning type of solidification.

An interesting example of these phenomena is illustrated in figure 20, based on MnS dispersions

in 303 (high-sulfur) stainless steel [15]. After conventional metallurgical practice, the MnS inclusions are on the order of 10 pm in size; however, after rapid solidification even with compositional

partitioning, the size is reduced to about 0.1 pm (figure 20a), and after massive solidification.

18
TEMPERATURE

I

i

I

! /^/c^dKAH^,

MSI

Al wt % Si Al Si

Figure 15 Schematic phase diagram of the Al-Si system showing the solidification path of the

rapidly solidified Al-6 wt.% Si droplet illustrated cm the right. The el ectron-transparent

micrograph indicates three sequentional modes of solidification. Region A is due to parti

tionless solidification during interfacial recalescence from to T^; region B continued

to solidify in a smooth-front mode with partitioning at the interface as suggested by solid

and liquid compositions moving along the dashed lines to the liquidus and solidus phase

boundaries; region C has solidified in the cellular mode due to onset of morphological

instability with silicon precipitation at the cell walls. The micrograph is from refer- ence [10].

there is a further refinement to about 0.01 pm (figure 20b). In the latter case, the MnS precipitation has taken place after the completion of solidification during the subsequent cooling to room

temperature. Under more extreme conditions, as shown in figure 20c, the MnS precipitation can be

suppressed entirely.

Such dispersions of relatively stable second-phase precipitates are very effective in pinning

grain boundaries at elevated temperatures, and so they play a significant role in retarding 9'^3‘'n growth during the hot consolidation of atomized powders as well as during later heat treatment. A relationship between the limiting grain diameter D which comes into i nterfacial-tension balance with

the dispersed precipitates is given by [16];

( 11 )
TEMPERATURE

Jc A LJ PARTITIONLESS

J

/C|_dT»AH|vi , SOLIDIFICATION

'O

Al wt % Si Al Si

Figure 16 Schematic phase diagram of the Al-Si system showing the solidification "path" of a mas- sively solidified Al-6 wt.% Si droplet illustrated on the right. Partitionless solidification occurred during recalescence from to T^. The micrograph is from reference [10],

20
54321 01 2345

DISTANCE, MICRONS

Figure 18 Photomicrograph and microsegregation in a supercooled powder particle of maraging 300 steel

alloy. (a) Photomicrograph at 250X, (b) distribution of Ni , Co, and Mo concentrations

measured at various points along the chosen microprobe path across a dendritic arm [14].

where d = mean diameter of the precipitates,

f = volume fraction of the precipitates, and

= ratio of the largest to the mean grain diameter.

It is evident, then, that the pinned grain size is smaller, the smaller the precipitate size and the

larger its volume fraction, both of which are favored by rapid solidification, and especially by

parti ti onl ess sol i di f i cati on.

21
a) DIAMETER = 100/Am

“ "0-25

K|^= 2 cm s"! K“l

Figure 19 Interfacial positions as a function of fraction solidified (g) for two sizes of atomized

aluminum droplets. Transition from concave to convex interfacial profiles is produced by

increasing droplet size at a given initial supercooling and also by increasing the interfacial kinetic coefficient [ 11 ].

However, the d/f^ ratio in eq. (11) tends to increase with increasing temperature because of

precipitate coalescence (Ostwald ripening) and resolution in the solid matrix. The kinetics of

coalescence can be expressed by [17]:

_o , 64 D C „ d-d _ s s _ ,,2 j. - 77—=

7^ a V t

0 9 R T p

g where d^ = mean precipitate diameter at zero time,

Dg = diffusivity (in the solid-solution matrix) of the rate-controlling species.

( 12 )

22
Figure 20 Electron transmission micrographs of centrifugally atomized high-sulfur 303 stainless

steel, illustrating (a) cellular, and (b and c) massive solidification, (a) MnS dispersion

in fee matrix; precipitate formed as part of solidification process. Dark-field image,

(b) MnS dispersion in bcc matrix; precipitate formed from solid solution after massive

solidification. Bright-field image, (c) Bcc matrix after massive solidification, but with

MnS precipitation suppressed. Bright-field image [15].

Cg = concentration (in the solid-solution matrix) of the rate-controlling species,

Rg = gas constant,

T = absolute temperature,

a = specific interfacial energy,

Vp = molar volume of the precipitate phase, and

t = time at temperature T.

The rate-controlling species, e.g. , M or X if the dispersed phase is taken as MX, is that for which

the product DC is the smallest. Combining eqs. (11) and (12) yields a t^/^ time dependence for grain

coarsening when controlled by particle pinning of the grai n boundaries. It turns out [18] that d is

dominated by the initjal pre_cipit_ate size d when d > 2 (d^ - d^)^/^, and is dominated by precipitate coalescence when d < 2(d^ - d^)^^^.

o o o

0 0 Concommitantly, f^ of the dispersed phase MX is maximized if the alloy composition in M and X corresponds to the stoichiometric ratio of MX [18]. This condition guarantees that the kinetics of

Ostwald ripening will be controlled by the slower diffusing species. The coarsening rate can be

further diminished by using second phases of higher stability, i.e., having still lower solubility

products relative to the solid-solution matrix; this serves to decrease the concentration C in

eq. (12).

23
An example of the resistance to austenitic grain growth after rapid solidification is given in

figure 21 for a 9Ni-4Co-0.2C steel [19]. Very little grain growth takes place in the rapidly solidified material up to 16 hours at 1100 °C, and even after coarsening sets in, the grain size remains

very much smaller than in the conventionally processed steel. Such unusual grain-growth characteristics have been observed in a number of steels [20].

The fracture toughness of some ultrahigh-strength steels can be significantly improved by austenizing at relatively high temperatures, but other adverse effects may result from the attendant grain

coarsening. On the other hand, with steels that are processed by rapid solidification, the resistance to grain growth should favor more attractive combinations of properties. An illustrative trend

is shown in figure 22 for an 0.25C-4Mo martensitic steel after tempering at 150 °C [21]. In this

instance, whether the steel is processed conventionally or by rapid solidification, both the sharpcrack and blunt-notch fracture toughness are enhanced by raising the austenitizing temperature.

However, the influence of the difference in grain size is clearly manifested when the toughness

testing is carried out at subzero temperatures, as plotted in figure 22c. The shelf energy and the

ductile condition are maintained to much lower temperatures, comparing the rapidly solidified to the

conventional materials, because of the small grain size retained in the former after high austenitizing treatments. J

The fine dispersions of very stable phases, made possible by rapid solidification, also play aft

beneficial role in oxidation resistance. For those alloys which tend to form protective oxide films®

at high temperatures, the stable second-phase precipitates seem to improve the adherence of surface®

films via a keying action due to fine-scale interfacial penetration around the uniformly distributed*

particles [22]. An example of this favorable effect is shown in figure 23 for rapidly solidified®

FesAl containing TiB 2 as the dispersed second phase [23]. ®

Of course, all of these second-phase considerations, when applied to larger volume fractions, 5: can lead to interesting possibilities with regard to dispersion strengthening. '

Figure 21 Grain-growth characteristics of 9Ni-4Co-0.2C steel (a) after conventional processing

(normal ingot and hot-rolling practice) and (b) after rapid solidification processing. The

latter involved centrifugal atomizing and consolidation by hot extrusion [19].

Closure

fundamental modes of solidification have been described according to the morphological

s a 1 1 y 0 the advancing growth front and the extent of solute partitioning at the interface. In tne regime of cellular and dendritic morphologies, even with no appreciable supercooling, rapid cooling through the solidification range modifies the microstructure by refining the scale of the

24
LUH

o

o

cc

- 30

“ 20

c\j COE

z

(f) LU

_l

D

o

AUSTENITIZING TEMPERATURE, ° C

30

20

10

C/)

LU

_l

D

o

TEST TEMPERATURE, °C

Figure 22 Fracture-toughness characteristics of conventionally processed and RSP 0.25C-4Mo steel;

water quenched from austenitizing temperatures shown and tempered at 150 °C for 1 hr. (a) and (b) Effect of austenitizing temperature on room- temperature properties, and

(c) effect of test temperature on fracture toughness after austenitizing at 1150 °C. Kg

is the plane-strain fracture toughness based on the maximum load vs crack-opening

displacement for a fatigue-precracked specimen in a slow-bend test [21].

microsegregated regions which result from solute partitioning; the accompanying dispersions of secondphase precipitates are also refined. Of further interest is the attainment of massive (partitionless) solidification via supercooling in which solute trapping is substantially complete. By this

means, extremely high supersaturations relative to very stable phases are obtained which, in turn, can be precipitated subsequently from the solid matrix. The required conditions for massive solidiL/q

fi cation are both thermodynamic and kinetic; the interfacial temperature must be below , and the

interfacial velocity must exceed the diffusive velocity to avoid solute partitioning.

25
OXIDATION AT 1095 °C

TIME, hr

Figure 23 Effect of RSP and a dispersed second phase (TiBg) on cyclic oxidation resistance of FegAlj at 1095 °C [23].

Stable second-phase dispersions resist coalescence on heating and so remain effective in pinning]

grain boundaries even at unusually high temperatures. Some consequences of these microstructural

features have been dicussed with respect to mechanical and oxidation behavior.

A proper definition of rapi d sol idi fi cation has still not been established. When dealing with

cellular or dendritic modes of solidification, it is found that increased rates of cooling from thej

melt are beneficial because they lead to enhanced compositional uniformity and second-phase refinement. In this sense, cooling faster than about 200 K s~^ has been regarded as a rapi d cool i ng rate < inasmuch as significant changes in microstructure can thereby be obtained [1]. Nevertheless, cooling'

rates of 10“* to 10® K s"^ are often sought for further benefits.

However, it is obvious that a cooling rate cannot be synonymous with a solidification rate. In connection with smooth-front solidification, as achieved by appropriate supercooling, the solidification rate can be identified with an interfacial velocity, at least for a given nucleation event. Under such circumstances, the interfacial velocity has further significance from the standpoint of

affecting the degree of solute trapping. In the limiting case, to reach the highest levels of compositional uniformity and of stable second-phase refinement, a solidification growth rate commensurate

with the diffusive velocity might be considered as rapid solidification , since this condition reflects

substantial amount of solute trapping. In the present work which focuses on the solidification of

atomized droplets, it appears that interfacial velocities greater than about 1 m s ^ might be regarded

as "rapid solidification." Solidification rates of this magnitude are obtained more readily by supercooling than by rapid cooling per se without supercooling.

Acknowl edgments

We wish to express our appreciation to several agencies who have sponsored various aspects of our I respective research programs on rapid solidification: the Office of Naval Research under Contract No. | N00014-81-K-0013 at the Massachusetts Institute of Technology; the National Bureau of Standards under

Contract No. NB83NAHA4004 at the Massachusetts Institute of Technology; the Defense Advanced Research i Projects Agency under DARPA Order Number 3751 at the National Bureau of Standards. Enlightening ^ discussions with Drs. J. W. Cahn, W. J. Boettinger, and S. R. Coriel 1 of the National Bureau of i Standards are gratefully acknowledged. j

26
I!

References

[ 1 ]

C 2]

[ 3 ]

[ 4]

" [ 5]

I [ 6 ]

C 7]

[ 8 ]

[ 9]

I [10]

! [ 11 ]

[ 12 ] [13]

[14]

[15]

[16]

[17]

[18]

[19]

[ 20 ]

[ 21 ]

[ 22 ] [23]

M. Cohen, B. H. Kear, and R. Mehrabian, Proc . 2nd Inti . Conf . Rapid Sol idification Processi ng : Principles and Technologies II , R. Mehrabian, B. H. Kear, and M. Cohen, eds. , Claitor's Publishing Division, Baton Rouge, p. 1 (1980).

S. R. Coriel 1 and D. Turnbull, Acta Metal 1. 2135 (1982).

M. C. Flemings, Solidification Processing , McGraw-Hill Book Co. , p. 60 (1974).

B. Chalmers, Physi cal Metal 1 urgy , John Wiley, New York, p. 258 (1959).

S. R. Coriell and R. F. Sekerka, Rapid Sol idification Processi ng : Principles and Technologies

II , R. Mehrabian, B. H. Kear, and M. Cohen, eds., Claitor's Publishing Division, Baton Rouge,

p. 35 (1980).

S. R. Coriell, private communication.

R. F. Sekerka, Crystal Growth: An Introduction, North-Hoi 1 and Amsterdam, p. 403 (1973).

M. J. Aziz, J. Appl. Phys. 1158 (1982).

W. J. Boettinger, S. R. Coriell, and R. F. Sekerka, these proceedings.

C. G. Levi and R. Mehrabian, Metall. Trans. A 13A , 13 (1982).

C. G. Levi and R. Mehrabian, Metall. Trans. A 13A , 221 (1982).

J. H. Perepezko, Y. Shiohara, J. S. Paik, and M. C. Flemings, these proceedings.

P. A. Joly and R. Mehrabian, J. of Mat. Sci. 9, 1446 (1974).

T. Z. Kattamis and R. Mehrabian, J. of Mat. Sci. 9, 1040 (1974).

T. F. Kelly, M. Cohen, and J. B. VanderSande, submitted to Metal 1 . Trans, for pubi ication.

T. Gladman, Proc. Roy. Soc. A294 , 298 (1966).

I. M. Lifshitz and V. V. Slyozov, J. Phys. Chem. Solids 19, 35 (1961); C. Wagner, Z. Electrochem. 65, 581 (1961).

G. B. Olson, H. C. Ling, J. S. Montgomery, and J. B. VanderSande, Rapidly Sol idified Amorphous

and Crystalline Alloys, B. H. Kear, B. C. Giessen, and M. Cohen, eds., North-Hol land Publishing,

NY, p. 355 (1982).

M. Suga, J. L. Gross, G. B. Olson, and J. B. VanderSande, Rapi d Sol idification Processi ng : Principles and Technologies II , R. Mehrabian, B. H. Kear, and M. Cohen, eds., Claitor's Publishing Division, Baton Rouge, p. 364 (1980).

G. B. Olson, M. Cohen, and G. J. Yurek, these proceedings.

P. M. Fleyshman, Fracture Toughness of Rapidly Solidified Martensitic Steels, S. M. Thesis, MIT,

Cambridge (1982).

G. J. Yurek, D. Eisen, and A. Garratt-Reed, Metall. Trans. A 13A , 473 (1982).

E. R. Slaughter, C. Adam, M. J. Blackburn, R. G. Bourdeau, and C. V. Law, unpublished research

(1982).

27
SOLIDIFICATION OF UNDERCOOLED LIQUIDS

+ J.H. Perepezko , Y. Shiohara , J.S. Paik and M.C. Flemings !

Department of Metallurgical and Mineral Engineering

University of Wisconsin-Madison

Madison, Wisconsin 53706

'^Department of Materials Science and Engineering

Massachusetts Institute of Technology

Cambridge, Massachusetts 02139

ABSTRACT

During rapid solidification processing (RSP) the amount of liquid undercooling is an important

factor in determining microstructural development by controlling phase selection during nucleation and

morphological evolution during crystal growth. While undercooling is an inherent feature of many

techniques of RSP, the deepest undercoolings and most controlled studies have been possible in care- fully prepared fine droplet samples. From past work and recent advances in studies of nucleation

kinetics it has become clear that the initiation of crystallization during RSP is governed. usual ly

by heterogeneous sites located at surfaces. With known nucleant sites, it has been possible to identify specific pathways of metastable phase formation and microstructural development in alloys.

These advances have allowed for a clearer assessment of the interplay between undercooling, cooling

rate and particle size statistics in structure formation. New approaches to the examination of growth

processes have been developed to follow the thermal behavior and morphology in small samples in the

period of rapid crystallization and recalescence. Based upon the new experimental information from

these studies, useful models can be developed for the overall solidification process to include nu- cleation behavior, thermodynamic constraints, thermal history, growth kinetics, solute redistribution

and resulting structures. From the refinement of knowledge concerning the underlying factors that

govern RSP a basis is emerging for an effective alloy design and processing strategy.

The main attention in rapid solidification (RSP) methods has been focused upon the attainment of

high cooling rates in the range of 10^ to 10^ °C/sec. With these methods small liquid volumes are quenched either in contact with a substrate as in melt spinning and directed energy beam processing or

in an inert gas as in atomization. However, the most important impact of RSP can be related directly

to the solidification structures and the new opportunities for structure control by innovative alloy

design and processing strategies. The wide range of reported microstructural variations encompass not only equilibrium phase mixtures with refined microstructural scale, but also novel microstructures

such as supersaturated solid solutions, metastable intermediate phases and amorphous solids.

The existence and understanding of novel microstructures representing non-equilibrium phases

requires a consideration of the level of undercooling at the onset of nucleation in addition to rapid

growth kinetics. In fact, the available free energy for non-equilibrium phase formation is directly

related to the amount of undercooling. With increasing levels of undercooling the product phase

selection involved in nucleation can be expanded to involve an increased variety and to allow competitive growth kinetics to play an important role in the evolution of different microstructural morphologies. Accordingly, the examination of solidification in undercooled liquids can contribute to the

basic understanding necessary for an effective microstructure control and alloy design in the optimization of RSP.

Experimental Techniques for High Undercooling

Even with liquid metals of high purity the attainment of a high level of undercooling is precluded

usually by the catalytic action of residual impurities or container walls. In order to observe extensive undercooling it is necessary to circumvent the catalytic effects of the active nucleants.

An effective experimental approach that may be applied to yield large undercooling prior to solidification involves the slow cooling (1 0-30°C/min) from the melt of a dispersion of stabilized fine metal

Introduction

28
droplets [1]. By starting with high purity metal and dispersing the liquid into a large number of

small droplets (<20y) within a suitable medium, the catalytic effects of active nucleants may be res- tricted to a small fraction of the total droplet population. A suitable means of carrying out the

nucleant isolation operation is illustrated schematically in fig. 1.

In considering the droplet approach to undercooling, a distinction should be made between the

several modifications of this technique which have been employed by a number of investigators in the

past. For example in the droplet substrate method [2], a small sample of liquid metal is placed on an inert glass substrate in a chamber with an inert or reducing atmosphere. With this approach the onset

of solidification is determined by visual observation of the change in surface reflectivity of a metal

droplet which is about 50-100y in diameter. In the entrained droplet technique [3] a bulk alloy

sample is equilibrated in a liquid-solid two phase field with the temperature and alloy composition

selected so that only a small fraction of the volume will be liquid in contact with a primary solid

solution phase, as a potential heterogeneous nucleant. In addition, several other procedures have

been developed for generating liquid metal droplet dispersions which exhibit a propensity for substantial undercooling. These droplet forming techniques involve exploding wires [4], shotting [5] and the

containerless solidification in drop tubes [6]. While the majority of droplet undercooling schemes

have a provision for some method of temperature measurement, the undercooling tendency of liquid droplets in several cases including all atomization methods is judged usually from microstructural observations.

With bulk samples, the removal of active nucleation sites usually involves a physical separation

of the sites from the melt or a chemical treatment either to incorporate or to deactivate possible

catalytic sites in the presence of an inorganic containment layer. These approaches which include a bulk sample encased in organic glasses [7] as well as a bulk melt levitated by induction [8] indicate

that some remarkably large levels of undercooling (i.e. 0.2 T^) may be achieved. The inorganic glass

encasing is believed to promote undercooling not only by isolating the liquid from possible catalytic

sites on container walls, but also by a possible scavenging action on nucleating centers that may be

distributed within the volume of the melt. It is worthwhile to note that in principle the undercooling

attained in properly conditioned bulk samples should approach the levels in droplet samples. Besides

these specific undercooling techniques, it is clear that atomization of liquids is a commercial RSP

method which incorporates an effective nucleant isolation operation as well as rapid quenching.

29
Nucleatlon Kinetics in Undercooled Liquids

In the present discussion the main interest in nucleation kinetics will be focused on the possible

relationships between undercooling and cooling rate during the solidification of a distribution of

liquid droplets. These relationships may be examined by nucleation rate calculations, but the computed

results are highly dependent on the specific values assigned to the kinetic rate parameters. Reliable

information is available for some parameters, however, in most cases the kinetic parameters are assigned values based upon theoretical models which require experimental verification. To minimize the

uncertainty associated with calculations based model dependent parameter values, it is of interest to examine the range of droplet nucleation kinetics in a general manner. With this approach it is possibl

to deduce several useful guiding relationships for rapid solidification behavior.

Nucleation of a crystalline phase in an undercooled liquid phase may be classified into two mechanistic types. When nucleation occurs without an association with any catalytic nucleation sites, it can be classified as homogeneous nucleation. In the vast majority of experimental conditions nucleation is initiated at some catalytic nucleation site as heterogeneous nucleation. For a heterogeneous

nucleation, an expression for the nucleation rate, J, has been developed on the basis of homogeneous

nucleation theory [9] as:

J = exp [-AG(n*)/kT] (1

)

a

where is proportional to the number of liquid atoms in contact with unit area of the catalytic

surface as well as the jumping frequency of atoms, AG(n*) is the activation barrier for nucleation and

kT has the usual meaning. For nuclei having the form of a spherical sector.

AG(n*) = (16/3) TTC^ f(0)/(AG^)^ ( 2 )

where a is the interfacial energy between the nucleus and the liquid, AGy is the driving free energy

for nucleation of unit volume of the product phase and f(0) = (2-3 cos 0 + cos^ 0)/4 is a contact

angle factor. In the description of nucleation kinetics during rapid liquid quenching, steady state

nucleation conditions may be considered in general, except for glass forming alloys [9].

Nucleation Rate Parameters

The nucleation rate is a relatively steep function of temperature with a steady state magnitude

determined principally by the value of the exponential terms in AG(n*) at the nucleation temperature

Tn and to a lesser amount by the prefactor. In evaluating the temperature dependence of J usually a constant value based on theory is taken for fia 6ven though the catalytic site may vary. Similarly,

little information is available to judge the catalytic potency, f(0) so that a range of values is

used in calculations. However, the most important parameters in determining J and hence the undercooling level are AG^ and a which have received continued experimental and theoretical study.

In pure metals, the recent measurements of the heat capacity for undercooled liquids [10] have

allowed for a direct evaluation of AGy. In general, the experimental values are represented to within

a few percent by the relation proposed by Turnbull [11,12].

AG„ = -AK, (AT/TJ = - AH^ (1-T^) (3)

where AH^ is the heat of fusion per unit volume of product phase, AT is the undercooling and T^ is the

melting point of the product phase. For alloys the value of AGy is determined from the composition of

the nucleating phase at T^^. The composition dependence of AGy can be expressed as

AG^ = [X^ Ay^ + (1-X^) Ay^]/v ' (4)

A B where Ay and Ay are the chemical potential differences between the undercooled liquid and crystal and

V is the molar volume for the crystalline phase at composition X/\. It is apparent from eq. 4 that AGy

is maximized when Ay^ = Ay^ so that AGy = Ay^/v. This condition has been proposed by Hillert [13] to assess the nucleus composition. However, the favored nucleus composition should yield a minimum value

of AG(n*). Only if v and a are not composition dependent can the maximum AGy value occur at the same

30
I

'xomposition as that for a minimum in AG(n*). Apparently this behavior may be approached in some alloys.

liFor example, recent calculations [14] of the composition dependence of the nucleation temperature,

I iTp(x) have reproduced the observed experimental trends [1] which reflect the composition dependence of

;;tne liquidus, but the generality of the calculated behavior requires further examination.

i In the study of nucleation one of the most important parameters is the solid-liquid interfacial

[jenergy. Since the nucleation rate is proportional to exp -(a)3 as shown in eqs. 1 and 2, it is very

ijsensitive to even slight errors in the a value [15]. A direct measurement of a is very difficult

yeven with a large uncertainties because of sensitivity to impurity effects. Usually values of o are

n calculated from classical nucleation theory by assuming that experimentally obtained maximum undercool

iling values are associated with homogeneous nucleation. Unless the maximum undercooling value is

ij proven to refer to homogeneous nucleation, the estimated a value is expected to underestimate the

{itrue value. There have been a number of theoretical approaches to study the structure and properties

iiof the solid-liquid interface on the basis of various model constructions. The models range from the

; limiting cases in which cf values are calculated from enthalpic bond breaking arguments [16] or derived

j solely from a basis of entropic considerations [17] to intermediate approaches where a weighted conitribution of enthalpic and entropic terms [18^20] is incorporated into the evaluation. At the present

I time, the theoretical values appear to underestimate a compared to values calculated from nucleation

'undercooling results [15]. Certainly, an improvement in the independent experimental determination

! of o would be invaluable in allowing for a clear evaluation of theoretical models and more reliable

j nucleation calculations.

|j Effect of Cooling Rate on Nucleation Kinetics

!. Since the onset of nucleation in a highly undercooled liquid is expected to be a sharp function of

i temperature it is useful to consider a limiting case for the effect of cooling rate on nucleation

: temperature. In general, for a given nucleation kinetics the critical condition to observe nucleation

! experimentally in a time, t, in a droplet of volume, v, and catalytic surface area, a, is given by

I Jvt :i 1 for homogeneous nucleation and by Jat ~ T for heterogeneous nucleation. From eq. 1, the

I heterogeneous nucleation time can be represented as a function of temperature as:

In t

= -ln[a (fl^)]+ 16 ^a3f(e)/[3 k (AH^)^ T^ (1-T^)^ T^] (4)

where a fig represents the heterogeneous nucleation site population associated with a droplet of volume

V.

j Furthermore, the rate dependence of the nucleation temperature which represents the shape of the

i T-T-T diagram associated with a specific nucleation kinetics is represented in general as

d X3k ^(1-T^)3

^ L(3VD j

where the possible temperature dependence of ^g, Ah^, a and f(9) has been neglected in order to ex- amine a limiting value and minimize the uncertainty of model dependent values. For homogeneous

nucleation f(9) = 1, but for a heterogeneous nucleation O<f(0)<l. Therefore, T-T-T diagrams for

heterogeneous nucleation can have different breadths representing a different temperature dependence

of the nucleation rate for different heterogeneous sites. In eq. 5, dTy-/d(lnt) becomes infinite at Ty = 1/3 as the nucleation time is minimized. Therefore, T^ = 1/3 is believed to estimate a lower

bound to the onset of sensible nucleation of a crystalline phase during slow cooling [21]. If the

undercooling limit of the droplet sample represents the solidification of the nucleant-free droplet

population, the crystallization temperature is determined either by the catalytic nature of the droplet surface coating or by the intervention of homogeneous nucleation kinetics. Also, during cooling

the operation of a given catalyst kinetics and the kinetic competition between different catalysts and

different product phase structures will be influenced strongly by the active catalyst distribution at

the undercooling limit. In order to reach this limiting undercooling the size refinement of droplets

becomes one of the most important factors. Further analysis on the distribution of active nucleants

among droplets then becomes necessary to understand the undercooling behavior of droplet samples.

Nucleation Kinetics and Size Distribution Statistics

As long as a single nucleation kinetics with a steep temperature dependence is operating the

effect of droplet size on the nucleation temperature is relatively small. In fact, if the size

distribution is the only factor associated with the nucleation temperature range during cooling, a

31
sharp nucleation peak over a narrow temperature range (<10°C) is expected even with a sample of a I

rather broad droplet size distribution. However, droplet emulsion samples of pure metals and alloys Hi

often show a broad crystallization peak over a wide temperature range (>20°C). In this case the H;

operation of a single nucleation kinetics is unlikely but the broad crystallization peak can be H

associated with multiple nucleation events continuously activated during the cooling process [21].

Morever, at a given cooling rate, nucleation of the crystalline phase in each droplet is associated B

with the catalytic potency of the most active nucleant. As a result, nucleation at different levels B;

of undercooling due to crystallization over a wide temperature range can result in the formation of Si

different structures in different droplets. Even if the same structure is nucleated at different ] levels of undercooling, this structure will not necessarily be retained completely in the final solid- yU

ification pattern. For example, the final product structure can develop variations due to a different '

thermal history associated with the recalescence period following nucleation at different undercooling - |, levels in separate droplets. •J j An example of the nucleation of different phases due to crystallization behavior over a broad Ji / undercooling range may be illustrated by considering the careful studies on Fe-Ni alloys reported by |l Cech [22]. In fine droplet samples, (<40y), Cech demonstrated the presence of droplets of a metastable i, bcc structure in addition to droplets of the equilibrium fee structure. Size distribution analysis

results showed that small droplets less than 13y are more likely nucleated at high undercooling since 'ji the major fraction had a metastable bcc structure. Since there were also droplets of the metastable ‘*

| structure within the large size grouping (d > 20y ), it is necessary to consider the mode of nucleant | distribution among droplets. When all possible nucleation sites which can be activated a lower levels

of undercooling than that necessary to retain the metastable phase are considered to be nucleants, a i Poisson distribution function may be introduced to describe the random distribution of these nucleants , j- among droplets [21,23]. In this case the fraction of droplets free from active nucleants is represented 'u by I'

X = exp [- m (v)] (6) {

I t where m(v) is the average number of active nucleants in each droplet of volume V. Depending on the ',11' distribution mechanism of nucleants during the atomization process m(v) is proportional either to the

volume or to the corresponding surface area of the droplet. Since a nucleant distribution proportional

to the droplet volume would result in a rapid decrease in the yield of metastable product with increasing droplet size, it appears that for the Fe-Ni case, a surface area dependent nucleant distribution is most probable. A further application of the Poisson distribution has been reported for the solidification of Pd-Si droplet samples [24]. In this case the droplet fraction which solidified continuously

as glass represents X in eq. 6 which can also be written in terms of droplet diameter, d, as X = exp

[-(d/do)'^]. When m(v) is proportional to the droplet surface area, n=2 for a heterogeneous surface

nucleation. For heterogeneous volume nucleation n=3 and for homogeneous nucleation n=4.6. As in other

droplet studies the solidification behavior of Pd-Si droplets was described most closely by a heterogeneous surface nucleation kinetics.

Metastable Phase Formation

Thermodynamic Features

A large number of examples of metastable phase formation have been demonstrated during rapid

quenching experiments [25]. While the possible reaction paths leading to metastable phase formation

in alloy systems are numerous a consideration of thermodynamic constraints can be useful in providing

guidelines and limiting bounds on the reaction sequences in undercooled liquids. For example the free

energy relationships that are relevant to the generation of solid solutions are illustrated in fig.

2. Two features are of interest in this case. For the formation of an equilibrium solid solution of

composition Cq at temperature Tf^ an overall free energy change of AG] is involved in reaction. However,

in the early nucleation stage of reaction when a small amount of a forms from the liquid the driving

free energy is maximized for a with a composition C] rather Cq. When the molar volume of solution does

not change appreciably with composition, the C] composition may be determined by constructing a tangent

to the a free energy curve which is parallel to a tangent to the liquid free energy. As noted previously, this consideration indicates that a with solute content different from Cq can develop initially

during nucleation [13]. Another feature of interest in fig. 2 concerns the bound set by the Tq dashed

curve between the liquidus and solidus [26], which represents the locus of equal free energy between

liquid and solid. Formation of a product with the same composition as the liquid is excluded at compositions beyond the Tq limit.

A thermodynamic consideration is of most value when the analysis also can provide an estimate of

potential metastable intermediate phases that can enter in the kinetic competition. For example, in

alloys the requirement that the free energy, /£, exhibit an overall decrease with reaction can lead

to several interesting possibilities for phase reaction sequences [26] as illustrated by one example in

32
TEMPERATURE

Fig. 2 Free energy relationship between undercooled liquid and a phase.

Fig. 3 Free energy relationship between under

cooled liquid and several crystalline

phases [26].

(a) (b)

Fig. 4 Schematic illustration of a possible reaction path for primary phase catalysis

of a metastable solid solution (a) and a metastable y intermediate phase (b).

33
fig. 3. In this case the equilibrium solidification of liquid of composition Cq would yield a twophase mixture of ct with composition Cq^ and B with composition Cg. However, if the nucleation kinetics are favorable, an intermediate reaction step can occur involving the formation of the metastable y phase with composition Cq and free energy reduction AG(L->-y )• The presence of y phase affects the

reaction sequence in that a phase cannot form since the free energy change AG(y^ ) is positive.

Only after the y phase transforms into 3 phase can the a phase form with a free energy reduction. ] When there are several metastable phases possible, the number of intermediate reaction steps can in- crease and present interesting opportunities for structure control during both solidification and subsequent solid state processing. One way of summarizing the potential for structure modification in

alloys is to construct a metastable phase diagram [25,27]. !

I Competitive Kinetics Anal ays is

Since the driving force for the formation of a metastable phase is related directly to the amount i of undercooling below the melting temperature of the phase, a high level of undercooling can expand the

product selection involved in nucleation to an increased variety of types of structures. This is

illustrated by the nucleation transition from fee to bcc in Fe-Ni alloys [22], stainless steels [28]

and tool steels [29]. Other examples have been reported in Pb-Bi, Bi-Cd and Pb-Sn alloys as well as

' Ga and Bi [1,10,21]. For Pb-Bi alloys a multiple phase selection transition has been observed in

droplets [1]. As the level of undercooling increases an equilibrium eutectic mixture is replaced by

a supersaturated (hep) phase which in turn is superceded by a metastable intermediate phase at the

highest undercooling. In addition for Pb-Sn, Sn-Ge and Cd-Sb alloys the operation of the nucleation

product transition has been related to the catalytic effect of specific heterogeneous sites [21,30].

One type of effective nucleation site is a primary solid solution which can participate in catalysis

reactions such as those illustrated in fig. 4 [21]. When the a phase is a potent catalyst, a continued

growth of a is possible during cooling to form a supersaturated a product as in fig. 4a. When a is

strongly catalytic in promoting the nucleation of a metastable intermediate y phase as indicated in ' fig. 4b, the formation of an equilibrium 3 phase may be bypassed as the liquid is undercooled below

the metastable y phase melting point. Thus, for an alloy that may exhibit metastable products, it ! is possible to select a particular phase by controlling the degree of undercooling prior to sol idifica- jj tion and by providing a high density of favorable catalytic sites.

In order to understand the kinetic competition between different phases during nucleation in an undercooled liquid, the comparison of thermodynamic stability may not be sufficient. In terms of

overall nucleation kinetics, the driving free energy is not the only parameter of importance, but also

nucleation site density and the magnitude of AG(n*) associated with a given phase are critical. As an illustration of the application of nucleation kinetics to reveal the relationships between cooling rate,

undercooling and phase selection, it is useful to consider the competition between two nucleation

kinetics during continuous cooling of a droplet in several different situations that are illustrated

in fig. 5. As an initial example, when two catalysts of different nucleation site density and potency

are in the same droplet (B and C), the undercooling limit of the droplet at T^l is determined by

catalyst C at cooling rate t]. At cooling rate Jz, however, the undercooling of the droplet increases

to T^2 by circumventing the catalytic effect of C. Thus at a continually increasing level of cooling

rate, there can be a significant improvement of undercooling as the catalytic effects of highly potent

nucleants are circumvented. During this process the controlling nucleation kinetics involves catalysts

of decreasing potency. When two different phases are competing with each other on the same catalytic

surface present in a highly undercooled liquid, according to nucleation theory, the magnitude of the

activation free energy barrier becomes the most dominant factor. Therefore, in order to favor nucleation of a metastable phase at a given cooling rate the activation energy barrier for the metastable

phase should be lower than that for the equilibrium phase. In terms of eq. 1 this condition may be

expressed as

T(9')

( ah ^'")2

( i - t ;)2

f(e)

( AH ^)2 ( 1 - T^)2

(7)

where a^-| and 0' are the interfacial energy and the contact angle between the metastable phase and the

catalytic surface, AH^^ is the heat of fusion per unit volume of the metastable phase and T^,' is the

relative nucleation temperature of the metastable phase. It has been shown in eq. 5 that the breadth

of the T-T-T diagram associated with different nucleation kinetics is proportional to (AHy )^Tni/<^“f (0)

When the melting temperatures of two different phases nucleated by the same catalysts are very close

and the metastable phase is dominant, the condition given by eq. 7 may be approximated by

34
1.0

1

0

Log(time)

Fig. 5 T-T-T diagram representing different nucleation kinetics that may occur during a continuous cooling in undercooled droplets.

(AH (AH

3

^ ~3 (2)

Olm f(0') ^ f(0)

In this case, the breadth of the T-T-T diagram for the nucleation of the metastable phase (A) is larger

than that for the equilibrium phase (B), but the noses of these two diagrams are placed at the same position as illustrated in fig. 5 schematically (A and B). Unless there are nucleants of higher

catalytic potency for nucleation of the equilibrium phase (for example C in fig. 5), the equilibrium

phase is always masked by nucleation of the metastable phase regardless of the cooling rate. In fact,

the droplet emulsion technique which disperses potent nucleants of small population into a minor fraction of the droplet sample can be viewed as eliminating the possible activation of nucleation kinetics

C in the major fraction of the sample. Alternatively, if the activation barrier for A is higher than

that for B, the equilibrium phase will nucleate at all cooling rates. In addition, it is also possible

for the nucleation of the equilibrium phase to be favored by an increase of the cooling rate if the

nucleation site density for the equilibrium phase (curve D) is greater than that for the metastable

phase. For these conditions an increase in cooling rate can result in a decreasing yield of a metastable product. In fact, a transition in nucleation from surface dependent to volume dependent kinetics

with increasing cooling rate such as C to D in fig. 5 can yield a large increase in rate and may con- tribute to a propensity for multiple nucleation as reported in quenched Al-Si alloy droplets [31].

However, other factors involving the specific thermal history of droplets are also likely to contribute

to the development of multiple nucleation.

In summary, the examination of the simple. competitive kinetics examples reveals several types of

general behavior. First, for a single operating kinetics an increase in cooling rate translates

usually into some increase in undercooling. With several competing kinetics the increase in undercooling with cooling rate may not be uniform. For several competing kinetics and possible product

structures, an increase in cooling rate can allow for nucleation of a metastable phase to be favored

or for an increase in the yield of a metastable phase product, but this is not a general rule. Depending on the density of catalytic sites, it is possible for an increase in cooling rate to result in no change in phase selection or to result in a decreasing yield of a given metastable phase in favor of

another metastable product or even an equilibrium phase product. From these examples illustrating the

possible synergistic effects between high undercooling and rapid quenching to yield unique product

structures, it is clear that attention to catalyst potency and site distribution as affected by surface

conditions during atomization can be as important as cooling rate in optimizing RSP treatments.

35
Heat Flow and Pure Metal Solidification

In solidification of small droplets, two extreme heat flow situations may arise. In the first fi one, the solidification front moves into a liquid at its melting point or above, and the latent

heat of fusion must be withdrawn through the solid. The second situation involves solidification

into a supercooled melt. Nucleation occurs only after substantial supercooling, and initial ri solidification then occurs primarily by flow of heat into the liquid itself, therby resulting in

f droplet recalescence. -

For solidification without undercooling, the Newtonian approximation is applicable when the |t; dimensionless Biot number, Bi , is lower than 0.1, where Bi = h rg/kp^. The Biot number is the

ratio of the surface conductance or heat transfer coefficient, h to the metal conductance(k^/rg) : _ ;i

kp, is the thermal conductivity of the metal and rg is the droplet radius. The heat transfer

coefficient is determined from the Nusselt number [32]; i|-

1/2 1/3 * I Nu = 2 h r /k. =2+0.6 Re Pr (9) ' J 0 T ; I'l where Re and Pr are Reynolds and Prandtl numbers respectively and k^ is the thermal conductivity

of the fluid. In the case of solidification of metal droplets in a quiescent liquid or gas,

the Nusselt number reduces to 2, and the Biot number is simplified to k^/k^. In the case of pure

tin droplets emulsified in oil (polyphenylether) , the Biot number is about 4.4 x 10“\ since

k^ is 0.112 kcal/m hr °C and kp^ is 25.3 kcal/m hr °C. Thus, the dominant resistance to heat flow

in this example occurs at the surface and no significant temperature gradients can develop inside

the droplet in absence of undercooling. Similarly in most if not all gas atomization processes,

heat transfer is limited by the solid-gas heat transfer resistance.

Thus, the simple Newtonian model is widely and correctly used for single phase cooling and

solidification without undercooling [33]. Table 1 shows cooling rates for aluminum gas atomized

droplets calculated in this way. Table 2 shows comparable cooling rates for atomized or _

emulsified metals in several liquid media. I

When supercooling occurs in atomized or emulsified droplets, it is not obvious that the

droplet will have a uniform temperature distribution during the recalescing period of solidification;

Table 1. Calculated Cooling Rates of Pure Aluminum at 660°C During Gas

Atomization (in Helium at 300°K): (°K/sec)

Rel ati ve Gas Velocity Particle Diameter (ym)

mach V (m s"M 10 100

0 0 7.21 X 10" 7.21 X 10"

0.1 97 1.00 X 10^ 1.56 X 10®

0.3 290 1.22 X 10" 2.26 X 10®

0.5 483 1.36 X 10" 2.69 X 10®

1.0 965 1.56 X 10" 3.53 X 10®

Table 2. Calculated Cooling Rates of Pure Metal Droplets (lOym dia.)

Cooled in Different Media at Low Reynolds Numbers

Metal and Quench Medium Cooling Rate at Equilibrium

Meltinq Point (°K/sec)

Aluminum in Helium gas at 27°C 7.21 X 10®

Aluminum in Argon gas at 27°C 9.68 X 10®

Aluminum in Liquid Gallium at 30°C 2.12 X 10®

Tin in Argon at 27°C 4.57 X 10®

Tin in Oil at 100°C 2.28 X 10®

Tin in 0il/CCl4(l/2) at -25°C 3.77 X 10® Iron in Argon at 27°C 1.15 X 10® Iron in Glass at 1000°C 3.99 X 10"

Iron in Molten Salt at 482°C 3.11 X 10®

I

36
I i.e., the temperature of the growing interface may rise faster than the average temperature of

I the droplet. In this case, the validity of the Newtonian approximation during recalescence will

depend not only on the Biot number, but also on interface velocity and interface morphology.

I

Levi and Mehrabian [34] have considered in some detail the solidification of undercooled

droplets of pure metals, assuming a planar (non-dendri ti c) interface and the interface kinetic

relationship suggested by Turnbull [35] and modified by Cahn [36] for the steady state continuous

growth of a planar front. Figure 6 is an example of results of their computer calculations for an aluminum droplet with a single nucleation event. For the conditions chosen, significant temperature

differences develop within the liquid during recalescence. Of course, as pointed out by Levi and

Mehrabian, the magnitude and longevity of internal temperature differences depend sensitively on

the number of nuclei growing, and on whether or not the interface is dendritic. In the case of

multiple nucleation, the scale of the heat transfer problem is reduced from that of the droplet

radius, rg, to that of half of the interparticle spacing d^x/2- In the limit of lacy dendrites

(or, dendrites that become unstable and break into individual particles), the heat transfer scale

is similarly reduced. Figure 7, also from Levi and Mehrabian, shows schematically how decreasing

d^x althers the "enthalpy path" during solidification. Decreasing d^x nnoves paths more toward the

top and right of the diagram and here recalescence becomes more rapid, more adiabatic, and more

Newtonian.

Figure 6. Thermal profiles for a recalescing

pure aluminum droplet, lOym dia.

Advancing liquid solid front shown in

insert which denotes fraction solid

and interface velocity(cm s“M» fi^om Levi and Mehrabian [34].

TEMPERATURE (K) 700 800 900

Figure 7. Effect of the ratio K/Bi on the

thermal history of undercooled

droplets for a fixed nucleation

temperature. Note that for typical

atomization conditions (k/Bi>100) the

external cooling is not able to slow

down recalescence to a significant

extent, from Levi and Mehrabian [34].

We do not yet have experimental measurements of growth velocities or interface morphology in

undercooled pure metal droplets. We can, however, turn to other studies where growth of crystals

into pure undercooled liquids has been measured, and where this growth is apparently dendritic.

Perhaps the highest measured crystal growth velocities are those of nickel and cobalt into their

pure undercooled melts, reported by Walker [37], Colligan and Bayles [38]. In their experiments,

columns of the bulk melt were undercooled by varying amounts up to 250-300°C and then seeded at one

end. The velocity was determined from the times at which the freely moving dendrite tip passed

thermal sensors sited along the column. Similar experiments have been performed by Glicksman et al [39] on succi nonitrile, and successfully analyzed by combining heat flow and interface stability

theory. By addition of interface kinetics to this analysis, Coriell and Turnbull [40] showed

excellent agreement of theory with experiment, for the undercooled nickel experiments described

above. These experimental studies, as well as interface stability theory suggest solidification of

undercooled pure metals (and alloys) should be dendritic. Nonetheless, there are indications that

undercooled alloy droplets at high undercooling solidify in a non-dendri tic fashion. Perhaps this

is because of transient effects, or perhaps because the wavwlength of the stability is long compared

37
with the droplet size [41].

Alloy Solidification

In an earlier section, the possibility has been described of achieving non-equilibrium phases

through undercooling (e.g. obtaining a BCC phase in a normally austenitic steel). And of course, glassy structures can also be obtained. In this section, we focus on structure and segregation

in alloys where the expected phases are formed (i.e. those obtained in usual solidification process).

A number of investigators have shown that solidification in bulk iron, nickel, and copper

alloys is dendritic at moderate undercoolings, with the dendrite structure becoming gradually a simple rod-like structure with increasing undercooling. Then, at a critical undercooling which

often seems to be about 100-200°C grain size drops abruptly; an example is shown in Figure 8 [42]. For undercoolings greater than the critical, the dendrite structure is completely absent, replaced

by more-or-less spherical "elements" which are the separate grains [7,42]. Both the dendrite arm

spacing in moderately undercooled specimens, and the size of the spherical elements are strongly

dependent on cooling rate after recalescence. Hence, it is clear that the structures we see are strongly affected by ripening and are not therefore the actual structures that grew during recalescence. The abrupt grain size reduction has been althernately explained by an acoustic nucleation

mechanism, and by a ripening or dendrite remelting mechanism.

Especially at intermediate to rapid rates of heat extraction during and/or immediately after

recalescence, a number of investigators have observed solute rich cores within the dendrites or

"elements" in undercooled alloys. Figure 9 is an example from work of Wu [43] showing high solute

cores in 5mm dia. Fe-25%Ni alloy that was undercooled 90°C. The solute rich cores certainly

formed during recalescence. Perhaps they grew in diffusionless fashion below the Tq, but the

probable remelting, coarsening, and solid state diffusion occuring during and after recalescence

did not completely eliminate these traces of the original structure.

Related structures have been obtained by Levi and Mehrabian in aluminum rapidly solidified

droplets. Homogeneous solute rich material is found in a portion (or all) of the droplets when

a given droplet is not completely homogeneous it contains also a solute depleted zone and dendritic

or cellular region as well. The homogeneous solute rich region must have formed at some temperature

below the thermodynamic Tq temperature, and grew in an apparently non-dendritic manner, at some

critical temperature during recalescence, solute began to be rejected, eventually rendering it

unstable and resulting in the cellular structure at the upper right.

AT, Degree of Undercooling (®F)

Figure 8. Grain size versus undercooling in

a bulk undercooled melt, from

Kattamis and Flemings [42].

Figure 9. Solute rich cores in undercooled

Fe-25%Ni alloy, AT=90°C, 700X,

from Wu [43].

With sufficient undercooling and sufficiently rapid heat extraction, supersaturated structures

and non-equilibrium phases can be retained to room temperature. Kelly, Vander Sande and Cohen [28]

have obtained a homogeneous (non-cellular), non equilibrium BCC structure in atomized 303 stainless

steel powder in diameter up to about 70 microns. Figure 10 shows similar homogeneous structures in

316 stainless steel obtained by Macisaac [44] in thesis work at MIT. These latter structures are

of both the equilibrium FCC phase and the non-equilibrium BCC phase; they were obtained in an emulsification of steel in glass. These structures also appear to have grown in non-dendritic manner.

38
5

i Figure 10. (a) Optical micrograph of cross-sectioned and etched 316 BCC and FCC particles

I after undercooling. Fee particles are light in color, (b) Backscattered

I electron micrograph of the same particles, from Macisaac [44].

1 The attributes of high undercooling and rapid heat extraction are optimized in emulsions

i when the emulsion itself is rapidly cooled by external means. One way to do this is shown in

I Figure 11 [45], where an emulsion of a metal in a carrier fluid is injected into a cooler miscible

fluid that is being vigorously agitated. This is a technique developed by M.G. Chu in his doctoral

thesis work at MIT. Figure 12 shows two structures of Sn-5%Pb alloy illustrating the importance of

this external cooling. When the alloy is emulsified in such a way that it can achieve an undercooling of 100°C or more at slow cooling rates, and then is rapidly cooled using the apparatus of

! Figure 11, a fully homogeneous, non-dendritic structure is obtained, at least at lOOOX, Figure 12a.

j This is in spite of the fact that the maximum equilibrium solid solubility of lead in tin is 2.5%.

When, however, these samples are more slowly cooled, or reheated, or held for an extended period of

time at room temperature, substantial precipitation of lead rich phase occurs. Figure 12b. In the

case of Chu's work, the effect of the quench was probably primarily in preventing solid state

decomposition. In other cases (e.g.) quenching liquid iron in molten salt. Table 2, heat extraction

rates might conceivably be increased enough by the cold salt to influence solidification kinetics.

Figure 11. Diagram of apparatus for rapid

quenching of molten emulsified

droplets, from Chu et al [45].

Figure 12. Microstructures of emulsified

droplets of Sn-5%Pb: (a) Rapidly

cooled to -25°C, (b) Rapidly cooled

and then held one month at room temperature lOOOX.

39
Quite complex solidification structures can be obtained in undercooled alloys when more than || one phase forms during recalescence. Chu has studied a wide variety of such structures in tin-lead j| alloys; the structures obtained depend on alloy, undercooling and rate of heat extraction during |' and after recalescence. with the particular emulsification technique he used, and a moderate - cooling rate (10°C/min), he found the nucleation temperatures plotted in the phase diagram in

Figure 13a. A tin rich phase forms at and a lead rich phase at Tp. Perepezko [21] previously

|| had found similar nucleation temperatures. Figure 13b, and by using a different emulsification |i medium, he was able to get the lower nucleation temperatures also shown in Figure 13b. iSolidification of hyper-eutectic tin-lead alloys was found to occur in the following sequence: |j (a) nucleation at Tp of a lead rich phase, (b) dendritic growth of this phase, first during recales- li cence and then during subsequent cooling Tn, where a tin rich phase nucleates, (c) growth of the tin :! rich phase during recalescence and during subsequent cooling. Note that the recalescence and solute if distribution during growth of the second phase can be expected to influence the first phase to form. || In addition the intervention of the formation of a metastable phase prior to the tin rich phase due

to catalysis by the lead rich phase [16] can also influence the overall solidification sequence and

microstructure. Figure 14 shows examples of solidification structures of such an alloy. Figure 14a ! is a strucutre that was undercooled and rapidly quenched. Lead rich dendrites (wliite) first formed h and tin rich phase (dark) later filled in within the interstices. Both the tin rich and lead rich 'jj phases shown substantial second phase with themselves, presumably from solid state precipitation. jij Figure 14b is a sample that was allowed to nucleate on slow cooling and then rapidly solidified. p- The dendritic structure formed as before, and the final structure is very little different. The j! sample allowed to cool slowly (10°C/min) however, shows massive ripening of the original lead rich

phase. Figure 14c. !'

Three different tin-lead structures of eutectic composition are shown in Figure 15. The first [ is comparable to the slowly cooled Sn-45%Pb alloy discussed above, in that it shows primary, ripened

lead rich phase (with entrapped and precipitated tin rich phase). Figure 15a. The second shows the T

same alloy rapidly cooled. Figure 15b. Apparently the rapid cooling suppressed the nucleation l!' temperature of the lead rich phase, Tp below that of the tin rich phase, Tn. The third structure,

Figure 15c, from work of Cooper, Anderson and Perepezko [46] is similar to 15b, but somewhat coarser. :|i In this case the nucleation temperature Tp was depressed below Tp by a modified emulsion; this

depression is shown in figure 13. The second two eutectic structures are similar to those obtained

in conventional rapid solidification processes, e.g. wheel casting [47]. li

Recalescence Measurements in Droplet Solidification

Ewasko [48] and Wu [43] in

separate theses at MIT are attempting

to gain some experimental understanding of the "solidification

paths" of undercooled alloys by

looking at thermal behavior and

structure of 5mm dia. metal droplets

undercooled in molten glass. Their

results to date illustrate well

how much we do not understand yet

about this process.

As one example, the rate of

recalescence of the droplet surface

increases dramatically with increasing undercooling, Figiure 16. We believe that this cannot be a heat flow caused phenomenon since

experiments on pure nickel show

orders of magnitude faster recalescence. Thus, a preliminary con- clusion is that these rates of

recalescence (at least at temperatures approaching and above the

calculated Tq) are limited by

solute transport. The Tq temperature in Figure 16, incidentally,

was calculated from a regular

sol ution model

Composition of lead (at“/o Pb) 20 40 60 80 Pb Sn

Composition of lead (at7*Pb) 20 40 60 80 Pb

Figure 13. Nucleation temperature in supercooled Sn-Pb

alloys. T is nucleation temperature of

tin rich pnase and Tp of lead rich phase.

Diagram on left is from Chu et al [45] and

on the right is from Cooper et al [46].

Exact amounts of undercooling depend on emulsification procedure. The two different

sets of curves in diagram on the right result

from the different emulsification procedures.

40
i| Figure 14. Microstructures of Sn-45%Pb alloy droplets, 7000X. (a) Rapidly quenched

i (b) Slowly cooled to just below and quenched (c) Slowly cooled to room

' temperature.

Figure 15. Microstructures of emulsified droplets of eutectic Sn-Pb alloy, (a) Slowly

cooled, moderate undercooling, (b) Undercooled and rapidly solidified,

(c) Slowly cooled, high undercooling.

The photomicrographs in Figure 17

provide some insight into factors influencing this recalescence, by showing

the very different structures at different

undercoolings. At moderate undercooling

where temperature is estimated to be

above the thermodynamic Tq temperature,

recalescence is very slow, and a two phase structure results that has a coarse dendritic appearance. Figure 17a.

Initial results show the core of this

"dendrite" is of approximately Cq

composition, the periphery is depleted,

and the interdendritic regions appear

to be somewhat above C^, possibly it

grew as a single phase initially

into solute redistribution at least

during later stages of growth. If

so, we would expect re-melting on

a fine scale within the dendrites

as it recalesces above some temperature approaching T^; the liquid

pools remelting could lead to the

finely distributed second phase

within the "dendrite".

Figure 16. Recalescence curves from 5mm

droplets of Ni-Sn eutectic alloy

at several undercoolings, from

Ewasko [48].

The sample at the intermediate undercooling in Figure 17 show a fine grain structure with a multiplicity of growth sites, each of which has a central region (two phase) that may well have

41
initially grown in a diffusionless manner below Tq and later partially remelted during recalescence.

The fine scale of this structure due to the multiplicity of growth sites would explain the more

rapid recalescence, even above T . This qualitative description of solidification behavior for

such an alloy is largely conjecture, but the approach suggests an experimental direction toward

an eventual fuller thermal and solidification description of under,cooled alloys.

Undercooling 70°C Undercooling 120°C Undercooling 380°C

Figure 17. Structure of 5mm droplets of

undercoolings, from Wu [43].

li-Sn eutectic alloy solidifed at 3 different

Concl usions

In examining the nucleation and solidification of undercooled liquid metals it is apparent that

the experimental and theoretical understanding has made good progress since the initial pioneering

work of Turnbull. From the present experience with droplets studies it has been possible to identify

some general guidelines for the interplay between process variables such as cooling rate, undercooling, alloy composition and particle size and the nucleation phase selection involved in RSP. With droplet samples and controlled undercooling conditions new, understanding has been gained into

the development of solidification microstructure and the influence of thermal history on the final

microstructural morphology. This area of research in the study of RSP, we believe, has a great

potential for the further development of nucleation and solidification theory. It also has the

potential for the development of new processes of engineering importance. In this regard it seems possible that we can someday learn to enhance undercooling with controlled surface treatments in

existing RSP methods, and perhaps, as well, find ways to bring other processes such as emulsification

to the point of engineering usefulness.

Acknowledgement

The support and the ARO (DAAG29-80-K-0068) for studies on solidification structures and the

NSF (DMR-79-1 5802) for work on nucleation catalysis kinetics at the University of Wisconsin-Madi son

is gratefully acknowledged. Portions of this work (those at MIT) were sponsored by NASA Grant

No. NSG 7645 , from Materials Processing & Space Division, NASA office, with Dr. Louis Testardi,

project monitor.

References

[1]

. J. H. Perepezko, in Proceedings of Rapid Solidification Processing: Principles and Technologies

J_I_, R. Mehrabian, B. H. Kear and M. Cohen eds., Claitor's Pub., Baton Rouge, La. 56 (1 980).

[2]

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[3]

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[4]

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[5]

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42
[6]

. L. L. Lacy, M. B. Robinson and T. J. Rathz, J. Crystal Growth 47 (1980).

[7]

. T. Z. Kattamis and M. C. Flemings, Trans. AIME 236 , 1523 (1966). [8]

[9] . S. Y. Shirashi and R. G. Ward, Canadian Met. Quaterly, 117 (1964).

[10]. D. Turnbull, J. Chem. Rhys., 411 (1952).

. J. H. Perepezko and J. S. Paik, in Proceedings of the 29th Midwest Solid State Conference:

Novel Materials and Techniques in Condensed Matter, G. Crabtree and P. Vashishta eds.. North

Holland, N.Y., 57 (1982).

[11]

. D. Turnbull, J. Chem. Phys., [[8, 198 (1951 ).

[12]

. C. V. Thompson and F. Spaepen, Acta. Met., 1855 (1979).

[13]

. M. Hillert, Acta Met., 763 (1953).

[14]

. C. V. Thomson and F. Spaepan, in Proceedings of the Materials Research Society Symposium : Materials Processing in the Reduced Gravity Environment of Space, G. E. Rindone eds.. North

Holland, N.Y., 603 (1982).

[15]

. J. S. Paik and J. H. Perepezko, to be published.

[16]

. A. S. Skapski, Acta Met., 576 (1956).

[17]

. F. Spaepen and R. B. Meyer, Scripta Met., J_0, 257 (1976).

[18]

. R. H. Ewing, Phil. Mag., 778, 1972.

[19]

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[20]

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[21]

. J. H. Perepezko and J. S. Paik, in Proceedings of the Material Research Society Symposium : Rapidly Solidified Amorphous and Crystalline Alloys , B. H. Kear, B. C. Giessen and M. Cohen

eds.. North Holland, N.Y., 49 (1982).

[22]

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[23]

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[24]

. A. J. Drehman and 0. Turnbull, Scripta Met., 1_5, 543 (1981 ).

[25]

. B. C. Giessen and R. H. Willens in Phase Diagrams: Materials Science and Technology Vol . Ill , A. M. Alper ed.. Academic Press, N.Y. , 104 (1970)

[26]

. J. C. Baker and J. W. Cahn, in Sol idification , ASM, Metals Park, Ohio, 23 (1971).

[27]

. J. L. Murray, to be published

[28]

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[29]

. C. M. Adam, Private Communication.

[30]

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[31]

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Company, Reading, MA(1973).

[33] N.J. Grant, in " RAPID SOLIDIFICATION PROCESSING-PRINCIPLES AND TECHNOLOGIES I ", Baton Rouge,

Claitor Publishing Division, 230(1978).

[34] C.G. Levi and R. Mehrabian, " HEAT FLOW DURING SOLIDIFICATION OF UNDERCOOLED METAL DROPLETS ", to be published.

43
[35] D. Turnbull, " THERMODYNAMICS IN PHYSICAL METALLURGY ", Americal Society for Metals, Metals

Park, 0H(1949).

[36] J.W. Cahn, W.B. Hilling and G.W. Sears, Acta. Met. , 12, 1421(1964).

[37] J.L. Walker, unpublished work presented in " PRINCIPLES OF SOLIDIFICATION ", B. Chalmers,

Chapter 4, John Wiley(1964).

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[40] S.R. Coriell and D. Turnbull, submitted Acta. Met. (1981).

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TECHNOLOGIES III. , to be published.

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,

75, 191 (1967).

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METALS VOL. 1 " 107, Sendai, Japan (1981).

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[48] Ricky Ewasko, S.M. Thesis, Department of Materials Science and Engineering, Massachusetts

Institute of Technology(1982).

44
MECHANISMS OF SEGREGATION-FREE SOLIDIFICATION

f

fi

W. J. Boettinger, S. R. Coriel 1

Metallurgy Division

National Bureau of Standards

Washington, DC 20234

and

R. F. Sekerka*

Carnegie-Mel 1 on University

Pittsburgh, PA 15213

ABSTRACT

A frequent microstructural feature of rapidly solidified alloys is the absence of

cellular or dendritic microsegregation, often accompanied by incorporation of solute in excess of the equilibrium solid solubility. Two solidification mechanisms can produce

microsegregation-fre'e crystalline solids: planar growth and partitionless solidification.

The solidification conditions required for planar growth with equilibrium partitioning

at the interface are well known. For growth at high velocities, capillarity provides a stabilization of the planar interface. This regime is referred to as absolute stability

and applies only when the net heat flow is toward the solid. Partitionless solidification

arises from the kinetics of interface motion. These kinetics can involve the trapping

of solute by the moving interface, causing the partition coefficient to deviate from the

equilibrium value. Solute trapping generally occurs at high interface velocities and

may be pronounced when the interface temperature is far below the liquidus. In the limit

of partitionless solidification, when the partition coefficient is unity, the interface

temperature is below the Tq temperature for the composition at the interface. Partitionless

solidification provides a mechanism for the formation of segregation-free crystalline

alloys when the net heat flow is toward the liquid or when the interface is not planar.

Recent work has combined a velocity dependent partition coefficient with interface stability

theory to predict the solidification conditions required for microsegregation-free solidification. Under some conditions, oscillatory instabilities are predicted. The relation

of the interface temperature to the liquidus, solidus and Tg temperatures as a function of

growth rate and considerations necessary to predict the conditions of microsegregation-free

solidification in concentrated alloys are discussed.

Introduction

Microsegregation-free crystalline alloys including those with solute in excess of equilibrium

solubility provide an ideal starting microstructure for the development of interesting mechanical

properties during subsequent thermomechanical processing [1]. In this paper, the following three

topics related to the solidification conditions required to produce such microstructures will be

described: the interface temperature during solute trapping; the interface stability theory; and

the conditions for microsegregation-free solidification of concentrated alloys. It will be assumed

! that the phase which is forming from the melt has been selected previously by nucleation kinetics or

I by competition among growing crystalline phases. Under this assumption, the primary role of nucleation is to allow the development of a large bulk undercooling of the liquid into which growth can

; occur quite rapidly. This paper focuses on the conditions at the solid-liquid interface, with the

[ realization that a full analysis of heat and solute transport must be conducted to form the basis

[j for a final choice of processing conditions.

Consultant, National Bureau of Standards

45
Interface Conditions

In 1971, Baker and Cahn [2] posed the solidification problem of a binary alloy using two response

functions. One choice for the response functions describes the interface temperature, Tj, and the

composition of the solid at the interface, These response functions can be written as follows:

where V is the local interface

T^r is a capillarity constant,

zero velocity, the functions f

Tj = f(V,CL )

Cs* = C|_*k(V,CL*)

T,^rK

velocity,

and K is

and k are tk

is the composition

e mean curvature of

very simply rel^ated to

liquidus and k(0,C. ) is

of the liquid at the sol id -liquid

(

1

)

( 2 )

the interface

the equation for the phase diagram

partition coefficient, k£, which can depend on compositionT

the phase diagram:

the equation for the

interface. Atj

f(0,CL*) is

equi1 bri umi

These response functions are constrained by thermodynamics. Fig. 1 shows the molar free

energy versus composition for a liquid and a solid phase at a fixed temperature. During solidification, the temperature and the compositions at the interface are constrained by must be a solid of

Fi g. 1 , and

net decrease jn the free

composition C. from liquid

energy per mole,

of composition

the fact that there!

aG, needed to form an infinitesimal amount ofj

[2]. The free energy change is shown inj

is given by

(3) i;

Under the usual assumption of local equilibrium at the solid-liquid interface, believed to be. valid j

| at slow rates of solidification, the chemical potentials of the liquid, , and solid, for ij

each species, i, must be equal. During rapid rates of solidification, this is not necessarily the ji i case and nonequilibrium conditions at the interface can exist. The term solute trapping is loosely ; applied to conditions when the partition coefficient deviates from the equilibrium value, but as originally defined, is restricted to the case when the chemical potential of the solute increases i'j during solidification. Baker and Cahn showed that the restriction that aG be negative could be [’M represented on a figure where, for constant temperature, the range of possible solid and liquid j

Fig. 1. Construction to show the Gibbs free energy per mole, aG, for

solidification of an infinitesimal amount of solid of com- position Cg from a liquid of composition C|_* under non- equilibrium conditions„for which the chemical potentials

and i

1-;

46
(a) (b)

I Fig. 2. Regions (cross-hatched) of thermodynamically allowed solid compositions that

I may be formed when solid solidifies from liquid of composition at various

I temperatures. The value of Tg is the highest temperature at which partitionless

I solidification of a liquid of given composition can occur. In (b), the Tg

i curve plunges and partitionless soldification is impossible for liquid of

] composition C|^ . (From reference [3]).

concentrations are shown. Fig. 2 shows an alternate representation of this restriction [3]. The

I cross-hatched jjegion shows the range of possible solid compositions that can form from a liquid of

! composition C|_ at vari^ous temperatures. When the temperature is equal to the liquidus temperature

I for the composition C|_ , the only possible solid composition is the equilibrium solid composition.

I For interface temperatures below the liquidus, the range of allowable solid composition expands. Of

I particular interest is the Tg curve which lies between the liquidus and and solidus. The Tg curve

; is the locus of compositions and temperatures for which the molar free energy of the liquid and solid

: phases are equal. When the temperature of the liquid at the interface reaches the Tg curve, the

j solid can have the same composition as the liquid, i.e., partitionless solidification. Fig. 2b shows

a case where the T^ curve plunges quite steeply. In this case, for the composition C|_ , partitionless solidification is not possible, i.e., the partition coefficient, k, cannot approach unity at high velocity.

: Many models for the dependence of the partition coefficient on velocity have been formulated

' [4-7]. This dependence on velocity constitutes one of the two response functions required to analyze

i rapid solidification problems. The model formulated by Baker [4] is quite general and includes a

I wide variety of possibilities. Fig. 3 shows the results of more recent theories which predict that

the partition coefficient changes monotonical ly from its equilibrium value to unity as the growth

velocity increases. In these models the interface partition coefficient is a function of a dimensionless velocity, SgV, where 3g is the ratio of a length scale, ag, normally related to the

interatomic distance and a diffusion coefficient, D. The diffusion coefficient in the various

models ranges between a liquid diffusion coefficient and a diffusion coefficient characteristic of

the interface itself. The functional form of the model proposed by Aziz [5] and Jackson, Gilmer and

Leamy [61 for continuous growth is given as

kp + 3„V

k(V) = (4)

where gg = ag/D. At a velocity of l/Bg, the partition coefficient is the average of the equilibrium partition coefficient and unity. If we choose a liquid diffusion coefficient typical of metals

(10“^ cm^/s) and the length scale to be 5 a , l/ 3

_ is 200 cm/s. However, the value of I /3 which

best fits [5] recent experimental data for dopecP Si is lO m/s; for Si , D = 5 x 1 0-4 cm^/s [ 8 ],

and hence ag = 50 A. It remains a challenge to design good experiments to determine the velocity

required for solute trapping in metallic alloys.

There are significant shortcomings in the present models of the dependence of the partition

coefficient on velocity. The models primarily treat dilute alloys so there is no explicit dependence

on composition. For the case of Fig. 2b, where the Tg curve plunges, the partition coefficient

cannot approach unity at high velocity. A second shortcoming of the present trapping models is that

the response function for the interface temperature is not treated.

47
Fig. 3. Curves showing the dependence of the interface partition

coefficient, k, on velocity according to the models of various

i nvesti gators.

Following the suggestion of Aziz [5] a model for the response function for the interface temperature can be formulated. For pure materials Turnbull et al. [9,10] have suggested that the rate at which atoms attach to the growing crystalline solid may only be limited by their rate of collision

with the solid. In such a case, the interface velocity is related to the interface temperature, Tj,

by an expression

V = fVod-exp (aG/RTi)} (5)

where f is the fraction of sites on the solid-liquid interface available for attachment, Vq is the

speed of sound in the liquid metal, AG is the molar free energy for solidification and R is the gas

constant. For most metals, f is thought to be close to unity and for velocities much less than the

speed of sound, this expression can be written as

V

= -V

0

( 6 )

For the solidification of disordered crystals, this collision limited model of solidification may

apply as well. This model will be employed to formulate a response function for the interface temperature for a binary alloy. The molar free energy for solidification is evaluated from Eq. (3) and,

in special cases, can be described simply in terms of undercooling. When solute trapping occurs a simple expression in terms of undercooling is not possible. For pure metals aG reduces to

AG (7)

where L is a latent heat and T^ is the melting point of the pure metal. For dilute alloys, when the

partition coefficient equals the equilibrium partition coefficient, k^, the thermodynamic driving

force for solidification is given by

48
AG =

( 8 ) -L(VTj

; Iwhere T|_ is the liquidus temperature for the concentration of the liquid at the interface

' jwhen the partition coefficient is identically equal to one, aG has a similar form,

AG =

-L (VT,)

( 9 )

where the undercooling (Tq-Tj) is referred to the Tq temperature for the interface composition.

i'lThe similarity of the expressions for the case where k equals one, i.e., partitionless solidification,

'and the expression for pure metals, is the origin of the often stated idea that the kinetics of

I partitionless solidification are quite rapid and similar to that for the solidification of a pure

Uetal. In the general case when the partition coefficient is not the equilibrium value, the thermo-

! dynamic driving force for solidification must be evaluated from the full expression given by Eq. (3).

I In this expression the chemical potentials are functions of concentration and temperature that may be

determined by modelling the phase diagram. This demonstrates the importance of modelling phase

' diagrams, even for applications to rapid solidification. If Eqs. (3 & 5) are combined with, for

i example, the trapping model of Aziz, Eq. (4), the interface temperature can be determined as a function of the interface velocity and the concentration of the liquid at the interface. A similar

I calculation for other models of solute trapping or interface temperature can be performed. Alternatively, part of the free energy for solidification can be dissipated by solute drag, as reported by

! Aziz in the current proceedings [11].

i

I Fig. 4 shows a composite plot of the two response functions obtained by using Eqs. (3-5) superI imposed on the phase diagram including the liquidus, solidus and Tq curve. The composition of the

j solid at the interface and the interface temperature are plotted along a curve parameter! zed by

interface velocity for a given fixed liquid concentration at the interface. The figure is based on

! thermodynamic data for Ag-Cu [12], a liquid conposition of 8.2 at% Cu, a value of of 10“^ s/cm

and for two values of V (infinity and 2 x 10^ cm/s). At zero velocity, the composition of the

solid lies on the solious curve. At intermediate velocities (~10 cm/s), the concentration of

the solid moves towards the concentration of the liquid with a small increase in the undercooling.

At high velocities the solid composition at the interface approaches the liquid concentration at

the interface near the Tq curve. The dashed curve shows the case where Vq is infinite. This

corresponds to the thermodynamic bound on the solid concentration at the interface given in the

Fig. 2. The dashed curve can also be obtained by letting aG = 0 in Eq. (3). The solid curve,

for a finite value of Vq, shows the interface temperature plunging quite rapidly. This analysis

provides, for the first time, a pair of response functions for the conditions at the liquid-solid

interface, thus permitting the analysis of rapid solidification problems when combined with solute

redistribution and heat flow analysis. Such curves are of interest during recalescence when the

velocity of the interface starts out at a high value and decreases due to the evolution of latent

heat. They are also required for a detailed analysis of the stability of the solid-liquid interface

during solute trapping.

Interface Stability

Interface stability theory is of interest because, during planar growth, solidification occurs

with no lateral microsegregation. The theory was first developed by Mullins and Sekerka [13] and

has been extended and discussed in more detail for rapid solidification by Coriell and Sekerka [14-

15]. The present paper is restricted to a discussion of the simple asymptotic form of the stability

theory relevant at high velocity. This asymptotic form is known as absolute stability and is given

by the expression

V

= mD(l -k^)C

( 10 )

where m is the liquidus slope, D is the liquid diffusion coefficient and C is the alloy composition.

Whenever V exceeds the value given in Eq. (10), a planar liquid-solid interface is stable. Absolute

stability of a planar interface occurs because of the influence of capillarity. At high growth

velocities, there is only limited time available so lateral solute segregation in the liquid can

49
L

Fig. 4. Interface temperature and solid compositions for solid that

forms at the indicated velocities from a liquid of fixed

composition at the interface. Thermodynamic parameters

are for Cu in Ag for a specific kinetic law with = 2 x 10^

cm/s (full curve) and Vq + “ (dashed curve).

only occur for perturbations of the solid-liquid interface with very short wavelengths. These ^ short wavelengths require such a large increase in the area of the interface that the perturbations*!

are retarded by capillary forces and the interface is stable.

4

There are several restrictions on the use of the simple asymptotic form given by Eq. (10). i First, rather trivially, this form is only valid if the partition coefficient is the equilibrium j partition coefficient. Modifications required when the partition coefficient depends on the velocity

will be discussed in a later section. A second restriction is that the net heat flow must be into

^ the solid, i .e.

,

*

G

k . G I + k G

L L s s

k, +k

L s

> 0 ( 11 )

Here ki_ and kg are the thermal conductivities of the liquid and solid respectively. Gi and Gg are

the temperature gradients in the liquid and solid respectively. In the case of solidification

following surface melting, the conduction of heat is almost completely into the solid and G* is

always positive. For quenching against a substrate or for atomization, the sign of G* is less

clear. What is important is the relative amount of heat flowing into the solid and flowing into the ,, liquid. For quenching against a substrate, bulk undercooling of the liquid can occur if nucleation i is difficult. In this case, during solidification heat can flow both into the solid and into the ; liquid to remove the latent heat evolved at the liquid-solid interface. An experimental observation

in which planar growth is stable, and G* is negative, would require a reexamination of the assumptions

50
i|;in the present stability theory. In this event two possible modifications [14] would be an allowance

I for a velocity of solidification that depends on time or an analysis of solidification in samples in

which the wavelength of the fastest growing instability is of the same order as the sample itself. A

|third restriction on the use of the absolute stability equation is that one must be far from the

pdified constitutional supercooling regime of the full stability theory. This condition will be

Isatisfied whenever the velocity

V >>

( k + k I s L

( 12 )

li ^provided also that the temperature gradient in the liquid is not too large [GL<<VL/(2ki_) ] as is

jusually the case for rapid solidification. Here, L is the latent heat per unit volume. This

restriction is more important for doped Si [15] than for metallic alloys due to the relatively high

'latent heat and diffusion coefficient for Si.

I In a recent set of experiments employing electron beam surface melting [16-17] the applicability

I

of the morphological stability theory to rapid solidification has been demonstrated. Fig. (5) shows

j the transition from cellular to cell-free structure which occurs in Ag-1 wt% Cu and Ag-5 wt% Cu alloys,

j For the former, the transition occurs between 10 and 20 cm/s, and for the latter, the transition

I occurs between 50 and 60 cm/s. Fig. 6 shows a summary of the experimental results for the two alloys.

The transition from cellular to cell-free material is seen to occur at velocities roughly a factor of

j2 less than the theoretical predictions. Such agreement is considered reasonable in view of the

I imprecise knowledge of the materials parameters required for the theory. A more complete discussion

jj of these results occurs elsewhere [16].

: i

: I

Fig. 5. a) Cellular and b) cell-free structures obtained in Ag-1 wt% Cu alloys rapidly solidified at 10 and 20 cm/s respectively,

c) Cellular and d) cell-free structures obtained in Ag-5 wt% Cu alloys rapidly solidified at 30 and 60 cm/s respectively [16],

TEM.

51
Fig. 6. Summary of structures obtained [16] by electron beam

surface melting and resolidification of Ag-1 and 5 wt% Cu alloys at various growth rates. The comparison of the

absolute stability theory with experiments is shown.

When the partition coefficient depends on the interface velocity, the stability theory must be

modified [18]. This modified analysis of stability suggests the possibilty of increased insta- 4 bility of a planar liquid-solid interface, attributed to a feature called the "solute pump mecha- 1 nism". Physically, the partition coefficient will vary across a perturbed solid-liquid interface j because of local differences in velocity. This leads to a variation of the solute rejection rate |: along the interface and produces lateral solute segregation without the necessity of lateral diffu- j. sion. Hence perturbations with wavelengths much longer than those normally encountered at high

| growth velocity can occur and capillarity is less effective in stabilizing such perturbations. The [ results of the theory are complex and will only be described qual itati vely here. It is found that Ij the actual value of the partition coefficient can be substituted into the absolute stability equation jj only under certain conditions. If the rate of change of the partition coefficient with velocity is |

relatively small, and/or the absolute value of the rate of change of the interface temperature with ; respect to velocity is relatively large, this simple substitution method is valid. On the other

j hand, if the rate of change of the partition coefficient with velocity is large and/or the absolute i value of the rate of change of the interface temperature with velocity is small, the simple substitu- I; tion is not valid. Clearly the details of the size of these derivatives depends critically on the

parameters 3g Vq described in the previous section. Fig. 7 shows a plot of the demarcation j| between stability and instability for Ag-Cu alloys for a choice of 3 g of 10“^ s/cm and for various j" values of \ij where yj = - 3V/aTj. On the right side of the figure is shown the variation of i the partition coefficient with velocity. On the left side is shown the demarcation between stability f and instability. Depending upon the value of yj, the stability criterion can change dramatical ly. j,

I It should be pointed out that the instability in the case of the "solute pump mechanism" is j oscillatory in nature. This arises because of the temporal phase difference between the local veloc- ( ity of the interface and the interface shape. Such oscillatory instabilities lead to microsegrega- | tion patterns that are three dimensional, whereas an ordinary microsegregation pattern is two dimen-

[ sional (in directions transverse to the growth direction). Such oscillatory instabilities have not been observed, and only future experiments will determine their importance. g

52
100

10

0 5 10 15 0.4 0.6 0.8 1.0 COMPOSITION (wt % Cu) INTERFACE PARTITION

COEFFICIENT, k

Fig. 7. Curves showing the onset of morphological instability for nonequilibrium

conditions for which the partition coefficient, k, varies with V as shown on the right. Velocity for instability is shown as a function of

the composition of bulk liquid for various values of the kinetic co- efficient \ij. For uT <. 1 , the transition to instability is non-oscil latory

in time and corresponds to a simple substitution of k(V) into the standard

criterion for instability. The units of uT are cm/(Ks).

The solidification conditions required to produce microsegregation-free crystalline structures

for concentrated alloys are generally associated with the temperature difference between the liquidus

curve and the Tq curve. For most terminal solid solutions, this difference increases with increasing

concentration. The usual idea is that, compared to dilute alloys, such alloys require more rapid

solidification or higher undercooling to produce microsegregation-free crystalline alloys. The

simple trapping models presented earlier in this paper involve no explicit concentration dependence.

Furthermore, a recent set of experiments conducted on Ag-rich Ag-Cu alloys [16-17], suggest that as

the composition approaches the eutectic composition, the growth velocity required to form microsegregation -free structures decreases. At 15 wt% Cu, an interface velocity of approximately 2 m/s is

required. At the eutectic composition (28 wt% Cu), a velocity of only 70 cm/s is required. Therefore,

in this alloy, it is actually easier to form the microsegregation-free structure for the very concentrated alloys. Such results have several possible origins and one possibility is related to the

dependence of the liquid diffusion coefficient on temperature, as discussed below.

The dependence of the liquid diffusion coefficient on temperature is generally neglected in

ordinary solidification theory. Such is the case because of the fact that solidification almost

always takes place very close to the liquidus temperature. Diffusion data is often described by an Arrhenius equation. For a melting point of lOOOK and an activation energy of 10 kcal/mol, an undercooling of 300K will reduce the diffusion coefficient by an order of magnitude.

The dependence of the diffusion coefficient on temperature can affect substantially various ideas

in solidification. In the following, it is assumed that the diffusion coefficient is only a function

of temperature. The first example involves solute redi stribution in the growth of concentrated

alloys. In particular it has been noted previously [3] that the dependence of the diffusion coefficient on temperature can lead to a maximum in eutectic growth rate of approximately 10 cm/s.

maximum velocity has origins that are identical with the nose of the C curve for eutectoid growth

observed in solid state transformations. Such a maximum eutectic growth rate is consistent with the

fact that eutectic spacings are rarely observed to be finer than 200 A. A second example may occur

in the stability theory. For planar growth with equilibrium at the solid-liquid interface, the

interface temperature is at the solidus for the average alloy concentration. In many concentrated

Concentrated Alloys

53
alloys, the solidus may easily lie 300 K below the liquidus. An order of magnitude change in the

diffusion coefficient can affect the prediction of the stability theory significantly. A third

example involves a very simple modification of the trapping model presented previously. The parameter 60 involves a diffusion coefficient. A simple approach to include composition dependence

in a trapping theory would be to calculate the diffusion coefficient at the Tq curve for various

compositions. For pure silver, the diffusion coefficient is 2 x 10~5 cm^/s at its melting point.

For the eutectic composition the Tq temperature is approximately 670 °C and the diffusion coefficient,

based on an activation energy of 10 kcal/mole, is 6 x 10"° cm^/s. This simple idea leads to the

conclusion that the growth velocity required for significant levels of solute trapping would be

reduced by a factor of 3.5 as the composition changes from pure Ag to the eutectic composition.

Such a trend is consistent with the observed data.

Concl usi on

In this paper, a response function for the temperature of the interface during solute trapping

has been discussed. This model for the temperature is consistent with thermodynamics and with

models of solute trapping and collision limited growth. Such interface response functions should

permit a rational approach to the solidification problem in a wide variety of conditions of heat flow

and solute redistribution. Secondly, circumstances where the absolute stability equation can be used

in lieu of the full stability theory have been addressed. Experimental results obtained by using

electron beam surface melting of Ag-Cu alloys are consistent with the stability theory. Thirdly, a few comments have been presented on the conditions for microsegregation-free solidification of

concentrated alloys. This area obviously remains ripe for experimental and theoretical investigation.

Acknowl edgment

This work was partially sponsored by the Defense Advanced Research Projects Agency. One of us

(RFS) received partial support from the National Science Foundation.

We would like to thank M. J. Aziz, J. W. Cahn, and D. Turnbull for helpful discussions.

References

[1] M. Cohen and R. Mehrabian, these proceedings.

[2] J. C. Baker and J. W. Cahn, in Sol idi fication (ASM, Metals Park, 1971) p. 23.

[3] W. J. Boettinger, in Rapidly Solidified Amorphous and Crystalline Alloys , ed by B. H. Kear,

B. C. Giessen and M. Cohen (Elsevier, 1982) p. 15.

[4] J. C. Baker, Interfacial Partitioning During Solidification, Ph.D. Thesis, MIT (1970) Chapter V. Also reported by J. W. Cahn, S. R. Coriel 1 and W. J. Boettinger, in Laser and Electron Beam

Processing of Materials , ed. by C. W. White and P. S. Peercy (Academic Press, 1 980)

p. 89-103.

[5] M. J. Aziz, J. Appl. Phys. 1 158 (1 982).

[6] K. A. Jackson, G. H. Gilmer, and H. J. Leamy, in: Laser and Electron Beam Processing of

Material

, ed. by C. W. White and P. S. Peercy (Academic, NY, 1980) p. 104.

[7] R. F. Wood, Phys. Rev. 2786 (1 982).

[8] See, e.g., J. Narayan, J. Crystal Growth^, 583 (1 982).

[9] F. Spaepen and D. Turnbull, in Rapidly Quenched Metals , ed by N. J. Grant and B. C. Giessen

(MIT Press, Cambridge, MA, 1976) p. 205.

[10] D. Turnbull and B. G. Bagley, in Treatise on Solid State Chemistry , ed. by N. B. Hannay

(Plenum, New York, 1975) Vol . 5, p. 513.

[11] M. J. Aziz, these proceedings.

[12] J. L. Murray, Met. Trans. to be published.

[13] W. W. Mullins and R. F. Sekerka, J. Appl. Phys. ^5, 444 (1964).

i- 54
[14] S. R. Coriell and R. F. Sekerka, in Rapid Solidification Processing Principles and Technologies

1

, ed. by R. Mehrabian, B. H. Kear, and M. Cohen (Claitor, Baton Rouge, LA, 1 980) p. 35.

[15] J. W. Cahn, S. R. Coriell, and W. J. Boettinger, in Laser and Electron Beam Processing of

Material

, ed. by C. W. White and P. S. Peercy (Academic Press , Rew York , 1 980) pT 8^

[16] W. J. Boettinger, R. J. Schaefer, F. S. Biancaniel lo, and D. Shechtman, Met. Trans. _A, to be

publ i shed.

[17] W. J. Boettinger, R. J. Schaefer, F. S. Biancaniello, and D. Shechtman, these proceedings.

[18] S. R. Coriell and R. F. Sekerka, J. Crystal Growth, to be published.

55
MICROSTRUCTURAL CHARACTERIZATION OF RAPIDLY SOLIDIFIED MATERIALS

H. L. Fraser* and J. B. Vander Sande**

*Department of Metallurgy and the Materials Research Laboratory

University of Illinois, Urbana, IL 61801

**Department of Materials Science and Engineering

Massachusetts Institute of Technology, Cambridge, MA 02139

Introduction

It is now well established that rapid solidification of alloys holds considerable advantage ' over conventional processing because of the fuller potential for microstructural control realizable i | through rapid solidification. With this new emphasis on microstructural control via rapid sol idif ica-l; j tion processing has come the need for detailed microstructural analysis. Also, with the reduction jit I in size of characteristic microstructural entities that occurs on rapid solidification, there developed a need for new experimental techniques that could probe both structure and composition with ew

much higher resolution than was previously available. Fortunately, the advent of scanning transmis- fn

sion electron microscopy (STEM) occurred virtually simultaneously with the developments of rapid t; solidification technologies. With STEM, the capability for compositional analysis by x-ray fluores- 7 cence spectroscopy with a spatial resolution on the order of 5 nm in electron transparent sections i has been realized. 1

>

[ Instrumentation (

The commercial instrument used in the part of this work conducted at MIT was a Vacuum Generator {j | HB5 field emission STEM [1]. Using a cold tungsten field emitter as an electron source and a two 1 : lens probe forming system, 2 nA of 100 kV electrons can be formed into a probe of approximately

1.5 nm diameter on the specimen. The vacuum in the specimen chamber is ~ 2 x 10 ^ torr, which mini- ij mizes the problem of specimen contamination. The instrument used in the part of this work conducted : at the University of Illinois was a Philips EM400T. The EM400T can operate as a conventional trans- 1

, mission electron microscope or as a STEM instrument, and has excellent selected area and convergent ' beam diffraction capabilities. An energy dispersive x-ray (EDX) analysis system as well as an elec- - tron energy loss spectrometer (EELS) are available on both the HB5 and the 400T. In each instrument, i the Si(Li) EDX detector and EELS system are coupled to a multi-channel analyzer and minicomputer '}! system for spectral manipulation. "jInterpretation of energy-dispersive x-ray data requires a great deal of care. The usual way to

|| consider the effects of x-ray production is to employ the sensitivity factor method due to Cliff and

Lorimer [2], embodied in the equation: P

C

C A

B

^a'^a 'b^b

where C^/Cg is the ratio of the concentrations of the elements A and B in a specimen; I^ and Ig are

*jij the integrated x-ray counts in chosen lines due to each element; and K. and Kp, are sensitivity ' A B |i factors. Often the ratio K^/Kg is used as a relative sensitivity factor denoted by K^g, relating j| the sensitivities of the two elements.

If the x-rays were generated in an infinitely thin specimen in free space, the measured x-ray 1 signals would correspond exactly to the generated x-ray intensities. Such, however, is not the case, j, and it is necessary to consider (a) the modification to the x-ray intensities as they travel through

adjacent material and selective absorption occurs, and (b) generation of spurious x-rays from remote

parts of the specimen and the microscope as a result of spurious radiation (electrons and x-rays)

generated in the microscope column or as a result of interaction of the main electron beam with the i specimen itself.

It has been shown that in the HB5 fitted with the so-called "virtual objective aperture," the j contribution of spurious radiation generated by the microscope is negligible compared with that , generated by the interaction of the beam with the specimen. Further, if the microscope is not so !;j equipped, this "hole count" can simply be recorded by passing the beam through a hole in the specimen,

and subtracting the resulting x-ray spectrum from that recorded while the beam is incident on a P
fracture in the specimen. The main disadvantage of this latter procedure is the large increase in

the background noise level.

For uniform thin foils, correction equations have been developed to allow for the effects of

selective absorption [3]. In such cases fluorescence is usually of less importance. However, in

multiphase materials the situation is not so simple. In general the x-ray path through each phase

is not known, and the distance from the x-ray source to different fluorescing phases is not well

characterized. Thus accurate correction of the x-ray intensities is not usually possible in such

cases. To some extent the problems may be overcome by the use of extraction replica techniques, but

these cannot be employed in cases where it is required to measure the compositions of, say, precipitates and matrix in immediately adjoining regions. In such instances it is of some assistance to

use specimen geometries in which the extent of interaction between the various phases is minimized.

Applications to Rapidly Solidified Fe-Based Alloys

Research at MIT has centered on microstructures and properties produced in rapidly solidified

steels. Of particular importance in this research has been the construction of composition profiles

across solidification-related microstructural features. Where grain boundaries or cell walls have

been analyzed, the analysis was performed by orienting the grain boundaries parallel to the electron

beam and then manually stepping the electron probe along a line perpendicular to the boundary.

After each point has been analyzed, the boundary was examined to ensure that no specimen drift had

occurred before repositioning the probe at the next data point. Whenever possible, boundaries were

selected that were perpendicular to the edge of the foil so that the step scan occurred in a region

of uniform thickness. Additionally, the line of the step scan was oriented perpendicular to the

line between the specimen and the x-ray detector.

303 Stainless Steel

A notable finding in the course of investigation into the microstructures of a rapidly solidified

austenitic 303 stainless steel is the presence of a distinct ferromagnetic component in the centrifugally atomized powders. It became evident that the ferromagnetic particles are the smallest of the

powder particle population. X-ray-diffraction analysis of the magnetically separated particles

shows that the nonferromagnetic powder is 100 percent fee. The ferromagnetic powder is predominately

bcc but has as much as 30 volume percent fee. Light microscopy of metal lographic sections of the

ferromagnetic powder using a "dark etch" on the bcc phase indicates that the smallest ferromagnetic

particles are 100 percent bcc. Some of the larger ferromagnetic particles have both fee and bcc

grains present.

The microstructure of the fee particles if found, in all cases, to be that of a cellular or dendritic solidification. The microstructures of the bcc particles, however, are found to be of two

types; cellular and noncellular. As described below, important differences exist between the fee

and bcc structures in regard to solute segregation patterns and resultant precipitate morphologies.

Rapidly Solidified Cellular FCC

The microstructure of a cellular fee powder particle is shown in figure 1 with a corresponding

composition profile. Precipitate particles 100 to 200 nm diameter reside at the cell walls of this

structure and delineate the cellular geometry. These cell-wall precipitates are observed to be

about the same size in both cellular fee and cellular bcc rapidly solidified structures. This

cell-wall type of precipitate will be referred to as the 100 nm precipitate. X-ray fluorescence

spectra indicate that these precipitate particles are nominally manganese sulfide.

With reference to figure 1, note that there is an increase in nickel content and corresponding

decrease in iron content associated with the cell walls in the microstructure. The chromium content

of these cells appears to be uniform though there is indication of a relatively slight increase at

the cell wall. Other profiles in these fee cellular structures have shown slight chromium enrichment,

coincident with more pronounced nickel enrichment at cell walls.

A notable aspect of the manganese variation along the profile is a decrease in concentration in

the vicinity of the cell wall. Since manganese must be segregating to the cell walls to form the

precipitates, this observed decrease in manganese must be due to manganese redistribution along the

cell wall during formation of the precipitates in the late stages of solidification.

Changes in silicon concentration within the structure appear counter-variant with those of

manganese. This may be evidence of a coupling of the partitioning behavior of manganese and silicon

during solidification. Similarly, changes in sulfur concentration in the structure appear covariant

with manganese.

57
COo

0.2

o

o

72

716967

18

16

13

I I

9

7

2.0

1.5

1.0

0.5 03

1

\

1—~—1 r

:

' \

\

"1 1 1 1 —1 z , *Fe

/

/

: /

:

\ V. -

- ..-

Si

*“—• •. ^

1 1 i 1

;

' ; 1

^ DISTANCE

2 10

FROM BOUNDARY

12( /ym)

Figure 1. Rapidly solidified cellular fee structure of 303

stainless steel. Montage of two scanning transmission electron microscopy annular dark-field

images and composition profile.

Note that compositional redistribution in the solid in this rapidly solidified structure is cal- culated to be on the scale of 8 nm or less, for example, for chromium and nickel using recently

measured diffusion data for stainless steels [4,5], It appears that compositional redistribution

in the solid does not significantly alter the observed composition profiles in these rapidly solidified powders. Though these variations are not extreme, this cellular fee structure of the rapidly

solidified powder does exhibit some partitioning of both the metal and metalloid solute species.

Rapidly Solidified Cellular BCC

A view of the cellular bcc structure is shown in figure 2. Cells appear delineated by sulfide

precipitates of 100 to 200 nm diameter, similar to the rapidly solidified cellular fee structure with

the 100 nm precipitates at the cell walls. The elongated cells in figure 2 are apparently longitudinal sections of cells. A notable feature of this structure is the presence of the small precipitates, 10 to 20 nm diameter, within the cell. These 10 nm precipitates are observed in the bcc

58
DISTANCE ifjm)

Figure 2. Rapidly solidified cellular bcc structure of 303

stainless steel. Scanning transmission electron

microscopy bright-field image and composition profile.

structures only. There appears to be a decrease in the volume fraction of 10 nm precipitates at the

cell walls commensurate with the appearance of the 10 nm precipitates in the cell interiors. In situ

x-ray fluorescence analysis of the 10 nm precipitates indicates that they are nominally manganese

sul fide.

The location of these 10 nm precipitates within the cells corresponds with the position in the

composition profile at which the concentrations of manganese and sulfur go through large variations.

These peaks in concentration are about 0.2 micron in spatial extent and correlate with localized

maxima in the number density of the 10 nm precipitates in the central region of the cell. It should be noted here that care has been taken to distinguish the observed increases in manganese and

sulfur concentrations as due to local increases in the matrix concentration and not simply due to the presence of a precipitate in the volume of material analyzed at a given point. The composition

profile of this rapidly solidified cellular bcc indicates that the structure is homogeneous with

respect to iron, chromium, nickel, and silicon.

59
Rapidly Solidified Noncell ular BCC

The term "nonce! 1 ular" is used here to mean that solidification occurred without the formation ,|

of a cellular liquid/solid interface morphology. Indications are that a smooth-front growth of the

solid took place without development of solute partitioning; hence, this solidification mode is termed massive solidification.

jjjj The microstructure of a noncellular bcc particle is shown in figure 3. There is no suggestion U

of a cellular solidification structure, especially as would be delineated by precipitates. The I matrix composition of this noncellular bcc structure is characterized by a lack of microsegregation. j|* The composition profiles of figure 3 indicates virtually complete homogeneity over a distance of

| nearly 5 microns. The most significant variation in composition observed here is that associated

with manganese. As pointed out earlier in connection with the cellular bcc structure, there appears

a periodic variation in manganese concentration along the profile. In the noncellular bcc structure, ji however, the period of the variation is about 2 microns as compared with about 0.2 micron in the

-J case of the cellular bcc. i|

0^

%

O

CL

o

CJ

DISTANCE (ym)

If

l;|

ifa

Figure 3. Rapidly solidified noncellular bcc structure of 303

stainless steel. Scanning transmission electron microscopy bright-field image and compositional profile.

60
A predominant feature of some bcc noncellular particles is a large number density of small

precipitates (10 to 20 nm diameter) throughout the entire microstructure. This feature is readily

seen in the image of figure 3. The morphology of these 10 nm precipitates is similar to that noted

in the intracellular region of the cellular bcc structure (fig. 2). Not all noncellular bcc particles

are found to have the 10 nm precipitates. This latter microstructure is the result of retention of

the solid solution that was produced during massive solidification. This precipitate-free noncellular

bcc microstructure was only observed in a sample of very small (20 ± 3 micron diameter) powder

particles. Since the cooling rate in the processing of these powders increases with decreasing

particle size, it appears that the supersaturated solid can be retained to room temperature as a result of greater cooling rates from solidification temperatures.

High Phosphorus Steel

A Fe-16Cr-10Ni-0. 4P steel also produced by centrifugal atomization has been under investigation.

As in the 303 stainless steel, there exists in these powders a fee and a bcc fraction with the

smaller powder particles being bcc. The range of solidification microstructures encountered can be

seen in figure 4. The intercellular phase observed in figure 4 is glassy [6]. Of particular interest

are the two powder particles which do not exhibit a dendritic structure. These observations exhibit

the same trends as seen in 303 stainless steel suggesting that, at the smallest (< 20 pm) powder

particle sizes massive solidification may be at play.

Figure 4. Rapidly solidified high phosphorus austenitic steel

powders exhibiting three different solidification

microstructures. All of the powders are 'v- 45 pm in

di ameter.

A1 Alloys

There are several potentially beneficial modifications to microstructure which may be effected

by rapid solidification processing (RSP). One of these is the possibility of metastable phase

formation. Two studies are currently underway at Illinois involving metastability in A1 alloys,

the first of which is directly aimed at developing an understanding of metastable phase formation in

Al-Ge alloys during rapid solidification. The second is part of a program attempting to improve the

elevated temperature properties of A1 alloys by developing a very fine scale dispersion of intermetallic compounds using RSP. In this study, metastable phases form as a result of rapid solidification,

and an example of this will be given.

In the first set of experiments, alloys from the Al-Ge system are used, because there is in the

literature a number of reports of metastable phase formation resulting from the application of

various means of rapid solidification [7-11]. In the present study, a new technique for the direct

i n situ observations of rapid solidification using transmission electron microscopy (TEM) has been

61
(ipv/ploDed where the incident electron beam is used as a local heat source to cause melting of

indiiiduai alloy powders [12], The powders are produced by electrohydrodynamic atomization (the EHD o?oJessr and wer^M^^^ by Phrasor Scientific Inc. and the National Bureau of Standards. J

,

As has been reported previously, such powders when supported on a thin carbon film may be i ntroduced , readily into a^TEM, and are often sufficiently thin for microstructural analysis [p]. When using f I

the in situ method for melting of powders, a range of cooling rates may be achieved; the slowest is produced by slowly defocusing the condenser lens. Intermediate rates of cooling may be effected by

deflecting the electron beam using electromagnetic coils, and the fastest rate of cooling is produced

by inserting an opague stop between the electron gun and sample.

For the case of Al-Ge powders of eutectic composition, the microstructure following iji situ , rapid solidification has been shown to vary widely depending on cooling rate [14]. For the slowest [ rate a metastable phase forms, the structure of which has been determined as hexagonal (space group | P 6 /mmm) with a = 1.42 nm and c = 0.74 nm [15]. At intermediate cooling rates, a refined distribution I of randomly oriented crystals of a”Al and p"Ge (solid solutions) was observed. These phases are those which are predicted by the equilibrium phase diagram. At the most rapid rate of cooling, a i microstructure is produced which consists of single crystal o(“Al , in either cellular or dendritic

form, surrounded by a Ge~rich amorphous phase. This microstructure is shown in figure 5, for the

j case of the cellular A1 , and figure 6 for the dendritic morphology. Included in both figures are I selected area diffraction patterns recorded, in each case, with the whole powder being irradiated;

j the single crystal nature of the crystalline phase is evident. Microdiffraction patterns of the

^ i nterdendri tic phase and carbon film are shown in figure 6 , and the former pattern is consistent

with the amorphous nature of the Ge-rich phase. Energy dispersive x-ray spectra show that this

amorphous phase is considerably enriched in Ge (~ 50 at.%). i

The mechanism of formation of the mixed crystal 1 i ne/amorphous structure has not been established.

However, it is possible to hypothesize that when the beam is occluded from the specimen, undercooling

of the molten droplet occurs to a temperature where a-Al nucleates. The a-Al phase will grow in size,

rejecting Ge-rich solute ahead of the advancing interface, until the liquid composition and temperature are such that a glassy metal forms. The glass transition temperature for hypereutectic compositions in this alloy system is, over a certain range, expected to increase with increasing Ge concen- ^ tration because of the difficulty in nucleating p-Ge. i

An interesting aspect of the present observations is that the metastable hexagonal structure

forms at a lower level of undercooling than that required to produce a mixture of a-Al and p-Ge

phases as predicted by the equilibrium phase diagram. This is the subject of further study. )

I

rr>

In the second set of experiments, metastable phases have been observed in Al-Fe alloys following

laser surface melting, and Al-Mo alloys which has been processed into ribbon using melt spinning. « ; In this latter case, an example will be given for melt spinning of Al-1.75 wt.% Mo. The microstruc- gJj ture of this material is shown in figure 7. A number of precipitates may be seen, the size of which Sv|

is typically ~ 0.1 pm. One of these precipitates is shown at a higher magnification in figure 8 , ? | and it can be seen that the phase is somewhat faulted. Selected area diffraction patterns were *

recorded with the incident electron beam parallel to a number of low index zone axes, but it was not*

possible to interpret these data on the basis of any of the reported phases in the Al-Mo system [16]. j

In an attempt to identify the second phase shown in figures 7 and 8 , the techniques of energy .w|;i dispersive x-ray analysis and convergent beam electron diffraction were used. The chemical composi-ch

tion of the phase was determined in this way and corresponds to AI 5 M0 ; for this stoichiometry, the ilj

previously reported crystal structure is hexagonal (space group P 63 ), and a = 4.89 A and c = 8.80 A. Convergent beam electron diffraction patterns with the electron beam parallel to the <100> and <111>B'

zone axes are given in figure 9. These various patterns may be interpreted using the tables in M.

reference [17]. Thus the projection diffraction and whole pattern symmetries for the <100> zone S||

axis patterns are both 4 mm, so that this diffraction group may be deduced as being either 4 mm or

4 mml|^. In the case of patterns produced with the beam parallel to the <111> zone axis, the projec-Bj

tion diffraction symmetry is 6 mm, and the bright field and whole pattern symmetries are both 3 m; the deduced diffraction group is G^mm^. Consideration of these possible diffraction groups leads to the conclusion that the point group of the precipitated phase is m3m. The space group may be deter- *! mined by projecting the reflections contained in the first and second order Laue zones onto the zero jr,

order zone of the <100> pattern. When this is done, a face centered reciprocal lattice results,

which corresponds to a body centered real lattice cell. There are two space groups that possess the

m3m point group and exhibit a body centered cell, namely Im3m and Ia3d. The intensities recorded inMthe convergent beam patterns is only consistent with the Im3m space group. The lattice parameter K:

o '

has been determined from the diffraction patterns, and is given by a = 5.39 ± 0.01 A. The occurrence*

of this phase, AI 5 M0 , possessing the space group Im3m, has not been reported previously in the 8

literature and is presumably a metastable phase formed on rapid quenching of the alloy. I
50 nm

Figure 5. Two phase (cellular, single crystal

interior plus amorphous shell) structure with

corresponding SAD pattern produced by interrupting the beam from a molten droplet with the

condenser aperture mechanism.

c

SAD

Figure 6. Larger two phase (dendritic, single

crystal interior plus amorphous shell)

with corresponding SAD pattern and microdiffraction patterns from the shell and

the carbon film. Structure was produced

by interrupting the beam from a molten

droplet with the condenser aperture

mechani sm.

63
Figure 7. TEM micrograph showing the precipitates and associated dislocation structure in Al-1.75 wt.% Mo ribbon. Beam direction

close to <101>.

50 nm

Figure 8. TEM micrograph of precipitate in Al-1.75 wt.% Mo

ribbon. This precipitate was used in the convergent beam diffraction studies. The electron

probe was focused in the unfaulted region in the

upper left corner.

64
Figure 9. Convergent beam electron diffraction patterns from the precipitate shown in

figure 5. (a) <100> zone axis, zero order Laue zone projection, (b) <100>

zone axis, whole pattern, (c) < 111> zone axis, bright field and zero order

Laue zone projection, (d) <111> whole pattern. Patterns recorded on a Philips EM400T operating at 120 kV.

Ti Alloys

Research in the area of RSP of Ti alloys has generally concentrated on the possibility of producing ultrafine dispersions of rare-earth oxides in an a-Ti matrix (e.g. [18]). Such a study is

currently underway at Illinois where it is intended to produce rapidly solidified powders of Ti and

Ti-8A1 (at.%) containing refined dispersions of Y2 O 3 , NdgOg, and ErgOs. Initially, the response of

the alloys Ti and Ti-8A1 with additions of elemental Y, Nd, and Er to rapid solidification is being

established. In this case, laser surface melting is being used as a convenient and efficient method

of obtaining small volumes of rapidly solidified material. An example of this study is given for

the case of Y containing alloys.

Figure 10 shows a backscattered electron image obtained in the scanning electron microscope of

the interface between as-arc melted and laser surface melted Ti-8A1-4Y (at.%). The as-arc melted

alloy consists of a relatively coarse dendritic microstructure, whereas a substantial refinement in

this microstructure has been effected in the laser affected region. A comparison between the microstructures of as-arc melted and laser surface melted materials is given in figure 11. The as-arc

melted alloy is seen to consist of a dispersion of second phase within an a-Ti matrix, where the

average size of the dispersoids is 0.1 pm. Following rapid solidification, the matrix consists of

laths of a' martensite, and there is a refined dispersion of second phase present, the average size

G5
Figure 10. The microstructure of an arc-melted Ti-8A1-4Y specimen which has

been subjected to laser surface melting. Backscattered electron

image recorded in scanning electron microscope.

of these precipitates being 120 A in diameter. It is necessary to determine the nature of these | particles, since in preparing the alloy, elemental additions of Y (nominal metallic purity of 99.95%)_

have been made. Such characterization has been performed, and the identity of these second phase I

dispersoids has been established [17]. f

O i

The identity of all particles with diameters > 400 A in both the as-arc melted and laser surface”'

melted samples has been determined. All these particles were found to have identical crystal structures; the nature of the smaller precipitates in the rapidly solidified samples has not been uniquely

O determined. For the precipitates with sizes > 400 A, selected area diffraction patterns corresponding i to a number of low index zone axes were recorded. It was not possible to interpret these patterns

in terms of one crystal structure, and so convergent beam patterns were recorded with the electron ! beam parallel to the <100> and <111> zone axes. These are shown in figure 12. The projection

diffraction, bright field and whole pattern symmetries for the <100> patterns are all 2mm. In this

case, the deduced diffraction group is either 2mm or 2mml|^. For the <111> patterns, the projection

diffraction symmetry is 6 , and the bright field and whole pattern symmetries are both 3; the deduced j diffraction group in this case is 6 ^. The only point group containing these deduced diffraction

groups is m3. By projecting the reflections contained in the higher order Laue zones onto the zero order zone, it is possible to deduce that the reciprocal lattice cell is face-centered so that the

unit cell of th real lattice is body-centered cubic. The space group can be determined by combining

the determinations of point group and cell type with information concerning kinematically forbidden

reflections. In this case, 011 reflections are absent from <001> patterns, and so it is possible to assign the space group Ia3. The lattice parameter may also be determined from the diffraction

O patterns and is given by a = 10.6 ± 0.05 A. Since previous work [20] has determined that Y2 O 3 has

the same space group and lattice parameter as that given above, it is reasonable to assume that the

second phases observed in the present study are, in fact, YgOg.

6G
a

III

b

Figure 11. Transmission electron micrographs of Ti-8A1-4Y alloy,

(a) as-arc melted, (b) laser surface melted.

67
Figure 12. Convergent beam electron diffraction patterns from the second phase precipitates found in

Ti-8A1-4Y alloy, (a) <100> zone axis, bright field and zero order Laue zone projection

(b) <100> zone axis, whole pattern, (c) <111> zone axis, bright field and zero order

Laue zone projection, (d) <111> zone axis, whole pattern. Patterns recorded on a Philips'

EM400T operating at 120 kV.

Summary

Examples of microstructural characterization of Fe-, A1-, and Ti-based alloys following rapid

solidification processing have been presented. Since the products of such processing are often

extremely fine, it is necessary to make use of the recently developed techniques of analytical

transmission electron microscopy. This is absolutely necessary, not only to aid in the understanding

of the mechanism of microstructural formation, but also in developing microstructure/property relationi

ships.

Acknowledgments

The MIT research was conducted by Dr. Thomas F. Kelly and Mr. Matthew Libera under ONR sponsorship (N00014-81-K-0013). Extensive use was made of central facilities provided through the Center

for Materials Science and Engineering (NSF-MRL).

u8
The Illinois research was conducted by Michael J. Kaufman, Fritz C. Grensing, and

Douglas G. Konitzer. The work on Al-Ge and the Ti alloys was sponsored by the U.S. Department of

Energy, Division of Materials Sciences, under Contract # DE AC02 76ER01198. The studies involving

Al-Mo were sponsored by the U.S. AFOSR, under Contract # AFOSR 82-0186, with Dr. Alan H. Rosenstein

as Technical Monitor. The electron microscopy studies were performed using the facilities of the

Center for the Microanalysis of Materials in the Materials Research Laboratory of the University of

111 i nois.

References

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2. G. Cliff and G. W. Lorimer, J. Microsc . 103 (1975) p. 203.

3. J. I. Goldstein, "Introduction to Analytical Electron Microscopy," (eds. J. J. Hren,

J. I. Goldstein, and D. C. Joy) Plenum Press, (1979).

4. R. A. Perkins, Metal 1 . Trans . 4 (1973) p. 1665.

5. R. A. Perkin, R. A. Podgett, and N. K. Tunala, Metal 1 . Trans . 4 (1973) p. 2535.

6. T. F. Kelly, G. B. Olson, and J. B. Vander Sande, "Rapidly Solidified Amorphous and Crystalline

Alloys," (eds. B. M. Kear, B. C. Giessen, M. Cohen), Elsevier Science Publ . Co. (1982).

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8. M. Laridjani and R. W. Cahn, Mat . Sci . Eng . 23 (1976) p. 125.

9. V. Kdster, Z. Metal Ik . 63 (1972) p. 472.

10. P. Ramachandrarao, M. G. Scott, and G. A. Chadwick, Phi 1 . Mag . 5 (1972) p. 961.

11. C. Suryanarayana and T. R. Anantharaman , J. Mater . Sci . 5 (1970) p. 992.

12. M. J. Kaufman and H. L. Fraser, Scripta Met . 17 (1983) p. 141.

13. C. G. Levi and R. Mehrabian, Metal 1 . Trans . 13A (1982) p. 13.

14. M. J. Kaufman and H. L. Fraser, Metal 1 . Trans . (1983) in press.

15. M. J. Kaufman and H. L. Fraser, Matls . Sci . and Eng . 57 (1983) p. 117.

16. Bulletin of Alloy Phase Diagrams, 1 (1980) p. 71.

17. B. F. Buxton, J. A. Eades, J. W. Steeds, and G. M. Rackham, Phil. Trans. Roy. Soc. A281 (1976)

p. 171.

18. K. K. Sankaran, S. M. L. Sastry, and P. S. Pao, Metal 1 . Trans . IIA (1980) p. 196.

19. D. G. Konitzer, B. C. Muddle, and H. L. Fraser, Metal 1 . Trans . (1983) in press.

20. M. G. Paton and E. N. Maslen, Acta Cryst . 19 (1965) p. 307.

69
ASPECTS OF SOLIDIFICATION IN THE RSR POWDER PROCESS

R. J. Patterson, II and C. M. Adam*

Pratt & Whitney Aircraft

P. 0. Box 2691

West Palm Beach, FI. 33402

ABSTRACT

Early work in rapid solidification suggested that the production of new structures,

metastable phases, and in particular the amorphous state, is a strong function of quench

rate. More recent work suggests that in addition to quench rate, nucleation plays a substantial role. We report here work related to nucleation in the glassy alloy

NiysBiySis and also a study of undercooling in rapidly solidified alloy steel powders.

The characterization of the glassy powders consisted of differential scanning

calorimetry (DSC) on precisely sized fractions from a number of powder runs where both

quench gas and other process parameters were varied. The results suggest that any of a number of nucleation modes may occur, including heterogeneous volume nucleation and

homogeneous volume nucleation. The possibility of other modes is also raised.

The results on the alloy steel powders were obtained by x-ray diffraction (XRD) and

transmission electron microscopy (TEM). Both the body centered cubic phase ferrite and

the face centered cubic phase austenite were observed. Preliminary evidence indicates

that the ferrite phase results from solidification of liquid droplets supercooled by a few hundred degrees, producing carbon supersaturation approaching 0.5 wt%, while the

austenitic powder results from the solidification of droplets which have experienced

lower levels of supercooling.

component liquid metals was recognized some three decades ago by Duwez [1]. Several conferences have

since provided increasing metallurgical data about the solidification behavior, phase stability, and

phase transformations of alloys solidified at cooling rates above lO^K/sec [2,3]. There exist several

One such device, developed by the Government Products Division of Pratt & Whitney Aircraft, has been : used for the past six years to produce quantities of powder quenched from the liquid at cooling rates j from about 10^ to 10®K/sec [4]. This device consists of a rapidly rotating disk onto which is directed!

a liquid metal stream, producing a normal distribution of liquid metal droplets ranging in size up to ; about 200 pm in diameter with a mean particle size of up to about 90 pm. Inert gas is used to provide

forced convective cooling at rates between 10^ and lO^K/sec, with the quench rate depending on the type

of gas used and the particle size produced.

The bulk of work published during the first decade or so of rapid solidification implied that the

controlling factor in the production of new phases, structures, and the amorphous state was the rate

of quenching of the melt. Early work at Pratt & Whitney on rapidly solidified powders suggested that

the situation must be more complex: Holiday et ai. [4] reported on two structures found in nickel-base j superalloy powders, a conventional dendritic structure and a previously unreported "microcrystalline" > structure. STEM analysis by Field and Fraser [5] suggested that the latter structure consisted of

discrete crystallites exhibiting lower levels of segregation than found in the former structure. The

former authors observed that both structures were found in different particles of any given particle

size fraction, though the occurrence of the microcrystalline form was found to depend on both alloy

chemistry and particle size, increasing rapidly as particle size decreased. ‘

*Now with Allied Corporation, Corporate Technology, Morristown, N.J. 07960.

INTRODUCTION

The possibility of producing new alloys with novel properties by rapidly solidifying multilaboratory scale devices which produce liquid cooling rates greater than lO^K/sec in bulk quantities.

70
This work plus unpublished data on powders produced from 99.97+ v^t% nickel led us [6] to conclude

that in addition to quench rate, the effects of nucleation must be considered. In any given size

fraction of the nickel powder, structures ranging from a very fine microcrystalline to a very coarse

dendritic structure were observed, as shown in Figure 1. This structural difference seemed difficult

to explain on the basis of quench rate alone since the quench rate of this powder process should be

quite uniform for any given particle size. The heat transfer work of Glickstein et ai. [7] also

suggested that recalescence should occur during nucleation and growth due to the release and conduction

of the latent heat of fusion; their model indicated that internal conduction should dominate over the

rate of heat removal from the droplet surface during solidification. These factors suggested that the occurrence of nucleation at varying levels of supercooling might give rise to the variety of microstructures which had been observed.

Figure 1. Scanning electron

micrograph of the -230+270 mesh

size fraction of RSR processed

99.97+ weight percent nickel

powder.

It thus seemed desirable to investigate the effect of process parameters on supercooling and

nucleation. Unfortunately, for many alloys this would imply the metal lographic examination of large

numbers of particles from each of many size fractions, in order to insure that conclusions drawn would

be statistically valid. Systems were sought for which other criteria might be used in order to avoid

this laborious effort; Two, studied here, seemed reasonable. First, in the case of metallic glasses,

the amorphous state would only be reached if nucleation and growth of the crystalline phases did not

occur, and one could then evaluate large numbers of particles simultaneously by DSC to determine the

variation of the heat released during crystallization of the glassy particles as a function of powder

particle size and of process parameters; earlier results have been presented elsewhere [8]. Another

system which which might be useful would be one in which a series of crystalline phases could form on solidification at various levels of supercooling, some of which might be magnetic. For such a system,

magnetic separation of each particle size fraction, coupled with x-ray diffraction, could then shed

light on the significance of alloy chemistry and operational parameters in achieving the various levels

of supercooling and resultant microstructures.

Ni The work reported here is thus separated into two sectionsiOne on the glass forming alloy

(atom percent) and the other on a series of alloy steels containing C,Cr,Mo, and V. In the

former case, the powders were precision sieved into 5 ym or 10 ym size fractions, which were then

subjected to DSC. The results suggest that a number of nucleation modes can occur, sometimes in

different size ranges of the same powder lot, and that run to run results can vary widely. For the

steels, the structures described resulted from a wider systematic study of the solidification behavior

of a wide compositional range of alloy steels [9]. Normal metal 1 ographic techniques suggested that in

all cases the crystalline solidification structure observed was cellular-dendritic, as shown in Figure

2, with dendrite arm spacings less than one micrometer indicating liquid quench rates on the order of

lO^K/sec or slightly higher. Subsequent XRD studies on these steels showed that in many cases the

powders had solidified as the body centered cubic phase ferrite (Figure 3a), or alternatively as the

face centered cubic phase austenite (Figure 3b), without any evidence of the metastable martensite

phase, and that the alloy preference for ferritic or austenitic solidification did not appear to be

easily related to chemistry. Also, some alloy compositions (e.g. Fe-1 .2C-19Cr-2Mo-lV and Fe-1.15C71
14.7Cr-4Mo-lV, wt%) showed almost equal quantities of austenitic and ferritic phase powder (Figure 3c).

solidification structure

Light etching powder

darker areas are ferritic.

Figure 2. Cellular-dendritic

of 14-19 wt% Cr alloy steels,

particles are austenitic; the

Since the austenitic powder is non-magnetic it was possible to separate the ferrite phase powder

particles magnetically and examine the distributions of the two phases independently. The non- magnetic austenite phase was found to be confined to the coarser size fractions, above about 80 ym diameter, while the magnetic ferrite phase was confined to particle diameters finer than 50 ym.

A further test of correlation between particle size and austenite/ferrite phase separation was obtained

by selectively removing the -25 ym fraction by sieving. This powder was found to be entirely ferritic, indicating that a strong correlation existed between particle size and phase distribution. This was later confirmed independently for a commercial 303 stainless steel (Fe-17Cr-8Ni-l .6Mn, wt%) converted

to powder by the P&WA process [10]. The structures found in both phases, shown in Figure 2 above, indicate direct solidification from the liquid rather than solid state transformation.

AMORPHOUS POWDERS

nr.v.+Ji'

Of’ amorphous powders [8,11] preliminary evidence has suggested that individual

^hn.it^inn

9^3ssy or crystalline. In our work, limited microscopy on particle diameters to about 100 ym has shown no particlesexhibiting both crystalline and glassy regions. For processes and this occurs, the size distribution of the amorphous particles can be determined by

!!

^ heat released upon crystallization (by DSC) as a function of powder particle size. Since

thP

Pt^oduce a fraction which has been determined to be totally amorphous by x-ray diffraction, the results we have obtained are relative to the most energetic fraction observed to date.

Arguments have been raised that the presence of crystalline particles

necessary for glass formation does not exist. However, since glassy parti size fractions, the quench rate would seem to be satisfactory and the more

events transpired which prevented the crystalline particles from reaching and subsequent growth due to the presence of motes or nucleation sites wou explanation than the presumption that the quench rate varies in some compl put not all) p

the smaller particles and some (but not most) of the coar Perepezp and Anderson [12] suggest that if heterogeneous volume nucleatio mode, then the probability x that a droplet of volume v will NOT contain a produced from a melt containing m sites per unit volume is given by a Pois

implies that the quench rate

cles can be found in the same pertinent question is: What

the glassy state. Nucleation

Id seem to be a more plausible

ex fashion, allowing most

ser particles to be glassy,

n is the dominant nucleation

nucleation site when it is son distribution of the form:

72
Figure

(a) Fe-

(a)

(b)

i X-ray diffraction patterns of ferritic and austenitic alloy steel powders,

L!l5C-14.7Cr-6Mo-lV; (b) Fe-1 .15C-14.7Cr-2Mo-lV; (c) Fe-1 . 15C-14 .7Cr-4Mo-lV

73
X = exp(-Mv) ( 1 )

This form is not explicit, since m is seldom known and must be back-calculated to fit the expression to the experimental data. A more useful form is obtained, following Drehman and Turnbull [11],

by converting the volume to a particle diameter and by realizing that at some point the product Mv

must be unity, at which point the expression (1) takes on the value X = 0.3679. We may thus normalize

the equation through the use of a particle diameter , which will be that diameter at which X = 0.36 1

X = exp{-(j )^} (2)

°o I

Unfortunately, the use of graphical analysis in Cartesian coordinates is also cumbersome. First, r Drehman and Turnbull point out that there may be other nucleation modes; they suggest that for hetero- I geneous surface nucleation the exponent on d should be 2.0, while for homogeneous volume nucleation J

they expect the exponent to be on the order of 4.6. The latter exponent results from considering that I

the time available for nucleation is inversely proportional to quench rate, which in turn depends I

roughly on particle diameter to the -1.6 power. 1

Further difficulties may also confound the analysis. First, one must expect a degree of uncer- tainty in the experimental data so that exact agreement between such data and expressions such as (2)

with exponents as suggested by Drehman and Turnbull is unlikely. Second, only three nucleation modes

^ i have been mentioned but others are not outside the realm of possibility. For example, the heat transfe a work of Glickstein et al. [7] implies that thermal gradients should exist within rapidly quenched ! droplets, resulting in a droplet surface which is colder than the interior, and which will thus also

begin to supercool sooner than the interior. This suggests that homogeneous surface nucleation may

also be possible. Third, beyond the normal uncertainties encountered with experimental data, "outliers I may be encountered which would be difficult to detect when fitting an exponential function in Cartesian k coordinates. Fourth, if a variety of motes having different catalytic activity exist in the melt, then I the mode of nucleation encountered in the data might vary with particle size fraction. i

To circumvent these difficulties, it is preferable to use a general expression in which the 1 exponent is to be determined from the data; I

X = exp{-(^ )y> (3) I “o

If we then take successive logarithms we arrive at;

'

ln(-ln X) = y ln(d) - y In(d^) (4)

Since this expression is linear

not only the best exponent, y, the line of best fit. This sti

different nucleation modes in d ical analysis; Data fitting the

log paper. Unfortunately, the

data when x is large. Thus, a relative energy of crystal lizat

the point. If other nearby poi

y will also change.

in In X and d, linear regression analysis may be performed to determine

for the data, but also the correlation coefficient, r, of the data to

11 does not solve the problem of faulty data points or the presence of

ifferent size fractions, but these may be readily determined by graphexpression (4) will be linear when -In x is plotted against d on log- ' price for this convenience is that this form puts heavy emphasis on the’:

small (1% or so) change in a large value of x (which we recall is the j ion determined by DSC) can cause a substantial shift in the location of

nts also shift (perhaps data are being remeasured), the slope (exponent)

Experimental Results

Data were obtained using a depending on absolute energy re using an ATM Model L3P sonic si were found to deviate from the

or as large as 10 ym in range.

DuPont 1090 thermal analysis system on small samples (3 to 15 mg, I_

lease) of powder which had been precisely sized into 5 or 10 ym fractions

fter equipped with precision el ectroformed sieves. A few of the sieves

stated size by a few microns, so a few fractions were as small as 3 ym

Powder was obtained from remelting one of three masterheats which had

74
1 been produced by vacuum induction melting lots of from 35 to 40 Kg, Each lot was then poured to form

i ingots weighing about 12Kg. The remelting of each ingot was done under vacuum in the RSR device; upon

j melting, the chamber was backfilled with the gas to be used for quenching and the melt temperature was

ji then increased to provide the necessary superheat. Argon was used as the quench gas for two of the

runs, while helium was used for the other three. The data obtained from DSC are presented graphically

j in Figures 4 (argon quench) and 5. Values of the slopes of the various segments, determined by linear

I regression, are presented in Table I along with the correlation coefficient determined for the data

I of each segment:

TABLE I Figure SI ope Correlation

Number y Coefficient, R 4A 1.01 0.9987

2.45 0.9990

4B 1.50 0.9973

2.75 0.9994

4,31 0.9999

5A 2.36 0.9995

5B 0.617 0.9972

0.992 0.9967

5C 2.94 0.9956

Two of the exponents can be correlated with the work of Drehman and Turnbull. The exponent for

Figure 5C is sufficiently close to 3.0 to suggest that in this instance, heterogeneous volume nucleation

has occurred. The exponent for the third segment of Figure 4B, at 4.31, is reasonably close to the

analytically determined 4.6 and sufficiently far above the 3.0 expected for heterogeneous volume

nucleation that homogeneous volume nucleation is likely the origin of the crystalline particles. At

2.75 the exponent of the second segment of Figure 4B may be sufficiently close to 3.0 to again

indicate that volumetric heterogeneous nucleation has occurred. However, exponent values between 2.0

and 3.0 may imply an interchange of heterogeneous surface nucleation (y=2.0) and heterogeneous volume

nucleation (y = 3.0); that is, a form of heterogeneous nucleation which is not site specific. This

may be the case for the values of 2.75 of Figure 4B and 2.36 of Figure 5A; although the latter value

may, with the uncertainties present in the data, be sufficiently close to 2.0 to indicate the heterogeneous surface nucleation observed by Drehman and Turnbull [11]. The latter situation may also apply

to the slope of 2.45 of Figure 4A.

Then we have the seemingly low values of 1.01, 1.50, 0.617 and 0.992 tabulated above for various

segments of Figures 4A,4B and 5B. Whether these represent new modes of nucleation or are the result of

a superposition of several modes remains to be determined, but some speculation may be warranted. For example, the dependence of the film coefficient, which relates the rate of heat removal from the surface

of a droplet, on diameter is of the order of d~°*® while the bulk quench rate dependence is of the order

of d~^-^. Along the lines of a surface to volume ratio, the ratio of film coefficient to quench rate

is on the order of d^-®. The value of 0.617 for a segment of Figure 5B thus suggests a dependence on film coefficient, while the value of 1.50 of a segment of Figure 4B is close to the 1.6 expected for

quench rate dependence.

It would thus appear that there are some operational regimes of the P&WA process where nucleation

may be controlled by macroscopic quench rate, however little correlation to run conditions has yet been

found. This is due in part to the number of operational parameters in the process, and also to the

fact that for us, this is a relatively new alloy; our past experience has indicated that as many as twenty runs may be needed to determine the conditions necessary for proper atomization. Presently,

operating variables have been selected not so much to study their effect on nucleation as to learn of their effect on atomization. In most of these runs we have obtained a log normal particle

size distribution which is not characteristic of proper atomization by this process, and the general

appearance, as observed via closed circuit TV and high speed film, suggests that a multiplicity of

atomization modes are occurring. It may be that this factor is confounding the data presented; surely, more work is needed to distinguish between competitive nucleation modes.

ALLOY STEEL POWDERS

There has been continuing controversy in metallurgical literature over the respective roles of

75
Figure 4. Variation in the relative volume fraction of amorphous powder, X, with particle size for

^ two argon quenched runs of Ni^^B^^^Sig (atom%) , plotted according to Equation 4. Data were obtained by^

differential scanning calorimetry at a scan rate of 50K/min. The width of the data bars represents : the nominal range between the precision sieves used for each size fraction. The slopes of the various , segments of the curves are presented in Table I and discussed in the text. j'

i

i

76
5

.0310 20 30 40 50

PARTICLE DIAMETER, MICROMETERS

60 80 100

Figure 5. Data for three helium quenched RSR powder runs of Ni^^B^^^Sig, obtained and plotted as in

the previous figure.

77
nucleation kinetics and subsequent growth kinetics during the solidification of metals. It has been

|

|j=,j generally considered that bulk liquid metals are incapable of massive supercooling because of the large,':'

number of heterogeneous catalytic nucleants normally present in commercial liquid metals, and that con-J tinuous dendritic growth at low supercooling levels is kinetically easier than repeated renucleation.i

However, the undercooling experiments of Cech [13] on Fe-Ni droplets, of Powell and Hogan on copper y , alloys [14], and more recently of Perepezko and Paik [15] on low melting point alloys has shown that K,|

in the absence of suitable heterogeneous nucleants, bulk ( Ig or more) quantities of liquid metal canMj!

be supercooled to 0.4 T„,, where T,„ is the normal pure metal thermodynamic freezing point. This work B.i

suggests that in general the degree of supercooling achieved before nucleation occurs increases as Jfjj

particle size decreases. Thus in our work it is reasonable to assume that supercooling levels in the Mr,

finer ferritic powders are greater than those achieved in the coarser austenitic powders, using the g!

same arguments as advanced by Perepezko and his co-workers. It is thus instructive to calculate the ^ undercoolings for ferrite and austenite nucleation and to speculate why higher supercooling might lead

to ferritic rather than austenitic solidification in the finer powder size fractions.

The iron-carbon equilibrium diagram can be viewed as a stable-metastable system in which either

6 ferrite or y austenite is stable, depending on the variables of composition and temperature. Cech

[13] showed this situation qualitatively; based on such phase diagrams, one expects solidification of

high-chromium (14-19 wt%) steels containing 1% carbon to be austenitic. The work of Kaufman and

Bernstein on phase diagram calculation allowed Kelly [10] to calculate the relative free energies of

the reactions Liq -* 6 and Liq -> y for a commercial grade 303 stainless steel which is typical of many

of the highly alloyed steels used in our investigation. This work indicated that the relative free

energy changes are very similar and that there is no overwhelming free energy growth advantage for the ; subsequent solidification of either ferrite or austenite. This observation then leads to the conclusiort'

that the presence of ferrite or austenite in rapidly solidified alloy steel powders is controlled by a

j critical nucleation energy, AG*, with some energetic advantage being enjoyed by ferrite at high super-«i

cooling levels and by austenite at lower levels, for high-chromium steels containing 1 wt% C. The

homogeneous nucleation rate, Jy, in an undercooled liquid may be expressed in simple form by the

Turnbull equation:

\ “v 3TIg7^T> ' — > (5)

where is a pre-factor of the order of 10 ^^ cm"^sec~\ a is the interfacial energy between the

nucleus and the undercooled liquid, AG^ is the driving force for nucleation of the product phase, k is]

the Boltzmann constant, and t is the absolute temperature. It is therefore useful to determine the

relative values of a and AG^ for austenite and ferrite and to speculate on how undercooling to different supercooling temperatures must affect the ferrite and austenite nucleation rates. Growth of ferrite

or austenite can then proceed according to the free energy driving forces calculated from KaufmanBernstein data. Since AG^ = AH AT/T^, we can substitute in Equation (5) for AG^ in terms for which

data are available, yielding:

J V

= exp{-

V ^ 16 TT g3 Tm^

3 3 AH^ AT^ kT

V

( 6 )

The nucleation event in a droplet can be observed [16] when the product J^vt = 1, where v is the

droplet volume and t is time of flight in the temperature range over which freezing occurs. The nuc- leation rates in droplets 30 ym and 100 ym in diameter differ by roughly one order of magnitude based

on their relative volumes, while the difference is roughly two orders of magnitude between 20 and 100

micrometer droplets. If we calculate the nucleation rate for a 20 ym ferrite droplet arbitrarily

based on the time spent in a BOOK temperature range below the equilibrium liquidus temperature, using

a nominal liquid cooling rate of lO^K/sec [17], then the nucleation rate for such ferrite droplets is:

J

V J_vt

3 (5x10-^)4tt(10)3x10-12

4.775x10

and

16 IT o3 Tn,^ ^ ln{^^^

3 AH 2 at2 kT

V 4.8x10

11

} = 53.70 (7)

Using the latent heat of fusion, AH^, of iron as 2.1 x lO^J/m^, the surface energy, o, as 0.2J/m2, and

78
V = 1812K and T = 1500K, the undercooling for ferrite nucleation becomes;

! m

ATa = {

16 TT (0.2)3 (1812)2

3(2.1x109)2(1.38x10-23)x1500x53.7

300K ( 8 )

'hoarser droplets cool at somewhat slower rates, e.g. 5xl0^K/sec for a 100 ym diameter particle. Over

|the same temperature range (500K), t == 10-^ sec, and the nucleation rate for austenitic droplets

becomes

J = 1 .6 X 10® cm~3s~^

V

and

16 Ti

3 AH 2AT^kT

V

= ln{-

1035

1.6 X 10®

-} = 61.7 : 9 )

Solving for AT yields a value of 260K. These calculations indicate that ferrite enjoys a nucleation

advantage at undercoolings of about 300K, but the available free energy data for subsequent growth

indicates that this advantage becomes suppressed by the faster growth of austenite at these undercoo mg

levels There is not sufficient precision in either the calculated undercoolings or in the KaufmanBernstein free energy data to be completely accurate in a predictive sense, in determining where austenite or ferrite nucleate in strongly undercooled liquid steel droplets, however the available data

does indicate that a nucleation advantage enjoyed by ferrite overcomes a free energy disadvantage for

subsequent growth. At droplet undercoolings of about 260-300K one would expect to find evidence of

both ferrite and austenite nucleation and subsequent growth, and indeed that is the case for powder

particles of intermediate size, i.e. between 50 and 100 ym diameter. In these powder particles both

ferrite and austenite nucleate, but it is apparent that austenite has a substantial growth advantage.

Carbon Supersaturation

The x-ray diffraction data shown in Figure 3 have been used to calculate accurate lattice parameters for rapidly solidified austenite and ferrite present in high chromium content alloy steels containing up to 1.2 wt% carbon. By using a standard Nelson-Riley extrapolation technique precme lattice

parameters have been determined for a number of steels. The distribution of austenite and ®

as a function of the particle size distribution obtained for one such alloy, Fe-1.15C-14.7Cr-4Mo-lV,

is shown in Figure 6.

150

Particle Size, n

Figure 6. Austenite-ferrite distribution in

Fe-1.15C-14.7Cr-4Mo-lV as a function of particle size.

79
Table II presents x-ray data obtained for two alloys: One (Fe-1 .15C-14.7Cr-4Mo-lV) with fine ferritic

and coarse austenitic powder and the other (Fe-1 .2C-19Cr-2Mo-lV) with a range of austenitic and ferritic powders for which lattice parameters varied systematically with powder size. For both ferritic

and austenitic powder the lattice parameters increased systematically with decreasing powder particle

size. We expect that this data reflects both increased undercooling and subsequently a faster solidification rate for finer powder particles, with a consequent increase in solute supersaturation.

I

TABLE II “i Lattice parameter data for austenite and ferrite phases in 14-19wt% Cr steels

A1 1 oy Particl

Lattice Parameter, nm Chemistry Diameter, ym Austenite Ferrite

Fe-1.15C-14.7Cr-4Mo-lV <45 0.2880

80 to 105 0.3623

Fe-1.2C-19Cr-2Mo-lV >180 0.2867 0.3617

150 to 180 0.2868 0.3617

105 to 150 0.2869 0.3616

75 to 105 0.2870 0.3617

45 to 75 0.2873 0.3617

38 to 45 0.2873 0.3619

<38 0.2872 0.3620

1

Lattice parameters determined for the alloys of Table II are understandably greater than the usual

lattice parameters for ferrite and austenite [18] indicating that the alloy solutes have expanded the

basic crystal lattices. The lattice parameter of austenite is related to alloy chemistry by the following relationship:

a (y) = 3.578 + 0.033%C + 0.0031%Mo + 0.0018%V + 0.0006%Cr

0

( 10 )

in weight percent, from which it is clear that interstitial carbon exercises the greatest effect on lattice parameter expansion. Lattice parameter determinations of the ferritic phase (Table II) yielded

ao(ct) values between 0.2867 and 0.2872 nm (Fe-1 .25C-19Cr-2Mo-lV) and 0.2880 nm (Fe-1 .150-14. 7Cr-4Mo-lV)

Using relationships determined for the lattice parameter of ferrite [19], ao(a) = 2.8664 + 0.0084% C (atom percent), solidification of the ferrite droplets thus resulted in substantial carbon supersaturation, varying from 0.4 to 0.6 wt% for the alloys shown in Table II. This calculated carbon supersaturation is imprecise due to lack of data about the magnitude of the ferrite lattice dilation by theHs

substitutional elements vanadium and molybdenum. Chromium produces an almost negligible change in

lattice parameter of ferrite [ 20 ] (0.005 nm/atom %) and molybdenum and vanadium appear to have an effect similar in magnitude to chromium in ferrite [ 21 ]. H|

The measured ferrite lattice parameters are due to supersaturated solid solutions of carbon in the||

6 ferrite phase field, for which maximum solid solubility is 0.09 wt% at 1811K. There is no readily ij available data relating 6 ferrite lattice parameters to alloy composition, however it is conceivable ’i that the measured values represent an extension of the stable solid solubility limit of carbon in 6 ferrite into the metastable region. Extrapolation of the 5 ferrite solidus line to a temperature some«|

300K below the ferrite melting point produces a carbon solubility of about 0.6 wt% in 6 ferrite. ThisBI

value is close to the carbon supersaturation calculated from the x-ray measurements reported in Table

Equilibrium solid solubility data for the Fe-Cr-C alloys [ 21 ] indicated that the maximum carbon i

solid solubility in austenite in Fe-12Cr alloys is 0.92 wt% at 1500K, and for Fe-17Cr alloys, 0.65 wt%5|i at 1533K. Using the experimental lattice parameters determined for the high-chromium alloys, the con- ’j centration of carbon in austenite was calculated at about 0.8%. Knowledge of the partition of carbon i between complex M 23 C 5 carbides and austenite is unavailable, however TEM showed that the volume fraction

of M 23C

0 present in Fe-1 .150-14. 7Cr-4Mo-lV austenitic powder was that expected from the balance of the

| alloy carbon content, e.g. 0.35 wt% C, producing about six volume percent M

23C

5 in the as-solidified

austenitic powder. The carbon content of the austenite is roughly that predicted by the equilibrium ; diagram with no extreme supersaturation, but there is a general trend apparent in Table II for the

finer size fractions of austenitic powder to contain more solute carbon than the coarser powder. j

80
SUMMARY

The results of preliminary studies of undercooling and nucleation in rapidly solidified powders

have been presented. Calorimetry on a glass-forming alloy has been used to evaluate nucleation. Suggested improvements to the analysis of data coupled with the data obtained suggest that a number of

nucleation modes may occur, including heterogeneous and homogeneous volume nucleation. The work presented on alloy steel powders implies that supercoolings of a few degrees in the coarser particles,

and a few hundred degrees in the finer particles, may be anticipated. In addition, extension of the

solid solubility limit of carbon in ferrite into the metastable region is suggested by the work on the

lattice parameters in the finer ferritic powder, while the work on the coarser austenitic powder suggests that the carbon content is roughly that predicted by the equilibrium diagram.

ACKNOWLEDGMENTS

We are indebted to Pratt & Whitney Aircraft for the support of the work on the metallic glass

powders, and to the Defense Advanced Research Projects Agency, the Materials Laboratory of the Air Force

I Wright Aeronautical Laboratories, and to the Office of Naval Research for their support of various asj pects of the work on alloy steel powders. In particular, we thank A.R. Cox, R.L. Athey and J. B. Moore

of P&WA, Drs. E.C. vanReuth and L.A. Jacobson of DARPA, A.M. Adair of AFWAL/ML, and Dr. Bruce MacDonald

of ONR for their interest and support.

The views and conclusions herein are those of the authors and should not be interpreted as necessarily representing the official policies, either expressed or implied, of the Defense Advanced Research

Projects Agency, the U.S. Air Force, the U.S. Navy, or the U.S. Government.

REFERENCES

[1]

. P. Duwez, R.H. Willens and W. Klement, J.App. Phys. ^ (1960) 1136, and Trans. Met. Soc. A.I.M.E.

227(1963) 362.

[2]

. N.J. Grant and B.C. Giessen, eds., Proc. Second International Conf. on Rapidly Quenched Metals,

Materials Sci. and Eng. V23 (1976); Proc. Third International Conf. on Rapidly Quenched Metals,

B. Cantor, ed.. The Metals Society, VI and 11(1978); Proc. Fourth International Conf. on Rapidly

Quenched Metals, H. Jones, ed . , Japanese Inst. Metals (1981).

[3]

. Proc. First and Second Int. Conf. on Rapid Solidification Processing: Principles and Technologies,

R. Mehrabian, B.H. Kear and M. Cohen, eds., Claitor's Publishing Division, Baton Rouge LA, 1978,

1980.

[4]

. P.R. Holiday, A.R. Cox and R.J. Patterson II, in Ref. [3] (1978) pp.246.

[5]

. R.D. Field and H.L. Fraser, ibid., 270.

[6]

. R.J. Patterson, II, unpublished. West Palm Beach FL June 1977 and La Jolla CA July 1977.

[7]

. M.R. Glickstein, R.J. Patterson, II and N.E. Shockey in Rapid Solidification Processing: Principles and Technologies , R. Mehrabian, B.H. Kear and M. Cohen, eds., Claitor's Publishing Division,

Baton Rouge, LA (1978) 46.

[8]

. R.J. Patterson ,I I , D.L. Ledwith and J.C. Dwyer, in Proc. ACS-AIME Powder Processing Symposium,

K.W. Lay and R.M. German, eds., TMS-AIME, Warrandale, PA (1982).

[9]

. C.M. Adam et ai., AFWAL Final Report, Contract F33615-76-C-5136, P&WA/GPD Report FR15325, September 1981.

[10]

. T. Kelly, Ph.D. Thesis, M.I.T., Cambridge, Massachusetts, 1981.

[11]

. A.J. Drehman and D. Turnbull, Scripta Met. 1^ (1981) 543.

[12]

. J.H. Perepezko and I.E. Anderson in Synthesis and Properties of Metastable Phases , E.S. Machlin

and T.J. Rowland, eds., TMS-AIME, Warrandale, PA (1980) 31.

[13]

. R.E. Cech, Trans. A.I.M.E. 206 (1956) 585.
[141, G.L.F. Powell and L.M. Hogan, Trans. Met. Soc. A.I.M.E. 242 (1968) 2133

(151. J.H.

B.C. Perepezko and J.S. Paik in Rapidly Solidified Amorphous and Crystalline Alloys, B.H. Kear Giessen and M. Cohen, eds., North-Hol land , N.Y., NY (1982) 49.

[16]. J.P. Hirth, Met. Trans. 9A (1978) 401.

[17] . C.G. Levi and R. Mehrabian, Met. Trans. IIB (1980) 21.

[18] . D.J. Dyson and B. Holmes, J. Iron and Steel Inst. 208 (1970) 469.

[19]. F.J. Fasiska and H. Wagenblast, Trans. Met. Soc. A.I.M.E. 239 (1967) 1818.

[20] . W.C. Leslie, Met. Trans. 3 (1972) 5.

[21]. F.B. Pickering, Physical Metallurgy and the Design of Steels, Applied Science, London, 1978.

d2
OSTWALD RIPENING OF RAPIDLY SOLIDIFIED SOLID-LIQUID MIXTURES

P. W. Voorhees* and M. E. Glicksman

Materials Engineering Department

Rensselaer Polytechnic Institute

Troy, NY 12181

ABSTRACT

A new theory of Ostwald ripening in two-phase mixtures has been developed which ex- plicitly accounts for the diffusional interaction between the dTspersed coarsening second

phase. This theory predicts the morphology of coarsening solid/liquid mixtures in terms

of time invariant distributions of interfacial curvature. The theory also predicts a dependence of the curvature distributions and ripening kinetics on the volume fraction

solidified. Experiments were performed to measure the response of an ultra-precise therma

probe immersed in a coarsening rapidly solidified solid/liquid mixture of high purity

white phosphorous over a wide range of fraction solids. Through this theory it is now possible to interpret the experimental results to gain a deeper insight into the nature o Ostwald ripening following rapid solidification.

Introduction

Ostwald ripening has long been viewed as a si ow morphol ogical relaxation process which follows

the initial phase transformation. The concept of ripening or coarsening as a kinetically slow

process no doubt stems from the extensive investigation of the effects of coarsening on the microstructure of slowly solidified castings. The work of Kattamis, et al., has established that in large

castings, where the local solidification times are on the order of many minutes or hours, ripening

plays a major role in determining the size scale of the as-sol idi fied microstructure [1].

By contrast, most rapid solidification processes are associated with interfacial

velocities and generally shorter local solidification times. The influence of coarsening of the

microstructure of rapidly solidified materials is commonly thought to be negligible because of a preconception of coarsening as a kinetically slow relaxation process.

A major impediment in the assessment of coarsening processes

the extremely short local solidification times which are commonly encountered. These br ef local

solidificatiL times--often as short as 10'^ s--preclude the use of standard experimental techniques,

such as quenches at different aging times for determining the role of coarsening in the morphological

rogressirof so?id-]iquid mixtures. As a result, there is little direct experimenta evidence

to indicate that coarsening plays any role in the mi crostructural development of rapi^

materials. A second impediment in the study of dendritic ripening in solid-liquid mixtures is the

lack of cogent theoretical guidance. This is due primarily to the d^f^^culty in construe^

theorv of coarsening in solid-liquid mixtures which realistically accounts for the diffusional inter

actions among regioL of the system with differing interfacial curvature and, consequently, differing

chemical potential.

To address these current deficiencies in our understanding of coarsening following rapid solidification, experimental measurements were made of the kinetics of coarsening in rapi ^

^

. solid-liquid mixtures and a realistic theory of microstructural relaxation in solid-liquid mixtures

was deve?oped which is capable of providing new insights into the dendrite ripening problem.

*Presently at the National Bureau of Standards, Washington, DC 20234

83
Experimental

s-i-i

A novel experimental procedure was employed to produce rapidly solidified mixtures with local

solidification times on the order of many hours. These long local solidification times were essential

in permitting the coarsening kinetics of a rapidly solidified solid-liquid mixture to be measured

experimentally. The following is a brief description of the apparatus and procedure used to measure , the kinetics of dendritic coarsening in rapidly solidified solid-liquid mixtures. A more detailed - j: description of the procedure and apparatus can be found elsewhere [ 2 ]. !

Measurements of the coarsening kinetics were made using rapidly solidified, high purity, white ; phosphorus (aP 4 ). aP 4 was chosen as the model material because of the relative ease in achieving

large supercooling. The supercooling tendency of 0P 4 enabled precise control of both the interfacial velocity and the volume fraction of solid transformed during solidification. The volume

fraction of solid, rapidly solidified under adiabatic conditions, is numerically equal to the normal- | ized supercooling, 44> = CaT/L where L is the latent heat of fusion, aT is the supercooling, and

C is the specific heat of the liquid [2]. The interfacial velocities achieved in aP4 were usually ,i on the order of 200 cm/sec, and thus were of the same magnitude as interfacial velocities commonly

reported for other rapid solidification processes. To preserve the adiabaticity of the two-phase |; mixture immediately following nucleation of the supercooled melt, the 30 ml bulb containing the j solid-liquid mixture was transferred immediately to a stirred bath carefully thermostatted approx- , imately 4mK below the bulk melting temperature (44. 1C). Placing the bulb in the thermostat greatly

reduced the heat loss from the bulb, thus preserving the sol id-to-1 iquid ratio in the two-phase

mixture for periods of many hours.

The equilibrium temperature of an interface is directly related to its curvature, k, via the

Gibbs-Thomson equation, viz..

T(k*) = T. (

1

)

where <* is the total dimensional curvature at a point on the interface, y is the solid-liquid

interfacial energy, T^ is the bulk melting point, and ASf is the entropy of fusion. The progress

of dendritic coarsening can, in principle, be measured in situ via a precision thermometer capable of

resolving temperature changes of 10“^K within the two-phase mixture.

Results

A typical log-temperature vs. log-time output from the thermistor is shown in Fig. 1. The slow

rise in thermistor temperature with time is consistent with increasing interfacial radii of curvature i linked to the ripening of the solid-liquid mixture. To determine the rate of Ostwald ripening at times too short to be directly measurable with the instrumentation used in this experiment, it is necessary to assume that the power-law kinetics shown at long times in Fig. 1 can be extrapolated *

TIME (sec)

i

I

Fig. 1. Liquidus temperature, T. , minus the ; thermistor temperature versus time (log-log f- scale) showing short-time response of the ]' thermistor (due to thermal inertia) and longtime response due to Ostwald ripening. The two ii kinetic processes are easily distinguished here, i as is the power-law character of the dendritic

coarsening.

84
' back toward shorter times. Theoretical justification for this back extrapolation is found in ripenling theory which asserts that power law coarsening kinetics exist as long as the system is at steadyI state. Theory also indicates that steady-state ripening should occur under most circumstances in a

j very short time compared to the duration of the steady-state coarsening process. Further discussion

I of this point can be found elsewhere [2]. The equation of the line shown in Fig. 1 is

l| T = Tl - Bt-^ (2)

ill

i

!j It has been shown theoretical ly that the temperature of a thermistor immersed in a coarsening solidliquid mixture is related to the average radius of curvature by [2],

I T = T

L Ay

AS^ <R> ’ ( 3 )

\

i where A is a constant of order unity and <R> is the average total radius of curvature of the solidI liquid interface.

Substituting eq. 3 into eq. 2 yields

<R> = . Ay

^ B AS^ ^ •

It is possible to compute the coarsening rate from eq. 4 as.

d<R> = .ft. _ Ayn n-1 dt “

B AS^ ^

( 4 )

( 5 )

The data in Fig. 1 yield values of B and n of 0.0486 and 0.476, respectively. The coarsening rates

and average radii of curvature of the solid-liquid mixure at shorter times can now be found by extrapolating the data of Fig. 1 backwards in time. Clearly, extrapolation to t = 0 is precluded because

eq. 5 diverges as t > 0. A reasonable estimate for the time corresponding to the beginning of steadystate coarsening in these experiments is 10"^ seconds. Inserting the materials constants of aP4 into eqs. 4 and 5 and taking A = 1 yields the average interfacial radius of curvature and coarsening

rates as functions of time, respectively, as shown in Fig. 2. A number of interesting observations

can be made on the basis of Fig. 2: 1) the average radius of curvature soon after solidification,

~ 10-3 s, is indeed small, but well above the atomic scale; 2) this small radius of curvature gives

rise to an extraordinari ly rapid coarsening rate. Our experimental measurements and subsequent

order-of -magnitude estimates of coarsening rates in rapidly solidified solid-liquid mixtures have

shown that significant coarsening of the structure can occur with local solidification times on the

order of 10“3 s. Further insights into the morphology and kinetics of dendritic coarsening in

solid-liquid mixtures can be obtained by considering the theory of Ostwald ripening.

Fig. 2. Average interfacial radius of curvature <R> and <R> versus time. These plots

correspond to the thermal response data shown

in Fig. 1, as interpreted through Equations 4 and 5 in the text.

85
Theory I

Developing a predictive theory of dendritic coarsening in solid-liquid mixtures which is based | upon detailed interfacial morphology is virtually impossible. We shall therefore consider the solidliquid mixture present following rapid solidification in a topological sense whereby only the statistical distribution of interfacial curvature associated with the mixture will be considered, and not i its detailed morphology. A distribution of interfacial curvature at some instant of time following j rapid solidification will be modeled in a statistical sense with a distribution of spherical solid ! particles having the same distribution of local equilibrium temperatures. Using this approach it is possible to gain considerable insight into the coarsening behavior of solid-liquid mixtures.

Recently, the authors have developed a fast algorithm for accurate numerical solution to the

multi -particle diffusion problem (MDP) [3]. The MDP has been employed to determine the temperature

field within a coarsening solid-liquid mixture, and as a result interparticle thermal interactions

are real i stical ly included in the theory. Interparticle interactions allow the volume fraction of

coarsening phase to play a role in determining the ripening kinetics and morphology. A detailed

description of the theory of the MDP can be found elsewhere [3]. The theory indicates that coarsening solid-liquid mixtures quickly reach a steady -state, wherein the inverse of the cube of the average

curvature grows linearly with time, i.e..

<<(t)>-3 - <k( 0)>-3 = Kt. ( 6 )

Here <<(0)> is the dimensionless average curvature at time t = 0, <<(t)> is the dimensionless

average curvature at time t, and K is the rate constant. The theory also shows that the distribution

of interfacial curvature, when normalized by the average curvature, is time invariant. The specific

predictions on the morphology and coarsening kinetics are discussed below.

Morphol ogy I As was mentioned earlier, under steady-state conditions the distribution of interfacial curvature

scaled to the average becomes time invariant. However, theory also indicates that the distribution

of interfacial curvature should depend on the volume fraction of solid in the two-phase mixture.

The volume fraction dependence of the distribution of curvature results from the diffusional interaction between regions of differing curvature. As the volume fraction of solid is increased, the

regions of differing interfacial curvature become closer, on average, which causes the thermal i. interactions among these regions to increase. The increasing interactions in turn produce a changej

in the interfacial morphology of the solid-liquid mixture. Since the volume fraction of solid is

equal to the normalized supercooling, the morphology of coarsening sol id -liquid mixtures should

depend on the degree of supercooling prior to nucleation. The changes in interfacial morphology with

volume fraction can be seen in the steady-state distribution of curvatures shown in Fig. 3. As is evident, the increased diffusional interaction with increasing volume fraction manifests itself as a broader and more symmetric distribution of interfacial curvature.

Kinetics of Morphological Relaxation

As expected intuitively, the increasing thermal interactions among regions of differing curvature'

as the volume fraction of solid increases cause an increase in the growth rate of the average \ interfacial radius of curvature. Such an increase in microstructural coarsening rates is expressed

by a change in the rate constant, K, appearing in Eq. 6. The functional dependence of K on volume

fraction is shown in Fig. 4. This figure illustrates that partially solidified mixtures with a volume fraction of solid, (or normalized undercooling) approaching 0.5 should coarsen at about four

times the rate as that of a mixture with a solid volume fraction near zero. The results shown in

Fig. 4, in conjunction with the rate equation, Eq. 6, can be combined to yield specific predictions

on the interfacial coarsening rates of solid-liquid mixtures. The coarsening rate for a given average’

curvature can be found from Eq. 6, viz.

<R> = ^ = K <[^>-2

dt 3

( 7 )

where <R> is the dimensionless average radius of curvature of the solid-liquid mixture. The inverse

squared dependence of the coarsening rate on the average radius of curvature indicates that mixtures

with a small average radius of curvature will coarsen at high rates. If the volume fraction dependence of K is included, then the coarsening rate behaves as is shown in Fig. 5. As was observed

OO
J!

i|

Fig. 3. Theoretical distributions of interfacial curvature. Distribution

functions are shown for the various volume fractions of solid studied by computer

simulation of the steady-state MDP.

experimentally, the theory predicts that a dendritic solid-liquid system with small radii of curvature will coarsen at high rates relative to the size scale of the microstructure. This can be seen

more clearly by denormalizing the variables shown in Fig. 5. For example, chosing a normalized

undercooling of 0.35 and <R> = 10^ yields a dimensionless coarsening rate of 4.2 x 10“^. Using the

materials constants of a-P 4 , <R> and <R> can be dimensional ized to give <R> = 6.51 y and <R> 10.4 y/s. These theoretical predictions further reinforce the idea that Ostwald ripening plays

an important role in the development of the morphology of rapidly solidified solid-liquid mixtures.

Conclusions

The following can be concluded on the basis of this paper:

1. An experimental procedure has been developed which permits investigation of the coarsening

behavior of rapidly solidified solid-liquid mixtures via direct thermometry.

2. The average interfacial radius at the beginning of steady -state dendritic coarsening following

rapid solidification is estimated to be a moderately small multiple of the capillary length.

87
i/i (O

o

°0q ^00 0,10 0.20 0.30 0.40 0.50 0.60 0,70 0.80 0.90 1.00 VOLUME FRACTION

Fig. 4. Dependence of the cubical rate constant K on volume fraction.

3. The large interfacial curvatures following rapid solidification give rise to rapid coarsening

rates relative to the size scale of the microstructure. Calculations indicate that a-P 4

in a solid-liquid mixture with an average radius of curvature of 0.1 y will double its radius in

approximately 3 x 10~4s.

4. Steady-state distributions of interfacial curvature can be calculated on the basis of Ostwald

ripening theory, and are dependent on the degree of normalized liquid supercooling.

5. At steady -state , the cube of the average radius of curvature grows linearly with time, with

the rate of growth being dependent on the normalized liquid supercooling.

6. Theoretical predictions of the coarsening rates of solid-liquid mixtures are in qualitative

agreement with experiment.

Acknowl edgment

The authors acknowledge the support provided by the National Science Foundation (NSF) through

the Division of Materials Research, under contract DMR-7721638. We are espeically grateful for the

special provision of a Research Creativity Award provided to one of us (MEG) by NSF.

88
Fiq. 5. Coarsening rate <R> versus average interfacial radius of

curvature <R> for dendritic mixtures supercooled initially an amount A4>.

1 .

2 .

3.

T.

P.

P.

Z. Katamis,

W. Voorhees

W. Voorhees

References

J. C. Coughlin, and M. C. Flemings, Trans. Met. Soc. AIME, 2^ (1 967) pp. 1504.

and M. E. Glicksman, J. Crystal Growth, to be submitted,

and M. E. Glicksman, Acta Met., to be submitted.

89
J. J, Richmond, J. H. Perepezko, S. E. LeBeau and K, P. Cooper

SOLIDIFICATION MICROSTRUCTURES OF UNDERCOOLED DROPLETS

Department of Metallurgical and Mineral Engineering

University of Wisconsin-Madison

Madison, Wisconsin 53706

^01 in Metals Research Labs

New Haven, Connecticut 06511

ABSTRACT

+

Following nucleation in an undercooled liquid the subsequent crystal growth will yield rapid

solidification conditions. Fine liquid droplets exhibit large undercoolings which can be measured B

accurately at a slow cooling rate and can be varied experimentally in a controlled manner. With thisW

approach the development of solidification microstructures as a function of melt undercooling and

composition has been examined in several eutectic alloy systems. In Pb-rich Pb-Sb alloys with com- positions extending to the eutectic a transition from a segregated, two-phase microstructure to a B

supersaturated solid solution has been identified with increasing undercooling level. These observa-B

tions have been related to the onset of hypercooling conditions during isenthalpic solidification. ^ 9

In hypereutectic Al-Fe alloys a transition from growth controlled features to nucleation controlled H

structures has been observed where formation of the metastable Al 5 Fe occurs in undercooled droplets ^

even in the presence of primary Al 3 Fe particles. For Pd-Cd alloys the usual eutectic pattern developed!

at modest undercooling is circumvented by a phase separation reaction as the initially homogeneous

alloy liquid enters a miscibility gap region before crystallization at high undercooling. These

observations highlight several solidification paths for microstructure development.

The emergence of Rapid Solidification Processing (RSP) as an innovative approach to materials

processing has been based to a large extent on favorable microstructural features [1-3]. As a broad

generalization of the microstructures associated with RSP there appear to be two categories of characteristics. In one case the influence of RSP treatment can be viewed in terms of the decrease in

local solidification time and the growth features of the principal microstructural elements. Another

category of microstructural development is the evolution of novel structures comprised of phase mixtures that are not associated with the equilibrium features of phase diagrams. The examination of

microstructure control in the case of novel structures requires a consideration of nucleation kinetics

and significant melt undercooling. In addition, for both general categories the influence of thermal ’ history should be incorporated in the analysis of microstructure. ‘

One approach to the examination of structure development during RSP involves the study of droplet

samples. The dispersal of a liquid into fine droplets has been demonstrated to yield very high undercooling levels which can be measured accurately [4]. This feature is a prerequisite for a detailed -j examination and understanding of the phase selection kinetics involved in nucleation controlled product

structures. In addition, with droplet samples and thermal cycling treatments it is possible to examine growth controlled microstructures and thermal history effects associated with RSP, In the present

work, the development of solidification microstructures has been examined in highly undercooled drople

samples of Pb-Sb, Al-Fe and Pb-Cd alloys.

Droplet samples were prepared by an emulsification process in which a high speed motor (30,000

rpm) attached with a shearing blade assembly disrupts a mixture of carrier fluid and liquid metal

under an inert atmosphere into a droplet dispersion. With a large number of droplets only a small

fraction of the drops may contain the most potent nucleant sites. The fine droplets (d<20y) are maintained independent by the use of thin surface coatings produced by chemical reaction. For metals

which melt below 500°C, a mixture of organic carrier fluid dissolved with organic or inorganic acid

surfactants and a liquid metal alloy, which is prepared with metals of 99.999% purity, is sheared at

Introduction

Experimental

90
IbO to 100°C above the liquidus temperature. Recently the droplet method has been extended to jallow the study of high melting temperature alloys such as Al-Fe with the use of molten salt carrier

i fluid mixtures. The crystallization temperature and kinetics can be measured accurately by either a

i differential thermal analysis (DTA) or a differential scanning calorimetry (DSC) technique. Following

thermal analysis, the phase selection and solidification microstructure were examined by applying an x-ray diffraction (XRD) technique and by metal lographic study of exterior surface and interior

droplet sections.

Results and Discussion

A large volume of bulk liquid can be considered to have a distribution of nucleating centers with

different catalytic potencies. The amount of undercooling at a given cooling rate is determined by the

effect of the most active sites. By progressively removing or isolating nucleation sites from a sample, an undercooling level will be observed which is characteristic of the most active nucleant still

remaining [5]. The importance of the degree of the droplet size refinement in promoting a high level

of liquid undercooling has been demonstrated clearly [4]. Based upon the current experience, a uniform

undercooling characteristic of a single crystallization exotherm far below the melting point is represented usually by droplets in the 5-20 ym size range. In addition it has been demonstrated that in^

fine droplet dispersions with a narrow size distribution, the extent of liquid undercooling is sensitive also to the surface coating used to maintain droplet independence [4].

Pb-Sb Alloy droplets

An example of uniform undercooling to a single crystallization exotherm is represented in fiQ*

1 for a Pb-11.4 a/o Sb alloy. The maximum undercooling nucleation temperature, T[\], occurs at 113°C

which is equivalent to an undercooling, AT, of 162°C below the alloy liquidus, T[_. The sensitivity

of the undercooling behavior to alternate nucleation sites is also illustrated in fig. 1 where the

cooling trace yielding final crystallization at Thet ri85°C) represents the catalysis of an a-Pb primary phase in promoting solidification.- In examining the influence of melt undercooling on microstructure development, alternate nucleation sites, that are generated by the application of different

surfactant treatments, were used as an effective approach to produce controlled undercooling levels.

A summary of the undercooling behavior for Pb-Sb alloys for compositions spanning the eutectic is

superimposed on the portion of the phase diagram plotted in fig. 2. As observed in other alloys [5j

the trend of undercooling with alloy composition exhibits a pattern which tends to follow roughly the

alloy liquidus.

The apparently simple behavior of the undercooling trends reported in fig. 2 however, does not reflect the variety of microstructural patterns developed in droplets following rapid solidification.

This variety is illustrated for a Pb 14.4 a/o Sb alloy by the microstructures given in fig. 3. At a low AT of 14°C the microstructure in fig. 3a a reveals coarse a dendrites with a complex regular eu- tectic within interdendritic regions [6] and would be represenative of a conven-tional processing

treatment. In droplet samples at an undercooling level of 111°C a strongly modified microstructure

is evident in fig. 3b. In this case the light, clear regions represent an a primary solid solution

which was supersaturated up to the alloy solute content while the remaining region represents a fine

scale two-phase mixture. The initial supersaturation of the a phase has been established by separate

XRD measurements on droplet samples immediately following thermal analysis. During sample preparation

to reveal interior droplet sections, the supersaturated a phase undergoes a precipi tation reaction

_

which produces a pattern of fine, dark Sb precipitates. Another stage of microstructure modification

is apparent in the micrograph of fig. 3c corresponding to droplet solidification at a AT of 162 C.

In this case, XRD measurements have been used to establish that the initial solidification structure

was essentially entirely a supersaturated a phase. Again, during sample preparation a fine scale

precipitation of the Sb phase developed in the single crystal droplets.

Based upon the measured undercooling values and the ini-tial droplet structures, it -is possible to develop an understanding of the observed microstructural modifications in terms of an adiaoatic solidification model [1,3]. When the nucleation temperature, T^, is close to Tl, a segregated structure can develop as in fig. 3a. It is possible for crystallization to yield a solid of the alloy composition

Cn as Tw is reduced below Tq which represents the onset of partitionless solidification [7] and is

also drawn schematically in fig. 2. Since the rate of crystal growth following nucleation in a highly

undercooled liquid is rapid, the latent heat will evolve at a rapid rate during recalescence. The

temperature rise during recalescence reduces the amount of interfacial undercooling and the growth

which for a solid of composition Cq can reduce to zero upon reaching Tq which appears to set a bound on partitionless solidification. Toward the end of recalescence and during final droplet solidification

the solid composition can change leading to a segregated structure and a mixed morphology (fig. 3bj.

With the adiabatic model, when the undercooling is sufficient so that the recalescence temperature rise

does not lead to the development of solute partitioning, solidification can occur entirely to a single

phase solid as shown in fig. 3c. The limiting undercooling for this process defines the onset of
400

DTA OF Pb-11.4 a/o Sb ALLOY

Fig. 1 DTA Thermogram for Pb-11.4 a/o Sb alloy

droplets.

Fig. 2 Summary of undercooling behavior in Pb-Sb

alloys. Different symbols refer to undercooling

trends produced by different droplet surface

coating treatments.

Fig. 3 Solidification microstructures developed in Pb-14.4 a/o Sb alloys at different undercooling

levels (a) AT = 14°C; (b) AT = 111°C; (c) AT = 162°C.

92
hypercooling, ATh which is given to first order by AT^ = AHj^/Cp where AH^ is the latent heat and Cp is the liquid heat capacity [3,4]. In Pb-Sb alloys ATh appears to be reduced with solute content from

about 157°C for pure Pb [4] to about 130°C for a Pb-14.4 a/o Sb alloy [6]. This is in agreement with

the solidification to a supersaturated solid solution in droplets at AT = 162°C and the formation of

partially segregated droplet structures at AT = 111°C.

In a number of hypoeutectic alloy compositions microstructural patterns similar to those reported

for the Pb-14.4 a/o Sb alloy have been observed. For alloy compositions near the eutectic however, a different pattern of microstructure was observed as illustrated in fig. 4 for a Pb-17.7 a/o Sb eutectic

alloy. At a AT of 19°C a complex regular eutectic pattern develops. Also, at undercoolings from

70°C to 130°C the eutectic pattern is replaced by a dendritic microstructure instead of single phase

a. This agrees with other rapid melt quenching work [8] where the limit for solubility extension

was approximately the eutectic. The transition from a eutectic to a dendritic morphology with increasing undercooling has been observed to occur also in hypereutectic alloys and appears to be related to the location of the coupled zone for eutectic growth [9]. The coupled zone region usually tends

to be skewed towards compositions high in the faceting component such as Sb. As a result, with increasing undercooling, a transition can occur from a eutectic to a dendritic structure. In addition,

measurements of the secondary dendrite arm spacing (DAS) reveal only a small decrease from 0.37y at

AT = 70°C to 0.23 p at AT = 130°C in the eutectic alloy which suggests that a significant structural

coarsening may control DAS at high undercooling [10]. In this case the application of DAS measurements

to characterize parameters such as solidification rate may not be reliable.

Undercooled Al-Fe Alloys

The Al-Fe alloy system exhibits a eutectic reaction at 1/8 w/o Fe and 655°C between A1 containing

0.052 w/o Fe and a binary Al-Fe intermetal 1 ic compound, Al 3 Fe. The Al-Al 3 Fe eutectic microstructure

at moderate freezing rates is an example of a faceted-nonfaceted irregular eutectic structure consisting of plates of the faceted Al 3 Fe phase in an aluminum matrix. A transition from the irregular

Al-Al 3 Fe structure to a fine fibrous metastable eutectic of Al-Al 5 Fe has been observed in chill casting [11] controlled direction solidification studies [12] and in rapid quenching experiments [13].

The potential for Al-Fe based high-strength PM alloys for use at elevated temperatures is receiving

continued interest [14].

The eutectic systems of the Al-rich end of the Al-Fe phase diagram have been examined using the

droplet-emulsion technique. Fig. 5 is a typical DTA thermogram for a hypereutectic droplet emulsion

sample of Al-9 w/o Fe. Upon heating the droplet sample (t = 30 K/min) an endothermic peak associated

with the melting of the metastable Al-Al 5 Fe eutectic, m-T^, occurs at 649°C corresponding to a eutectic composition of the Al-AlgFe at 3.8 w/o Fe by a simple linear extrapolation of the solidus

and liquidus of the equilibrium phase diagram. Upon further heating the thermal signal returns to

the baseline value at a T[_ of 875°C. During the subsequent cooling of the sample (T = 30K/min) a small exothermic peak which indicates the crystallization of primary AlgFe is observed at 850°C. Any

further crystallization and/or growth of the primary phase is not easily distinguished from the normal

baseline behavior. Upon cooling below the eutectic temperature a pair of broad crystallization

exotherms occurs with peak temperatures of 640°C and 565°C respectively. The DTA thermogram suggests

that even in the presence of Al 3 Fe primary particles substantial undercoolings of at least 85°C below

the eutectic isotherm can be achieved. The relatively low catalytic activity of the Al 3 Fe primary

phase can be substantiated by reheating the sample above the eutectic temperature to equilibrate a liquid-Al 3 Fe primary solid two-phase mixture within each droplet. Subsequently, upon cooling the

samples, the crystal 1 ization behavior observed was identical to that obtained in the previous cycle

during cooling from above Ti_. The droplet microstructure obtained after the thermal cycling in fig. 5 is shown in the SEM micrograph in fig. 6. The microstructure consists of a primary Al 3 Fe plate which

is surrounded by a metastable Al-Al6Fe eutectic mixture. The undercooling and microstructure characteristics that are presented in figs. 5 and 6 reveal that a distribution of nucleants can be active in

a droplet sample. These observations suggest that the range of microstructural variations that have

been reported in aluminum base systems in rapid quenching studies may be the result of variations in

initial liquid undercooling in addition to variations of cooling rate.

Pb-Cd Alloys

Often it is implicitly assumed in solidification that an initial alloy liquid is compositional ly

homogeneous. During conventional processing this characteristic may be assessed by examining the

equilibrium phase diagram. However, during RSP which can involve liquid undercoolings up to about

0.3 Tpi, it may be necessary to reconsider the assumption of initial melt homogeneity. For example,

at high undercooling the alloy liquid may enter into the domain of a metastable miscibility gap. The

role of initial melt constitution on microstructural development in rapid solidification has been

examined in the Pb-Cd system which exhibits an equilibrium eutectic reaction. However, for hypereutectic alloys, DTA studies on undercooled droplets reveal that prior to crystallization a phase

separation can occur in the liquid. The consequences of phase separation on the solidification

93
Fig. 4 Solidification microstructures developed

in Pb-17.7 a/o Sb alloys at different undercooling

levels. A transition from eutectic (a) [AT=19°C]

to a dendritic surface morphology (b) [AT=130°C]

is illustrated.

Fig. 6 Solidification microstructure developed

in Al-9 w/o Fe droplets.

Fig. 5 Undercooling behavior of Al-9 w/o Fe droplets.

Fig. 7 Microstructure developed in Pb-40 w/o Cd

droplets in which phase separation proceeded

solidification.

94
microstructure are illustrated strikingly in fig. 7. The interior droplet section reveals a Pb-rich

region (light) surrounding a Cd-rich core (dark) which is a substantially different morphology than

that developed during usual eutectic solidification. For Pb-Cd alloys the tendency for liquid phase

separation may be judged by the shallow Cd liquidus so that the undercooling required to achieve

liquid demixing is not large [15]. In general the likelihood of a metastable miscibility gap may not

be readily apparent from an equilibrium liquidus without a thermodynamic analysis, but this feature

can be encountered more often at large undercooling and for conditions involved in glass formation.

Summary

The controlled undercooling in fine droplet samples provides an effective experimental approach

to the study of RSP. In Pb-Sb alloys the capability to vary melt undercooling has allowed for the

study of a transition from a segregated two-phase morphology to a single phase solidification microstructure. This microstructural transition has been related to the onset of hypercooling in terms of

rapid solidification under adiabatic conditions. Similar studies on eutectic Pb-Sb alloys have revealed only a weak dependence of DAS on undercooling which may be the result of structural coarsening

effects. In droplets of Al-Fe alloys a large undercooling and solidification of a metastable Al 0 Fe phase has been established even in the presence of primary Al;jFe nucleants. The microstructural effects of melt inhomogeneity due to a metastable phase separation have been demonstrated in undercooled

Pb-Cd alloys. These observations highlight the use of undercooled droplet samples to examine paths

for microstructure development during RSP.

Acknowledgement

The authors are grateful for the support of the U.S. Army Research Office (DAAG29-80-K-0068) . We

are indebted to W. J. Boettinger, G. J. Hildeman and C. M. Adam for discussions and assistance.

References

[1]. R. Mehrabian, "Rapid Solidification Processing-Principles and Technologies", R. Mehrabian,

B. H. Kear and M. Cohen (eds.), p. 9, Claitor Pub. Div., Baton Rouge, LA, (1978).

[2]

. M. Cohen, B. H. Kear and R. Mehrabian, "Rapid Solidification Processing - Principles and

Technologies II", R. Mehrabian, B. H. Kear and M. Cohen (eds.), p. 1, Claitor Pub. Div.,

Baton Rouge, LA, (1980).

[3]

. C. G. Levi and R. Mehrabian, Met. Trans. 13A 13, 221 (1982).

[4]

. J. H. Perepezko, ref. 2 p. 56.

[5]

. J. H. Perepezko and J. S. Paik, "Rapidly Solidified Amorphous and Crystalline Alloys", B. H. Kear, B. C. Giessen, and M. Cohen (eds.) p. 49, Elsevier Pub. Co. N.Y. (1982).

[6]

. J. J. Richmond and J. H. Perepezko (to be published).

[7]

. J. C. Baker and J. W. Cahn, "Solidification" p. 23, ASM, Metals Park, OH (1971).

[8]

. C. Barromee-Gautier , B. C. Giessen and N. J. Grant, J. Chem. Phys. 1905 (1968).

[9]

. W. Kurz and D. J. Fisher, Int. Met. Rev. 5-6 177 (1979).

[10]

. P. W. Voorhees and M. E. Glicksman, ref. 5 p. 33.

[11]

. L. Backerud, Jernkont. Ann. 152 109 (1968).

[12]

. R. Hughes and H. Jones, Jnl . of Mat. Sci. 1J_ 1781 (1976).

[13]

. H. Jones, Mater. Sci. Eng. ^ 1 (1969).

[14]

. C. M. Adam and R. G. Bordeau, ref. 2 p. 246.

[15]

. J. H. Perepezko and W. J. Boettinger (to be published).

95
MICROSTRUCTURAL STUDIES ON SUPERCOOLED SUB-MICROMETER POWDER PARTICLES

D. Shechtman , S. D. Ridder^ and R. Mehrabian^

*Center for Materials Research, The Johns Hopkins University, Baltimore, MD 21218

'Center for Materials Science, National Bureau of Standards, Washington, DC 20234

ABSTRACT Pj

Various Al-Cu alloy powders were made by the Microparticle Processor and their |l microstructure studied by transmission electron microscopy. The powder particles which

sizes range from 3 nm to a few microns demonstrate different structural characteristics

^ ranging from amorphous to ordered crystalline intermetal 1 ic depending upon composition

and particle size. Extended solid solubility was observed in all the alloys examined.

^ Some of the results obtained from Al-Cu alloy powders will be shown and discussed. The ^ cooling rates needed to suppress crystallization in some alloys will be calculated. r 1

Introduction

I' The use of electrohydrodynamics (EHD) in the atomization of metal alloys, in an apparatus L hereafter referred to as the micro-particle processor (MPP)^ has provided a means of producing small

quantities of submicrometer size metal particles. These small powders are, in general, transparent

to 100 Kev electrons, used in conventional TEM analysis allowing a three-dimensional view of the

entire solidification structure. This paper will discuss some aspects of the microstructures observed

in aluminum-copper alloy powders produced by this process. Of particular interest here is the

formation of metastable solid solutions in many of these powders and powders showing a transition

from partitionless (massive) solidification to a segregated solidification mode as described in an earlier paper by Levi and Mehrabian [1] and in this volume by Cohen and Mehrabian [2]. The influence ' of particle size and composition on microstructure is discussed, especially as it relates to the

apparent bulk supercooling achieved prior to nucleation and the solid/liquid interfacial velocity || during growth and recalescence.

Experimental Procedure

Al-Cu binary alloy powders were produced using the microparticle processor (MPP) at NBS. A detailed description of this apparatus is given in these proceedings [3]. Droplet generation in the ,i- processor essentially involves the application of a very intense electric field on a conductive ijj liquid at the tip of a capillary nozzle (emitter). The alloys contained 0.5, 1, 2, 3, 4.5 and i;| 6 wt.% Cu. The production parameters were 14 kV_extraction voltage with an argon feed pressure of lij 50 mm Hg X cm 2 against a vacuum of 10 ^mm Hg cm The flight distance to the cellulose acetate

tape collector (target) was 150 mm. Specimens for transmission electron microscopy were prepared by

depositing a thin carbon layer on the monolayer of particles captured on the target tape. The tape ii was subsequently dissolved in acetone. Specimens thus made were mounted on 400 mesh grids and ,! studied in a scanning transmission electron microscope (STEM) equipped with a nondispersive x-ray Ij fluorescence analysis system. The study of composition variation across a powder particle involves iil several experimental difficulties, including the proper geometrical positioning of the particle J relative to the grid base, the electron beam and the detector. Appropriate adjustments were made to ; compensate for these so as to produce minimum experimental error. Other difficulties are inherent to i the spherical shape and size of the individual particles. These require a correction procedure that

will take into account the specific electron beam location on the particle, i.e., the cross-section

of the particle at that point and its location relative to the x-ray detector. Such correction

procedure was not done, but an attempt was made to measure x-ray intensities generated from points of ji equal thickness. This may account for some inaccuracies in the composition analysis. | Finally, iji si tu heat treatment of individual powder particles was performed under controlled

electron beam in the STEM and the sequence of second phase precipitation from the supersaturated J solid solutions was recorded. S

^Disclaimer: Certain trade name and company products are mentioned in the text in order to adequately specify the experimental procedure and equipment used. In no case does identifica- , tion imply a recommendation or endorsement by the National Bureau of Standards, nor does it imply that products are necessarily the best available for the purpose. This

processor uses electrohydrodynamic atomization technique to break a molten metallic

droplet situated at the tip of an extractor tube into smaller droplets, which then i cool by radiation and solidify.

»

95
Results and Discussion

Most of the Al-Cu particles studied were in the size range of 3 nm to 2 pm. In general,

larger particles did not solidify in flight. Furthermore, particles larger than 2 pm are not

transparent to 100 Kev electrons. Representative microstructures of individual powder particles are

shown in Figure 1. Three kinds of microstructures were observed. The first is a segregation-free

particle which solidified with no partitioning at the solid/liquid interface. Figure la. The second

particle solidified in a more-or-less conventional cellular solidification mode. Figure Ic. The

powder particle in Figure lb exhibits a mixed solidification mode with both a segregation-free and

cellularly segregated regions. These types of microstructures have previously been reported in fine

powders of Al-Si and Al-Cu alloys [1,2].

While the structures in Figure 1 are not restricted to any particular size range or copper

content, some general trends were established. At the lower end of the size range (below 0.5 pm)

and compositions below ''' 2 wt.% Cu most of the particles solidified without apparent lateral segregation. With increasing particle size and/or copper content the microstructures in Figure lb and Ic become predominant. No attempt was made to deduce a statistical correlation between size, composition and mode of solidification.

These observations can be interpreted in terms of bulk supercooling prior to nucleation and

interfacial temperatures and velocities during solidification. Partitionless solidification requires

an interfacial temperature below the T^ temperature of the alloy and an interfacial velocity in excess of the diffusive velocity of solute in the liquid in front of the interface. The latter

conditions can only be achieved through substantial bulk supercoolings prior to solidification. In the absence of such supercoolings, the rate of external heat extraction limits achievable solid/liquid

interfacial velocities to below that required for partitionless solidification [2,4].

Figure 2 shows the proposed solidification "path" of a droplet with no discernible segregation.

L/tt First, the thermodynamic restriction on partitionless solidification is denoted by the T^ curve

on the binary phase diagram. Partitionless solidification of a crystalline solid can only occur

while interfacial temperatures are below this curve. On the other hand, kinetic considerations

require interfacial velocities in excess of the diffusive velocity of the solute in the liquid. It

has been shown that velocities in excess of several meters per second are achievable if the droplet

supercools to a temperature Tj^ prior to nucleation. The supercooling required for the complete

partitionless solidification of the powder particle in Figure 2 has previously been defined by the

following equation (1):

C dT > AH^ (1)

p - M

^ ^

L /(V where T^, is some temperature below T^ (2).

The solidification "path" is then that depicted by the arrows in Figure 2. After nucleation

the evolving heat of fusion increases the interfacial temperature. The growth velocity, while

decreasing with rising interfacial temperature, remains fast enough to avoid solute partitioning even

at the end of solidification at T^.

The mixed mode of solidification observed in Figure lb was studied by establishing the composition profile across an Al-4.5% Cu powder particle. The results are shown in Figure 3. The points

on the electron micrograph indicate locations where the microanalysis was performed. The data are plotted on the right-hand side of the figure as a function of distance from the interface delineating

the plain-front solidified region from the cellular. The most important finding was the depletion

of copper at this interface. This observation verifies earlier hypothesis of solute depletion

between partitionless and cellularly solidified structures [1,2,4].

Figure 4 shows the phase diagram and the postulated solidification path of the powder particle

with a mixed mode of solidification. The lower region of the powder underwent partitionless solidification while both interface temperature and average droplet temperature recalesced to T^ along

the vertical "path" shown. The solute poor region of the particle shown on the right-side of this

figure, which is documented in Figure 3, was due to solute partitioning as the interfacial velocity

slowed down. The solid and liquid interfacial compositions are indicated by the dashed lines. Solute

97
93
HYPERCOOLED

O

o

CO

o

CM

O

O

O

o <

3UniVd3dV\l31

99

WT

%

Cu

WT

%

Cu

Figure

2.

Schematic

phase

diagram

of

the

Al-Cu

system

showing

the

solidification

"path"

of

the

rapidly

solidified

Al-3

wt.%

Cu

droplet

illustrated

on

the

right.

Partitionless

solidification

occurred

during

recalescence

from

T.,

to

T^.
100

Figure

3

Copper

concentration

as

a

function

of

distance

from

the

interface

between

the

planar

(partitionless)

and

cellular

(segregated)

regions

in

a

particle

containing

an

average

of

4.5

wt.%

Cu.
/cpdT <

AH

Oo

COo

CVJ

O

I- oo

3dniVU3dVV31

Q) r— U

<T5 n3 <D

•1- M- SZ M- 3 US- 4-> O O 0) m 0) -C -P 05 -P ^ -p S- c C <D . OJ •!- o to -P C 4-> r— C 5 T- C QJ 03 O

•1- -C

LD "D -P 05

O . Q) JC. C -P ^ +j o^+J 1 U •«- fO > 0) 1— •«- -C O 3 < a. o) E "O

O) -C c

0) TD -P -iC

C ^ •!- C o o

-p 0) 5 o •r- P •!- -P 4-> M- m c -P •c- 03 o o •>- to 05

l/l •!- -PO 0) = 0) -P P CL S- ^ T3 (0 fO E 05

-P

OU CL

O 03

(t) E •>-

u to CL +- >» = C *1- -Q T3 r- O “O •1- OJ C "O O S- o -P CU O" CU

•r- fO

O 5 -p

+->

U «/)

O 1— 03

(T3 •!- I—

(j <4- to 1— “O

•r- •(— to

O c sz

M- -a <D «4- fO -P r” “O »— C to o 5

•1-

O

O n3 1— to *f- 5 r— CU O -P o *o

to r~ •!-

C to o

03 -P

O E u •>- S- Vx: -P 03 -P s4-> C Q. 0) •(- 03

OJ s- 5 •-

O) 3 O o c O" -P to “O r~

•f- <U •!- 0) r— 5 to OJ ^ 3 CU O 3 -P C U

-C

U "O •r— to a> • -P fO

s- to o c E ^ t- o c

<u hu •«- •P >, o

U1 +J -P

C -O

>, . -rO CU

(/) E Z <-•.-

03 1— +-> «4- 3 05 S- 03 o -P E U “O 1 s- X O •f— CU s- r— < Q) M- •r— O

s- “O 1/5 CU 4-> CU 03 -C 5 u r— C

•p

C o c O o o r— 03 to •r— M- U O) o “O CU to M- 03 03 -C CU O S- -p E +-> p- r— •r— 03 03 03 CU -P t—

S- iu T3 to •r- CT-P • 03 O 03 -Q

03 to S- E r— 03

•r- 3 o. P o f— 03 05 -P 03 to

S- c C -C c

CU 05‘*- o P •r— to O S- S- 03 -P S- 3 M- • ^ r—

-C CUu t: 1 to 03 CL r— 03 03 U

Q. E to

C

C •1—

uO i CU 03 05 • r- S- C •!- 1— O

+J o O -P a. r—

03 S- 'C- •o o E >^-P U 03 03 CU o u o r— -C a.

CU •— •f— to su r— 03 03 o

uo 03 03 > 3 -O E

01

S101
3dniVU3dl/\l31

102

Figure

5

Schematic

phase

diagram

of

Al-Cu

system

and

electron-transparent

micrograph

of

a

rapidly

solidified

Al-4.5

wt.%

Cu

submicron

droplet.

The

particle

solidified

with

no

initial

bulk

supercooling.
103

Figure

6

Electron

micrographs

of

as-solidified

and

heat

treated

particles

of

Al-6

wt.%

Cu
rejected ahead of the interface resulted in morphological instability and in the cellular solidific^.

tion noted in the upper portion of the powder particle.

Figure 5 shows a powder particle that solidified in a more-or-less conventional manner without V!

any bulk supercooling. In this case solidification was initiated at or close to the liquidus temperature. Interfacial breakdown early in the growth process lead to cellular solidification with n

lateral segregation. The solidification "path" in this figure is that usually expected in slower ff, conventional sol idi f icaton processes with equilibrium at the solid/liquid interface. ftj

Some of the powder particles were heat treated by the electron beam in the electron microscope,

to study precipitation from the supersaturated solid solution. Results obtained are shown in Figure 6 The particle has 6 wt.% Cu and its microstructure contains cells in about 50 percent of its volume. .. When heat treated, two kinds of precipitates appear. The 0' precipitates which appear in the parti- *’ tionlessly solidified volume are thin platelets which form on the {100} planes. Their diameters vary

between 50 nm and 400 nm. The other kind, probably 0 precipitates, forms from the cell boundaries

which are richer in Cu. These precipitates too, form on the {100} planes. They are round, 25 nm thick, and about 75 nm in diameter.

Acknowledgment

This work was sponsored by the Defense Advanced Research Projects Agency.

References

1. C. G. Levi and R. Mehrabian, Met. Trans. 13A , 13 (1982).

2. M. Cohen and R. Mehrabian, in these proceedings.

3. J. Perel , J. F. Mahoney, S. Tayler, L. Shanfield, and G. Levi, in these

proceedi ngs.

4. R. Mehrabian, International Metals Reviews 27(4), 185 (1982).

.!: ,N )'

rl

'.tJ

104
AN EXPERIMENTAL TECHNIQUE FOR DETERMINING THE DEGREE OF

PARTICLE MELTING IN A PLASMA JET OPERATING AT LOW PRESSURE

R. W. Smith

General Electric Gas Turbine Division

Schenectady, NY 12345

M. Paliwal and D. Apelian

Drexel University

Philadelphia, PA 19104

ABSTRACT

Low pressure plasma deposition, a recent advance in plasma spray metallizing, is now

being developed for high performance materials applications. A study of particle melting

and solidification in the plasma jet, operating at low pressure, will allow one to identify the effect of the process variables and material properties on the resultant structures.

In order to study and model the particle melting and solidification phenomena, experimental

techniques must be developed to measure the progression of the particle melting that is

occurring. In tandem, the experimental results are being utilized in the development of a model describing particle melting in ^apid Solidification by Hasma Reposition (RSPD).

Two methods of determining the degree of powder melting in a plasma jet have been

evaluated. The first method intercepts the particle path in the jet with a glass slide and

the other method samples the jet with a nozzle, quenching the powder before it impacts the

walls of the collector. The splats on the glass slide are studied metal lographical ly using

optical and scanning electron microscopes. Changes in splat appearances are documented for

the alloys Fe-20 wt% Mn, and Hastelloy Alloy X. The plasma sprayed powders trapped by the

collector are evaluated for surface changes via SEM and depth of melting via TEM.

1 . Introduction

The plasma spray process has been used as a method of producing rapidly solidified materials since

1963; reported structures are typical of RSR processing (i.e. amorphous phases, metastable phases,

extended solid solubility...) [1-3]. However, due to the entrainment of oxygen and nitrogen in the

deposits and the high process heat flux causing deposit recrystallization, the plasma process has not

previously been used extensively to produce RSR structures for high performance applications. The

recent development of low pressure plasma spraying [4] and its subsequent use in the deposition of

superalloys [5] has brought the plasma spray process into the forefront in RSR processing and has been

named Rapid Solidification by Rlasma Deposition (RSPD).

The addition of a low pressure environment to the plasma spray process has introduced several

advantages, making possible deposits of highly alloyed materials which can be used for high performance

applications. Some of these advantages are: i) inert environment; ii) pre-heat/cleaning ability of

plasma transferred arc; iii) process flexibility for spraying alloy particles having different melting

ranges, and iv) deposition size/uniformity. Work is now underway to evaluate and optimize this plasma

deposition process for producing high quality superalloy deposits which can be used as preformed RSR

materials or near net shape structures. This is particularly beneficial in alloy systems which are difficult to cast and forge.

In order to successfully and repetitively produce deposits of good quality, the basic controlling

mechanisms of the RSPD process must be understood. Variables which affect the deposit structure (presented by others in this conference [6-8]) include: i) plasma jet enthalpy and velocity (determined

by gun design, power level, and arc gas); ii) powder characteristics; iii) powder injection and iv)

substrate variables (distance, temperature, thermal properties and geometry). The objective of our

study has been to isolate the processes of melting and solidification through experimental techniques

and to study each process separately. The experimental results are used to identify the mechanisms of

105
melting and solidification during RSPC processing. These results are utilized in the development of

models describing particle melting and solidification during RSPD.

The predominant requirement for the production of dense high-strength deposits via RSPD is that

a large fraction of the particles be in a molten state when they impinge on and impact the substrate

(or the previously deposited particles...). Isolating the melting and solidification phenomena in thej

plasma spray process has been and still is a difficult task. The degree of particle melting, the mod^

of powder particle melting, the final splat thickness, the particle cooling rate, and the quenched

microstructures can be masked by the recrystallization caused by the extensive heat flux from the

plasma jet. Recrystallization of the deposits, especially in simple alloys, can be quite extensive

due to the low contamination in the low pressure plasma deposits. In addition to deposit recrystal 1 i-J zation, the other confounding variables in the process are the non-uni form powder sizes, non-unifonn

heat flux to the particles, and non-uniform powder trajectories. Experimental techni.,ues used to isolate the melting and solidification phenomena in the RSPD process are subsequently discussed.

2 . Experimen tal Proced ure

Two different plasma jets were utilized in this study; i) the I^a 3 plasma jet was an EPI Ar/ne

jet at 80 kw, Argon lib liters/min. Helium 40 liters/min, and an EPI03CA-80 anode; ii) the Ma 2.4

plasma jet was an Erl Ar/He ( 125/30 liters/min ) at 72 kw and an EPI03CA-109 anode. Both jets opera^

ted in a low pressure, 40 torr, argon atmosphere and incorporated internal anode powuer injection.

Glass microscope slides were positioned at various distances from the plasma gun exit, low powdei^

feed rates (<o0 gram/min) v/ere set, and the plasma gun was rapidly (>30 cm/sec) traversed over the

slides. The slides intercepted varying numbers of powder particles, but all the deposits were less

than a single particle thick without major overlaying of particles. The glass slides were then viewed}

via both optical and scanning electron microscopy (SEM). The powders studied were argon atomized

Fe-2U wt» t'in (averages of lb and 29um; classified), Hastelloy Alloy X (-400, -325+400, -270+325 and

-lOC+270 mesh), and Rene 00 (-400 mesh). Sampling distances ranged from 20-40 cm downstream from the

nozzle exit.

Another technique to sample the powders from the plasma jets was a powaer collector whose aperture

could be positioned in the plasma jet. The collector was a 45 mm ID by 95 cm long stainless steel J tube topped with a 90° stainless steel cone with a 3 mm aperture port drilled in the center, ana a a

stainless steel collector cup at the bottom. This collector was inserted into the jet at 40 cm axial

j| distance to collect powders over a 5 min deposition time. During this time powders deposited onto the.* cone surface and a small fraction entered the collector aperture and cooled during flight oefore - impacting the collector bottom 95 cm downstream. Hastelloy X, Rene 80, Fe-20 wt% Mn and Fe-20 wt% Cu ' powders were sprayed and collected by the collector at 300 and 400 mm, using the two plasma jets.

Evaluation of these small samples required a modified Field and Fraser technique [9]. The powders

v/ere spread thinly onto a replica tape and vapor deposited with a thin silver layer. These collected

powder sai.iples as vyell as the "virgin" atomized powder were then evaluated for surface morphology via

SEM. The tape was then sandwiched betv/een two washers, plated with 0.1 mm of nickel, the replica

tape dissolved off, mounted in a metal lographic mount, and finally diamond polished and etched. This

technique provides good edge retention and allows a good view of the cross sections. The metallographic evaluation was accomplished using both the SEM and transinission electron microscopy on carbon

replicas. The carbon replica technique has some shortcomings caused by the etch "grooves" around the

edge of the powder particles, but good resolution of powder particle interiors were obtained.

"Sweet-spot" (where the plasma gun is not traversed) deposits onto substrates were also made

to study the microstructural features (unmelted particles, splat thickness, porosity, etc.) of

actual built-up deposit. These deposits were made onto 15 cm x 15 cm plates positioned under the tvio different plasma jets, 30-40 cm downstream. Using spraying times of 1-2 minutes, 2-4 mm thick

deposits were made on preheated substrates (800-900°C).

3. Resul ts and Discussion

Glass Slide Experiments

The results of the glass slide experiments are used to study the morphological changes of individual particles which have interacted with the plasma jet and subsequently impacted the glass substrate. Particle morphologies for Hastelloy X powder plasma sprayed using the Ma 2.4 gun are shown

in Figure 1, Figure la shows the splatting of a partially melted particle which was in a "slushy"

state when it impacted. Examples of particle fracturing along the melted interdendritic region in a

105

I
partially melted particle, and the splatting of well melted particles can be seen in Figure lb. A large number of spherical particles much smaller than the starting original powder can be seen in

the micrographs. These smaller spherical particles result both from the spheriodization of the edges

of the splats as well as in-situ atomization of the particles and their subsequent resolidification

within the plasma jet. The in-situ atomization of the particles is more conclusively determined from

the results of the powder collector which are discussed below. In addition. Figure lb shows solidified

streaks caused by the branching of molten splats on impact.

The glass slide experiments can also be used to qualitatively assess the degree of particle

melting achieved using different gun to substrate distances, plasma gun conditions or the starting

particle size. Optical micrographs comparing melting of two different powder sizes for Fe-20 wt% Mn plasma sprayed using the Ma 3 condition are shown in Figure 2a (average particle size 29 ym) and 2b

(average particle size 16 ym). More uniform melting achieved using a finer size powder is evident

when comparing the two micrographs. The white circular areas seen in the micrographs are the well

splatted particles whereas the dark black areas represent the unmelted and partially melted particles.

(a) (b)

Figure 1: SEM photographs of Hastelloy X powders showing varying amount of melting and splatting upon

impact on a glass slide substrate.

(a) (b)

Figure 2: Optical photographs of Fe-20 wt% Mn powders plasma sprayed onto glass slides using the

EPI Ma 3 gun. a) Average particle size 29 ym, b) Average particle size. 16 ym.

107
Particle Col lector

Figures 3a to 3f are SEM micrographs of "virgin" and plasma sprayed powders of Fe-20 wt% Cu,

Fe-20 wtX Mn, Hastelloy X and Rene 80. Figure 3a is a SEM photomicrograph of the surface morphology

of the virgin Rene 80 powders. It can be seen that the argon gas atomized powders are mostly spherical

and quite regular in shape. A few particles which have been impacted or are splatted with another

particle, during atomization, can be seen. Figures 3b and 3c show the general distribution of morphol- . ogies of the plasma sprayed Rene 80 and Hastelloy X powders, respectively. When compared to the

virgin powders. Figure 3a, the plasma sprayed powders of both alloys show extensive amounts of fine

powders, agglomerated fine particles, fractured powder particles, and various other shapes which have

been formed in the plasma jet or during collection. Figure 3c shows an example of how a coarser screen fraction of powders has been "atomized" into a distribution of powder that has a large volume fraction

of particles below lOym. The probable origin of these fine particles could be i) interdendritic

fracturing of powders and their subsequent spherioidization into smaller particles, ii) detachment of \ the liquid layer from surface melted particles, and iii) secondary atomization of molten droplets

into smaller particles. Figures 3d through 3f are examples of the fracturing and spheroidization that i has occurred in Hastelloy X, Fe-20 wt% Mn and Fe-20 wt% Cu powders sprayed with the higher enthalpy jj Ma 2.4 plasma jet. Figure 3e shows that the Fe-20 wt% Mn powder, having the narrowest melting range,

j; was more irregular and less spheroidized than the wider melting range powder, Fe-20 wt% Cu, shown in ; Figure 3f. Comparison of surface morphologies of plasma sprayed and collected particles using this

[ evaluation technique has also shown that the Ma 2.4, higher enthalpy plasma jet, results in better

j and more uniform melting of the injected powder particles than the Ma 3 jet.

j Figures 4a, 4b, and 4c show polished cross-sections of the virgin Fe-20 wt% Mn, Rene 80 and

Hastelloy X, respectively. The SEM photomicrographs of the virgin powders show that Fe-20 wt% Mn and Rene 80 have solidified dendri tical ly and the Hastelloy X powders solidified cellularly. Figures

5a and 5b are cross-sectioned microstructures of collected plasma sprayed Rene 80 and Figure 5c is a cross section of a Hastelloy X powder. Microstructural features of the collected powders further

augment and confirm previously discussed observations such as the fracture of particles along melted

interdendritic regions.

Figures 6a and 6b illustrate that the collector technique may be trapping a large fraction of i completely melted particles; therefore not measuring the progress of the melting process. These

figures are carbon replica TEM photomicrographs of an as atomized and a plasma sprayed powder particle,

respectively. The details of the carbon replica evaluation technique show these particles to differ

in dendrite arm spacing and interdendritic MC carbide size. The finer structure of the plasma sprayed

powder particle indicates that it has cooled faster than the atomized powder. Moreover, the smaller j- , dendrite arm spacing suggests that this irregularly shaped particle was at one time completely melted.!

j Because the particle was completely melted, all information regarding the progression of melting has f

been lost. This shortcoming for studying melting of powders can be resolved by using shorter collection

distances to catch the particles earlier in flight. Modifications are being made to the collector to ^ include a high temperature ceramic nose cone. I j

The Fe-20 wt% Cu and Fe-20 wt% Mn powder cross sections have not been shown here because their 1

atomized and subsequent plasma sprayed dendritic structures showed no significant differences, suppres^

sing any conclusions about the progress of melting. Because of this observation, it is believed these!

powders melted before collection and then cool at a rate similar to the rates of cooling in the I I atomization process. To be able to distinguish the microstructures of the plasma sprayed and collected

powders from the miscrostructure of the as atomized powders, the starting powder will be pre-annealed

to eliminate the initial segregation structure. jP

Deposi ts I

Microstructures of "sweet spot" deposits made using Fe-20 wt% Mn and Rene 80 powders are shown I

in Figure 7. The Fe-20 wt% Mn powder (average size 29)jm) deposited using the Ma 3 gun are essentially!

unirolted but heated to a temperature sufficient to cause deformation of the particles upon impact I (Figure 7a). In contrast, the same powder plasma sprayed using the higher enthalpy Ma 2.4 gun exhibitsV

a structure where the particles are well melted and splat upon impact (Figure 7b). Rene 80 powder S deposits. Figure 7c, consist of splats of melted particles, partially melted particles, and fragments t of larger powder particles. In these deposits, particles impact on top of one another, bonding, inter-j

diffusing, and recrystallizing. These deposits supplement the structural knowledge obtained from

the glass slide studies, particularly in terms of observing the individual particle morphologies

and their interaction (overlayering) as the deposit builds up. The deposits can thus be used to measure the degree of particle melting and splatting as a function of radial position in the plasma

jet.

108
109

Figure

3:

SEM

micrographs

of

"virgin"

and

plasma

sprayed

powders,

a)

"Virgin"

Rene

80

powder.

Plasma

sprayed

powders;

b)

Rene
no

Figure

5:

Microstructure

of

plasma

sprayed

powders,

a)

and

b)

Rene

80

powder,

c)

Hastelloy X

powder.
(a) (b)

Figure 6: Replica TEM photomicrographs of Rene 80 powder, a) As received and b) plasma sprayed.

Summary

It has been shown that both the glass slide and the powder collector experiments can successfully

be used to study the effects of the plasma jet processing variables and those of material parameters

(particle size and physical characteristics) on the extent of particle melting.

The results also indicate that the melting of powder particles inthe plasma jet is not a classical homogeneous sphere melting problem. A large number of spherical particles smaller than the

starting original virgin powder have been observed on the inserted glass slide substrates as well as from the particle collector. These smaller particles are generated from particles fracturing along

melted interdendritic regions; melting and spheroidization of the fractured particles; and/or

secondary atomization of the melted droplets within the plasma jet.

The powder collector experimental technigue is a very useful tool for trapping droplets in flight

and for examining the morphological changes experienced by the parffcles such as fracturing and

spheroidization. However, monitoring the progress of powder melting (via the study of the microstructural features of the collected powders) is limited by the capability of differentiating between the

resolidified features of the collected particles and the solidification structures of the original

powder. In these cases, the injected particles should be treated to erase the features of the initial

atomization process - i.e., inject pre-annealed or amorphous particles.

Acknowledgements

The research reported herein was sponsored by DARPA Contract F33615-81-C-5156 and the authors

would like to thank Col. L. Jacobsen and Air Force Project Monitor Dr. H.C. Gegel for their support.

Ill
(c)

References

[1] M. Moss: Acta. Met., 1968, Vol . 16, No. 3, pp. 321-326.

[2] K.D. Krishnanand and R.W. Cahn: "Properties of Plasma Sprayed Aluminum-Copper Alloys" in

Rapidly Quenched Metals , Eds. N.J. Grant and B.C. Giessen, pp. 67-75, MIT Press, MA, 1976.

[3] R.W. Cahn: "Rapid Solidification by Plasma Spraying" in Rapid Solidification Processing:

Principles and Technologies I, Eds. R. Mehrabian, B.H. Kear and M. Cohen, pp 129-139,

Cl ai tor's Publishing Division, Baton Rouge, LA, 1978.

[4] E. Muehlberger: "A High Energy Plasma Coating Process" in Proceedings of the 7th International

Thermal Spray Conference , pp. 245-256, Welding Institute, Abington, Cambridge, U.K., 1974.

[5] M.R. Jackson et al . : J. of Metals, 1981, Vol. 33, No. 11, pp. 23-27.

[6] R.W. Smith, D.V, Rigney and J.R. Rairden: "Plasma Deposition Effects on Structure and

Properties of RSPD Superalloys", Third Conference on Rapid Solidification Processing,

Gaithersburg, MD, Dec. 1982.

[7] M.R. Jackson, R.W. Smashey and L.G. Peterson: "Structure and Properties of RSPD Materials",

Third Conference on Rapid Solidification Processing, Gaithersburg, MD, Dec. 1982.

[8] T.L. Cheeks et al . : "Temperature Effects on Plasma Deposition of Ni-Base Materials ", Third

Conference on Rapid Solidification Processing, Gaithersburg, MD, Dec. 1982.

[9] R.D. Field and H.L. Fraser: Met. Trans., 1978, Vol. 9A, No. 1, pp. 131-134.

112
AN ATOMISTIC MODEL OF SOLUTE TRAPPING

Michael J. Aziz

Division of Applied Sciences

9 Oxford Street

Harvard University

Cambridge, MA 02138

ABSTRACT

An atomistic model for nonequilibrium solute segregation during rapid solidification

is presented in terms of a single unknown parameter, the interfacial diffusivity D-. The

results for stepwise growth and for continuous growth at the same interface velocity differ

quantitatively but exhibit the same qualitative features. A transition from equilibrium

segregation to complete solute trapping occurs over approximately an order of magnitude in

growth velocity, as the interface speed surpasses the maximum speed with which solute atoms

can diffuse across the interface to remain ahead of the growing crystal. This diffusive

speed is given by D^/x, where X is the interatomic spacing, and is typically of the order

10 meters per second. A solute drag term and an intrinsic interfacial mobility term are

included in an equation which predicts the growth velocity of a planar interface.

Steady-state solutions are presented for the Ag-Cu system.

INTRODUCTION

Experiments have shown [1,2] that the chemical potential of the minor component of a binary

alloy often increases during rapid solidification. This result rules out concepts of local

equilibrium and independence of atomic species. In this paper, a simple model is described which

predicts suppressed segregation as the solidification velocity becomes faster than the diffusive

speed of solute across the interface.

Conventional modelling of interface kinetics uses the same atomic motions to describe

crystallization as those to describe interdiffusion. In such a case, the growth speed u can never

exceed the maximum diffusive speed D^/x, where D- is the diffusion coefficient at the interface and x is the interatomic distance. However, the collision-limited growth model developed by Turnbull and

coworkers [3-10], which has recently been confirmed by Broughton et. al . [11] indicates that for

metallic systems such a scaling is invalid and that u , the speed of sound, should be the only limit

to the crystal growth rate. Consequently the growth speed can easily exceed the diffusive speed by a large factor^ and when it does so in alloys we expect suppressed solute segregation and a diffusionless transformation.

The reason that crystal growth in metals should not scale with diffusion is illustrated in

Fig. 1. A crystallization event consists of an atom shifting its position a small distance to move

from a potential well in the liquid structure to a well on the crystal lattice. Since no bonds must

be broken, the activation barrier for this reaction ought to be quite small*, in the Turnbull

collision-limited growth model this barrier is zero. A diffusive jump, on the other hand, is a different atomic process in this system. Position shifts of a fraction of an interatomic distance do

not accomplish the solute-solvent redistribution which is necessary to avoid solute incorporation

into the crystal. A solute atom must accomplish three things to avoid crystallization; It must

escape from its potential well, as shown by the dashed arrow in Fig. 1, as the energy of the well

increases and the well moves to the lattice site. It must push solvent atoms out of the way. And it

must be replaced by a solvent atom on the lattice site, or else it will be pushed right back to where

it came from.

Atoms in the interface region tend to interdiffuse in order to establish equilibrium

compositions on either side of the interface. As the interface passes, solute-solvent redistribution

only occurs during a limited time. Before the interface arrives, the driving force for

redistribution is zero; afterwards, atomic mobilities are too low to permit significant

interdiffusion. When u << D^/x, there is plenty of time for the atoms at the

113
interface to equilibrate. As the growth velocity becomes greater, solute and solvent atoms have less

time while they are in the interface region to interdiffuse and maintain local equilibrium. When

u >> D/X, the atoms have no time to interdiffuse; the interface passes and locks the liquid

composition into the bulk of the solid before the atoms have a chance to react.

Fig. 1 Choreography of Solute Trapping.

In order to avoid incorporation into the

crystal (solid arrow), the "B" atom must:

1) ESCAPE from its potential well (dashed

arrow) as the well gains energy and moves

to the lattice site (solid arrow),

2) PUSH "A" atoms out of the way,

3) BE REPLACED with an "A" atom on the

lattice site.

coefficient with dimensionless velocity,

dilute solution limit. Dashed lines;

stepwise growth model (equation 1). Solid

lines: continuous growth model (equation

2). ko is k for a dilute solution; i.e.

kg at the melting point of pure solvent.

SOLUTE TRAPPING MODEL

The mathematical derivation and extensive discussion of the solute trapping model are published

elsewhere [12]. A summary is given here of the assumptions that go into the modelling.

Stepwise Growth Model . If growth occurs by the rapid lateral passage of steps of height X, then

the average time between the passage of steps is Z= x/u, where u is the growth velocity. The time

it takes solvent atoms to shift to their lattice positions is negligible when compared to f , and the

shifts can be treated as virtually instantaneous. The interface is thus assumed to be planar and

executing instantaneous jumps of length x separated by a time interval T= x/u. The jump of the

interface is assumed to incorporate one liquid monolayer, including all solute atoms , into the

crystal lattice. Solute diffuses back into the liquid during the interval diffusive-style rate equations give the composition in the solid monolayer at the interface as a function of time. (The

composition of the liquid at the interface is assumed to be constant and given by the solution to the

diffusion equation in the bulk of the liquid.) Solute remaining in the solid monolayer at t = is

permanently buried into the bulk of the solid, where diffusion is negligible, by the advance of the

interface across the next monolayer. The result for a dilute solution of B in A is

kB(u.T) = kg(T) + [l-kg(T)] exp(-D./uX), (1)

where T is the interface temperature, the equilibrium segregation coefficient k (T) is the ratio of

the solidus composition to the liquidus composition at the undercooled temperature T, and the

interface diffusivity D^ is the only fitting parameter in the model. The result for the

stepwise-growth model is plotted in Fig. 2, note that kg->k when u << D^/x and kg->l when u >> D^. /X.

114
Continuous Growth Model . If continuous solidification occurs, the impurity atom is being

dragged toward its lattice site during a time period of order t. The details are complex and

unknown, but in the limit of strict steady-state fluxes on a microscopic scale, the problem

simplifies. The fluxes and concentrations become time-independent and the same diffusive-style rate

equations as before yield, for a dilute solution of B in A,

kg(u,T) = [(uX/D.) + kg(T)]/[(uX/D.) + 1] (2)

This function is also plotted in Fig. 2.

GROWTH VELOCITY

Equations (1) and (2) for solute trapping require information, either theoretical or experimental, for the velocity-undercooling function. In many cases the assumption of no interfacial

undercooling is a useful one, but the assumption becomes unrealistic at high velocities and solute

concentrations. A dissipation-theory treatment of the interfacial undercooling problem for alloys is

used below.

For pure systems, the process is easily understood. In equilibrium, the crystallization

reaction and the backward (melting) reaction are equal in magnitude. When there is a finite driving

force for crystallization, the former exeeds the latter by a Boltzmann factor exp(Ay/RT); the growth

velocity is given by their difference. The kinetics are lumped into an intrinsic

growth-velocity-limit, Uo(T), which is the speed with which the interface would propagate at infinite

driving force (i.e. when the back reaction disappears). The growth velocity is then written:

u = Uo(T)[l-exp(Au/RT) ]. This becomes the familiar u = M aG at small driving forces. Inverting the

above expression, we have Ay = RT ln(l-u/uo). The left-hand side is the driving free energy; the

right-hand side can be thought of as the free energy dissipated by the irreversible crystallization

reaction at the interface. Thus the velocity-driving force relation can be written AG^p = AGp, where

aGqp is the driving energy, expressed as free energy per mole, and AGp, the "interface^friction

energy" due to a finite interfacial mobility, is the free energy dissipated per mole of material

solidified. We say that the available driving force is dissipated by the irreversible processes at the interface.

For alloys, we can write AG^p = z (3)

where the are the losses due to all of the dissipative processes going on simultaneously. In the steady state, this must hold: the driving free energy does no work on the surroundings and so it

must be entirely dissipated by irreversible processes during the transformation. If we can calculate

the dissipation and the driving force as functions of velocity, we have an equation which predicts

the growth velocity. Ignoring heat conduction, the two dissipative processes at the interface are the aforementioned crystallization and redistribution reactions. The dissipative terms associated

with them are, respectively, the "interface friction energy" AGp and the "solute drag" aG^ [13,14].

By assuming interface friction to be the same for alloys as for^pure systems, we can calculate these

terms [15]. The results are AGp = RT ln(l-u/uo) and, for the dilute solution continuous growth model

(eq'n 2), aG^ = X^RT[(l-kg)/kg]^ln[ko/kg(T) ]. The driving force is given by [14]

aGqf = X^Aug^+ (l-X^)Auy^;^ the growtn velocity equation becomes

AGgp(u) = AGp (

) + aGq(u) (4) It turns out that one of the dissipative terms is usually negligible compared to the other; the

transition occurs over a very small velocity range. Consequently the equation can be simplified away

from the transition. For e-beam melting experiments on Ag(Cu), with interface velocities in the

centimeter-per-second to meter-per-second range, aG^ dominates most of the time; for laser annealing

of doped Si with meter-per-second velocities, AGp almost always dominates. During a diffusionless

transformation in any system, aG^ vanishes.

The dilute solution approximation is invalid when the mole fraction of solute is no longer

negligible or when activity coefficients are no longer constants. When the appropriate corrections

to the fluxes and driving forces in the continuous growth model are made [15], the results are

kg = [(uX/D.)+>io]/[(uX/D.)+l-(l-Xo)X^.], (5)

AGg = X^[(l-kg)/kg]RT ln[kg/k^/io], (6)

where Ko(X

3

,Xj^. ,T) = exp[-A( ug-y;)/RT], u'(X,T) = u(X,T) - RT In X = u°(T) + RT In ^(X,T), X^. i: Am

the mole fraction in the liquid at the interface, k. = X /X . , A = solute, B = solvent, and a

- "solid minus liquid". The implicit assumption has Been made^that at the interface the transition

state behaves in an ideal manner. In practice, temperature and composition dependence may show up in

Di.

115
RESULTS

Equations 4, 5, and 6 were solved numerically; some of the results follow. For the Ag-Cu

system, the thermodynamic model of J. Murray [16] was used to obtain the chemical potential of Ag and

of Cu across the phase diagram. The following parameters were used: X = 2A, Uo = 340000 cm/s,

D- = .00122 exp[-(10 kcal /mole)/RT] cm^/s. The results are plotted as interface temperature and

cimpositions on the phase diagram in Fig. 3; we see that for a given composition X . in the liquid

at the interface, the solid composition increases and the interface temperature dropi as the velocity

is increased. In a steady-state e-beam melting experiment [17], the solid composition is constrained

to be equal to the bulk liquid composition. The velocity is independently controlled; the dependent

variables are X • and T. The way they behave is shown in Fig. 4. As the velocity is increased^ the

liquid composition at the interface decreases due to solute trapping, and the interface temperature

actually rises, trying to follow the liquidus. When the To line is reached, a diffusionless

transformation becomes thermodynamically possible. However, diffusionless solidification does not commence until some distance beyond the To line, when the velocity is great enough to suppress the

redistribution reaction. Fig. 5 illustrates that in the model for the Ag(Cu) system, two effects

facilitate trapping at high concentrations. The interface diffusivity drops by a factor of two as we

go from the melting point of silver to the eutectic temperature, thus making the redistributive

reaction more difficult while not sensibly affecting the growth mobility. In addition, the driving

force for redistribution is smaller at high concentrations.

Acknowledgements . The author is indebted to E. Nygren and K.F. Kelton for assistance with the

computer calculations; to J.A. Murray for the thermodynamic modelling of the Ag-Cu and Si-Bi

systems; to J.W. Cahn, S.R. Coriell, W.J. Boettinger, and F. Spaepen for much helpful discussion;

and to D. Turnbull for guidance. This work was supported in part by an Allied Corp. Fellowship Grant

and by National Science Foundation Grant DMR-79-23597.

Fig. 3 Interface temperature and solid

composition for fixed X

^ in Ag(Cu). For a given liquid composition'x . at the

interface, the solid composition increases

and the interface temperature drops as the

velocity is increased.

MOLE FRACTION COPPER Fig. 4 Interface temperature and liquid

composition at interface for fixed X^.

When the To line is crossed the

diffusionless transformation becomes

thermodynamically possible, but it does not occur until some distance beyond.

Fig. 5 Segregation Coefficient

vs. Velocity for Dilute and

Concentrated Alloys of Cu in

Ag. In this system, two effects facilitate trapping at higher concentrations. The

interface diffusivity drops and

the driving force for

redistribution drops.

116
REFERENCES

[1] J.C. Baker and J.W. Cahn, in Solidification (ASM, Metals Park, Ohio, 1970).

[2] C.W. White, S.R. Wilson, B.R. Appleton, and F.W. Young Jr., J. Appl . Phys. 738 (1980).

[3] D. Turnbull, Contemp. Phys. 1^, 473 (1969).

[4] D. Turnbull, in Sol idi fication (ASM, Metals Park, Ohio, 1970).

[5] D. Turnbull, J. de Physique-Col loque C-4, p. 1 (1974).

[6] 0. Turnbull and B. G. Bagley, in Treatise on Solid State Chemistry , ed. N. B. Hannay (Plenum,

New York, 1975), vol . 5.

[7] F. Spaepen and D. Turnbull, in Rapidly Quenched Metals , 2nd International Conf.,

ed. N.J. Grant and B.C. Giessen (MIT Press, Cambridge, Mass. 1976).

[8] P. Chaudhari and 0. Turnbull, Science 199 , 11 (1978).

[9] F. Spaepen and D. Turnbull, in Laser and Electron Beam Processing of Semiconductor

Structures, ed. J. M. Poate and J. W. Mayer (Academic, New York 1981).

[10] D. Turnbull Met. Trans, m, 693 (1981).

[11] J.Q. Broughton, G.H. Gilmer, and K.A. Jackson, Phys. Rev. Lett. 1496 (1982).

[12] M.J. Aziz, J. Appl. Phys. 1158 (1982).

[13] J.W. Cahn, Acta Metal 1. 10, 789 (1962).

[14] M. Hillert and B. Sundman, Acta Metal 1. 11 (1977).

[15] M.J. Aziz, to be published.

[16] J.A. Murray, to be published.

[17] W.J. Boettinger, these proceedings.

117
TEMPERATURE EFFECTS ON PLASMA DEPOSITION OF Ni-BASE MATERIALS

T.L. Cheeks

Cornell University

Ithaca, New York 14853

M.E. Glicksman

Rensselaer Polytechnic Institute

Troy, New York 12180

M.R. Jackson and E.L. Hall General Electric C operate Research and Development

Schenectady, New York 12301

ABSTRACT

Temperature of deposition has been varied from -500C to the melting range for Ni, NiAl,

NiCr and a superalloy, IN-100, in structures produced by rapid solidification plasma deposition.

This was accomplished by controlling the spacing between the plasma gun and the substrate on which deposition occurred. Considerable difference in microstructural response was noted for the

different alloys. Much of this difference is attributable to grain growth occurring in situ, due to continued heating by the pdasma beam for the less heavily alloyed compositions. During

deposition at a given distance, a continuous change in temporature measured in the substrate was

observed, due to the evolution of the latent heat of solidification of porticles impacting the

substrate. The rate of change in tempjerature is a function of the arrival rate of molten particles

and the distance the substrate is from the pdasma gun.

Introduction

Plasma spx’aying has been used to produce coatings on substrates to imp>rove resistance to wear or aggressive environments (1,2). The process involves melting individual p>owder particles, propelling the

molten droplets toward the substrate, and solidifying the droplets at the surface. Because of spx’eading of the

dropdet at the surface, very great heat extraction rates are atteuned, and rapid solidification occurs (3,4). When deposition is performed in a low-p>ressure environment, coating microstructure is significantly impx*oved

in density and cleanliness comp>ared to deposition performed in air or under a protective gas shroud (5-7). Low-pressure deposition of structural alloys, rather than coating compositions, to produce free-standing

bodies has been termed rapid solidification pdasma deposition (RSPD) (8,9). The microstructure of such

deposits 8U*e typically fine-grained as-sp>rayed, chemically homogeneous, and consolidated to nearly full density. One of the major features of this rapid solidification technique is the ability to produce structurally

sound components without the need of a sep>arate consolidation step (10).

Although the RSPD process is a single step operation, conceptually it can be considered as a series of unit operations. These operations include injection of powder into the plasma arc or jet, transfer of heat and

momentum from the gases to melt and transport the perticles, and impact and solidification of the droplet

against the substrates (11).

The p>resent study treats the last of these unit operations, addressing the interrelationship of substrate

surface temperature and RSPD microstructure in Ni-based aUoy deposits. Since structures and properties of

heat treated deposits may depend on the as-deposited microstructure, it is important to determine if the

initial structure can be controlled, and if it can, by what procedure. Substrate temperature might be

expected to affect microstructure by influencing the rate of heat extraction and the degree of droplet

spreading before solidification begins. The fine scale at which these microstructural changes may occur

requires observation by transmission electron microscopy.

Exporimental Procedure

Temporature measurements on substrate surfaces during exposure to the plasma were obtained using

.050 cm diameter Pt-Rh thermocouple wires that were threaded through previously drilled holes to the front

and back surfaces of either Cu substrates of varying thickness or 10 x 10 x 1.25 cm stainless steel substrates.

118
The thermocouples were spot welded to the front surface and insulated with a both-ends-open ceramic sleeve

which was locked into the hole by upsetting the steel surface by hammer peening. The substrates and

thermocouples were attached to a fixture which provided stability for the substrate in the plasma stream and

also served as a means of connecting substrate thermocouples to other connections leading to the outside of

the 150 cm diameter low-pressure chamber. Measurements were recorded using computer-aided devices.

Atomized powders of Ni-7.5 wt% Al, Ni-13.5 wt% Cr and a Ni-based superalloy (Ni-15 wt% Co, 10 wt%

Cr, 3 wt% Mo, 5.5 wt% Al, 4.5 wt% Ti, .04 wt% Zr, .015 wt% B, .05 wt% C), as weU as pure Cu, were

prescreened to -400 mesh (37/cm or smaller) before spraying. Powder was weighed before filling the powder

feeder and unsprayed powder remaining was weighed after spraying in order to determine feed rates and

powder usage. A measure of the substrate weight before and after spraying then allowed a determination of

deposition yield.

Data from the temperature measurements were taken for gun-to-substrate distances (g-t-s-d) of 20, 25,

30, 37.5 and 45 cm. The measurements began after the process parameters such as voltage, amperage,

pressure and gas flow rate were stabilized with the plasma plume off the substrate. Deposition conditions

used in all experiments were 52V, 1300A, Ar-He plasma in an EPI, Inc. plasma gun, using a 20° powderinjection angle. Initially the chamber was pumped to 0.5 torr, the arc was ignited, and the vacuum pump was

throttled to maintain a pressure of 60 torr due to the plasma gases. With the plasma plume centered on the

substrate, temperature measurements were record^ once a second and continued until the surface

temperature rise decreased to 1 K/sec. This temperature was denoted "temperature before deposition of

powder." Once this increase was achieved, powder injection into the plasma gases began and continued for 55 seconds for most experiments. This corresponded to a powder mass flow of approximately 12 kg/h. The

temperature achieved the instant deposition ceased was denoted "temperature after deposition of powder." A

temperature vs. time plot was generated using computer-aided graphics to compare the different results.

After deposition the substrates were sliced into sections which encompassed the "sweet-spot," an area

of the substrate upon which the deposits were the most homogeneous and uniform, and which varied in diameter from 2 to 5 cm, as g-t-s-d was varied. The sections were subjected to both optical and transmission

electron microscopy (TEM), as well as density measurements.

Optical samples were cut, polished and etched using Marble's Reagent. TEM samples were cut and

ground to .0075-.0125 cm thickness and mechanically punched into 3 cm diameter disks. The disks were

electrochemically polished using a solution of perchloric acid and methanol.

Results and Discussion

Temperature

Typical time-temperature curves are shown in Figure 1 for deposition of Cu on Cu substrates. The

curves show a form characteristic of aU depositions made, whether Cu or Ni alloys deposited on Cu or

stainless steel substrates. As the plasma initially heated the substrate, a rate of temperature rise at the

exposed surface of ~i0 K/sec was observed. Eventually the substrate was hot enough that radiative losses to the nearly-black body chamber walls became significant. The rate of temperature rise then fell rapidly. For

the purposes of this experiment, an effective steady state, heating from the plasma nearly equal to radiative

loss, was defined when the rate of temperature rise feU to 1 K/sec. At that point, powder was injected into

the plasma gases for melting and transport to the substrate. Again, a rapid rise at the substrate surface was

Figure 1. Temperature change with time for Cu deposited on Cu for variation in a) substrate thickness or

b) deposit thickness.

119
observed. Heat transfer now included not only the plasma gas contribution, but also contributions from

heated solid (unmelted) particles and from the latent heat evolved by solidification of melted particles. The

rate of temperature rise measured at the substrate surface again fell as the exposed (deposit) surface grew

away from the substrate surface and the total mass being heated increased. At the termination of deposition,

the plstsma was extinguished and the substrate-deposit composite cooled due to continued radiative loss.

As seen in Figure la, substrate steady-state temperature varies with substrate thickness, although the

larger effect is on time to reach steady state. Substrate temperature is also dependent on the total amount

of deposition. For very thick deposits, thicker than shown in Figure lb, the substrate may show a decrease in temperature during deposition, as the system reaches a new equilibrium between heat input and output and

the exposed surface moves away from the substrate surface.

The effect of gun-to-substrate distance (g-t-s-d) is shown in Figure 2. Steady-state temperature is a

stronger function of g-t-s-d than is the temperature achieved during deposition. It can be assumed that the

heat input from latent heat is essentially constant over the range of g^^s-d. The fraction of that heat lost by

radiation will decrease with increasing g-t-s-d because of the lower T^ losses from less plasma gas heating.

As g-t-s-d is increased, not only steady-state temperature decreased, but the time to reach steady state is increased significantly (Figure 3).

The steady state temperatures are actual surface measurements, while the substrate temperatures

during deposition are being measured further and further from the deposit surface, as noted earlier. For Cu

substrates, front-to-back substrate surface gradients are -^10 C/cm, while for Ni and stainless steel substrates, gradients of as much as 75 C/cm were observed. At any g-t-s-d, Cu substrates were hotter, but

the differences were <50C. Heat capacities are nearly equivalent for these substrate materials, but Cu has

about five times the thermal conductivity of stainless steel and Ni. Even for the largest gradient observed in the substrates, the difference between substrate surface temperature and deposit surface temperature is likely to be small. The deposits of this study were typically ^^.5 cm thick, so that the deposit surface is expected to within 40C of the measured substrate temperature.

Figure 2. The effect of gun-to-substrate

distanceon maximum temperature

achieved before and after deposition.

AVERAGE TIME (SEC)

Figure 3. Time to reach maximum temperature before deposition for various

g-t-s-d.

Microstructure

The deposits consist of a composite containing small fractions of unmelted particles interspersed in a melted and rapidly solidified matrix. Microstructure of the rapidly solidified regions is shown in Figure 4 for both longitudinal and transverse sections. The thin foil in the longitudinal section is parallel to the plane of

the substrate surface, while the transverse section is perpendicular. Thus, the spheric^ liquid droplets spread

on impact (12) and solidify rapidly. Grains are equiaxed columnar in nature. For the powders used, the

largest liquid droplets shoidd be 37yt<m in diameter. The maximum measured splat thicknesses of l-3yw.m

indicate a spreading to the order of lOOy/m prior to solidification, assuming no mass loss due to splashing of the droplet.

Temperature gradients can be approximated from the temperature measured and the splat thicknesses.

For the heavily alloyed Ni-base material, with a liquidus temperature of 1355C, a maximum temperature

120
Figure 4. Longitudinal and transverse TEM views in RSPD Ni superalloy.

during deposition of 1164C ^d splat thickness of 3 leads to a temperature gradient of (1355C1164C)/.0003 cm, or 6.4 x 10 K/cm, while for a maximum temperature of 731C, the gradient is 2.1 x 10”K/cm. A lower limit of growth rates can also be estimated from the measured splat thicknesses. For a powder-feed rate of -^20 kg/h, assuming aU particles are 37yum in diameter, approximately a million particles per second enter the plasma. For the degree of sf^eading noted above, any point on the substrate is covered

about 350 times/sec. For 3yUm splat thicknesses, this corresponds to a solidification rate of at least 0.1 cm/sec, if each splat solidifies just as the next one spreads. Assuming solidification occurs well before

spreading of the next droplet, the solidification rate is greater than 0.1 cm/sec.

From this upper estimate of particle solidification times and measured substrate temperatures, cooling

rates can be estimated. For 350 particles/sec, a solidification time of ^^.0029 sec is obtained^ Again, for a

liquidus of 1355C and a maximum substrate temperature of 1164C, a celling rate of >6.6 x lO^K/sec results,

while for a maximum temperature of 731C, a cooling rate of >2.2 x lO^K/sec results. The degree to which

these cooling rates underestimate the actual rates is related to the overestimate of solidification time.

7 2 For the temperature gradients, G, and solidification rates, R, a G/R ratio of 0.7 to 2.2 x 10 K/sec/cm

results. Since a ratio of ^5 x lO^K/sec/cm ^ is sufficient to achieve planar growth (13,14), the columnar,

equiaxed grains in Figure 4 are, as expected, free of any dendritic or cellular solidification features.

Although the substrate temperature range is quite broad, it results in small ranges in G and therefore in G/R. As would be expected, the variation in splat thickness is also quite small, less than a factor of 2, between the lowest and hipest substrate temperatures. The variation in deposit temperature from beginning

to end of deposition is smaller than the variations due to g-t-s-d, so similarly the variation in splat thickness

from first deposition to last is insignificant.

Other features of the microstructure were alloy dependent, such as fraction of particles unmelted and

degree of microstructure modification after solidification. The simple alloys, Ni, Ni-13.5 Cr and Ni-7.2 A1

have liquidus temperatures in excess of 1400C while the heavily alloyed composition has a liquidus of 1355C.

There were considerably more unmelted particles in the materials with the higher liquidus temperatures, in spite of a plasma gas nearly 10 times hotter than the metal liquidus. This was manifested in deposit density

as well. When compared to cast standards, the materials with large fractions of unmelted particles were ^93% as dense as their cast counterparts, while the heavily alloyed material was fully dense (Figure 5).

Comparing the temperatures measured with the optical micrographs of Ni-7.5 wt% A1 and Ni-13.5 wt96

Cr revealed structures that had experienced an in situ anneal. This was evident from the enlarged grains ("6

yum) that are normally not apparent in plasma sE»*ayed materials ('^.^Ajn) (Figure 6). The annealing was

attributed to the heat of the plasma which in the 20 cm g-t-s-d case caused melting of the deposit during

deposition, resulting in a cratered contour. The contours in general, became broader and more uniform as the

g-t-s-d increased. The grain growth due to the anneal was attributed to the low-aUoy content and lack of

grain boundary inhibitors such as carbides, as compared to a chemistry more typical of a superalloy. The true microstructure of these alloys could not be correlated to the temperatures measured. However, for the case

of the melted sample there was evidence that a temperature gradient existed between the substrate surface

(T 1325C) and the liquid deposit surface (1400-1450C) through a deposit thickness of M cm.

121
Figure 5. Deposit density for various g^t-s-d

for Ni-Al, Ni-Cr and Ni superalloy.

Figure 6. Grain size in as-deposited material

as a function of substrate temperature.

Summary

Deposition temperature, as estimated from the substrate temperature, is dependent on distance from

the pdasma source, substrate thickness (mass), and powder-deposition rate. Heating is due in part to the

plasma gases and in part to the heat involved during re-solidification.

Spreading of liquid droplets varies less than a factor of 2 (in thickness of the spread disc) fro^ about 2.5

j^^m to 4/Am, for the range of cooling rates achievable in RSPD. Cooling rates are ~'10 K/sec, and

temperature gradient, solidification rate ratios (G/R) are such that planar solidification occurs.

Degree of particle melting is strongly dependent on liquidus temperature. For simpler alloy

chemistries, in situ annealing can occur for the higher deposition temperatures. This is noted by a loss in remnant splat structure and a growth in grain size.

Acknowledgement

The helpful technical discussions with J.R. Rairden, M.F.X. Gigliotti, D. Lee, F. Zaverl and A.M. Ritter

(aU of General Electric Company) are gratefully acknowledged. Considerable experimental skills were

provided by W.A. Seaman, S.F. Rutkowski and J.A. Resue. R.W. Kopp was responsible for designing

temperature measurement equipment.

References

(1) T.A. Roseberry and F.W. Boulger, "A Plasma Flame Spray Handbook," Rpt. No. MT-043, Naval Sea

Systems Command, Dept, of the Navy, March 1977.

(2) J.R. Rairden, "A Study of the Properties of Hi^-Velocity Arc Plasma Sprayed NiCrAlYC Coatings"

Proc. 9th Int. Thermal Spray Conf., The Hague, p. 329, 1980.

(3) R.W. Cahn, "Rapid Solidification by Plasma Spraying," Proc. Int. Conf. on Rapid Solidification

Processing, Reston, Va., p. 129, 1977.

(4) H. Herman, "Metastable Phase Produced by Plasma Spraying," Fall Meeting Program, TMS-AIME,

Pittsburgh, p. 22, Oct. 1980.

(5) R.W. Smith, "Mechanical Properties of a Low-Pressure Plasma Applied CoCrAlY Coating," Proc. Int. Conf. on Metallurgical Coatings, San Francisco, 1981.

(6) J.R. Rairden, M.R. Jackson, M.F.X. Gigliotti, M.F. Henry, J.R. Ross, W.A. Seaman, D.A. Woodford and

S.W. Yang, "Coatings for Balanced Environmental/M echanical Behavior," Proc. Int. Conf. on HighTemperature Corrosion, NACE, San Diego, p. 72, 1981.

(7) E. Muehlberger, Proc. 7th Int. Thermal Spray Conf., London, p. 245, 1973.

122
(8) M.R. Jackson, J.R. Rairden, J.S. Smith and R.W. Smith, "Production of Metallurgical Structures by

Rapid Solidification Plasma Deposition," J. of Metals, p. 23, Nov. 1981.

(9) M.R. Jackson, R.W. Smashey and L.G. Peterson," The Mechanical Behavior of RSPD Material," this conference proceedings.

(10) J.R. Rairden, M.R. Jackson, J.R. Ross and W.A. Seaman, "Plasma Sprayed Alloys Considered for

Advanced Concept for a Gas Turbine Bucket," Thin Solid Films, 73, p. 37, 1980.

(11) M.G. Benz, General Electric Co., Schenectady, NY, private communication, 1981.

(12) J. Madejski, "Solidification of Droplet on Cold Surface," Int. J. of Heat and Mass Transfer, 1^, p. 1009.

(13) R. Mehrabian, "Relationship of Heat Flow to Structure in Rapid Solidification Processing," Proc. Int. Conf. on Rapid Solidification Processing, Reston, Va., p. 9, 1977.

(14) H. Matyja, B.C. Giessen and N.J. Grant, "The Effect of Cooling Rate on the Dendrite Spacing in Splat

Cooled Aluminum Alloys, J. of the Institute of Metals, p. 30, 1968.

123
SUPERSATURATED SOLID SOLUTIONS BY MASSIVE

SOLIDIFICATION: OBSERVATIONS AND IMPLICATIONS*

Thomas F. Kelly, Morris Cohen, and John B. Vander Sande

Department of Materials Science and Engineering

Massachusetts Institute of Technology

Cambridge, Massachusetts 02139

Abstract

Massive solidification structures have been observed in centrifugal ly

atomized and rapidly solidified bcc powder particles of a 303 stainless

steel. These structures are solid solutions of uniform composition throughout. A most important characteristic here is that the as-grown solid

solution is highly supersaturated. Normally, large manganese sulfide precipitates form during cellular or dendritic solidification of this steel.

Since the solubility of the liquid state is generally much greater

than that of the solid state, massive solidification to supersaturated

solid solutions offers great potential for development of new alloy

compositions and alloy design concepts. These possibilities and some processing conditions that favor massive solidification are discussed.

* Sponsored by the Office of Naval Research, Contract Number

N00014-81-K-0013, NR653-032.

Introduction

303 stainless steel**is an austenitic steel with approximately one volume percent sulfide

precipitate due to an intentionally high sulfur content. Table 1. Owing to the low solubility

of the solid iron-base matrix for sulfur, partitioning of sulfur during solidification is severe. As a consequence, in conventional ingot processing sulfides, principally of manganese,

precipitate from the enriched liquid at dendrite cell walls during the latter stages of

sol idifi cation.

As the solidification rate increases, it is found that the size of these intercellular

precipitates decreases as the size of the cells decreases. At the most extreme solidification

rates, solute partitioning is kinetically suppressed and homogeneous solid solutions are produced in a massive solidification. These solid solutions are highly supersaturated with

precipitate forming elements even at the moment of solidification which suggests novel

processing paths to precipitate-strengthened alloys.

Table 1

Fe Cr Ni

bal

17.3 8.68

- Composition of 303

Mn Cu 3i

1.60 0.78 0.66

Stainless Steel in

Mo S Co

0.37 0.34 0.71

weight percent.

C N P

0.034 0.032 0.022

**The 303 stainless steel was kindly supplied by Carpenter Technologies Corporation, Reading, Pennsylvania, which had originally requested the rapid solidification processing run.

124
Microstructural Observations

Light metallographic examination of ingot-processed (cooling rate < 10"^ K/s) 303 stainless

steel reveals large, 10 to 50 micron diameter manganese sulfide precipitates strung out in the

rolling direction. Figure la. These are the remnants, after hot rolling, of the more nearly

spherical precipitates formed by liquid-phase precipitation at dendrite cell walls in the ingot.

Chill casting of this steel (cooling rate 10^ K/s) reduces the precipitate size to 1 micron

diameter.

Centrifugal atomization with forced convective cooling* (cooling rate 10^ to 10® K/s) has been

used to produce rapidly solidified powder of 303 stainless steel with powder particle sizes in the

range 5 to 200 micron diameter. Interestingly, a body-centered cubic (bcc) crystallization phase

of the 303 stainless steel nucleates and grows in the smaller powder particles as a result of

changing relative nucleation rates of fee and bcc with increasing liquid supercooling [1]. Facecentered cubic (fee) is the usual crystallization phase of this steel at small liquid supercoolings

and is found in the larger powder particles.

The structure of the fee powder particles is invariably found to be cellular with sulfide

precipitates at the cell walls. These precipitates are about 100 nanometer diameter as a result of

rapid solidification but are still formed by a liquid-phase precipitation process as occurs in

ingot processing.

It is in the microstructures of the bcc powders that deviations from this solidification

scenario are observed. The extreme example of this is a homogeneous [2] solid solution produced by

massive (partitionless) solidification with a planar liquid/solid interface. Figure Ic. In this

process, the solid grows with the same composition as the liquid. The solid solution of

Figure 1C apparently decomposes in some cases, during continuous cooling to ambient, to the

structure of Figure Id. Here, lOnm diameter manganese sulfides have precipitated throughout

the structure from the supersaturated solid solution produced during massive sol ification.

An intermediate microstructure between the cellular fee and the noncellular bcc structures

is a cellular bcc structure. Figure le. Larger, 100 nanometer diameter sulfides that precipitate from the liquid are located at the cell walls of the structure. There is, as well, a massive solidification zone in the central region of each of the cells. Smaller 10 nanometer

diameter sulfides have precipitated from the supersaturated solid solution that results from

this localized massive sol ification.

Preliminary studies indicate that the supersaturated solid solutions of Figure Ic may

decompose to ultra-fine, 1 to 2 nanometer sulfide precipitates under controlled precipitation

heat treatments at 400°C.

Discussion

Supersaturation

It is important to know whether the solid solution that forms during solidification is

supersaturated with respect to manganese and sulfur at solidification temperatures. If the

solid solution is not supersaturated, then this is simply the case of a high temperature

single phase of high solubility that has been accessed by massive solidification. An important distinction must be made between massive solidification processes that produce

homogeneous undersaturated solid solutions and those that produce homogeneous supersaturated

solid solutions. Though there might still be some advantages to massive solidification in

the former case, the real importance of massive solidification as a processing route lies in

the ability to produce supersaturated solid solutions that cannot be produced otherwise.

Turkdogan et. al [3] have determined the solubility product of manganese and sulfur in

fee iron. By their results, the equilibrium solubility produce of manganese and sulfur at

the melting point of iron, 1800K, is [Mn (wt%)][S(wt%)]r = 0.0076. The composition product

of manganese and surfur in 303 stainless seel is 0.544 wnich is 72 times this fcc-iron

equilibrium solubility product at 1800K.

* The rapidly solidified powder was kindly processed by Pratt and Whitney Aircraft Group,

Government Products Division, United Technologies Corporation, West Palm Beach, Florida 33402.

125
BC C

(3

Q0® mi

Figure 1 - Process states of 303 Stainless

Steel

(a) light micrograph of longitudinal

section of hot-rolled ingot;

(b) annular dark-field STEM image of

rapidly solidified fee;

(c) and (d) bright-field STEM image

of rapidly solidified noncell ular

bcc

(e) bright-field STEM image of rapidly

solidified cellular bcc.

1^6
Complete solid solutions are observed only in the bcc structures. Although the solubility

of sulfur may be greater in bcc than in fee, this is a small effect compared with the large

amount of manganese and sulfur in 303 stainless steel. Furthermore, a one hour heat treatment

of some rapidly solidified 303 stainless steel at 1600K produced only a small amount of precipitate coarsening: the average precipitate size increased from 10 to 200 nanometer diameter

to 0.5 to 1.0 micron diameter. This indicates a low solubility product of the precipitateforming elements, manganese and sulfur, in this steel. Accordingly, it does appear that a bcc solid solution that is highly supersaturated with respect to manganese and surfur forms

during massive solidification of 303 stainless steel as a bcc structure.

Hypercooling: Fee vs. Bcc

It is not coincidence that massive solidification is observed only in the bcc structure.

There are some compelling reasons why bcc solidification in this steel would be massive more readily than fee solidification. This difference in solidification behavior between the two structures is associated with the degree of supercooling of the liquid prior to nucleation of

the solid.

The onset of massive solidification can be associated with the approach to hypercooling,

which is where the specific supercooling goes to unity:

^AT

AH

L->S

1 .

L“^bcc For 303 stainless steel, the enthalpy of fusion of fee, AH , is calculated to be 19% less

than the enthalpy of fusion of fee, AHL’^'*"cc [i]. This means that, for equivalent supercoolings,

AT, the specific supercooling of the liquid with respect to bcc.

ah'-"''“

is greater than the specific supercooling of the liquid with respect to fee.

P ahL^^"

Cp is the heat capacity of the liquid at constant pressure. For an alloy where the equilibrium

melting temperature of bcc is not too much lower than the equilibrium melting temperature of fee,

as for 303 stainless steel, the liquid will be hypercool ed with respect to bcc at a higher

temperature than fee. Bcc should then be more prone to solidify massively.

Implications and Applications of Massive Solidification

Massive solidification, if it is achieved, offers a means of producing solid solutions in

alloy systems in which it may not otherwise be possible to obtain such solid solutions. This

could have far-reaching implications for alloy design.

Consider an alloy consisting of a base metal plus elemental additions that will form a precipitate in the base metal. Solidification of this alloy in a cellular or dendritic mode

might produce a relatively coarse second-phase morphology. If a phase regime of high solubility

for the precipitate were to exist at high temperature for this hypothetical alloy, then

solution! zing of the alloy could be carried out. The penalty for this might be excessive grain

growth during the solutionizing treatment. If, however, there is no high-temperature regime of

high solubility for the precipitate, then there is no possibility of solutionizing the alloy via

this type of heat treatment.

On the other hand, processing of such an alloy to obtain massive solidification would offer

some advantages. Having achieved the solid solution during solidification, a precipitation heat

treatment could be carried out at a low enough temperature that grain growth is negligible.

Since the solubility of the precipitate phase is very low, the precipitate dispersion will be

very resistant to coarsening even at high temperatures. This fine precipitate dispersion will,

in turn, prevent grain coarsening at the high temperatures. Thus, an unusual precipitationstrengthened fine-grain material would result.

127
An alloy-design approach that utilizes massive solidification may be especially desirable

for the design of materials with a large volume fraction of the second phase. With rapid

solidification processing, alloy designs would not be limited to those that can be successfully

processed by conventional means. New precipitate/metal matrix combinations should be tried on

the basis of their potential effectiveness if the solid solution can be achieved by massive

solidification. Ideally, the supersaturated solution would be retained to room temperature to allow control of subsequent precipitation processes. Optimal precipitate volume fractions,

sizes, and dispersions could then be achieved with this type of processing.

Conclusions

Massive solidification of bcc structures, which occurs in the smaller (,^ 70 micron diameter)

droplets of 303 atomized stainless steel, results in a supersaturated solid solution with respect

to the precipitate forming elements, manganese and sulfur. This supersaturated solid solution

cannot be produced by conventional solid-state processing of this alloy.

As material s-processing route, the use of massive solidification to produce supersaturated

solid solutions appears promising as a means of structural control in a wide range of material

systems.

Acknowledgements

Sponsored by the Office of Naval Research, Contract Number N00014-81-K-0013, NR 653-032.

References

1. Thomas F. Kelly, Morris Cohen, and John B. Vander Sande, "Rapid Solidification of

Atomized Stainless Steel: Part 2 - Origins of FCC and BCC Solidification Phases,"

submitted to Metal 1. Trans. A , 1982.

2. Thomas F. Kelly, Morris Cohen and John B. Vander Sande, "Rapid Solidification of

Atomized Stainless Steel: Part 1 - Depressed Partitioning and Supersaturated

Solid Solutions in FCC and BCC Solidification Phases," submitted to Metal 1

Trans. A, 1982.

3 . E. T. Turkdogan, S. Ignatowicz and J. Pearson, J. Iron and Steel Inst. , 1955,

vol . 180, pp. 349-354.
SOLUTE REDISTRIBUTION DURING SOLIDIFICATION

OF SUPERCOOLED ALLOYS

S. D. Peteves, D. D. McDevitt and G. J. Abbaschian

Department of Materials Science and Engineering

University of Florida

Gainesville, FL 32611

ABSTRACT

During some rapid solidification processes considerable supercooling is achieved such

that nucleation takes place below the solidus temperature. The solidification in such cases

is initially partitionless, resulting in the formation of a solid with the composition of the

liquid. A simple model based on the heat of fusion, heat capacity, and the supercooling has

been proposed to calculate the fraction that solidifies without segregation. Calculations

based on this model are compared with STEM measurements of the composition profile in supercooled Fe-25%Ni structures. The results indicate qualitative, but not quantitative, agreement

between the two. The discrepancy is explained based on the non-uniformity of temperature, and

interfacial curvature and kinetic effects. A model for the associated heat transfer equations

is given, which indicates that solidification is kinetically controlled at the beginning of

recal escence , but becomes thermally controlled as recalescence progresses.

Introducti on

The conventional techniques of obtaining rapidly solidified structures are generally based upon

increasing the cooling rate during solidification by improving the heat extraction efficiency between

the casting and the heat sink, and/or by reducing the volume of material being solidified. However,

rapid solidification can also result from supercooling prior to nucleation, followed by rapid heat re- moval. With substantial supercooling, the initial solidification rate is large as the heat of fusion

is absorbed by the supercooled liquid. In fact, a supercooling exists beyond which the liquid can solidify completely without the need for transferal of heat to the surroundings. Substantial supercoolings before nucleation have been obtained by melting high purity metals within molten glass [1,2],

within emulsions [3,4], or by using a levitation technique [5,6]. The latter technique has been re- cently used [5] to study the effects of supercooling and cooling rate on the solidification structure

of Fe-25%Ni samples. The samples, approximately 0.5 cm in diameter, were supercooled as much as 270 K and were solidified in water, molten lead, or a ceramic mold. The sub-grain microstructure was found

to fall into three categories: dendritic, spherical, or a mixture of the two. All three morphologies

were observed when the samples had solidified with a supercooling less than 175 K. At larger supercoolings, however, the spherical morphology predominated. In addition to the spherical morphology,

other segregation patterns were noted at large supercoolings [6].

Numerous studies have been done to relate solute distribution and microstructure to the amount of

supercooling prior to nucleation [7-lD]. The most widely used model, proposed by Kattamis [7], is based on the relationship between the heat of fusion, heat capacities, and the supercooling. The

model indicates that when a liquid is supercooled below the solidus temperature prior to nucleation,

the fraction that solidifies segregationless is approximately:

where C_ is the average heat capacity of the solid and liquid, is the heat of fusion, T^ is the

solidus^temeprature, and T^ is the nucleation temperature. As the temperature recalesces above the

solidus, partitioning begins at the interface and solidification continues with the solid decreasing

in solute until the maximum recalescence temperature is reached. Beyond this point the solidification

continues similar to conventional solidification processes with no supercooling. The equations relating the composition profile to the fraction solid during and after recalescence are given by:

- ^s)^f 0 ^s

= ^0 C m ^°^^S<^S<^SR

P s

129
c 1 s 'sR

- fsR>

Ik-1

(1 - ^sr)

for < 1

where C-p is the solid composition at the maximum recalescence temperature and where is the fraction solid at maximum recalescence, m^ is the solidus slope, and C^p is the solid composition at the

maximum recalescence.

The present paper investigates the applicability of the above mentioned models to the solidification of supercooled Fe-25%Ni alloys.

Experimental Procedure

The experimental procedure consisted of utilizing electromagnetic levitation to levitate and melt

one gram samples of Fe-25%Ni [5]. The samples were melted while levitated and sunppcooled to a desired temperature where they were dropped in a molten lead bath held at 650 K. The quenched samples

were then analysed metal lographical ly and by a JOEL 200 scanning transmission electron microscope.

For the latter study, the samples were sliced by a diamond saw followed by electrochemical thinning in

a solution of 90% acetic - 10% perchloric acit at 280 K. The composition profile in the thinned samples were obtained by using the electron microscope.

Results

Similar to previous studies [5,6], the microstructure was found to fall into three major categories: dendritic, spherical, or a combination of the two. At supercoolings less than 175 K, all three

morphologies were frequently observed while at supercoolings exceeding this value, the spherical morphology tended to predominate. The metal 1 ographi c examination further revealed the presence of a fine

dendritic segregation pattern within the larger dendrites, as shown in Figure 1. Surrounding the solute-rich "primary" dendrites is a lightly colored region which gradually darkens toward the edge of

the larger dendrite arms. In general, there is a single primary dendrite arm within each larger arm. However, several primary arms were observed within a larger arm in some areas.

Figure 1. The dendritic morphology

etched to reveal fine "primary"

dendrites within the dendritic arms.

The sample was supercooled 175 K before quenching in molten lead.

Samples having the spherical segregation pattern exhibited similar features, containing soluterich cores, as shown in Figure 2. However, the solute-rich primary cores are not fuTly developed dendrites but have a simple cross shape. The crosses are surrounded by a lightly colored, solute-poor

region, darkening toward the edge of the spherical elements.

Figure 2. a) Brightfield microscopy

of the spherical morphology showing

dark colored solute-rich crosses in

the center of each element surrounded

by a solute-poor region which gradually

darkens toward the edge, b) crosses revealed using Nomarski interference

contrast.

130
A scanning transmission micrograph of a spherical element, similar to the ones shown in Figure 2,

is shown in Figure 3. The compositional profile along a horizontal line passing through the center

Figure 3. STEM micrograph of spherical

element similar to those in Figure 2b.

of the cell was obtained using an energy dispersive spectrometer. The ratio of the intensities of the

iron and nickel K x-ray peaks was used for compositional measurements. The line scan data of a sample supercooled 8Q K below the solidus line, is plotted as a function of distance from the center of

the cell in Figure 4. The compositional profile based on the model discussed earlier is also shown in

the figure.

Figure 4. Model and measured compositional

profiles: a) and b) Scheil equation and

Kattamis' model, respecti vely, converted to

spherical geometry, c) Line scan of measured

x-ray intensity ratios converted into

concentration.

Di scussi on

The measured compositional profiles and the results of metallographic examination clearly indicate the existence of a solute rich core within the cast elements. This region is surrounded by a re- gion of lower concentration beyond which the concentration increases to a maximum at the periphery of

the cell. The measured compositional profile qualitatively agrees with that of the model but not quantitatively. The experimental results show the region of the partitionless solidification to be

much smaller than that of the model, and the measured minimum concentration to be higher than the theoretical one. These discrepancies are believed to be due to the assumptions or simplification which

were used in the model. These include: i) adiabatic solidification during recalescence; ii) negligible solid state diffusion; iii) no remelting during recalescence; iv) partitioning at T^; v) no curvature effects; and vi ) uniform temperature throughout the sample during recalescence.

131
The assumptions of adiabatic solidification seems to be justified considering the short recalescence time and relatively large volume of the samples about 0.5 cm in diameter. Furthermore, the effect of non-adiabatic conditions will be to increase the extent of partitionless solidification,

rather than to reduce it. The solid state diffusion can also be assumed to be negligible because of

the relatively short solidification time. The local solidification is approximately 0.007 sec for a sample supercooled 80 K below solidus and quenched in lead [5]. The smaller size of the central core,

as compared with the model, may be due to remelting of compositional ly unstable solid above the solidus temperature. However, for remelting to take place at the surface of a growing solid, the interfacial movement must cease completely before retreating. The continuous cooling of the sample during

this period can lower the temperature enough to elminate remelting.

Thermodynamic considerations based upon the Tq concept, detailed by Baker and Cahn [11] indicate

that partitionless solidification may continue above the solidus until the temperature reaches T^.

The T curve is where the free energies of the liquid and solid phases are equal. Above T^, solidification must occur by partitioning if the free energy is to be decreased. If partitionless solidification continues past the solidus to the Tq curve, the volume fraction of the solute-rich core would be

larger than the presented model. However, the actual temperature of the partitionless-partitioning

demarcation is debatable.

The assumption of the planar interface is not justified when the interface breaks down and forms

a dendritic morphology. The dendrites during partitionless solidification are caused by the thermal

field, since the solid is growing into a supercooled liquid. According to Gibbs-Tompson effect, the

interfacial curvature then shifts the phase diagram downward, leading to an earlier transition into

partitioning.

For the sample shown in Figure 1, the size of the primary dendrites is less than 10"^ cm. Although, it is rather difficult to determine the tip radius during recal escence, a rough estimate of

the radius can be made from the dendritic growth models relating the tip radius to the bath sugercooling and growth rate [12]. For 8fl K supercooling, the tip radius is estimated to be 10“^ - 10"^ cm.

The corresponding depression of the solidus and liquidus temperatures, as well as Tq, is in the range

of 3-3f1 K. As a result, the extent of the partitionless solidification will become considerably less

than that predicted from the model. It should be noted that a 30 K temperature depression will reduce

f° to a value close to the observed results.

Non-uniformity of the temperature within the sample during recalescence is another major factor

that could drastically alter the value f°. If the solidification process is controlled by the interfacial kinetics, the interface temperature will remain close to that of the bulk sample. However,

when the process is heat transfer controlled (i.e., the interfacial kinetics are faster than the rate

of the thermal diffusion), the interface temperature will be appreciably higher than the bulk. In this case the interface temperature will eventually rise to the melting point, even though the bulk

liquid still remains supercooled. Partitioning in the latter case takes place much sooner than the

former case. The former case is used in the derivation discussed earlier.

Numerical techniques have been recently used to calculate temperature profiles during solidification of atomized aluminum powders [13]. The results indicate that substantial temperature differences

are developed inside the droplets during recalescence. Obtaining an analytical solution for the asso- ciated heat transfer equations during recalescence is rather difficult. Heat transfer solutions of

supercooled melts to date have been usually concerned with isothermal solidification of a solid at its

melting point into the supercooled liquid. The interfacial temperature is also assumed to be fixed at the melting point and without any kinetic effect. A solution based on these conditions was given by

Frank [lA] for radially symmetric phase growth controlled by diffusion. The solution, however, is not applicable for transitional heat transfer problem during recalescence.

For a simplified case of a spherical solid growing into a relatively large supercooled bath, a quasi-stationary solution for the heat transfer equations can be given by modifying the analysis given

by Frank [13]. Assuming the heat transfer to be by conduction only, the heat capacity of solid to be

negligible compared with the heat of fusion, and the solid to have a uniform temperature equal to the

interface temperature, the temperature profile in the liquid ahead of the interface is;

T(r,t) = T|^ + S(T,-T,) 1/2

r4 at 1/2

erfc (-

S 1/2

2 (at) 1/2 )}

erfc(|)

1/2 - 1/2 where S = r a t , dimensionless "reduced radius. is a root of
exp (|s^) [exp (^S^) - |

^ S {1-erf (|s)}] = 2 (^-^---) Cp

in which r is the radial distance in the liquid, t is the time, T^- is the

interface temperature, Tm is the nucleation temperature, a is thermal

diffusivity, C is specific heat, and is heat of fusion. The interface

temperature can have any value between the nucleation temperature and the

melting point.

The above equations were used to calculate the heat transfer controlled interfacial velocity as a function of radius at various interfacial supercoolings. The results are plotted in Figure 5, where

the y-axis gives the interfacial velocity while the x-axis gives the radius of the growing solid.

Figure 5. Interface velocity as a function of the ra- dius of solid, growing at various constant interfacial

supercoolings: Solid and

dashed curves show heat flow

and kinetically controlled

growth respectively. The

dotted curve is the demarcation line betwen the two

growth modes.

The above mentioned calculations are based on the assumption that the interfacial kinetics is infinite. In reality, however, the interface kinetics is limited and it is a function of the material

parameters, the supercooling, the orientation and the nature of the interface. Therefore, the heat

transfer equations must be coupled with the kinetic equations. There are only limited number of accurate measurements on the solidification kinetics of metals, and these are mostly at low supercoolings

[15]. However, at high enough suprcool i ngs , the rate equations are expected [16] to be a linear function of the interfacial supercooling given by

gnH.AT.

V

= hRT^

[ where 3 is a correction factor for the structure and orientation, D is the liquid di ffusiti vity, R is

the gas constant, and h is the distance by which a position on the interface advances when an atom is

added or removed. The coefficient 3 is expected to be of the order of ten for symmetrical molecules

and atoms. Using the available data for iron and nickel and 3=10, the coefficient of the linearity is about 2 cm per second per degree of the interfacial supercooling for Fe-25%Ni alloys. The corresponding kinetically controlled solidification rates at various supercoolings are shown by the horizontal

dashed lines in Figure 5.

The preceding heat flow and kinetics calculations indicate the limiting cases for the solidification processes. For a given interfacial supercooling, when the radius of the solid is small, the

kinetically controlled rate is slower than the heat transfer controlled one. Thus the solidification

will be kinetically controlled. The situation is reversed however, as the radius of the sphere be- comes large and the process becomes thermally controlled. However, it should be noted that as solidification continues, the interface temperature also rises. This in turn reduces the kinetically con- trolled rates while it increases the heat transfer rate. Therefore, the various regions shown in

Figure 5, can be divided to two zones separated by the dotted curve and marked by I and II. In zone

I, where the radius of the solid is small and/or the interfacial supercooling is high, the solidification is kinetically controlled. In zone II, where the fraction solid is larger and/or most of the interfacial supercooling is lost, the solidification process is heat transfer controlled. It is in zone

II, that the interface temperature rises rapidly to the melting point and partitioning begins even though the bulk liquid is still supercooled. It should be noted that the time necessary for the so- lidification to change from kinetically controlled mode to the heat transfer controlled mode is rather

short. For example at a 40 K interface supercooling, the time for the transition is estimated at

2 X 10-^ sec.

133
Summary

The solute distribution profile in highly supercooled Fe-25%Ni was m^^^'-red by using STEM. The

results quanti ati vely agreed with the model proposed by Kattamis [7], but not quantitatively. The

disagreement is believe to be due to certain assumptions made in the theoretical calculations; namely

uniformity of temperature during recalescence, no curvature effects, and no interface kinetics. An approximate analytical solution has been given for the associated heat transfer equations. The re- sults are discussed in conjunction with kinetically controlled solidification, which qualitatively

indicate the solidification is kinetically controlled at the beginning of recalescence, but becomes

thermally controlled as recalescence progresses.

Acknowl edgements

The project was sponsosred by the National Science Foundation, under grant no. nMR-8202724 to the

University of Florida.

References

1. T. Z. Kattamis, and M. C. Flemings, "Solidification of Highly Undercooled Castings," Modern

Casting , pp. 97-104, July, (1976).

2. J. L. Walker, "Physical Chemistry of Process Metallurgy, Part I," pp. 845, Interscience, NY,

(1961).

3. n. Turnbull, "Kinetics of Solidification of Supercooled Liquid Mercury Proplets," J. Chem. Phys. , vol. 20, No. 3, pp. 411-24, (1952).

4. J. H. Perepezko, "Crystallization of Undercooled Liquid Droplets," Conf. Proc. 2nd International

Conf. on Rapid Solidification Processing, Reston, VA, March 23-6, (1980).

5. 0. J. Abbaschian, and M.C. Flemings, "Supercooling and Structure of Levitated Fe-Ni Alloys," to

be published in Met. Trans.

6. 0. 0. McOevitt, and G. J. Abbaschian, "Microstructure and Solute Distribution in Highly ; Supercooled Fe-Ni Alloys," to be published in Microstructural Science, vol. 11, 1983.

!' 7. T. Z. Kattamis, "Redistribution of Solute in Highly Undercooled Iron-Nickel Alloy,"

Metallkunde, pp. 856-60, (1970). P

I, 8. T. Z. Kattamis, "Solute Redistribution in Undercooled Alloys," Ph.D. Thesis, MIT, (1965). j:r 9 H. Biloni ^nd B. Chalmers, "Predendritic Solidification," Trans. Met. Soc. AIME, vol. 236, pp. . 615-34, (1966).

10. S. P. Iyer and W. V. Youdelis, "Distribution of Bismusth in Primary Sn Crystals of Supercooled

Sn-Ri Alloy Ingot," J. of Inst, of Metals, vol. 101, pp. 218-20, (1973).

1

. 11. J. C. Baker and J. W. Cahn, "Thermodynamics of Solidification," in Solidification, an ASM | publication, (1971). !

r 12. R. Trivedi, "Growth of Dendritic Needles from a Supercooled Melt," Acta. Met., vol. 18, March,

'

pp. 287-96, (1970).

13. C. G. Levi and R. Mehrabian, "Heat Flow During Rapid Solidification of Undercooled Metal

Droplets," Met. Trans. A , vol. 13A, pp. 221-34, (1982).

14. F. C. Frank, "Rapidly Symmetric Phase Growth Controlled by Diffusion," Proc. Roy. Soc. A, 201,

pp. 586-99, (1950).

15. G. J. Abbaschian, S. F. Ravitz, "Rate Equations for Dislocation-Free and Dislocation-Assisted

Growth of Gallium, J. Crystal Growth , 44, pp. 453-66, (1975).

1 16. J. W. Cahn, W. B. Hillig and G. W. Sears, "The Molecular Mechanism of Solidification," Acta.

Met. , vol. 12, pp. 1421-39, (1964). «

134
MICROSTRUCTURE-PROPERTY RELATIONS AND THEIR ROLE IN RSP MATERIALS

J. C. Williams and J. P. Hirth

Cat iiegie-Mellon University

Pittsburgh, PA 15213

Ohio State University

Columbus, OH 43210

ABSTRACT

The prospect of creating microstructures which are truly optimum for a given application has moved a step closer to reality with the development of Rapid Solidification Processing (RSP) techniques. This paper discusses in general some of the possibilities for

creating optimum microstructures using RSP and considers in greater detail the selection

of microstructures for combined strength and toughness. Finally, a series of unresolved

issues regarding microstructure-property relations are identified and the role that these

issues play in restraining the selection of optimum microstructures is discussed.

Introduction

It has become accepted in recent years that optimum performance of a material in a given application can be paced by any of a variety of properties, depending on the details of the loading and service requirements placed on it. In simple cases, where weight is critical, the component may be sized

using strength as a primary criterion whereas in many other cases, where fracture resistance is critical, toughness or fatigue-related properties may fix the section size long before strength becomes an

issue. Many other comparable situations exist but these are too numerous to describe here. It is

clear, however, that the first step in optimizing the performance of a structural component is identification of the pacing property which dominates the design process. As already mentioned, this property may be structure-sensitive such as strength, toughness or fatigue resistance, or it may be one

of the so-called structure-insensitive properties such as modulus or density. Irregardless of which

of these properties pace the performance of a component, metallurgists have made considerable progress

in learning how to enhance the values of each of these properties. Unfortunately, a material cannot

be altered so as to maximize a given property without due consideration of the concurrent effect that

such alterations may have on other properties.

Among the better known examples of the trade-offs which result from improving one property at the

expense of another is the inverse relationship between strength and toughness. Many other comparable

trade-offs exist in practice and it is these which often limit the extent to which a particular property can be maximized. Frequently these trade-offs and limitations are the result of the metallurgist's

inability to simultaneously control the different microstructural features which affect the various

individual properties. If it were possible to synthesize optimum microstructures using a metallurgical equivalent of the molecular engineering methods used so successfully by developers of polymeric

materials, then some of these trade-offs might be lessened or eliminated.

One possibility for expanding the range of microstructures and alloy compositions which can be obtained in metallic materials is the use of rapid solidification processing (RSP) techniques. In particular, RSP permits fabrication of alloys which contain either greater concentrations of solute or

solutes which are difficult to deal with using ingot metallurgy. Thus, RSP opens new horizons in terms

of alloy composition. In addition, RSP offers the potential for creating finer and even novel microstructures which cannot be achieved in alloys fabricated from ingots. These opportunities thus raise

an important question: "If we could create any microstructure, what would we specify for a particular

application?" It turns out that the answers to this question are more complicated than might be anticipated, especially when fracture-related properties are considered.

135
The purpose of this paper is to discuss the question of specifying optimum microstructures. To do this, we will address only two properties in any detail; strength and toughness. The discussion

of these two properties will be followed by a short summary of key issues which face the metallurgist

before true microstructure synthesis can be attempted with confidence. These key issues could also be

identified as pacing issues or barriers to success in the implementation of RSP in a broader applica-:

tions arena. At any rate, we hope that the remainder of this paper will illustrate an approach to considering microstructure-property relationships in a more fundamental way than has been traditionally done in alloy development programs. We also hope that some food for thought can be found herein,

even for those who do not share our enthusiasm regarding the opportunities which RSP brings to metallurgical practice.

Strength

If one considers the possibilities for improving the strength of metals and alloys, there are basically three general types of strengthening mechanisms which can be employed. These are: solid

solution, particle and boundary hardening. Most articles on strengthening mechanisms list more hardening mechanisms but, after some scrutiny, all of the mechanisms can be grouped in the above three

types. From the standpoint of microstructural design or selection, this is a convenient way to con- sider the possible strengthening contributions since each is readily identifiable and distinct from a 1

microstructural viewpoint. The ensuing discussion will thus examine these three types of hardening,

” particularly in light of some of the effects that each has on other characteristics or properties. It

is not the intent of the discussion to present a rigorous analysis or review of strengthening mechan- , isms but, rather, to examine the role of microstructure on strength with an eye toward providing some M

insight into optimizing microstructures of RSP materials with respect to strength. Readers interested!

in a more rigorous review of strengthening, per se are referred to any of a number of books and reviews]

on this subject[l-6] . S

Boundary Hardening

Boundary hardening refers to the strengthening contribution due to the presence of boundaries

which act as obstacles to dislocation motion; these may be grain boundaries, interphase boundaries in

duplex or microduplex alloys, or sub-boundaries such as cell walls in a polygonized structure. In general, reducing the grain size or slip length of a material leads to increased strength, with these

strength increases being proportional to d"V2 where d is the grain diameter. Smaller grains also

usually lead to simultaneous improvements in ductility and toughness. However, very fine grains

(n. l-5ym) may also reduce the resistance to shear localization, thus an optimum grain size no doubt

exists for each material. The exact value for optimum grain size depends in a complex way on a vari- i ety of factors, including intrinsic work hardening, preferred orientation, slip character and others.

Nevertheless, it seems certain that refining the grain size of most structural materials to values in M

the 1-lOpm range can lead to simultaneous increases in strength, toughness and other properties such

as fatigue strength, although fatigue crack propagation rates may increase and environmentally- m

induced cracking may become more severe. <1

RSP has been used to introduce a uniform distribution of small second phase particles which act J

as intrinsic grain refining agents[7]. The advantage of RSP in this case is that the particles can be

kept small enough that they do not appear to participate in the initiation of microvoids during due- , tile fracture, yet still stabilize the grain size by preventing growth. As a result, the grain size ; refinement can be achieved without the deleterious effects associated with large inclusions or second

phase particles.

When the grain size is refined to diameters less than Ipm, then the strengthening from substructure becomes more important[8] . As a result, this seems to set the low end size limit for grain

size strengthening. However, the stabilization of substructure by fine particles is also a useful

concept and the use of RSP in producing a uniform distribution of fine particles which aid in stabilizing a substructure with a small cell size seems attractive. This approach has not been extensively f

studied in RSP materials, but similar structures have been produced in mechanically alloyed materials

such as ODS Ni-Cr alloys[9]. In these materials the thoria particles are very effective in stabiliz- J ing the substructure. j

Thus, while it is still difficult to use first principles to quantitatively specify an optimum

grain size, it is clear that grain diameters in the 1-lOym range are attractive and potentially beneficial. Such grain sizes are indeed rare in conventional ingot metallurgy products. As a result, the ' use of RSP to develop very fine grained materials appears very attractive. The use of very small, ' stable cells (< lym) for strengthening also appears attractive, especially if RSP can be used to produce a uniform dispersion of particles which will pin and stabilize these cell walls. In the final

analysis, boundary hardening is an attractive means of improving the strength because it typically im- .

136
proves ductility and toughness also. When employing boundary strengthening, however, the earlier

comment on fatigue crack growth rates must be borne in mind. Other property compromises may also

exist, for example, intragranular environmental cracking is usually more pronounced in fine grained

materials.

Particle Hardening

Particle hardening refers to increases in strength due either to dispersoids or to precipitates.

The distinction between these is that the former are typically by-passed by mobile dislocations

whereas the latter are sheared. This difference has a profound effect on the work hardening characteristics of the material which, in turn, affect the ability of a material to resist localized deformation and plastic instability which leads to premature fracture. Alloys which are hardened by shearable precipitates generally exhibit decreasing strain hardening with increasing strength. As will be

discussed below in the section on Toughness, this reduction in strain hardening tends to lead to earlier onset of plastic instability and thus to reduced toughness. The precipitates which form in

many of these alloys during aging have a very uniform size and a small spacing as shown in the examples in Figure 1. Since these particles are usually aligned with the matrix via an orientation re- lation they tend to be readily sheared by matrix dislocations. Extreme examples of this are found

in Al-Li alloys where the coherent 6' phase is sheared, leading to intense bands of shear localization and to the premature initiation of cracks along these bands[10]. Essentially all alloys which

contain shearable precipitates exhibit this tendency and the differences between alloy systems tends

to be one of degree rather than of kind.

Alloys which contain non-shearable particles (dispersoids) tend to exhibit increased strain hardening compared to the equivalent solid solution matrix. The reasons for this are the rapid accumulation of dislocation debris around the particles during the initial stages of straining. The dispersoids also are effective in producing particle strengthening via the Orowan process. In such cases

the initial strength varies as the inverse particle spacing. Thus, a fine, uniform distribution of

dispersoids is a highly desirable means of producing increased strength without accelerating the onset of strain localization. RSP has been used successfully to produce this type of microstructure

in A1 alloyslll] and in other alloy systems [7 ,1 2] , An example of these dispersoids in Al-Mn alloys

is shown in Figure 2. These dispersoids are Al6Mn[13] but a variety of other particles have been

shown to exist in Al-Fe-Ce and in other alloys[ll].

The size of the dispersoid is critical to the performance of these materials for several reasons.

First, if the volume fraction is fixed by solubility limits or other considerations, then a uniformly

small dispersoid size is necessary in order to achieve small spacings and high strengths. Second,

the size of the dispersoid is related to the ease with which it can act as a site for void formation

as discussed in greater detail below. Third, there is a size below which particles do not appear to act as void initiation sites. This size has been calculated to be 14nm[14] for the case of uniform

strain distribution in the matrix. If the strain is concentrated locally due to inhomogeneous slip or because cross-slip is difficult, then this becomes an underestimate. Also, elongated particles

appear to be more prone to decohesion and thus act as more frequent void nuclei. Despite the details

of the calculation, it is helpful to recognize that a particle size limit exists below which the particles contribute to strength but do not act as void nuclei for ductile fracture. RSP has the potential to produce a uniform distribution of small particles and thus may be very beneficial as a means

of producing dispersion hardened materials which are both strong and tough. There is another con- sideration regarding dispersoid size and strength. That is the existence of a minimum size below

which the dispersoid is readily by-passed by local climb or cross-slip. This size is estimated to be

in the range of a few nm. Particles which are smaller than this do not contribute to the flow stress.

Thus, the optimum particle size range for dispersion strengthening appears to be bounded on both

sides; the upper size bound is fixed by the onset of void nucleation at particles while the lower

size bound is fixed by the ease of cross-slip or climb around the particles. These sizes have been

subjected to calculation and within the uncertainties of these calculations because of necessary simplifying assumptions the optimum dispersoid size range lies between a few nm and tens of nm. This is

indeed a fairly tightly defined range. As a result, RSP may be very useful in providing the processing control necessary to produce this.

Solid Solution Hardening

Solid solution hardening is a well-known and useful means of raising the flow stress. However,

solid solution additions can also alter the deformation characteristics of the matrix. For example,

in face-centered cubic materials the slip character can change from wavy to planar if solutes which

reduce the stacking fault energy are added. Such solute additions not only raise the friction stress

but also increase the strain hardening rate. Other solute additions can create short range order

which increases the initial flow stress but tend to also cause strain localization and reduced work

137
Figure 1. TEM micrographs showing fine, uniformly distributed precipitates typical of

most age hardening alloys.

a) 0" in Al-4%Cu, bright field b) u-phase in Ti-12%Mo, dark field

micrograph (photo courtesy of micrograph using u-phase reflecS. Dumolt). tion.

Figure 2. TEM micrographs showing AlgMn dispersoids which have formed in a Al-10%Mn

powder metallurgy alloy (Ref. 58).

a) hot pressed at 400°C b) hot pressed at 500°C.

138
hardening [6] . Thus the choice of solutes to strengthen alloys must be carefully selected if the in- creases in strength are to be achieved without sacrificing the ductility and toughness of the alloy.

There are a number of rules regarding the strengthening contribution of solutes based on size difference, modulus difference, valence, chemical nature, etc.,[cf 3,4,14] but these rules only provide

insight regarding the expected change in stress. They do not consider, in any direct way, the effect

of these solutes on slip character and work hardening rate. It is this latter factor which becomes

important when striving for simultaneous improvements in strength and toughness.

Additivity of Strengthening Contributions

The use of two or more types of strengthening contributions is commonplace in commercial alloys.

However, the methods for summing the contributions from various hardening mechanisms are not welldeveloped, as is mentioned later in the key issues section. On the one hand, work on dilute solid

solutions containing Si02 particles showed that the increased friction stress due to the solid solution additions added linearly to the strengthening due to the Si02 particles [1 5] . However, as the

solute concentration is increased or as solutes which affect slip character are added, the attendant

creation of stronger or more intense dislocation pileups tends to alter the effectiveness of particle

or boundary hardening contributions. In such cases the linear superpositioning of individual hardening contributions can lead to either an underestimate or an overestimate of the true combined

strengtheni ng.

One last example of the compexity of strengthening which cannot be dealt with analytically may be

useful. Figure 3 shows the microstructure of a Ni-Al-Mo-W alloy which contains at least three distinct hardening precipitates in addition to the solid solution strengthened matrix. Figure 3(a) is

a bright field electron micrograph which shows a high volume fraction of cuboidal y' precipitates

separated by a continuous matrix which shows mottled contrast. Selected area diffraction patterns,

such as the one shown in Figure 3(b), show that two more ordered phases exist in addition to y'.

These have DO22 and Pt 2Mo prototype structures. Dark field images of the appropriate diffraction

spots (Figs. 3(c) and (d)) show that these additional phases lie in the matrix regions between the y'

particles. The point of showing this example is to illustrate the complexity which can arise in con- centrated alloys. The current situation with regard to combining strengthening contributions from

three coexisting types of particles is such that analysis is not currently possible. Many other ex- amples such as this exist in practice.

In view of the foregoing brief discussion, it is clear that improved insight into the effects of

multiple hardening contributions on strength must await development of improved analytically methods

for combining the individual contributions. Computer simulation studies coupled with detailed calculations and carefully planned experiments are needed here. The use of RSP to produce more uniform

microstructures may be helpful in this regard.

Toughness

The discussion of toughness which follows considers three central elements which affect the performance of a material with regard to toughness. These are plastic instability, ductile fracture and

brittle fracture. The factors which affect each of these elements will be discussed below.

Plastic Instability

It has long been known that the condition for macroscopic necking in a tensile test is that

^ ^ a Equation 1

For larger strain rates both rate effects and adiabatic heating can be important and the condition for

stable flow becomes [17].

de +

9e £,T

9-g

31 ue T,e

dl ne +

3T

dT ^ ade

e,e Equation 2

In either case, an increase in work hardening rate (3a 3e) clearly enhances the resistance to plastic

instabil ity.

More recently, it has become known that ductile failures often occur as a consequence of plastic

instability in the form of localized shear bands initiating either at the surface or in the bulk[18].

Continuum mechanics theories have been developed for the particular case of a yield surface with a

139
Figure 3. TEM micrographs showing the complex microstructure which exists in a Ni-AI-Mo-W alloy made from RSR powder (Ref. 59).

a) Bright field micrograph showing b) Selected area electron diffraction

large cuboidal y' and mottled pattern showing y, y' , D022> Ni 2Mo contrast in matrix regions. reflections.

c) Dark field micrograph showing

Ni 3 Mo (DO22 ) precipitates.

d) Dark field micrograph showing

Ni 2Mo (Pt2 Mo type) precipitates.

140
corner[20] »and these give good to fair agreement, respectively, with experiments revealing surface

[19,20] or bulk[21,22] initiated instabilities. To illustrate these newer ideas, we present the

critical condition for the onset of surface instability in the form of surface roughness which subsequently initiates shear bands propagating into the interior. The condition is expressed in terms

of the critical strain in plane strain tension as ^ = e* [1 - exp (- 2e*)] Equation 3 "s

where e* is the critical maximum principal strain and Ej and E5 are the tangent and secant moduli defined in Figure 4. Through the dependence on Ej = (da/de) for slow strain rate tests, again the material is stabilized against local instability by an increased work hardening rate.

Thus, so-called diffuse hardening mechanisms which tend to increase the yield or flow stress but

not the work hardening rate are detrimental to ductile fracture in that they increase E$ while leaving

Ej constant, thus enhancing the tendency for plastic instability. Diffuse mechanisms include solid

solution hardening, precipitation hardening, and grain size hardening. Contrariwise, dispersion hardening promotes multiple slip and large work hardening[2] , provided the dispersoid is impenetrable by

dislocations. Hence, if impenetrable and if resistant to cracking or decohesion, dispersion hardening should increase the resistance to plastic instability. Therefore, as also mentioned in other sections of this paper, an ideal microstructure for strength and toughness should involve dispersion

hardening as the primary hardening mechanism with the diffuse mechanisms as secondary ones.

Once localized shear instability ensues, the detailed mechanism of failure is not clearly understood, but failure is rapid once instability begins. Thus, avoidance of the critical condition for

onset of plastic instability is the critical factor for design purposes. Failure can occur by shearing to a chisel-point fracture[23] , by void formation and link-up under the influence of incompatibility stresses induced by the instabi 1 i ty[24] , or by transition from a mixed-mode I-II shear crack

to an unstable Mode I brittle crack which propagates by hole growth, by cleavage, or by a mixed mechanism. Conditions which favor the onset of localized shear instability, and hence ones that should

be avoided, include plane-strain plastic constraint and low work-hardening rate/yield stress ratio.

Ductile Fracture

Beginning with the early work of McCl intock [25] , it has become clear that the resistance to ductile fracture is improved by increasing the resistance to particle decohesion or cracking. Since

either of these events leads to void growth, whether for smooth or notched specimens, ductile fracture resistance is improved when the cohesion of the particle is above a minimum value. Theories re- lating to particle decohesion in simple[26,27] or more complex[28,29 ] ways are available and have been

quite successful in explaining some[30] but not all [28], experimental results. If the cohesive

strength of a particle is below a certain size-dependent critical value, most of the plastic strain to failure is involved with void growth and there is not much gain in fracture toughness with an increase

in cohesive strength. Above this critical value, failure is rapid once void nucleation occurs be- cause the voids enhance shear localization and essentially limit plastic strain prior to failure.

Thus, above the critical value, fracture toughness is enhanced directly by an increase in cohesive

strength[31 ]

Once hole nucleation occurs, failure can occur by necking of the ligaments. Figure 5(a) or by

shear localization between the ligaments. Figure 5(b)[16,21]. In either case, the amount of plastic

work required for crack propagation is limited to a volume of the order of the void spacing. Figure 6. Hence, a decrease in void spacing should enhance toughness as is observed in fractographis correlations of dimple size and toughness. Enhanced resistance of the matrix to plastic instability should

enhance toughness in the case where ligaments fail by shear localization, so the matrix properties

should follow the design precepts of the previous section.

Decohesion or cracking of a particle becomes more likely for large particles[14,l8,32] , particle

with large aspect ratios[14 ,18,33] , and particles with low intrinsic cohesive strengths[14,31 ,33]

The cohesive strength for the reversible formation of a decohesive nucleus is

^coh = Vrt + Tmat - Tint Equation 4 where the surface energies are expressed per unit area and Ypart, Tmat and Tint respectively, the

particle-void, matrix-void and particle-matrix interfacial energies. Values for Ypart and Ymat are available at high but not low temperature for some cases but values of Yint are sparse. Hence, data

or theory is needed before Ycoh can be used as a design parameter. One can crudely say that semi141
conducting particles such as carbides should have lower values of Tint than insulators such as oxides,

a factor which would tend to increase Ycoh 3nd help resist decohesion.

There is some theoretical indication that the critical stress for decohesion should vary inversely with the square root of particle diameter[14,34] . These models are upper bounds for the stress and

a more detailed theory is needed which includes the influence of slip inhomogeneity or dislocation

pileups in nucleating decohesion. Nevertheless, they provide a guideline in agreement with experiment

that decreasing the size of a dispersoid decreases its decohesion tendency. There is a limit roughly

of the order of lOnm diameter below which the particles become effectively penetrable because dislocations can cross-slip or cl imb around the particles[2] . Again, a more detailed theory is needed to better define this limit.

The particle asymmetry effect is associated with increased stresses favoring decohesion arising

from slip inhomogeneity and dislocation pileups. An array of edge dislocation pileups produces larger

stress concentrations than a single pileup, and a large aspect ratio particle would be more likely to sustain more than one pileup along its larger dimensions.

Brittle Fracture

Some progress in understanding brittle fracture and relating it to microstructure has been

achieved, beginning with the work of Ritchie, Knott and Rice[35]. These authors considered that the

critical event was the spreading of a crack from a carbide at a grain boundary one or two grain diameters ahead of a crack tip to link with the main crack. A critical combination of intensified stress

and carbide size (crack size) is needed, as illustrated in Figure 7, to initiate matrix cracking. The

calculation has been modified recently Evans and Hutchinson[36 ] who consider a statistical distribution of carbide sizes (crack lengths). In both calculations, the carbides are considered to be at grain boundaries and to be precracked. Since incompatibility and pileup effects are large at boundaries, and since carbides are often larger there because of preferential grain boundary nucleation,

the assumption seems plausible. The resulting prediction for the critical stress intensity for cleavage crack initiation is [36],

Y (N+D/4

^-l)/2^(N-2)/4(^^)l/4

Equation 5

where Yq is the effective surface energy for crack propagation, Y is the yield stress, r is the average

carbide radius, f is the volume fraction of carbides on grain boundaries, and d is the grain diameter.

The factor N is the reciprocal of the work hardening coefficient n in a power law description

Typically N>3 so all exponents are positive in Eq. 5.

The factors in Eq. 5 can be understood qualitatively as follows. An increase in Y shifts the

stress to a larger value at a given position. Figure 7, and favors fracture. An increase in op obviously represents increased resistance to cracking. An increase in r means an increase in local stress intensity at the edge of the carbide crack which favors its propagation into the matrix. An increase in

f increases the probability of a large carbide at a given position and hence favors fracture. The

weak dependence on d arises from the need for a number of cracks spaced at about the grain diameter

along the crack tip to propagate locally to lead to overall crack advance. The model agrees well with

data of Curry and Knott[37].

The above models are lower bound models in the sense that they underestimate Kjq if the carbide

particles are not precracked as assumed. The likelihood of carbide cracking or decohesion is related

to the magnitude of internal stress concentrations produced by localized slip or dislocation pileups.

These factors have been considered quantitatively[38] but they can be understood in qualitative terms. Pileups should be decreased, and fracture resistance promoted, by: a decrease in grain size provided

that it and not dispersions controls the wear slip path; increased stacking fault energy which promotes pileup blunting by cross-slip; a large number of available slip systems; also promoting blunting;

decreased hardening by solutes or precipitates which would tend to prevent operation of secondary dislocation sources to relieve pileup stresses; decreased dispersoid spacing, which bounds pileup lengths

in the dispersion hardening limit; and decreased particle size, which decreases the likelihood of decohesion or cracking in the limit for reasons already discussed. Particle shape anisotropy or greatly

asymmetric grains promote pileup stress concentrations and fracture: edge dislocations formed against

such asymmetric barriers interact in such a way that the pileup stress concentration is increased by

a factor of d/2,, where d is the pileup length (short grain dimension) and 2, is the pileup (slip band)

spacing.

Another ameliorating factor for the models mentioned above is the possibility that a crack in a carbide may not propagate in a brittle manner into the matrix. At the grain size level, added energy

142
6

Figure 4. Schematic stress-strain curve showing

tangent and secant moduli, Ey and Eg,

respectively.

0 .

Figure 5. Schematic showing ductile fracture

processes which occur in crack tip

region

a) hole formation ahead of crack tip

b) shear localization connecting

holes and crack tip.

Figure 6. Schematic showing relationship between

plastic zone volume and inclusion

spacing, D, ahead of crack tip.

Figure 7. Schematic showing the stress distribution as a function of distance ahead

of a crack tip.

143
must be supplied to create cleavage steps or tear ridges as the cleavage plane changes orientation

upon crossing a grain boundary. At a finer scale, a microcrack in, say, a particle must propagate

j into the matrix locally rather than blunting. The statistics of the probability of this occurrence as'

a function of particle size distribution were considered by Evans and Hutchinson[36] , as discussed ; previously. There are some other factors at the microstructural level that we now consider briefly

for the case of quenched and tempered steel as an example. ’

The orientation relationship between austenite and ferrite is (211 )aj | (001 )cm; (001 ]^| | [lOOJcm*

[111 lal I

Thus, the primary cleavage plane of (OOl)cm is inclined at 35“ or 6b“ to the

54“ to the {106}(j planes. Thus, a carbide particle within a grain is not favorably oriented for

propagation of a crack into ferrite. Neither are such particles favorably oriented for propagation

from ferrite into carbide, a factor that could be important in crack growth. At a grain boundary,

the carbides are oriented with respect to one grain to reduce surface energy. However, they are ran- domly oriented with respect to the second grain. Thus, some fraction of grain boundary carbides are likely to be oriented for cleavage crack propagation and grain boundary carbides can be more detrimental, but on a statistical distribution basis. If recrystallization or grain growth is imposed

subsequent to carbide precipitation as in "process-annealing", the orientation relation would be lost

for many of the carbides, and bulk carbides could become favorably oriented for cleavage propagation.

Thus, in the brittle fracture regime one might be cautious about attributing changes in behavior to grain size changes alone if the changes were accomplished in the above manner.

Carbides have larger volumes per iron atom than the matrix and smaller thermal expansion coeffi- , cients. Thus, residual stresses in the particles from precipitation and cooling should be compressive i in nature and would tend to reduce the brittle fracture tendency for radial cracks. Circumferential ' tensile stresses would exist though, and could promote fracture for precipitates with salient corners. *

If a crack nucleus starts to grow in a particle or grain boundary, it will gain inertia. This f dynamic effect would tend to promote crack propagation into the matrix, but calculations are needed

to determine the magnitude of this effect. If a particle decoheres, this effect should be diminished

because of the non-planar trajectory of the growing crack nucleus.

As mentioned previously, detailed consideration of the notion of microstructural synthesis brings

into focus a number of inadequacies in our current understanding of strengthening and fracture at the

fundamental level. We have elected to present these as "Key Issues" since we feel progress toward

elucidating these will substantially enable true microstructural synthesis.

1. Effect of dispersion size on properties.

a. Yield strength: Orowan theory and experiment verify that yield strength should

increase as particle spacing decreases (hence as particle size decreases for a fixed volume fraction) provided the particles are impenetrable by dislocations [2]

The trend is reversed only below a oarticle size where dislocations can readily

cross-slip or climb around particles. Rough estimates indicate that this critical

size is lOnm, but experiment and theory are needed to verify this.

b. Ductile fracture: Since ductile fracture (initiation or stable crack growth) occurs by particle decohesion (or, less-often, particle cracking) and finer

void spacings enhance Mode I (normal) or mixed Mode I-II (shear) failure[40],

it has been suggested that increased interface cohesion of particles would help

and that a critical size existed, below which decohesion would not occur[41].

Recent work [36] suggests that increased cohesion helps provided that the cohesion

exceeds a critical value which depends on volume fraction, but otherwise that

changes in cohesion have little effect in the sense that void growth tends to control fracture.

Also, the estimate of the critical size for decohesion of ^ 14nm[14] is a lower

bound if the local plastic strain is uniform but is an overestimate if plastic

strain can be concentrated in dislocation pileups. Experiments to follow up on these ideas, some new, are called for.

c. Brittle fracture: For years, the ductile-brittle transition has been little

understood in terms of mechanism. Ritchie, Knott and Rice[35] focused attention on the possibility of a critical stress being achieved for steel at a grain diameter or so ahead of a crack where a cracked carbide particle could act

as a nucleus for cleavage. Yet, many Kjq-T plots show an abrupt change at the

144
DBTT[42] and not a smooth one as predicted by the model, and estimates of the

critical distance vary from 1 or 2 to 10 grain diameters [43 ] . Possibilities

include, R-curve type crack resistance with microcracks extending over grains

but not propagating further[44], the critical event then being linking of a set of these and dynamic pop-in of a crack from a particle to the matrix; or extension of pre-existing cracks at particles into single grains, with the

cracks eventually linking up to form an unstable cleavage crack. Unresolved

issues are the requirements to propagate a cleavage crack across a boundary

where the crack orientation changes (requiring added crack surface area or linking shear zones); the local events needed for particle cracking and whether

these, at the dislocation level, determine the ductile brittle transition

temperature; whether microcracks exist in a zone near the crack initiation point

as suggested by an R-curve-type model; and whether crack initiation indeed oc- curs ahead of the precracked region[45]. Also, a statistical distribution of

particle sizes and strengths could mean that the critical event occurs from one

to ten grain diameters ahead of a crack [36], tying the mechanics of the Ritchie

et al model to a perhaps more realistic microstructural mechanism. Some of

these new ideas are evidently amenable to experimental test.

2. Effects of grain size.

A decrease in grain size always increases yield strength or flow stress, an effect which is

large provided that the size exceeds a critical value of ^ lym where substructure becomes a more important strengthening mechanism[46 ] . Empirical correlations for steels also indicate that decreasing

grain size is the only strengthening mechanism that also increases toughness[39] . However, this trend

is also limited by the decreased work hardening rate and increased tendency for strain localization

associated with it. Thus, an optimum grain size exists for a given material.

3. Dispersion hardening.

Dispersion hardening is attractive compared to other hardening mechanisms because: dispersoids pin grain boundaries and aid in achieving fine grain or substructure sizes (item 2); they

provide large increases in yield or flow stress (proportional to the inverse particle spacing) while

they are impenetrable and non-by-passable; and they give large values of work hardening rate which

provides resistance to plastic instability. The problems with dispersions are those mentioned in

item 1, concerning their role as crack or void sources.

4. Diffuse hardening mechanisms.

Argon[47] classified as "diffuse" hardening mechanisms those which increase yield stress;

do not affect slip character or da/dT; have little or no positive effect or a negative effect on work

hardening; and as a consequence decrease the resistance to plastic instability. These include grain

size hardening (item 2), solid solution hardening, and the various forms of precipitation hardening.

Because of the decreased resistance to plastic instability, these mechanisms are not as attractive as primary strengthening mechanisms as they are as secondary mechanisms in conjunction with dispersion

hardening. This holds true in particular for RSP materials.

5. Work hardening.

Despite the theoretical and experimental activity since the 1950's, which provide guidelines

on work hardening in terms of rough average dislocation properties [48 ,49 ] , we still have little knowledge of the detailed mechanism of work hardening, particularly in polycrystals or multiphase systems

and at high strains.

6. J and tearing modulus as a function of yield stress.

Continuum mechanics models predict that if the critical crack opening displacement for crack

nucleation and propagation is constant, that the value of Jic should increase with yield strength

while the value of the tearing modulus dJ/da should decrease with increasing yield strength[27,48]

There are data to support the latter effect, but some measurements indicate an inverse relation of

JlC and yield strength[49] . Evidently, the experiments also conflict with the mechanics prediction

(which has some support in a compilation by Paris [52], although the data is so scattered that the

opposite trend would be possible) that materials with large J have small dJ/da and vice-versa. This

suggests a possible role of microstructure leading to a dependence of the critical crack opening displacement 6ic on yield strength.

145
7. Additivity of strengthening mechanisms.

In the absence of detailed studies or predictions, most treatments consider that multiple

mechanisms contribute linearly to strength. For the special case of point obstacles and simple dislocation line tension, the rule instead should be the square root of the sum of the squares of the

individual contributions[53 ] . Work is clearly needed to provide better guidelines in this important

subject area. ^

8. Dislocation mechanisms.

A major problem in complex structures in relating microstructural mechanism to macroscopic

properties such as flow stress is the statistical averaging problem. To first order, single disloca-

tion models for Orowan hardening, precipitation hardening, etc., are known. However, averaging must

be done over 10^^ segments, each experiencing differing resolved internal stresses arising from

interactions with other segments [54] . In the absence of guidelines, simple averaging is assumed, for 1

example in the relation of strain rate to dislocation density and dislocation velocity. Some work on i

a limited number of segments, but nevertheless distributed realistically in three dimensions, is

needed to indicate the accuracy of the simple averaging process. To second order, more refined single,

dislocation calculations, using anisotropic elasticity, nonlinear inhomogeneity effects for second- 1 phase particles, and dropping the simple line tension model, are needed to better define the forcedistance curves necessary in the dislocation rate theories.

9. Limiting volume fraction of dispersoid.

, Of importance in RSP alloy design is the issue of the limiting volume fraction of dispersion. i Where strength alone is the dominant design criterion or where interface cohesion prevents void formation, properties improve with decreased spacing (or increased volume fraction at constant size). ; For particles poorly wet by the matrix in the capillarity sense, a limit should be the percolation i limit of 18-25 percent, and some work on steels tends to confirm this limit[55]. On the other hand, 1 when the particles are wet by the matrix, as for tungsten carbide in cobalt, the material has useful

properties. Wetability increases with decreased surface energy of the particle-matrix interface, or, equivalently with increased interface cohesion. Work is needed to establish alloy design guidelines.

10. Plastic instability.

Some work on the initiation of plastic instability, particularly at surfaces [21 ] is in quite

good agreement with continuum mechanics predictions of critical strains [20]. Yet hydrogen charging

has been shown to drastically lower these strains (by about a factor of two)[21]. Work is needed to determine whether strong local defects such as sharp micro-notches could produce such changes in the

continuum predictions, and to determine whether the role of hydrogen is to produce such defects.

11. Cyclic crack growth.

In cyclic loading, wavy slip alloys appear superior in low cycle fatigue, where crack initiation is important, but poorer in da/dN - AK crack growth compared to planar slip alloys. This is

clearly a microstructural effect and deserves further study.

12. Transformation induced plasticity.

In a number of alloys, resistance to plastic instability is provided by martensite transformation as a function of strain. The J resistance curve does not appear to be optimized at the

same conditions of strain rate, temperature and composition as the necking resistance in tension[56].

An important issue is whether martensite forms continuously during the deformation process and influences necking resistance throughout the uniform strain increment or whether it preferentially forms

locally at incipient necks. In other words, is the TRIP effect a local or global one. If local, the

models for transformation toughening of ceramics by Evans and Hutchinson[36] could apply as well for TRIP materials in the plastic zone at crack tips. J

13. Multi axial stress and strain effects. 1

Work on multiaxial stress effects is only beginning. Certainly from both a continuum and

microstructural viewpoint[23] , void formation and the propensity for plastic instability are influ- 1

enced by hydrostatic tension. Also, the strength-differential effect reflects multiaxial stress influences. Similarly, strain path has been shown to be an important factor influencing strain to failure in sheet steel and aluminum. There is little guideline on how microstructural variables relatej

to these effects.

146
14. Design needs.

Alloy design should not be focused on a single property. The potential spectrum of property needs for given designs should be considered. In many cases secondary properties such as buckling

resistance or density may be limiting. A systems approach to materials and design seems to be the

optimum method for improving overall structural efficiency.

15. Other effects.

Other properties are quite important and could receive equal attention with the yieldtoughness properties emphasized in the above enumeration. Environmental interactions can lead to severe degradation and could be limiting in RSP structures. Fatigue properties, particularly initiation, relate perhaps more closely to events at the crystal plasticity-microstructure level than the

yield-toughness properties. Optimization of yield-toughness behavior almost certainly would not give

optimum corrosion or fatigue resistance.

16. Production of precipitation hardening materials by RSP.

It is fairly clear that materials which are precipitation hardened in the absence of dispersoids or very small grains have a tendency for strain localization. They thus do not typically

have particularly attractive ductility. Accordingly, if RSP is used to make materials which are precipitation hardening then it appears to be advantageous to take advantage of this technology to introduce either grain refinement or dispersoids in addition to the intrinsic precipitates which form

during aging.

Concl usions

In view of the foregoing discussion we draw the following conclusions.

1. RSP has the potential to fill an important need in the area of microstructural

synthesis

2. Our ability to specify optimum microstructures is not yet fully developed but with

respect to improving strength and toughness several trends can be identified as follows

a. grain size and dispersion hardening are the most attractive

strengthening mechanisms because they tend to increase strength

and work hardening simultaneously

b. factors which promote strain localization and/or plastic instability must be avoided

c. a relatively narrow size range exists over which dispersoids

can provide increased strength without impairing toughness

d. other strengthening mechanisms can be used as secondary means

for improving strength provided their effect does not override

the dispersoids with regard to plastic instability.

3. A number of areas of microstructure-property relations are presently not well -enough

understood to permit microstructure design from first principles.

4. RSP can make a significant short term contribution to improving materials by providing a method for preparing more uniform, finer microstructures in conventional

alloys.

5. RSP can make an even bigger long term impact on improving materials when and if we

can specify optimum microstructures with more precision.

Acknowledgments

The authors gratefully acknowledge useful discussions with other members of the DARPA Materials

Research Council. In particular, those with Professors Hutchinson, Evans and Budiansky, and with

Dr. Jacobson, formerly of DARPA. The authors, also, acknowledge with gratitude the permission of

several collaborators to cite the results of their work prior to publication. Included in the group

of collaborators would be Professors Hutchinson and Evans as well as Drs. Terlinde, Martin and Paris.

147
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149
THE EFFECT OF RAPID SOLIDIFICATION ON THE MICROSTRUCTURES

AND PROPERTIES OF ALUMINUM POWDER METALLURGY ALLOYS

J.R. Pickens

Martin Marietta Laboratories

1450 South Rolling Road

Baltimore, MD 21227

and

E.A. Starke, Jr.

Department of Materials Science

University of Vi rginia

Charlottesville, VA

ABSTRACT

Rapid solidification (RS) applied in conjunction with powder metallurgy (P/M) can

be used to improve the microstructures of alloys by refining grain and constituent

particle size, increasing solubility limits, and enabling the use of alloying elements

not practical by ingot metallurgy (I/M).

RS has been particularly successful in producing improved aluminum P/M alloys,

resulting in materials having high room-temperature strength, high specific modulus,

and improved elevated temperature properties. This paper reviews the role of the RS microstructure in determining properties such as room and elevated temperature strength,

fatigue resistance, toughness, creep resistance, stiffness, and resistance to environmental attack.

I. INTRODUCTION

Rapid solidification (RS) and powder metallurgy (P/M) consolidation are presently being

used in aluminum alloy development programs directed towards developing new alloys that combine

high strength, fracture toughness, and corrosion resistance with improved thermal stability,

reduced density, and greater stiffness. Much of this work has been reviewed recently [1,2]

and subsequent research in the RS aluminum P/M field has been so extensive, that a complete

state-of-the-art review is an impractical task. Consequently, these authors will not review

the work of many ongoing investigations, but will describe in detail the relationships between

the microstructure and properties of RS aluminum alloys based on current investigations to which they have access. Since the P/M alloys of interest for structural applications fall into

three classes, i.e., high strength, low density/high modulus, and high temperature alloys, we

will discuss them in these groupings.

II. HIGH-STRENGTH P/M ALUMINUM ALLOYS

The strengthening mechanisms most often employed for aluminum alloys include those associated

with solid solutions, precipitates and dispersions, and grain and subgrain boundary effects [3,4].

Rapid solidification may be used to decrease the grain size, extend solid solubility, refine the

intermetal 1 ic particle size, and increase chemical homogeneity. Most of the commercial highstrength P/M aluminum alloys are based on the composition of I/M 7XXX series alloys (A1 -Zn-Mg-Cu)

with relatively small additions of low-solubility transition elements such as Co, Fe, or Ni [2].

A representative alloy of this group that has received much attention is X7091 (nominal

composition in weight percent, 6.5 Zn, 2.5 Mg, 1.5 Cu, 0.4 Co, 0.05 Fe, and 0.06 Si). Typical

production procedures are as follows: A melt of the desired chemical composition is air atomized

with solidification rates in the 10^-10°°C/s range to produce irregularly shaped powder particles,

sieved to obtain a desired particle-size distribution, isostatical ly pressed at room temperature

to ~ 75% theoretical density, encapsulated in an aluminum can, evacuated at high temperature,

hot pressed to 100% theoretical density, and extruded. Subsequent processing is typical of that

used for I/M products.

150
A. Microstructure

The non-axi symmetri c geometry of the extrusion die coupled with the cylindrical nature of

the hot compacted P/M billet creates non-uniform metal flow during extrusion. The changing

nature of flow between the edge and the center of an extruded X7091 plate is shown in Fig. 1 [5]. Metal flow at the center is similar to that which occurs during hot rolling; it gradually

deviates to a pseudo-radial nature at the edge similar to that which occurs during axisymmetric

extrusion. The closer spacing of the flow contours revealed by etching the short transverse

section of the plate corresponds to greater strains [6]. Consequently, variations in the

texture, degree of recrystallization, and grain shape are expected within the plate; Walker

and Starke [5] have characterized these parameters for extruded X7091 plate.

The texture measured from the interior of the plate is represented by the pole figures in

Fig. 2. The texture near the edge of the plate. Fig. 2, (a) and (b), contains strong deformation

(110) [111] and recrystallization (100)[001] components often observed in axisymmetric extrusions

of aluminum alloys. The partial radial metal flow creates a pseudo-fiber texture indicated by a strong [111] peak in the extrusion direction and spreading of the (100) along the fiber orientation. The radial nature of the recrystallization and deformation components of texture are represented by the broken and solid lines in Fig. 2(c). Figure 2(d) contains strong deformation

(100)[112] and recrystallization (100)[001] components characteristic of duplex sheet textures.

The surface texture is quite different from the interior texture and this is especially true

near the center of the plate. The (100) pole figure from the center surface. Fig. 3(a), has a strong deformation component spread over a wide range in the transverse direction and no recrystallization component. However, a recrystallization component is present in the (100) pole

figure from the center interior. Fig. 3(b).

The absence of a strong (100)[001] recrystallization component in the surface texture.

Fig. 3(a), suggests an unrecrystal 1 ized structure, and the presence of a strong (100)[001]

component in the interior texture. Fig. 3(b), suggests a partially recrystal 1 i zed structure.

Optical metallography. Fig. 4, appears to support the x-ray results, although it is difficult to distinguish recrystal 1 i zati on from unrecrystal 1 i zed regions in fine-grained P/M products.

Apparently the higher surface temperature, resulting from die friction and large deformation

strains, allows extensive dynamic recovery and the formation of a substructure that is stable

during subsequent annealing operations. Conversely, the lower temperature experienced in the

interior does not allow extensive dynamic recovery, and partial recrystallization occurs during

solutionizing treatments. Although an inhomogeneity in grain structure does exist in the P/M

product, the degree of inhomogeneity appears to be much smaller than normally found in recrystallized or partially recrystallized I/M products. A fine distribution of oxides, originating

from the atomized particulate, work in conjunction with small (< 0.4 ym) dispersoids. Fig. 5, to inhibit subgrain and grain boundary migration and prevent the large grains commonly found in

I/M products from developing.

The small i ntermetal 1 i c particles that form during solidification and fabrication of the

P/M billet do not fracture or coarsen during primary processing as do the large constituent

phases present in I/M products. SEM observation of bromine-etched specimens revealed a uniform

distribution of dispersoids from the edge to the center of the plate. Fig. 6. Mean particle

spacing measurements give no significant statistical variation in interparticle spacing. The

dispersoids were identified by x-ray diffraction energy dispersive analysis as [Co,Fe] 2 Alg.

Most dispersoids were less than 0.4 pm in diameter although some primary i ntermetal 1 ics, ~ 2 pm

in diameter, were uniformly distributed throughout the extruded plate. No large Al 7 Cu 2 Fe constituents, normally found in I/M 7XXX alloys, were detected. The precipitates that form

during aging are the same as those in I/M 7XXX alloys. For the T7E69 aging condition of this

study (an overaged temper) the matrix precipitates are metastable n'. Equilibrium n (AlCuMgZn)

precipitates are associated with subgrain and grain boundaries. Fig. 7.

Alloy X7090 is similar in composition to X7091, but is more highly alloyed with Zn, contains

less Cu, and has four times the Co content (nominal composition in weight percent: 8.0 Zn, 2.5 Mg,

1.0 Cu, 1.6 Co, 0.05 Fe, 0.06 Si). The microstructures of X7090 and X7091 are compared in Fig. 8 and are basically similar, with the exception of the larger size and the larger volume fraction

of cobalt-containing particles in the X7090.

151
B. Mechanical Properties

Tensile tests were conducted at six locations across an X7091 extrusion to evaluate the

effect of the microstructural variations on tensile properties. No significant variation was detected, and the fine grain size, particle size, and lack of coarse constituents result in

better combinations of tensile properties than observed for I/M 7XXX alloys in similar tempers.

Table 1.

TABLE I

Average Tensile Properties of P/M X7091 and Selected I/M Alloys

Alloy

X7091-

T7E69

7175-

T73511

7178-

T76511

7050-

T76511

YS (MPa) 571 462 484 514

UTS (MPa) 658 534 554 576

RA (%) 28.6 — — 31 Elongation 13 14 10 13

The strain-life behavior of P/M X7091 is compared to that of a conventionally processed I/M

of similar composition in Fig. 9. Although cyclic deformation was very homogeneous for the P/M

product. Fig. 10, no improvement in low cycle fatigue (LCF) life, normally obtained when reducing

grain size [7], was observed. Careful examination of the surface of the P/M LCF specimens re- vealed that cracks initiated at grain boundaries, whereas they initiated at slip bands in I/M

samples. The unexpectedly poor LCF behavior of the fine-grained P/M X7091 is most likely due to the presence of oxides at the grain boundaries. This has been confirmed by recent studies of

Hirose and Fine [8] and Kuo and Starke [9], which will be described in more detail in the next

section.

C . Microstructure and Mechanical Property Modification by ITMT

Kuo and Starke [9] and Chang and Starke [10] have recently studied the effect of microstructure and texture on the monotonic and cyclic properties of X7091-T651 rolled plate. Various

structures were developed using conventional I/M, P/M, and intermediate thermal mechanical treatments (ITMT). As mentioned previously, P/M produces a finer grain structure and particle size

distribution than I/M. Intermediate thermal mechanical treatments produce a recrystallized

grain structure with a weak texture, compared to the unrecrystal 1 i zed grain structure and sharp

texture normally obtained with conventional processing. For the sake of discussion, the following

designations will be used: Conventionally processed I/M (IMCP); I/M ITMT to produce a large

grain size (IMIP)-L and small grain size (IMIP)-S; conventionally processed P/M (PMCP); and P/M

ITMT to produce a fine, recrystallized grain structure (PMIP).

Representative microstructures are shown in Fig. 11 and the basic microstructural data

obtained from the optical micrographs are summarized in Table II. Particles larger than 0.1 ym were observed on bromine-etched surfaces by SEM and were coarser in the I/M products; higher

volume fractions were observed for the I/M materials, Table III. Intermediate thermal mechanical

treatments generally produced a more uniform distribution of particles than their conventionally

processed counterparts. The monotonic properties, which were comparable for the different

products, are listed in Table IV; they generally have higher yield strengths than those observed

for conventional I/M 7XXX alloys.

152
TABLE II

Basic Microstructural Data of the X7091 Materials

Materi al Ll X Li T X Lst

(ym)-5 (ym)

^pl -1 i n DR

{%)

IMCP 170 X 61 X 20 81 0.76 20-40

IMIP-L 139 X 109 X 76 130 0.33 100

IMIP-S 11.5 X 10.7 X 9.2 13.0 0.13 100

PMCP 3.2 X 3.4 X 2.0 3.5 0.22 7 PMIP 31.2 X 30.4 X 9.6 25.9 0.53 100

L = mean grain intercept length

L = longitudinal direction

LT = long-transverse direction

^lpl_lin = orientation factor of grain shape

(e.g., 0 for a sphere, 1 for a sheet)

ST = short-transverse direction

Dgq = diameter of a sphere of equivalent

vol ume

DR = degree of recrystallization

TABLE III

Total Volume Fraction of the (Co,Fe) 2Al 9 and Cu-rich Particles

with Sizes Larger than 0.1 ym in the X7091 Materials

Volume Fraction

(%)

IMCP IMIP-L IMIP-S PMCP PMIP

Transverse Plane 2.17 2.23 2.69 1.55 1.52

Longitudinal Plane 2.56 1.98 2.77 1.74 1.71

Rolling Plane 2.36 2.12 2.20 1.52 1.42

Average 2.36 . 2.11 2.55 1.60 1.55

TABLE IV

Mechanical Properties of the X7091 Materials

El asti

Modulus

(GPa)

Yield

St rength

0.2% offset

(MPa)

ELONG

(%)

^0

4n —

A

(%)

Strai

Hardening

Exp., n

Ductility

Coeff

£f ' (%)

Fati gue

Ductility

Exp., c

Strai n Hardening

Exp., n'

IMCP 71 549 15.9 40 0.053 92 -0.89 0.071

IMIP-L 543 15.2 37 0.063 16 -0.62 0.075

IMIP-S 545 12.0 42 0.069 51 -0.74 0.076

PMCP 558 14.8 27 0.055 25 -0.74 0.090

PMIP 497 17.7 40 0.065 122 -0.84 0.106

153
The IMCP failed by transgranular shear, Fig. 12(a), and both the IMIP-L and IMIP-S failed

by a mixture of transgranular and intergranular modes. Fig. 12, (b) and (c). The lower ductility

highly textured PMCP material failed intergranularly. Fig. 12(d). The fracture features appeared

as pulled-out flat grains, corresponding to the shape of the flattened powder particles and

indicating fracture along particulate boundaries that had high oxide concentrations. The ITMT

of the P/M material changed the fracture mode from intergranular to a mixed mode for the PMIP

Fig. 12(e). The recrystallization process associated with the ITMT moved the grain boundaries

away from the oxides and reduced the occurrence of intergranular fracture.

As mentioned previously, reduction in grain size generally decreases strain localization

and increases resistance to fatigue crack initiation (FCI) [7]. This was observed for all mater- ials except PMCP, Fig. 13. The PMCP showed the poorest LCF life, even though it had a microstructure conducive for homogeneous deformation. The modified P/M microstructure obtained

by ITMT had the best FCI resistance of the materials investigated in this study. The slip in the IMCP was retarded by coarse particle stringers, and fatigue cracks were initiated at slip band particle intersections or at particles lying along slip bands. Fig. 14(a). Fatigue cracks nucleated at slip bands and grain boundaries in the IMIP-L, Fig. 14(b), most likely due to the material s long slip distance and slip compatibility across grain boundaries. Since P/M refined grain structure and homogenized slip, FCI in PMCP was associated with other detrimental features such as oxide clusters along particulate interfaces. Fig. 14(c). However, the ITMT moved the

’ grain boundaries away from the oxides, and fatigue cracks were nucleated at the slip bands on free surraces. Mg. 14(d), thus accounting for the improvement in FCI resistance.

j

Ai. experimental materials are shown in Fig. 15, by plots of a da/dN versus ' AK. A plateau connects the threshold and Paris regions on the FCGR curves from IQ-^ to IQ-^ m/cycle. The FCP resistance was found to be best for the IMCP and IMIP-L (AK^.u ~ 3.0 MPa m^/^1 ~ 2.5 MPa ml/^), and worst for the PMCp'Snd IMIP-S

P/M ^ Several studiP^^Fn 1 ^ 1 ^^’h°^®

reported similar results for extruded X7091

have Suggested that the FCGR's of aluminum alloys decrease with

n?l J'" attributed to the dislocation slip being more reversible

closure [Is]

^ accumulated [14], or to a greater degree of fatigue crack

rninhilati combination

X7091 produced the optimum microstructure, i.e., that which gives the best of mechanical properties. The important features include a small recrys^^lized

at "ralJ boJnSIriL!'"'

distribution, a weak texture, and a iow concentraUon oj Ixides

Properti es Relating to Envi ronmental Interactions

menta[^\Mck'''L\''pJ'th effect on resistance to various types of environprovidefb^RS fineJ microstructural refinement

intermetalfics that segregation. In addition, the fine dispersion of

consolidation can altpr^th

RS P/M, and the dispersion of oxides accompanying P/M

gen eib“Juemel!t'’(5Fl®®"r^!.^^°''TM‘' d^^rning the susceptibility of RS aluminum alloys to hydroheavtTr ntoirHE‘ry^7 SCC has been shown io

undertlken^Telucidate t^rslrmechanL!

Al-Zn-Mri/M^lMvPMvMvP.'w^^ tJ® numerous investigators to provide evidence that SCC in

Mode I (tension than in Mode^lfl Mnt

i greater susceptibility of I/M 7075 to SCC in susceot i bi 1 i tv whpn tha h ^ tiplane shear, i*e., torsion) and demonstrated changes in the

Mode'^a JML'arstress^^^^^^^^^^^

inhibitor. As, was added to the environment! In ciently for HE, whereas in Mode III ^thpr^

^ driving force for hydrogen to concentrate suffisolely by dissolution Conseaiipntl v th

^ tnaxiality and SCC is believed to proceed

is operating dSrinq see .J interpreted these findings as evidence that HE

in the SCC mechanism.

* proposed that HE was more dominant than dissolution

|

solution'^deJeloped'b^^Spro^^^^ technique for P/M X7091 in the following teft uy ^prowis et al. [20], referred to as chromate-inhibited brine:

154

I
0.6M NaCI + 0.02M Na2Cr207 + 0.07M Na2C2H502 + HC2H5O2, pH = 4.

They found results similar to those Green et al . [18] found for 7075. Figure 16 shows that

time-to-fail ure (ttf) at a given normalized stress intensity (Kj/Kjc or Kin/Kmc) ‘'s significantly less in Mode I than in Mode III. This shows that RS Al alloy X7091 is susceptible to HE

as are its conventional counterparts and furthermore, that the SCO mechanism is similar in the

two-alloy types in that crack advance occurs both by dissolution and HE.

To assess the relative contributions to crack advance during SCO, Pickens et al . [19]

measured the plateau crack velocity (vp) in Mode I using double cantilever-beam specimens and

a compliance calibration. They also measured crack velocity in Mode III [21] but, because the

Mode III apparatus did not include a torque-angular displacement measurement system, crack ve- locity could not be precisely measured. However, the total crack advance measured from the Mode

III see fracture surface (after fracture) was divided by the ttf to give the average crack velocity (v) in Mode III. Because ttf includes both initiation and propagation times, strict comparison of this velocity with Vp in Mode I cannot be done. However, v computed for a specimen

loaded to 95% of Kijiqj where initiation time is expected to be short, was 0.28 x 10”° m/hr, in

comparison to a Vp in Mode I of 3.3 x 10”° m/hr. Thus, the crack advance solely attributed to dissolution (i.e., in Mode III) is an order of magnitude less than that attributed to dissolution and HE (i.e., in Mode I), which suggests that the mechanism of SCO in X7091 is heavily

dependent upon crack advance by HE. Thus, microstructural features that trap hydrogen, or serve

as recombination sites for H2, should improve SCO resistance.

Pickens et al . [19] showed that resistance to SCO is enhanced by RS P/M processing and

there is a lower Vp for X7091 than I/M alloy 7075 in chromate-inhibited brine. Fig. 17.

Pickens et al . [21] have also examined susceptibility to SCC of X7091 and X7090 in several

tempers using double cantilever-beam (DCB) specimens in chromate-inhibited brine. The DCB speci- men geometry enables measurement of the SCC threshold stress intensity (Kj 5 qq) as well as

Vp. They found that susceptibility in each alloy decreases in the order: underaged temper

(UT) > peak aged temper (PA) > overaged temper (OA). Figure 18 shows the v-K curves for each

alloy in the UT and PA tempers. For X7090, Kj 5qq

is lower and Vp is higher in the underaged

temper than in the peak aged temper, demonstrating greater susceptibility in the UT. Several

specimens showed double plateaus which may be caused by corrosion product wedging or may reflect

a change in fracture mechanism in different stress intensity regimes. No cracking was observed

in the OA tempers for X7090 (T7E71) or X7091 (T7E69). This is significant because each alloy

affords a major strength advantage (~ 100 MPa for X7091, greater for X7090) over comparable

7XXX I/M alloys in the SCC resistant temper (e.g., 7075T73). The stress-corrosion fracture

path in both X7090 and X7091 was found to be largely intergranul ar (Fig. 19) in this work [21]

although Swanson et al . [22] have observed transgranul ar SCC in X7090.

Free Mg has been shown to segregate on the GBs of Al-Zn-Mg I/M alloys [24,25] during solution

heat treatment. Viswanadham et al . [26] proposed that this free Mg serves as a means of concentrating hydrogen by forming a Mg-H complex and thereby playing an important role in the SCC

mechanism. Pickens et al . [27] found that on high-purity Al-Zn-Mg I/M alloys, this Mg segregation

generally decreases with aging time, which is consistent with the decrease in SCC susceptibility

found by Speidel [23] on Al-Zn-Mg-Cu alloys. Furthermore, Pickens et al . [27] proposed that

the decrease in SCC susceptibility with aging time may in part be caused by the decreasing trend

in Mg segregation.

Figure 20 is an Auger depth profile of an X7090 specimen fractured under ultra-high vacuum

at 35°C after pre-exposure to liquid gallium (Gai_) to induce intergranular failure [28]. The

Auger beam spreads over many GBs and collects information from GB precipitates and elemental

segregants on the GB. Extreme Mg segregation is observed that is much greater than the Zn segregation, and no Cu segregation was found, which indicates that much of the Mg is not bound in

precipitates but exists freely. This is also similar to the relative concentrations of Mg and

Zn segregation found on I/M grain boundaries [24].

Other studies [28-29], which used the aforementioned Auger technique, measured GB segregation of Mg for I/M 7175 and RS P/M X7091 extruded bar. The magnesium-to-aluminum peak-to-peak

height ratio from the Auger spectra (Mg/Al) was used as a measure of GB segregation. Mg/Al was significantly lower for the RS P/M alloys as shown in Table V. Thus, RS P/M processing decreases

GB segregation of Mg, which probably is largely caused by the finer grain size in the RS P/M

alloy. The X7091 has a nearly equiaxed grain size of 2.0 x 1.6 x 1.6 ym in the L, LT, and ST

155
directions, respectively, compared with ~ 170 x 17 x 16 ym grain size in the I/M 7175. Therefore, the RS P/M alloy has much greater GB area competing for segregants.

TABLE V

Grain Boundary Composition Measured By Auger Electron Spectroscopy of LT Specimens Fractured

Intergranularly at 35°C, Under Ultra-High Vacuum after Pre-Exposure to Liquid Gallium

Alloy Temper Mg/Al

7175 T 6 1.9

1.5

X7091 T6E192 + Ihr at 157“C 0.24

0.33

X7090 T6511 0.33

A study of see initiation in X7091 and 1/M 7075 revealed that the surface oxide films -- an alloy's first defense against environmental attack — are similar in composition and thickness

for each alloy [19]. The "thermal" films that form during solution heat treatment of each alloy

are magnesia and are 600a thick, whereas the "natural" films that form after polishing are alumina and are 50a thick [29]. Therefore, RS does not alter the composition and thickness of

the protective surface oxide film. However, RS does eliminate the coarse iron-containing

constituent particles that are present in I/M 7XXX alloys; much of the iron goes into the

[Co,Fe] 2

Al 9 particles. Figure 21 shows a polished, unstressed I/M 7075 specimen that was exposed to distilled water, a mild SCO environment. Attack is apparent at the constituent/matrix

boundary, indicating that these particles serve as SCO initiation sites.

Lyle et al . [30] speculated that the presence of oxides in the RS P/M microstructure

decreases SCO susceptibility by changing the electrochemical nature of the GBs. Figure 22 shows

a specimen of X7091 that was solutionized and water quenched to reveal the oxides. They clearly

serve to pin GBs, thereby refining grain size, and also make the GB surface more irregular,

causing an intergranular SCO path to be more circuitous.

A major benefit of RS is that fine dispersions of low-solubility alloying elements can be

achieved that are not practical by I/M because these additions segregate coarsely using conventional solidification rates. Cobalt, one such addition, has been particularly successful and

was suspected early in RS P/M aluminum alloy development of enhancing SCC resistance [31]. Otto

[32] subsequently provided strong evidence of its beneficial effect. Christodoulou et al

[33] studied the mechanism of SCC in X7091-type alloys containing 0, 0.4, and 0.8 wt% Co and

having similar compositions and grain sizes. When bolt-loaded and precracked DCB specimens were alternately immersed in chromate-inhibited brine, SCC susceptibility clearly decreased with Co

content (see Fig. 23). Both and Vp reflect this trend. To understand the mechanism

by which Co decreases SCC susceptibility, a specimen of X7091 (0.4 wt% Co) was pre-exposed to salt water at room temperature for 30 min. Attack occurred in the region adjacent to the C02AI 9 particles (see Fig. 24), suggesting that the C02AI 9 is cathodic with respect to the grain

boundary region. This could have a beneficial effect on SCC resistance because this galvanic

action could serve to blunt a propagation SCC crack. Furthermore, these cathodic particles may

serve as sites for hydrogen recombination, thereby catalyzing hydrogen egress from the alloy

during stress-corrosion. This possibility is currently being explored.

Liquid metal embrittlement (LME) is a type of environmental attack studied as much for its

insights into mechanisms of bonding and fracture, as for its prevention. The susceptibilities

to LME of RS P/M alloys X7091 and I/M 7175, aged to similar strengths, were compared in liquid

mercury (Hg|_) and in a liquid mercury - 1.5 at% gallium solution (Hg + 1.5 Ga)L [29]. Specimens

were wetted by placing a 2-ml drop of liquid metal on top of a HF layer used to remove the surface

oxide films. It was quite apparent that the RS alloy was more difficult to wet. The specimens

were fractured in tension at 25°C, 5 or 60 minutes after wetting, at a constant cross-head

speed of 0.084 cm/min. All specimens fractured with no measureable ductility, so the decrease

in fracture stress is the sole means of assessing susceptibility to LME.

The X7091 was less susceptible to LME in each environment (see Table VI). In addition,

susceptibility was generally greater with (Hg + 1.5 Ga)[_ because of gallium's well-known ability

156
to penetrate and weaken aluminum grain boundaries [34,35,36]. With Hg^, susceptibility did not

vary significantly with pretest exposure time (tp) for X7091, but it did increase greatly for

7175. Fracture under the influence of Hgi_, (Hg + 1.5Ga)L, or GaL (Fig. 25) was invariably

intergranular. Fractographi c studies revealed that HgL dissolves the grain boundary region,

particularly the precipitates on the boundaries, so a GB corrosion reaction occurs simultaneously

with adsorpti on-LME*. Furthermore, the greater GB segregation (Table V) in I/M 7175 than in RS

P/M X7091 leads to greater corrosion by the Hg[_. Nevertheless, the susceptibility to adsorptionLME is also greater in the I/M 7175, which is consistent with the observations of Ichinose and

Oouchi [37] who found that susceptibility to adsorption-LME generally decreases with decreasing

grain size.

TABLE VI

Effect of Pretest Exposure Time (tg) and Liquid Metal Environment on Fracture

Stress, op> at 25°C (LT Orientation) for an RS and a Conventional Aluminum Alloy [29]

Alloy Heat Treatment (mm)

Liquid

Metal

Fracture

stress, ap

(MPa + standard

deviation)

X7091 Peak hardness [29] -- None 586

X7091 Peak hardness 5 Hg 135 + 13

X7091 Peak hardness 60 Hg 131 + 20

X7091 Peak hardness 5 Hg + 1.5 Ga 119 + 12

X7091 Peak hardness 60 Hg + 1.5 Ga 94 + 25

7175 T6 — None 545

7175 T6 5 Hg 80 + 15

7175 T6 60 Hg 15+9

7175 T6 5 Hg + 1.5 Ga 68 + 17

7175 T6 60 Hg + 1.5 Ga 22+6

Thus, RS for aluminum alloys generally decreases susceptibility to a variety of environmental

effects ranging from SCC-hydrogen embrittlement to liquid metal embrittlement. The improvement

offered by RS often is caused by the refinement of grain size, but also is affected by the reduced

segregation and unique dispersions made possible by RS and P/M consolidation.

III. LOW DENSITY, HIGH MODULUS P/M ALUMINUM ALLOYS

Lithium is the lightest metallic element and, with the exception of beryllium, which has

manufacturing and health-related problems, is the only metal that improves both the modulus and

density when alloyed with aluminum. Fig. 26. Each wt% lithium added to an aluminum alloy reduces

the density approximately 3% and increases the elastic modulus approximately 6%, for lithium

additions up to 4% [38]. Moreover, Al-Li alloys containing more than about 2 wt% Li can be

age-hardened due to the precipitation of the metastable, coherent, and ordered s' (Al 3 Li)

phase [39]. However, upon plastic deformation these A1 3

Li precipitates are sheared by glide

dislocations resulting in strain localization and concomitant fracture at low macroscopic

strains [40]. Consequently, research on Al-Li alloys has focused on improving tensile ductility

and fracture toughness [41].

* Adsorption-LME is the decrease in fracture stress or strain of a solid metal that occurs during

its fracture in a liquid metal environment, without the requirement of diffusion or strong chemical

interaction between the two metals.

157
Previous studies have shown that strain localization due to shearing of coherent precipitates can be reduced by decreasing the grain size and by adding small incoherent dispersoids

[42]. The development of Al-Li alloys using rapid solidification and powder metallurgy

processing is considered a promising approach for achieving a smaller grain size and high-volume

fraction of small dispersoids when compared with conventional ingot metallurgy [43], Many

compositions have been examined [44]; however, those containing large amounts of dispersoidforming elements have not produced the desired ductility.

The most promising alloys are based on the Al-Li-Cu system. Copper has been shown to be

very effective in increasing the strength of Al-Li alloys. The main strengthening precipitates that form during aging of Al-Li-Cu alloys containing 2-3% Li and 2-3% Cu are 6* (Al 3 Li),

Ti (Al 2 CuLi), and o' (AI 2 CU) [45]. The T^ and o' precipitates have interfacial strains

that may aid in homogenizing deformation. The equilibrium precipitates T

2 (Al 5 CuLi 3 ) and 6 (AlLi) may also form during aging [46,47]. A composition of A1 -3Li -2Cu-0.2Zr was selected in

a recent alloy development program [44,48]. Three percent lithium was chosen to maximize the

modulus and density effects and eliminate the fabrication problems associated with higher lithiumcontent alloys. Two percent copper prior to aging, which approaches the solubility that can be

obtained in an Al-3Li alloy, was selected for strength considerations. Zirconium was added to inhibit recrystallization. A comparison of the tensile properties of this alloy produced by

three methods--I/M, P/M-splat, and P/M-atomi zed--i s given in Table VII. Although the strengthductility relationship is considerably better than observed for Al-Li binary alloys [49], it

is still somewhat below that desired for commercial utilization.

TABLE VII

Comparison of Tensile Properties of I/M, P/M-Splat, and P/M-Atomized

A1 -3Li -2Cu-0.2Zr Alloy Artificially Aged to Peak Strength

A1 1 oy Yield Strength,

0.2% offset

Tensile Strength, onts

(MPa)

Strain to Fracture, cf

(%)

I/M 503 561 4.0

P/M-spl at 490 546 4.2

P/M-atomi zed 443 562 7.3

The low ductility of this alloy was due to strain localization in the preci pitate-f ree zones

(PFZ's) along the subgrain boundaries. The unrecrystal 1 i zed structure has equiaxed subgrains

that develop during hot deformation. Preferential precipitation occurs along the subgrain boundaries resulting in the PFZ's, Fig. 27(a). The PFZ's are much weaker than the matrix and many

are favorably oriented for slip. Deformation concentrates in these regions and low ductility

ensues. Fig. 27(b).

Attempts were made to harden the PFZ by solid solution strengthening and additional precipitation, and to minimize subgrain boundary precipitation by alloy modification and TMT's [44].

Because magnesium is a potent solid solution strengthener in aluminum, it was selected as an additive and one of the alloys had the nominal composition A1 -3Cu-2Li -lMg-0.2Zr. The age-hardened

licrostructure of this alloy is very complex and contains T^ (Al 2 CuLi)-, 6* (Al 2 CuMg)-, and

A1 2 MgLi -strengtheni ng precipitates [50]. Some 6' (Al 3 Li) may also form; however, due to the presence of Al 3 Zr, it is difficult to detect [48].

Although the equilibrium To and S phases were detected on the boundaries of the elongated

unrecrystal 1 i zed grains, no preferential precipitation and associated PFZ's were observed at subgrain boundaries in TMT-processed material. Fig. 28. The strength/ductility relationship of

this alloy. Fig. 29, exceeds that of I/M 7075-T76. The high ductility is associated with homogeneous deformation and a transgranul ar fracture mode as shown in the TEM's (Fig. 30) from

recent high-voltage electron microscope (HVEM) in situ deformation studies [51].

158
IV. HIGH-TEMPERATURE P/M ALUMINUM ALLOYS

In the 1950's, attempts were made to produce an aluminum alloy that would be stable at high

temperatures by utilizing sintered alloy powder, SAP, containing a high concentration of fine

oxide particles [52,53]. Although this product was stable to ~ 800° K, it had limited room- temperature strength, poor reproducibility of mechanical properties, and a tendency toward

embrittlement during sustained loading at elevated temperatures [54]. The common aluminum

alloys currently being used by the aircraft industry for high-temperature applications are based

on the Al-Cu system and include 2219 (Al-Cu-Mn) and 2618 (A1 -Cu-Mg-Fe-Ni ) . The transition elements form i ntermetal 1 i c phases that are stable at elevated temperatures and hinder grain-boundary

sliding. However, both alloys derive their strength from coherent and partially coherent intermediate precipitates that coarsen rapidly and transform to equilibrium precipitates at temperatures above 450°K. The rate of coarsening is controlled either by atom transfer across the

precipitate/matrix interface, or by volume or bulk diffusion. Low values of interfacial energy,

solid solubility, and solute diffusivity are desirable for low coarsening rates [55,56].

The solid solubility and diffusion parameters for a selected number of binary additions

to aluminum are given in Table VIII. The alloying elements normally used for precipitation

hardening, i.e., Cu, Mg, and Zn, have high solubilities and diffusion rates. However, the

dispersoid-forming transition elements and the rare earth Ce have low solubilities and low diffusion rates. A high-volume fraction of small dispersoids formed by these additions can give

adequate high-temperature strength. The dispersoids can also slow down all three contributors

to steady-state creep; grain boundary sliding, diffusional or "Nabarro-Herri ng" creep, and

dislocation creep [57]. Large additions of the transition elements can significantly increase

the elastic modulus of aluminum. Fig. 26(a), making the dispersion-hardened alloys especially

attractive for stiffness-critical high-temperature applications. Unfortunately, the limited

solid solubility makes it impossible to obtain a fine dispersion of closely spaced incoherent

particles using normal ingot-casting procedures. The inherently slow ingot solidification

rates enhance the precipitation of coarse particles that lower ductility and fracture toughness

val ues.

TABLE VIII.

Diffusion Parameters and Solid Solubilities of Binary Additions to Aluminum*

Solute Maximum Solid Diffusion Coefficient, Dg Activation Energy

Solubility (cm^/s) (k cal/g atom)

(wt%)

Ce <0.05 1.9 X 10-6 26 .6 Co <0.02 1.1 X 10-6^ 19 .9 Cr 0.72 3.01 X 10"^ 15 .4 Cu 5.70 0.15 30 .2 Fe 0.05 4.1 X 10-9 13 .9 Mg 17.4 0.12 28 .6 Mn 1.4 0.22

IQ-^

28 .8

Ni 0.05 2.9 X 15 .7

Ti 0.24 5 X 10 -7 25 .8

V 0.37 6.05 X 10-' 19 .6 Zn 82.2 11.6 27 .8

* Data from Hansen, 1958 and Bishop and Fletcher, 1972.

Rapid solidification and P/M consolidation are currently being used in aluminum

development programs aimed at producing elevated temperature properties competitive

of titanium [58,59,60]. Solute element additions having high liqud solubility, low

lity, and low diffusion rates are being used either individually or in combination,

include Co, Fe, Ni , Ce, and Cr. Although strengthening theory predicts an increase

with decreasing interparticle spacing, and therefore with volume fraction (Vf) for a

-alloy

with those

solid solubiThese

in strength

constant

159
particle size, ductility also decreases with thereby placing restrictions on solute concentrations. Argon and Im [61] have shown that when the volume fraction of particles approaches some

critical value, interaction effects come into play and interfacial stresses are enhanced due to overlap of the plastic zones around the particles. This results in interface separation at very

low strains [62].

The principal characteristics that dominate the microstructure and properties of dispersionhardened aluminum alloys fabricated from rapidly solidified particulate arise from the size

distribution of the particulate and the quality of the interparticulate bonding. A typical

microstructure of an Al-Fe-Ni-Co splat-quenched flake is shown in Fig. 31 [63], and contains

both coarse and fine precipitates analogous to the microstructure of RS Al-Fe alloys studied by

Jones [64]. Jones suggested that the amount of coarse precipitates could be reduced by increasing

the solidification rates and/or reducing the solute concentration. Although the bimodal structure

normally exists within one splat particle, the relative amounts of coarse and fine structure

depend on the splat particulate size. A larger amount of the coarse dispersoids appear in the

large particulates because they have a lower degree of undercooling during solidification. Since

a range of particulate sizes are used for consolidation, the finished product contains a bimodal

particle-size distribution similar to those shown in Fig. 32 for Al-Fe-Ni-Co extruded plate.

The fine structure occupied approximately 95% of the volume of the alloys studied by Mullins

[63] and appeared to control the flow stress. Although the yield strength could be correlated

with the interparticle spacing of the small particle regions. Fig. 33, no one-to-one relationship

existed with calculated values using the Orowan equation. This is probably due to the additional

strength contribution associated with the high dislocation densities introduced during compaction

and primary processing of the wrought product. Nondeformable dispersoids generate a large number

of geometrically necessary dislocations [57], which contribute to the strength when the final

product is unrecrystallized, as for the case described here. The coarse structure occupied less

than 5% of the extruded plate but most likely controlled the ductility. Cracks were first

formed at the large particles. Fig. 34, and final fracture occurred by the coalescence of microvoids associated with the dispersoids. Other research has also shown that even when small particles far outnumber large ones, the larger particles crack first and at lower strains than the

smaller particles [65,66]. The difference in strength-ductility relationship of the two alloys

of Fig. 34 may be related to more large particles in the higher solute alloy.

As mentioned earlier, Jones suggested that the formation of coarse particles may be suppressed by increasing the degree of undercooling. Because this often is impractical, other

approaches must be taken. It is well known that multiple alloy additions can affect precipitation

in many aluminum systems [67]. In certain cases, compositional modification can increase the

nucleation rate and decrease both the size and coarsening rates of the precipitates, probably by

altering interfacial energies and diffusion processes. This has been accomplished for the Al-Fe

alloys by additions of Ce. Rapidly solidified Al-Fe binary alloys have microstructures containing

both coarse and fine dispersoids [64] similar to those shown in Fig. 31. The coarse precipitates

are absent in Al-Fe-Ce alloys. Fig. 35 [68].

Al-8Fe-3.4Ce is a new aluminum alloy that shows promise for high-temperature applications

[68]. Figure 36 compares its yield strength at various temperatures with the program goal and

the currently used high-temperature alloy 2219-T851. The decrease in strength with temperature

is related to both recovery and coarsening processes and occurs at a lower rate than that for

2219-T851 at temperatures up to 505°K. The real advantage of the di spersoid-strengthened

material, however, is in its creep resistance. Fig. 37. Note that the creep rate of age-hardened

2219 and 2124 increases rapidly with time, whereas that of the Al-8Fe-3.4Ce alloy seems to level

off. The accelerated creep of 2219 and 2124 is associated with extensive overaging that occurs

at 505°K. However, coarsening of the dispersoid-strengthened Al-8Fe-3.4Ce alloy is unlikely

at 505°K and the initial creep strain is probably associated with recovery processes. Once

these are complete, additional creep strain is minimal.

This study clearly demonstrates the feasibility of producing aluminum alloys with usable

mechanical properties and creep resistance in the 505-615°K range. However, more work needs to

be directed toward improving ductility. The elongation to failure at room temperature is 5% for

the Al-Fe-Ce alloy. Argon and Im [61] suggest that some improvement may be obtained by controlled

deformation processing, i.e., keeping the increments of strain between annealing operations below

the critical strain for dispersoid fracture and interface separation.

160
V. SUMMARY AND CONCLUSIONS

Rapid solidification combined with powder metallurgy consolidation enables the production

of aluminum alloys which often have superior mechanical and corrosion properties with respect to conventional ingot metallurgy alloys. The alloys developed for elevated temperature service

have perhaps the greatest advantages over conventional alloys. The superior properties result

from increased alloying limits, mi crostructural refinement of grain and constituent particle

sizes, reduced harmful segregation effects, and the ability to finely disperse i ntermetal 1 i

phases and oxides in RS P/M alloys. These unique abilities of RS P/M allow the alloy designer

to combine various microstructural improvements to tailor an alloy to specific requi rements

As a result, alloy development efforts have been successfully producing promising alloys for

applications requiring high room-temperature strength, high modulus and low density, and improved

elevated temperature properties.

ACKNOWLEDGEMENTS

Much of the data discussed was generated in projects supported by

agencies: AFOSR, ARO, DARPA, and NSF. The authors are grateful to G. M.J. Koczak, R.E. Crooks, V.W.C. Kuo, J.A. Walker, J.W. Mullins, J.A.S

Chri stodoul ou for providing information described in this paper.

the following government

Hildeman, R.E. Sanders,

Green and especially L.

161
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Design of New Aluminum Alloys," in Proc. 6th Int. Conf. on the Strength of Metals and

Alloys, Melbourne, Australia, Aug. 1982.

[5] Walker, J.A. and Starke, E.A., Jr. (1982) "The Microstructure and Properties of an Extruded P/M X7091 Plate," (submitted to Powder Metallurgy).

[6] Pearson, C. and Parkins, R. (1960) in The Extrusion of Metals , New York: John Wiley and

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Metals Park, OH: ASM, p. 205.

[8] Hirose, S. and Fine, M. (1982) "Fatigue Crack Initiation and Propagation in High Strength

Aluminum P/M Alloys," Metal 1 . Trans. , ^ (in press).

[9] Kuo, Victor W.C. and Starke, E.A., Jr. (1982) "The Effect of ITMT's and P/M Processing on the Microstructure and Mechanical Properties of X7091 Alloy", Metall. Trans., A (in press).

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on Fatigue Properties of I/M X7091 Aluminum Alloy," paper presented at the Fall Meeting of

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[13] Fine, M.E. (1980) Metall. Trans., A, llA , p. 365.

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by the Aeronautical Res. Inst, of Sweden and the Swedish Plastics and Rubber Inst., Stockholm, Sweden.

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Metal 1 urgy , M.J. Koczak and G.J. Hildeman, eds., Warrendale, PA: TMS AIME.

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and G.J. Hildeman, eds., Warrendale, PA: TMS-AIME.

[23] Speidel, M.O. (1971) in The Theory of Stress Corrosion Cracking in Alloys , J.C. Scully,

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[24] Green, J.A.S., Viswanadham, R.K., Sun, T.S., and Montague, W.G. (1977) in Corros. /77, Int. Corros. Forum , Paper 17, p. 71/1, San Francisco, CA: Natl. Assoc. Corros. Eng.

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unpublished research.

[29] Pickens, J.R., Precht, W., and Westwood, A.R.C. (1982) "Embrittlement of P/M X7091 and I/M

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MD, to be published.

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E.A. Starke, Jr., eds., Warrendale, PA: AIME, p. 205.

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Technical Report #AFML-TR-78-102, July 1978. [44] Palmer, I.G., Lewis, R.E., and Crooks, D.D. (1981) in Aluminum Lithium Alloys , T.H. Sanders,

Jr. and E.A. Starke, Jr., eds., Warrendale, PA: TMS-AIME, p. 241.

[45] Silcock, J.M. (1959/60) J. Inst. Met. p. 357.

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[48] Lin, F.S., Chakrabortty , S.B., and Starke, E.A., J7. (1982) Metall. Trans. 13A , p. 401.

[49] Sanders, T.H., Jr. and Starke, E.A., Jr. (1982) Acta Metall. p. 927.

[50] Crooks, R.E. (1982) The Influence of Microstructure on the Ductility of an A1 -Cu-Li -Mg-Zr

A1 1 oy , Ph.D. Thesis, Georgia Institute of Technology, Atlanta, GA. [51] Crooks, R.E., Kenik E., and Starke, E.A., Jr. (1982) University of Virginia, Charlottesville,

VA, unpublished research.

[52] Irmann, R. (1952) Metallurgia p. 125.

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Technical Report for Period Apri 1 -September , 1981, Lockheed-Cal i forni a Company.

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[56] Wagner, C., Z. Elektrochem , Vol 65, 1961, p. 581.

[57] Ashby, M.F. (1968) in Oxide Dispersion Strengthening , G.S. Ansell, T.D. Cooper, and F.V. Linel, eds., NY: Gordon & Breach, p. 143. [58] Wald, G.G. (1981) NASA Contractor Report 165676, May 1981.

[59] Griffith, W., IBID [19].

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[61] Argon, A.S. and Im, J. (1975) Metall. Trans. p. 839.

[62] Tien, J.K. (1975) in Fundamental Aspects of Structural Alloy Design , R.I. Jafee and B.A. Wilcox, eds., NY: Plenum Press, p. 363

[63] Mullins, J.W. (1982) The Microstructure and Tensile Properties of RSP Al-Fe-Ni-Co Alloys , M.S. Thesis, Georgia Institute of Technology, Atlanta, GA. [64] Jones, H. (1969-70) Mater. Sci. Eng. p. 1. [65] Palmer, I.G. and Smith, G.C. (1968) in Oxide Dispersion Strengthening, G.S. Ansell, T.D. Cooper, and F.V. Linel, eds., NY: Gordon & Breach, p. 253.

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163
ED ST

Figure 1. Optical macrograph of the cross section of the extruded P/M plate

showing flow lines.

ED

Figure 3. The (100) pole figures from

the center of the P/M extrusion

showing (a) the absence of the

recrystallization component as well as broadening of the deformation component near the surface,

and (b) the typical duplex sheet

texture at the interior indicating

partial recrystallization.

Figure 5. TEM showing oxides along grain and

subgrain boundaries and (Co,Fe) 2Al 9 dispersoids.

Figure 2. The duplex pseudo-axisymmetrlc.texture from the interior of the edge

(a), (b); and the duplex sheet

texture from the interior of the

center (c), (d), of the P/M extruded

plate. ED represents the extrusion

direction, T the traverse, and S the

short traverse directions.

Figure 4. Optical micrographs from center of

extrusion showing (a) unrecrystallized sub-structure near the

surface; and (b) larger, more

inhomogeneous, partially recrystallized grain structure at the

interior.

Figure 6. SEM of bromine-etched surface of ST section of the extrusion showing the

uniform distribution of (Co,Fe) 2Al 9 dispersoids.

164
PLASTIC

STRAIN

AMPLITUDE

,ACp/2.

Figure 7. TEM showing uniform distribution of

metastable n' precipitates in the

matrix and equilibrium p along a grain boundary.

Figure 9. Comparison of strain-life behavior

for X7091 and I/M alloy of similar

composition.

Figure 8 . Microstructure of (a) X7091

(slightly overaged); and (b) X7090

T6511 showing greater number of,

and somewhat larger, C02 AI 9 particles in the X7090.

Figure 10. TEM showing homogeneous deformation

in P/M extrusion cycled to failure

at AEp/2 = 0.0007.

Figure 11. Optical metallographs

(a) IMCP, (b) IMIP-L,

(c) IMIP-S, (d) PMCP,

and (e) PMIP.

165
!

i

Figure 13. LCF Cof f in-Manson plots for crack

initiation in dry air, showing

variations in FCI resistance.

Figure 14. SEM's showing the surface markings

and FCI sites of the LCF specimens,

Aex/2 1/3%: A) IMCP, = 250 cycles; (b) IMIP, = 180

cycles; (c) PMCP, = 150 cycles;

and (d) PMIP, = 290 cycles.

Effect of loading mode on SCC of P/M x7(Nl and I/M 7075

I

Figure 15. FCGR's of the experimental alloys

in dry air with R = 0.1, 10 Hz. Figure 16. Effect of loading on SCC of P/M

X7091.

166
Figure 17. Stress corrosion crack velocity vs Kj for X7091 and 7075.

Figure 19. SCO path in X7091 showing largely

intergranular fracture.

Kj (ksi'/In.l

Figure 18. Stress corrosion crack velocity vs Kj for X7091 and X7090.

SPUHERING TIME (min)

(Sputtered at 31 A/min) Figure 20. Auger depth profile of intergranular

fracture surface of X7090.

Figure 21. SCC initiation at constituent

particle in 7075.

Figure 22. Oxide particles in RS P/M alloy

X7091 pin grain boundaries and

make an intergranular path more

circuitous (from Ref. 31).

167
Figure 23. Stress corrosion crack velocity (v) vs Kj for different cobalt contents

in X7091-type alloys.

Figure 25. Fracture of X7091 under liquid gallium is intergranular.

Figure 27. TEM showing (a) PFZ's along subgrains, and (b) intersubgranular

fracture of an age-hardened A1-3L1-

2Cu-0.2Zr alloy (from Ref. 48).

Figure 24. Localized attack around a coarse

C02AI 9 particle (Indicated by arrow)

in X7091 after 30-min exposure to 3.5% NaCl solution.

Figure 26. (a) The effect of solute additions

on the modulus of aluminum alloys,

and (b) the effect of solute additions on the density of aluminum

alloys (from Ref. 58).

168
Figure 28. TEM's of Al-3Cu-2Li-lMg-0. 2Zr showing uniformly distributed precipitates and a lack of PFZ ' s (from

Ref. 50).

Figure 29. Yield stress and reduction of area versus aging time after various TMT (stretch) (from Ref. 50).

Figure 30. TEM's taken during in situ deformation studies in HVEM, showing homogeneous deformation and transgranular fracture of aged Al-3Cu-2LilMg-0. 2Zr alloy (from Ref. 51).

Figure 31. TEM of rapidly solidified Al-3. 3Fe2. 3Ni-4. 6Co splat-quenched flake

showing regions of coarse and

fine dispersoids (from Ref. 63).

Figure 32. SEM’s of bromine-etched surfaces of

extruded Al-3. 3Fe-2. 3Ni-4. 6Co plate

made from rapidly solidified flake

showing coarse and fine dispersoids

associated with large and small

flake particulate (from Ref. 63).

169

%

REDUCTION

OF

AREA
YIELD

strength.

MPa

1

4ftO -

z-o

zU 300

AI-3 3 F«-2 3 NI-4.6 Co • O

AI-4 3 F.-6 Ni-5 Co O

D

) \

N

\

O \

- Q- - -oo20

IB

16

1412

10

e

64 H 2

2 6 3 0

Figure 33. The effect of interparticle spacing

on the yield strength of Al-Fe-Ni-Co

plate made from P/M consolidation

of rapidly solidified particulate

(from Ref. 63).

TEMPERATURE. K

Figure 36. Comparison of yield strength temperature behavior of Al-8Fe-3.4Ce forging made from rapidly solidified

particulate with two I/M 2XXX

alloys. Tests conducted after

1000-hr exposure at temperature

(from Ref. 68).

Figure 34. SEM showing crack nucleation

associated with large dispersoids

in an Al-3.3Fe-2.3Nl-4.6Co tensile

specimens (from Ref. 63).

Figure 35. TEM of a rapidly soldified Al-8Fe3.4Ce alloy (from Ref. 68).

Figure 37. Comparison of creep strain versus exposure time of Al-8Fe-3.4Ce forging made from rapidly solidified

particulate with two I/M 2XXX

alloys (from Ref. 68).

170
STRUCTURE-PROPERTY RELATIONSHIPS IN RSP STEELS

G.B. Olson, M. Cohen, and G.J. Yurek

MIT, Cambridge, MA 02139

ABSTRACT

Rapid solidification of austenitic and martensitic steels via centrifugal atomization is found

to produce significant refinement of inclusion phases. The thermal stability of the fine inclusions

makes them effective grain-boundary pinning particles in hot-consolidated materials, accounting for

unusual resistance to grain coarsening at high austenitizing temperatures. This allows the improvement of sharp-crack fracture toughness by high austenitizing treatments without the reduction in blunt-notch toughness and tensile ductility which may otherwise accompany excessive grain coarsening. Experiments on a 4Mo-0.25C martensitic steel demonstrate that high austenitizing significantly improves the Kjc toughness of both RSP and conventionally-processed materials, but the

finer-grained RSP material has a 60°C lower Charpy ductile-brittle transition temperature. The

fine-inclusion dispersion and grain-coarsening resistance also enhances the oxidation resistance of

an RSP 303 austenitic stainless steel; this effect is attributed to promotion of oxide scale formation and mechanical keying of the scale by the fine grain size and dispersed inclusions.

Introducti on

Rapid solidification processing (RSP) allows the achievement of a wide variety of novel

metastable structures, but the range of structures which can be practically exploited is greatly

narrowed by the available consolidation techniques through which rapidly solidified powders, flakes,

and ribbons are converted into bulk materials. While the development of lower-temperature con- solidation methods will expand this range, certain special aspects of structural metastability which

survive the high temperatures of currently reliable consolidation processes form the basis of an approach to mi crostructural control which is ready for practical exploitation today. Investigations

of structure/property relationships in rapidly solidified steels produced by centrifugal atomization

and consolidated by hot extrusion indicate that the most practical benefit of rapid solidification

is the fine dispersion of stable compounds normally regarded as "inclusion" phases. After

discussing the role of inclusion refinement in the mi crostructural development of these steels we propose to illustrate its potential for the improvement of material properties, using examples of fracture toughness in high-strength martensitic steels and oxidation resistance in an austenitic

stainless steel.

Inclusion Refinement and Grain-Coarsening Resistance

The refinement of inclusion phases via rapid solidification has been demonstrated in a study by

scanning transmission electron microscopy (STEM) of a centri fugal ly-atomi zed high-sulfur 303

stainless steel (Table 1) [1]. While coarse powder particles which solidified in a

TABLE 1 Steel Compositions (w/o)

_C _S Si Mo ^ 1 12. Cu Ji

303 Stai nless 0..06 1.60 0,.028 0.34 0.62 8.68 17.,31 0.37 - 0.15 0.78 0.032

4Mo -0.25C 0,.26 0.51 0,.005 0.007 0.05 - 4.10 - - - 0.001

M2 Matri x 0,.51 0.22 0,.015 0.006 0.19 _ 4.,68 2.72 1.07 2.02 - - 0.011

171
cellular/dendritic tnode produced 0.1 ym intercellular manganese sulfides as a part of the solidification process, smaller powder particles which solidified in a massive (segregationless) mode produced 100 A sulfides, the latter being precipitated from the solid state. After hot consolidation

of -140 mesh (<105 ym) powder by 10:1 extrusion at 900°C, an average sulfide size of ~0.1 ym was maintained. This is in contrast to the 10-100 ym sulfide stringers normally obtained by conventional ingot processing.

As was reviewed recently in the context of alloy design guidelines for rapid solidification

[2], the low solubility of inclusion phases, which causes them to precipitate early (and in coarse

form) during conventional solidification, provides the basis for unusual coarsening resistance once they are finely dispersed (metastably) via rapid solidification. The rate of inclusion-particle

coarsening is proportional to the matrix concentration of the rate-controlling species (the species

with the smallest product of diffusivity and matrix concentration), and this matrix concentration is controlled by the inclusion-phase solubility product. For example in the rapidly solidified 303

stainless steel, the ability of the manganese sulfides to maintain during hot consolidation a particle size two to three orders of magnitude smaller than in conventionally processed material arises

from the low matrix sulfur concentration associated with the low sulfide solubility in austenite.

The refinement of inclusion phases down to sizes of 0.1 ym and below renders the particles less

harmful as initiating defects in fracture and fatigue processes, and also allows them to play a new

and benefical role in the control of microstructure. At inclusion volume fractions of the order of

0.1%, refinement to this size level provides a strong grain-boundary pinning force which results in unusual resistance to grain coarsening at high austenitizing temperatures [3]. Figure 1 summarizes

the high-temperature austenitic grain-coarsening behavior (1 hour treatments) of a 4Mo-0.25C martensitic steel (Table 1), comparing the grain size of conventionally processed material with that of

hot extruded (9:1 reduction at 900°C) centri fugal ly-atomi zed RSP material of two powder sizes, -200

mesh (<75ym) and -140+200 mesh (75-105 ym) [4]. Dissolution of carbides above 1025°C leads to rapid

grain coarsening of the conventional steel, whereas the additional pinning effect of the dispersed

inclusions causes greatly retarded grain coarsening, and this is slower in the material produced

from the finer powder. While there is evidence of enhanced pinning in the region of the prior

powder-particle boundaries [3], which may contribute to the powder size dependence of the graincoarsening behavior, a comparison of hot-extruded and hot-pressed RSP powders of a 9Ni-4Co steel [5]

revealed substantial grain-boundary pinning within the powder-particle domains as well. Hence,

finer dispersions of inclusions formed by solidification at higher supercoolings may also contribute

to enhanced grain-boundary pinning in material formed from finer powders.

Fig. 1 Comparison of austenitic grain

size of conventional and RSP 4MO-0.25C steel after 1-hour

austenitizing treatments [4].

Austenitizing tempera ture,°C

Fig. 2 Comparison of austenitic grain

size of conventional and RSP

M2 matrix steel after 1-hour

austenitizing treatments [6].

172
Figure 2 shows a similar comparison of the grain-coarsening behavior of a matrix tool steel

(VASCO-MA) derived from the matrix composition of M2 (Table 1) [6]. Again the rapid coarsening of

the conventional material marks the dissolution of alloy carbides, while the RSP material

(consolidated by blank die compaction and 4:1 extrusion at 900°C) exhibits greatly retarded

coarsening; in this case the coarsening is less sensitive to powder size. Results of a STEM analysis of the dispersed phases remaining after austenitizing at temperatures above the carbidedissolution temperature are summarized in Table 2 [6]. Light-element analysis was permitted by the

use of an ultrathin-window X-ray detector. Though sulfides make a significant contribution to the

boundary-pinning particles at 1180°C, at still higher temperatures the predominant particles are silicates which occasionally possess substantial sulfur contents. Table 2 lists the observed size

range of each inclusion phase as well as the total average particle diameter 3 and estimated total

volume fraction f at each temperature.

Under conditions of particle-pinning-controlled grain coarsening, the limiting grain diameter

at which the total pinning force balances the driving force for grain growth is expressed by [7]:

D

JL ^ fA A\

6 f ^2 ( 1 )

where Z is a grain-size distribution factor defined as the ratio of the largest to the average grain

diameter. The high-temperature grain-coarsening behavior of the RSP material illustrated by Fig. 2 is in qualitative agreement with eqn. 1 in that the observed D is proportional to the 3/f determined from the data of Table 2. Adopting an empirical Z parameter of 1.5<Z<2.0 based on previous

studies of austenite grain coarsening [8] predicts a grain size smaller than that observed by a factor of 2 to 5. Thus, the observed particles are more than sufficient to account for the observed

coarsening resistance. The discrepency may lie within the accuracy of the particle volume-fraction

estimates in as much as these are subject to the uncertainty of foil thickness and degree of both

particle-image overlap and truncation in the thin-foil STEM observations.

TABLE 2

Austeniti zi ng Temperature (1 hr.)

Dispersed Phases in RSP M2-Matrix Steel [6]

Stable Particles d,pm

1180°C MnS 0.1-0.

VS 0. 1-0.5

(Mn,V)S 0.15-0.5

SiO^ 0. 1-0.5 d = 0.18pm

MgO'SiO^ 0.2-0.

f « 3.7x10"^

MnS'MgO^'SiO 0.2-0.

VS*SiO,

2 0.15-0.6

1220°C SiO^ 0.1-0.

MgO'SiO^ 0.2-0.

d = 0.19pm

VS'SiO^ 0.2-0.

f - 2.5x10"^

MnS’SiO^ 0.2-0.

1260°C SiO^ 0. 1-0.5

d = 0.23pm

MgO* SiO^ 0.2-0.65

f « 2.2x10"^

VS'SiO^ 0.2-0.

173
While further refinement of inclusion size would be desirable, the observed dispersion of inclusion particles and associated high-temperature grain coarsening demonstrate that phases which

normally play a deleterious role can be made to perform a beneficial function in microstructural

control through rapid solidification.

Mechanical Properties

Provided stable boundary-pinning dispersions can be kept sufficiently fine so that they themselves do not limit fracture toughness, significant toughness improvements should be expected from

the proper exploitation of the grain-coarsening resistance of RSP steels. It is well established

that the sharp-crack fracture toughness (Kic) of medium-carbon ultrahi gh-strength steels, such as 4340 and 300M, can be substantially enhanced by the use of high austenitizing temperatures up to 1200“C [9-11]. This toughening effect can be attributed primarily to the dissolution of soluble

void-initiating particles [11,12]. However, even though Kjc is increased , the blunt-notch toughness

(e.g. Charpy energy) and tensile ductility (reduction in area) may decrease due to the excessive

grain coarsening which normally accompanies such high austenitizing treatments [12,13]. A recent investigation of the toughening effect in conventional 4340 [5] reveals that the improvement of Kjc

with high austenitizing treatment persists through the first stage of tempering (e-carbide

precipitation) but is lost in the second stage (retained austenite decomposition) due to increased

intergranular fracture in the coarse-grained material. On the other hand, the grain-coarsening

resistance of RSP steels offers the prospect of using the beneficial solutionizing effect of high

austenitizing treatments without the deleterious effects of excessive grain coarsening.

The influence of high austenitizing treatments on the fracture toughness of the 4Mo-0.25C steel

of Fig. 1 was examined using the 75-105 pm powder size RSP material [5]. Room-temperature mechanical properties of the conventional and RSP steels tempered 2 hour at 150°C after austenitizing at 1025°C and 1150°C are summarized in Table 3. Due to some decarburization during rapid solidification processing, the RSP material was found to contain 0.03C less than the conventional steel of

Table 1. This accounts for the slightly lower hardness and strength level indicated in Table 3. The influence of austenitizing temperature on the room-temperature fracture properties of both

materials is depicted in Fig. 3. Sharp-crack fracture toughness, given as Kq in Table 3 and Fig. 3, was measured on standard pre-cracked single-edge-notched specimens of 0.41 in. thickness. The

higher Kq values in Table 3 fall just outside the range for valid plane-strain Kjq values for this

size specimen. In both the conventional and RSP materials, this steel exhibits a significant

increase in both sharp-crack Kq toughness and blunt-notch C\j energy with high austenitizing treatment. While the conventional material shows a substantial drop in tensile ductility (%RA) attending

the high-temperature grain coarsening, the coarsening-resistant RSP material is not appreciably

affected. Although the room-temperature Charpy energies of the two materials are equivalent after

1150°C austenitizing, the advantage of the finer-grain RSP steel shows up clearly in its superior

ductile-brittle transition behavior illustrated in Fig. 4. While shelf energies are similar, the onset of the transition is lowered by ~100°C in the RSP material, and the midpoint transition tem- perature is decreased by 60°C. This dramatic difference is too large to be accounted for by the

small carbon and strength differences between the two materials. Hence, the grain-coarsening

resistance of the RSP steel allows practical use of the toughening effects of high austenitizing

treatment without the adverse effects of grain coarsening on the tensile ductility and ductilebrittle transition behavior. More fundamentally, the results affirm that the influence of high

austenitizing treatment on sharp-crack toughness is independent of grain size while the blunt-notch

toughness and tensile ductility are grain-size sensitive.

The density of the consolidated RSP material was found to be 99.9% of full density, and

metal lography revealed occasional voids which were several microns in size. It is reasonable to expect better fracture properties after more complete consolidation. A similar level of porosity

was found in the M2 matrix steel of Fig. 2. Although this steel did not undergo any significant

improvement in toughness with high austenitizing treatments, the finer grain RSP material showed

some advantage in both Kig and Cy energy relative to the conventional material after high tem- perature austenitizing [6J. Efforts to obtain toughness improvements in similarly processed RSP 4340 and 300M steels have so far been hampered by higher levels of porosity and powder contamination

which severely limit the fracture properties in these high strength materials [5]. While the results

of the 4MO-0.25C steel demonstrate the potential of RSP steels, it is evident that consistent and

reliable methods of powder consolidation must be perfected to achieve reproducible results.

Oxidation Behavior

The fine inclusion dispersion and grain-coarsening resistance of RSP steels can also improve

oxidation behavior, as demonstrated by cyclic oxidation kinetic experiments on the RSP 303

stainless steel of Table 1 [14]. The rate of weight change with 20 hour cyclic exposures to pure

174
Kq Cv RA

(a) (b) (c)

Fig. 3 Influence of austenitizing temperature on room-temperature fracture properties of 4MO-0.25C

steel, 150°C temper; (a) sharp-crack Kq toughness, (b) Charpy impact energy, (c) tensile

reduction in area [5].

^ ^ ^ ^

-100 0 100

T,°C

Fig. 4 Charpy ductile-brittle transition curves for RSP and conventional 4MO-0.25C steel after

1150°C austenitizing, 150°C temper [5].

TABLE 3

Mechanical Properties of 4Mo--0.25C Steel [5]

Aust. Temp. Hardness > Rc '^y» UTS, ksi %RA CVN,ft-lb Kg,k!

Conv. 1025°C 45.8 195 247 53 11.5 68

1150°C 46.3 197 246 42 22.6 99

RSP 1025°C 46.0 183 227 56 15.3 73

1150°C 44.7 181 229 51 22.4 95

175
oxygen at 900“C was compared with that of chemically-equivalent conventionally processed 304

stainless steel. Both materials were initially annealed for 6 hours at 1000°C, producing a grain

size of 170 uni in the conventional 304 steel and 5.6 pm in the RSP 303 steel. STEM analysis of the

RSP material after the 1000°C anneal revealed manganese sulfides of 0.2 to 0.5 pm diameter.

Fig. 5 Kinetics of cyclic oxidation of Fig. 6. Kinetics of cyclic oxidation

conventional 304 stainless of RSP 303 stainless steel [14].

steel [14].

The weight-change results obtained for three specimens of the conventional 304 steel are shown

in Fig. 5. The kinetics of oxidation are quite variable, indicating both rapid weight gains and

losses due to oxide spallation. Results for the identically treated RSP 303 steel are given in Fig.

6. The RSP steel demonstrated excellent oxidation resistance, showing consistent weight gains with

no visual evidence of oxide-scale spallation. SEM analysis of the oxide films indicated that the

superior oxidation resistance was due to the formation and growth of protective Cr203 and Si 02 scales, which were promoted by the sulfide dispersion and enhanced grain-boundary transport of Cr to the oxide due to the stable, fine grain size. There was also evidence of increased adherence of the

scale by the formation of intrusions of Si 02 from the external scale into the alloy, which formed

around the sulfides and along the closely-spaced grain boundaries, and which acted to key the scale

mechanically to the base alloy. The overall oxidation kinetics were equivalent to a conventional

310 stainless steel containing considerably higher Cr and Ni . These results suggest that RSP inclusion refinement and grain-coarsening resistance can be employed either to improve the oxidation

behavior of high-Cr alloys or to obtain equivalent behavior at lower Cr contents.

Conclusions

A primary benefit of RSP in hot consolidated steels is the achievement of fine dispersions of

stable inclusion phases. The low solubility of these phases in the austenite imparts an unusual

coarsening resistance to the dispersed particles such that they act as effective grain-boundary

pinning dispersions at very high austenitizing temperatures. As demonstrated for a 4Mo-0.25C mar- tensitic steel, the resulting grain-coarsening resistance allows substantial improvement of sharpcrack fracture toughness through the use of high austenitizing treatment without the reduction of

blunt-notch toughness and tensile ductility that usually accompanies excessive grain coarsening.

While both the RSP and conventional steels show an improvement of Kjc with high austenitizing, the

finer-grained RSP material displays a 60°C lower ductile-brittle transition temperature. Obtaining

consistent properties, however, requires reproducible methods of achieving full consolidation of the

RSP powders. The fine inclusion dispersions and grain-coarsening resistance also provide improved

oxidation resistance in RSP stainless steels through enhanced grain-boundary transport of Cr to the

protective oxide together with mechanical keying of the oxide at the dispersed inclusions and

closely-spaced grain boundaries.

Acknowledgements

This research is sponsored by the Office of Naval Research under Contract No. N00014-81-K-0013

and by fellowships from the AMAX Foundation. The martensitic steels were supplied by the Climax

176
Molybdenum Company, converted to powder by Pratt & Whitney Aircraft, Government Products Division,

and consolidated at Wri ght-Patterson Air Force Materials Laboratories. Results reported here are based on thesis research at MIT by J.L. Goss, C.Y. Hsu, and P.M. Fleyshman, and an undergraduate

research project by D. Eisen.

References

[1] T.F. Kelly and J.B. Vander Sande, Proc. 2nd Inti. Conf. Rapid Solidification Processing:

Principles and Technologies , R. Mehrabian, B.H. Kear, and M. Cohen, eds, (Claitor's, Baton

Rouge, 1980) p. 100.

[2] G.B. Olson, H.C. Ling, J.S. Montgomery, J.B. Vander Sande, and M. Cohen, in Rapidly Solidified

Amorphous and Crystalline Alloys, B. Giessen, B.H. Kear, and M. Cohen, eds. (North-Hol land, NY,

1982) p. 355.

[3] M. Suga, J.L. Goss, G.B. Olson, and J.B. Vander Sande, Proc. 2nd. Inti. Conf. Rapid

Solidification Processing: Principles and Technologies , R. Mehrabian, B.H. Kear, and M. Cohen,

eds. (Claitor's, Baton Rouge, 1980) p. 364.

[4] J.L. Goss, Austenitic Grain Growth and Precipitation Strengthening in Rapidly-Solidified

Alloys , M.S. Thesis, MIT, Cambridge, MA, Sept. 1980.

[5] P.M. Fleyshman, Fracture Toughness of Rapidly Solidified Martensitic Steels , M.S. Thesis, MIT,

Cambridge, MA, Sept. 1982.

[6] C.Y. Hsu, doctoral research in progress, MIT, Cambridge, MA.

[7] T. Gladman, Proc. Roy. Soc. A (1966) 294 , p. 298.

[8] T. Gladman and F.B. Pickering, J. Iron Steel Inst . (1967) 205, p. 653.

[9] G.Y. Lai, W.E. Wood, R.A. Clark, V.F. Zackay, and E.R. Parker, Met. Trans . (1974)^, p. 1663.

[10] W.E. Wood, Eng. Fract. Mech. (1975) 7_, p. 219.

[11] J.L. Youngblood and M.R. Raghavan, Met. Trans. A (1977) 8A , p. 1439.

[12] R.O. Ritchie and R.M. Horn, Met. Trans. A (1978) 9A , p. 331.

[13] R.O. Ritchie, B. Francis, and W.L. Server, Met. Trans. A (1976) 7A, p. 831.

[14] G.J. Yurek, 0. Eisen, and A. Garratt-Reed, Met. Trans. A (1982) 13A , p. 473.

177
1

STRUCTURE AND PROPERTIES OF RSP COPPER-BASED ALLOYS

I.E. Anderson, R.A. Masumura, B.B. Rath and C.L. Void

Physical Metallurgy Branch, Material Science and Technology Division

Naval Research Laboratory, Washington, DC 20375

1

ABSTRACT

The benefits of rapid solidification processing (RSP) for fabrication of Cu-Mn based

alloys for acoustic damping applications were examined. When these alloys are processed by

conventional casting or hot working techniques, the products exhibit either heavily-cored

dendritic or coarse banded microstructures . In this investigation pre-alloyed RSP powders were consolidated by hot extrusion to produce high damping alloys with Increased

homogeneity. The compositions selected for rapid solidification processing correspond to commercial damping alloys containing 58Cu-40Mn-2Al and 37Cu-55Mn-4Al-3Fe-lNl , in weight

percent. The microstructures and properties of the powder extrusions were compared to the

ingot metallurgy products. These results permit an assessment of the potential advantages

of RSP and advanced consolidation techniques for damping alloy development.

i

INTRODUCTION

The damping of elastic stress waves in solids in the acoustic frequency range (20 to 5000 hz) can

be used to control unwanted noise and vibration. The decrease in amplitude with time can be the

result of several different types of loss mechanisms. One such mechanism involving magneto-elastic

interaction has been commercially exploited. During the past several years, this development of the

basic Cu-Mn alloy has lead to two types of high damping materials [1]. The first is a cast alloy,

known as "Sonoston" with composition, 55Mn-37Cu-4Al-3Fe-lNi (wt.%) developed by Stone Manganese

Marine, Ltd. The other is a wrought alloy, known as "Incramute" with composition, 58Cu-40Mn-2Al

(wt.%) developed by the International Copper Research Association. Based upon the Cu-Mn binary

system, damping in these alloys is dependent upon the interaction between an elastic wave and

microtwin domain walls formed by the antiferromagnetic Mn-rich regions when given a specific thermal

treatment [2-4]. The physical properties of these two alloys, "Sonoston" (Mn-rich) and "Incramute"

(Cu-rich) have been characterized and documented. These compositions were selected to examine the

effect of RSP on acoustic damping and to compare the RSP alloys with those fabricated by ingot

metallurgy processes.

Powder metallurgical methods have been employed in two previous studies of the Cu-Mn system

[5,6]. In both cases, elemental powders were consolidated by die compaction and sintering. After

sintering, the densities ranged from 90 to 95% of the theoretical values indicating a high degree of

porosity which is known to improve damping at the expense of strength and toughness. To obviate these

difficulties, the present approach was to produce rapidly solidified powders from pre-alloyed stock

and to consolidate the particles by extrusion at 50-150°C below the solidus. Using pre-alloyed RSP

process has the advantage of producing an homogeneous and fine grained microstructure which should

lead to improved corrosion resistance, ease of fabrication and increased damping properties.

I

POWDER CHARACTERIZATION

Ingots of the Cu-rich and Mn-rich alloys were made with compositions corresponding to their

commercial counterparts. Rapidly solidified alloy powders were generated by centrifugal atomization

and quenched in flowing helium gas at the Pratt and Whitney facility at West Palm Beach. Particle

cooling rates were estimated to be near 10^°C/sec. [7]. Chemical analyses of the two powders yielded

compositions to within 1% of that required.

178
The powders were screened by the producer to remove debris and non-typlcal particles larger than

80 mesh before consolidation. Particle size analysis showed that the powder samples closely followed

a log-normal distribution with a maximum size of about 200 pm. The mean particle diameter for the

Mn-rich alloy was 46 pm; larger than the Cu-rich alloy powder, whose mean diameter was 29 pm. These

distributions appear to match the typical characteristics of a centrifugally atomized powder [8].

Optical metallography of individual powder particle sections permitted an examination of the

solidification structure of the individual particles. Figure 1 shows primarily dendritic

solidification with a secondary dendrite arm spacing of about 0.5 pm for the Mn-rich alloy particles.

Similarly, the Cu-rich alloy also exhibited a fine microstructure that appears to be cellular [8]

,

with a cell size of about 1 pm. Internal porosity was observed in some of the Cu-rich and Mn-rich

alloy particles. These types of pores may result from entrapped quenching gas or gas dissolved in the

molten alloy which then form internal bubbles on solidification as shown in Figure 1. The impingement

of molten particles on solidified particles may cause "splats" as can be seen in Figure 1.

X-ray diffraction measurements of both alloy powder samples showed a retained fee structure at room temperature. The ternary Cu-Mn-Al isothermal section for 20°C [8] shows a yCfcc) + o-Mn two phase mixture for the Cu-rich alloy. The Mn-rich alloys contains a significant amount of Fe and Ni (4

wt.% total) in addition to the Cu, Mn and Al, and the equilibrium phases are only known empirically.

Due to line broadening and weak high-angle reflections, the lattice parameters of the fee phases could

not be calculated accurately. Generally, the unit cell of the Mn-rich alloy appears to be slightly larger than that of the Cu-rich alloy which is consistent with the similar atomic volumes

of Cu and Mn.

EXTRUSION

The powders were consolidated at the Experimental Metals Processing Laboratory at Wright-Patterson AFB. The screened powders were placed originally into stainless steel containers,

evacuated, and sealed under vacuum. The cylindrical containers were pre-heated to 800° C for 2 hours

prior to extrusion. The Cu-rich powder sample was used in the initial extrusion attempt at 800°C with

a cross section reduction of 12:1 and a final desired diameter of 1". Unfortunately the stainless

steel can ruptured, exposed the billet to the atmosphere, and resulted in an irregular extrusion with

a reduced diameter. As a consequence, the Mn-rich powder sample was recanned in a copper container

with improved co-reduction properties and a successful extrusion was produced at 800° C with 12:1

ratio. The measured densities relative to the commercial alloys were 99.6% for the Mn-rich and 99.1%

for the Cu-rich alloys. The microstructures of the as-extruded alloys are shown in Figures 2 and 3. Both extruded alloys show an equiaxed structure with grain sizes of about 10 iim and about 1-2 pm for

the Cu-rich and Mn-rich alloys, respectively.

HEAT TREATMENT

The first step in the thermal schedule for the powder extrusions was an homogenization anneal.

Since the y~phase (fee) is required for damping in Cu-Mn based alloys, the high temperature anneal was intended to minimize any microsegregation in the as-extruded microstructure. Encapsulated Cu-rich

alloy samples were annealed at 800° C for one hour and water quenched. This is consistent with the

recommended treatment for the commercial wrought product. The homogenization anneal promoted abnormal

grain growth in the Cu-rich alloy microstructure, resulting in large grains, about 500 pm in diameter,

with occasional clusters of small grains, about 20 pm in diameter, as shown in Figure 2. The grain

interiors showed some evidence of dispersed stringers, presumably oxide, but appeared to be free of

precipitates while some grain boundary regions exhibited isolated precipitate particles. X-ray

diffraction measurements indicated that the structure was primarily single phase fee.

The Mn-rich alloy responded to the high temperature anneal by forming a duplex microstructure. An annealing temperature of 900°C (0.95 Tsolidus^ chosen. This is similar to the normalized

temperature selected for the Cu-rich alloy anneal. Solidus temperatures were measured by differential

thermal analysis with powder samples, where TgQ^idus (Cu-rich) = 870°C and Tsolidus (Mn-rich) = 955°C.

Encapsulated Mn-rich samples were annealed at 900° C for 1 hour and water quenched. The resulting

duplex microstructure, shown in Figure 3, consists of dark constrast phase regions of about 8-10 pm in

diameter in a light contrast matrix. The dark phase appears as equiaxed particles and phase

boundaries or precipitates are not apparent in the matrix.

Because X-ray diffraction measurements of the annealed and quenched structures of the Mn-rich

alloy revealed primarily fee reflections, other microanalytic techniques were utilized to identify the

phase regions. Electron channelling patterns revealed the fee structure of the dark phase regions. A

(110) pole pattern of the "particle" phase region shown in relief can be seen in Figure 4. The erose

surface condition of the Interparticle region produced by electropolishing prevented the determination

179
RSP:

15 pm

As Extruded

50 ym

RSP: Homogenized

(794 “C, 1 hr)

Water Quenched

50 ym

WROUGHT: Homogenized

(800“C, 1 hr)

Water Quenched

Figure 2. Microstructure of Cu-rich alloy.

180
I 1 20 ym

RSP: As Extruded

I-20 ym

RSP: Homogenized

(900 °C, 1 hr)

Water Quenched

Figure 3. Microstructure of Mn-rich alloy

60 ym

CAST: Homogenized

(843 °C, 1 hr)

Air Cooled

Back Scattered Image (110) fee Channelling Pattern from

Grain A

CHEMISTRY (wt%)

Mn Cu A1 Fe Ni

"A" 54.8 36.4 2.8 4.7 1.3

"B" 46.0 46.7 3.8 2.2 1.3

nominal 55.0 37.0 4.0 3.0 1.0

Figure 4. Crystallography and Chemistry of Mn-rich alloy.

181
of the matrix phase structure, denoted as "B" in Figure 4, by electron channelling. Further

development of thin foil preparation methods is planned to identify the phases with electron

diffraction. The results of an electron microprobe analysis of the phase regions, are given Figure 4. The particle phase composition. A, is similar to the initial alloy composition, although lower in A1 and higher in Fe. A significant composition shift between the particle phase and the matrix, B, was observed in the microprobe results. The apparent matrix composition suggests that other phases may exist in the microstructure which have not been resolved by this method. The cast microstructure [11] was described as essentially a two-phase mixture of a y-phase (fee) and a 3-phase (bcc). A third

phase, an inter-metallic compound, Cu3 Mn2 Al, was also Indicated in the report as associated with the 3-phase. While the heavily cored microstructure of the casting is distinct from the equiaxed, duplex

structure of the extrusion, the identification of the phases [11] is consistent with the microanalytical results of this study.

The aging treatment selected for these alloys was based on the binary Cu-Mn alloy behavior [10]

and commercial recommendations. An aging treatment is required to establish an optimum morphology for acoustic damping. The desired morphology contains a coarsened "tweed-like" structure of localized

Cu-rich zones, separated by less than 1000 angstroms, in an antiferromagnetic Mn-rich matrix within

each grain. The Mn-rich matrix should consist of randomly oriented microdoraains. Acoustic energy

dissipation is related to elastic wave interaction with the domain walls where the wall motion is impeded by the Cu-rich zones. Implicit in this mlcrostructural model is that the aging treatment

should be terminated prior to significant precipitation of the oe-Mn phase, which would deplete the antiferromagnetic species.

The response of hardness and damping to aging treatments is summarized in Table I. The acoustic

damping measurements were performed using the band-width technique of a cantilever beam in forced resonance [12] . This method for determining the loss factor, a measure of the damping, was used

because it can cover a wide range (~100 to 4000 hz) of frequencies. Values presented in Table I are averaged over the frequency spectrum. Based upon X-ray diffraction results, o-Mn precipitation occurs after an increase in the hardness of the alloy [10]. Thus hardness can be used to monitor, at least

in a qualitative sense, the ot-Mn precipitation and, hence, to select an aging treatment. Table I lists

the loss factor and hardness (DPH) , as a function of aging time. It is clearly evident that the initial rise in hardness (0 to 3 minutes) is the critical aging time in contrast to the recommended

aging times of 8 hours for the Cu-rich alloy and 4 hours for the Mn-rich alloy [11,13]. The

accelerated aging response of the Mn-rich alloy may be caused by the enhanced solute migration

kinetics of fine grained materials due to grain boundary diffusion. In the Cu-rich alloy, the rapid

or-Mn precipitation may be related to the high density of the oxide inclusions acting as nucleation

sites, as evident in the longitudinal section of Figure 2.

The comments in Table I on the mlcrostructural evolution during aging are based on the results of an optical metallography study. A series of four micrographs, shown in Figure 5 was selected to illustrate the aging sequence of the Cu-rich alloy. In that figure, after 10 minutes of aging, the grain boundaries are decorated with precipitates, presumably ot-Mn. Precipitation of ot-Mn occurs in the grains after 180 minutes and continues well beyond 600 minutes.

For the Mn-rich alloy, three major aging steps are Illustrated in Figure 6. Higher magnification

was used to show the detail of the evolution of additional phases during the later stages of aging.

The mlcrostructural features of the Mn-rich alloy that was aged for 5 minutes, shown in Figure 6, appear unchanged from the as-quenched state in Figure 3. Solute migration and the presence of an unidentified, fine-scale phase mixture between the grains can be observed after 30 minutes of aging.

The overaged condition, represented by the micrograph of the 1080 min. aged sample, is characterized

by a continuous grain boundary phase and uniform contrast grains. The apparent light-contrast zone which surrounds the overaged grains may be an etching artifact.

DISCUSSION

Rapid solidification processing of high damping Cu-Mn based alloys results in general

mlcrostructural refinement and increased alloy homogeneity. The microstructures of the conventionally

processed versions of these alloys are often highly segregated, which can contribute to stress corrosion cracking and low fracture toughness. A marked improvement in these properties is expected

to result [7] from the increase in homogeneity realized by the RSP approach used in this study.

Consolidated specimens of the RSP powders displayed an accelerated response to aging treatments.

In contrast to more typical RSP applications where a continuous precipitate arrangement is desirable,

the optimum morphology of a Cu-Mn damping alloy should be developed upon aging prior to precipitation

from a metastable solid solution [10]. Some factors that contribute to the rapid aging effect in the extruded samples are the small grain size in the Mn-rich alloy, and the presence of fine inclusions in

182
Figure 5

5 minutes 10 minutes

.- ’'*

‘‘ f ^ 4. \ t-.f' { *? * ~

«, V , ^

.1 T', V’ ‘ - / ^ V-* ^

. *. . f ^ ^

'

s. ^ . . V

>/

J.'

.' V., ' 'W

. ’w

600 minutes

20 ym

Cu-rich alloy homogenized at 800 for 1 hour and aged at 400 °C.

183
5 minutes 30 minutes 1080 minutes

Figure 6. Mn-rich alloy homogenized at 900 °C for 1 hour and aged at 450 °C.

TABLE I: AGING RESPONSE RESULTS

Cu-rich alloy (Homogenized 800 "C, 1 hr; Aging Temperature 400 °C) Aging Time

—4

Loss Factor x 10 Hardness Gomments

(rain) (DPH)

0 5.2 140 Pronounced Grain Growth

3 — 155

5 8.2 151 Apparent a-Mn in Grain Boundaries

10 — 152

20 2.8 159

300 — 168 General Precipitation

Mn-rich alloy (Homogenized 900 °G, 1 hr; Aging Temperature 450 °G) Aging Time

—4

Loss Factor x 10 Hardness Comments

(rain) (DPH)

0 4.3 159 Duplex Microstructure

3 6.0 178

5 4.5 187

15 — 192 Solute Migration

30 — 191

45 — 189 Apparent a-Mn in Grain Boundaries

200 — 200 General Precipitation

184
the Cu-rich alloy. To establish convenient peak aging treatments and to enhance the thermal stability

of the RSP alloys, the over-aging reaction associated with ot-Mn precipitation must be suppressed. The

results suggest a strategy to achieve this goal that includes high temperature annealing to encourage

grain growth, improved consolidation processing to reduce the inclusion content, and alternative alloy

selection to stabilize a single phase microstructure during the high temperature anneal. Clearly, the

design of the commercial alloys used in this feasibility study was restricted by conventional ingot

metallurgy practice. However, the enlarged composition range accessible to an RSP approach can allow

alloy selection based primarily on damping property improvement.

CONCLUSIONS

Relative to the conventionally processed materials, the microstructure of the rapidly solidified

and extrusion processed alloys have significantly increased homogeneity. The fine grained structure

of RSP extruded products can contribute to a rapid aging response which may limit thermal stability.

However, the results suggest a strategy to achieve a homogenous microstructure and an extended service

life in RSP high damping Cu-Mn alloys. The RSP approach can be used to select improved alloy

compositions without many of the limitations of conventional processing techniques.

ACKNOWLEDGEMENTS

We wish to thank Dr. R.A. Meussner of this Laboratory for his valuable work in preparing the

micrographs and D. A. Gerard of Michigan Technological University for his assistance during the

preliminary phases of this investigation.

REFERENCES

1. L. Me. Schetkey and J. Perkins, Machine Design , 50(1978) , 202.

2. J.A. Hedley, Metal Science Journal , _2(1968) , 129.

3. R.J. Goodwin, Metal Science Journal , _2(1968), 121.

4. D. Birchon, D.E. Bromley, and D. Healey Metal Science Journal , _2(1968) , 41.

5. J.L. Holman and L.A. Neumier, Modern Developments in Powder

Metallurgy , 10(1976) , 259.

6. P.W. Taubenblat, W.E. Smith, and R.L. Bieniek, Modern Developments

in Powder Metallurgy , 10(1976) , 241.

7. M. Cohen, B.H. Rear, and R. Mehrabian, Proc . of the Second Int’l . Conf . on Rapid Solidification Processing , (l980) , pp 1-23.

8. J.E. Smugeresky, Met . Trans . , 13A , (1982) , 1535.

9. Y.A. Chang, J.P. Neumann, A. Mikula and D. Goldberg, "Phase Diagrams and

Thermodynamic Properties of Ternary Copper-Metal Systems,"

INCRA Monograph VI, The Metallurgy of Copper, International

Copper Research Association, Inc., New York, (1979).

10. J.M. Vitek and H. Warliraont, Metal Science , 10(1976) , 7.

11. J.M. Langham, Stone Manganese Marine Limited Tech. Paper, no. 6 (1968).

12. R. Plunkett, "Measurement of Damping" in Structural Damping , (J.E. Ruzicka, editor), ASME, New

York, 1959.

13. INCRAMUTE Technical Data Sheet, International Copper Research Association, Inc.

185
FUNDAMENTALS OF RAPID SOLIDIFICATION

PROCESSING IN NICKEL-BASE SUPERALLOYS

E.H. Aigeltinger*, A.F. Giamei** and D.M. Dimlduk***

*PratC & Whitney Aircraft, West Palm Beach, FL 33402

**United Technologies Research Center, E. Hartford, CT 06108

***AFWAL/MLLM/Wright-Patterson Air Force Base, OH 45433

ABSTRACT

A general overview will be given of the benefit currently derived from Rapid Solidification

Processing (RSP) of nickel-base superalloys and what may be expected in the future. The

discussion is focussed (<n powder RSP as this is the predominant method of production of

RS nickel-base superalloys today. Some microstructural features of nickel-base superalloy powder, which can be produced by RSP, and of consolidated material are presented.

Two categories of nickel-base superalloys are discussed: those processed at temperatures

higher than the y' solvus temperature and intended for high temperature use (above about

1800°F) and those processed at temperatures lower than the y' solvus temperature and

intended for low temperature use (lower than about 1500°F) . The influence of RSP on the

microstructure of both types of materials is discussed and possible avenues for further

RSP exploitation of nickel-base superalloys are presented.

I. INTRODUCTION AND HISTORICAL PERSPECTIVE

The application of Rapid Solidification Processing (RSP) to nickel-base superalloys has been interesting, innovative and technologically effective. As the hardener content, volume fraction y' and

strength levels rose during the fifties and early sixties, superalloys underwent serious decreases

in ingot homogeneity and workability. Detrimental macrosegregation occurred in large ingots with low

cooling rates and long freezing times. Homogenization was out of the question based on the adverse

economics of smoothing out fluctuations of relatively slow diffusers over distances of 0.01-0. 10m [1].

The alloys also became more difficult to heat treat with only a small gap between the solution and

incipient melting temperatures. At the same time, alloy costs began to rise and turbine disks were

designed to higher levels of complexity, which increased the demands on workability.

The dilemma was solved by the sequential develojment of superalloy Powder Metallurgy (P/M) and hot, isothermal, low strain rate forming (GATORIZING ® ) [2]. The P/M approach started with a homogeneous

liquid alloy, produced by Vacuum Induction Melting (VIM), which was then atomized using a gas pressure differential. The powder was canned and consolidated by extrusion (and more recently by HIP) below

the y' solvus temperature. The "blow" of a homogeneous melt with a cooling rate of lO^-lO^oc/S, vs.

10 °C/S for an ingot, combined with a powder particle volume of 10“13m3^ an ingot volume of Im^,

led to elimination of macrosegregation due to solute convection. Any residual microsegregation (due

to dentritic growth) was eliminated during hot working and/or heat treatment.

The real benefit derived from P/M of nickel-base superalloys resulted from the 1-5 /im grain size

produced during powder consolidation by extrusion or HIP. This fine grain size imparted superplastic

behavior to these alloys in that the strain rate sensitivity was high, allowing uniform flow and

good fill of complex dies [3]. This coupled with low superplastic flow stresses made forging

(GATORIZING) of large turbine d isks feasible with presses of reasonable tonnage capacity.

At the same time, material utilization was improved by the application of sonic or near-net shape

processing which reduced the required amount of input material. P/M was a technological success for

turbine disk application.

In the late seventies, interest turned to even higher cooling rates, namely 10‘^-10^°C/S in processes

such as Rapid Solidification Rate [4] and 10^-10^°C/S for Laserglaze [5]. In the first process

rotary (centrifugal) atomization is carried out in flowing helium and in the second process a thin

surface layer is first melted and then resolidified. In these processes convective gas and/or con- ductive cooling accelerate cooling of the melt. The objective was to obtain high undercooling

186
resulting in extended solid solubility and cellular or amorphous structures. Potential applications

of RSP include low strategic metal content superalloys, higher temperature turbine blade alloys, related, specialized fabrication schemes [6] and hardened surface layers to resist erosion. To date

the turbine blade alloys have shown significant promise and will be considered in some detail here.

The turbine material of interest was [111] textured and based on the NiAlMo ternary system. The

RSP procedure plus extrusion for texture control plus conditioning heat treatment plus directional

recrystallization (DR) led to a columnar-grained or single crystal material with a strong [111] axial

texture. It was clear that the final texture and microstructure were strongly dependent upon both

alloy composition and starting RSP microstructure.

In the following sections some microstructural features of these alloys, from powder particles to material for blade application, will be presented. Features associated with the outstanding high

temperature properties will be emphasized.

II. MICROSTRUCTURAL FEATURES

Powder particles produced by rotary disk atomization have either a dendritic or a microcrystalline

morphology, as shown in the particle cross-sections in Figures 1(a) and 1(b), respectively. The

microcrystalline morphology is found in the smaller, more rapidly cooled particles and may be produced by homogeneous nucleation in the supercooled liquid. Concentration profiles across interdendritic regions and cell boundaries are shown in Figure 2 [7]

Segregation is very low in the

microcrystalline particles indicating they are solidified at an extremely rapid rate. Segregation

in the dendritic particles is considerably greater, however, it is substantially less than that found

in cast materials. The reduced segregation results in a significant increase in incipient melting

point for consolidated powder compared to cast material as shown for several alloys in Figure 3. This is important for subsequent processing as shown below.

For high temperature application transverse grain boundaries in the consolidated RSP alloys can be

eliminated in the solid state by directional recrystallization. This consists of passing the material

through a furnace hot zone with a temperature profile illustrated schematically in Figure 4.

a b

I lOfitn

Figure 1: Dendritic (a) and microcrystalline (b) powder particle cross-sections

The DR process is driven by reduction in grain boundary energy (area) associated with the finegrained consolidated powder structure. The grain boundaries are pinned by y' until the y' solvus

temperature is reached and y' dissolves. At this temperature the grain boundaries become mobile and

one or several grains across the sample cross-section grow to a single crystal or columnar-grain

structure, respectively as the material moves into the temperature gradient. Figure 4 illustrates

why the RSP produced increase in incipient melting temperature is beneficial and sometimes essential

187
for DR processing. The higher incipient melting point case allows a larger difference between the

y' solvus and incipient melting temperatures through which the material must pass in order to achiev<

DR. A higher Incipient melting point also results in a higher thermal gradient at the y' solvus

temperature which generally produces faster DR rates and lower defect content (twins and retained

grains) in the DR'd material.

MICROCRYSTALLINE DENDRITIC

0.2 ym

Figure 2: Concentration profiles (weight percent) across interdendritic and cell interfaces

in dendritic and microcrystalline powder particles, respectively.

Any features of the as-consolidated microstructure which influence grain boundary motion can in- ^

fluence DR response in terms of DR rate, resulting grain orientation or defect content of the result- | ing material. For example, initial grain size, primary y' geometry and dissolution kinetics, primarj^

carbide and a (BCC) precipitate geometry, etc. all influence DR response. These features are all

affected to some extent by RSP and subsequent TMP prior to DR processing.

2500 1

200 Al-W

Figure 3: Comparison of RSR and cast incipient melting temperatures for several Ni-base

superalloys.

188
7' Solvus Temperature

Figure 4: Illustration of temperature profiles and DR conditions for two different

incipient melting - y' solvus temperature "windows".

A representative as-consolidated micro structure is shown in Figure 5. Powder consolidation is typically done by extrusion at temperatures about 50°F below the y' solvus temperature. The resulting

structure is generally fine grained, the structure shown in Figure 5 having a grain size of 1-2 pm. The featureless grains in Figure 5 are y' while the remaining grains have a y matrix with a fine,

cooling y' precipitate. The latter grains consisted of y phase at the extrusion temperature and the

y' precipitate formed on cooling from this temperature. This fine-grained structure provides the

driving force for grain boundary motion in the DR thermal gradient, resulting in the columnar or single crystal grain structures produced by DR.

Figure 5: RSR, NlAlMo-base superalloy, as-extruded structure shown in TEM bright field.

If the crystals produced by this RSP-DR, solid-state processing are compared with those produced

from the melt by directional solidification (DS), at least two microstructural features are found

which are significantly different. Shrinkage voids formed during DS are not found in crystals produced by DR processing. These voids may weaken a cast material as they provide sites for initiation

of cracks, especially under fatigue and/or creep conditions. Another microstructural difference is

that DS crystals contain an array of low angle boundaries. Figures 6(a) and 6(b) while crystals

189
produced by DR are defect-free in this regard. Figure 6(c). The substructure density of cast crystals

is a function of both crystal orientation and growth conditions.

Figure 6: Cast, (a), (b) and DR, (c) microstructures after solution heat treatment.

Another difference between crystals produced by DS and DR processing is that of which crystallographic

orientations can be readily produced by these processes. At the present time the [112] orientation

of the NiAlMo-base alloys is of interest for its desirable balance of mechanical properties (see

Mechanical Properties Section). This orientation can be obtained by DR processing while it cannot be

produced by current DS practice because rotation away from [112] occurs during DS crystal growth.

This demonstrates how RSP-P/M processing can be used to achieve homogeneous, fine-grained material

essential for subsequent processing to a desired structural morphology.

The NiAlMo-base alloys, in the DR'd grain morphology, have stress-rupture capability about 150° better

than that of MAR-M200 at high temperatures (see Mechanical Properties Section). The unique compositional requirement for these alloys is saturation in Mo at a temperature above the intended application temperature. This ensures Mo supersaturation at the application temperature, resulting in a number of important microstructural effects, as described in the following.

Mo partitions strongly to the y matrix phase in preference to the y' precipitate phase. A concentration of about 20-25 atomic percent Mo is reached in the y matrix which expands the y lattice so that at saturation the y-y' lattice mismatch, AaQ, is about 1%, with the lattice of y larger than

that of y'This is the reverse of the sign of misfit in most superalloys. The lattice parameters

of the Ni-Mo binary and the approximate lattice parameter of y' are shown in Figure 7 as a function

of Mo content [8] . The lattice parameter of y' does not change significantly with Mo content while

that of y increases dramatically as shown. The two are equal at some intermediate Mo content. As

Mo increases to saturation, Aao increases to its maximum value of about 1%.

This is shown experimentally in Figure 8 where Aao is plotted as a function of saturation level for

a number of alloys. The as-solutioned data show a continuous increase in Aao to about 1% with increasing saturation level, because Mo is quenched in solution in both the y and y' phases on air

cooling from the solution heat treatment temperature of 2400°F. The standard heat treatment consists of an additional four hours at 1975°F after the 2400°F solution heat treatment. This allows

Mo present in excess of saturation to precipitate leaving the Mo concentration in solution constant

for supersaturated compositions. After such an equilibration heat treatment, Aao increases with Mo content until saturation is reached and then remains nearly constant with further increase in Mo content

This relatively large value of y-y' lattice mismatch for a Ni-base superalloy is accommodated at the

Y~y' interface by formation of an array of interfacial dislocations, as shown in Figure 9 which should

190
provide a substantial barrier to dislocation motion during deformation. It is expected that these

interfacial dislocations are associated with the high strength of the Mo-saturated NiAlMo-base

superalloys.

Figure 7: y and y' lattice constants as a function of Mo content.

Saturation

Increasing Level

of Saturation

I I O-S/im Figure 8: Lattice mismatch, Aag, as a Figure 9: Bright-field TEM photomicrograph

function of alloy saturation level. showing interfacial dislocations in NiAlMo

alloy after solution heat treatment for two hours at 1316°C (2400°F) followed by 10 hours

at 816°C (1500OF)

Despite the relatively high y-y' interfacial dislocation content, and hence interfacial energy, the

y' precipitate structure exhibits excellent thermal stability. This is in contrast to the poor y'

thermal stability and associated decrease in strength generally found in Ni-base superalloys when

Aao deviates from zero by more than about 0.1 percent [9]

A possible basis for this excellent thermal stability is the influence of Mo on the diffusivity of A1

in the y matrix phase. Figure 10 shows the diffusivity of A1 in binary Ni-Cr and Ni-Mo alloys as a function of Cr and Mo content. The curves show that both Mo and Cr produce a decrease in A1 diffusivity but that Mo has a much stronger effect than Cr. Mo at saturation in Ni produces an order of magnitude decrease in A1 diffusivity which is equivalent to about a 250°F increase in temperature capability in terms of resistance to y' coarsening at Mo saturation compared to that at low Mo content.

191
This reduced Al diffusivity may account for the excellent diffusion controlled creep resistance of

the NlAlMo-base superalloys through both improved Y' thermal stability and reduced dislocation mobility.

Figure 10: Diffusivity of Al in Ni-Mo and Ni-Cr binary alloys as a function of Mo or

Cr content at 1093°C (2000°F) and 1149°C (2100°F) [10]

Additional evidence which supports a low value of Al diffusivity in the Mo-saturated alloys can be

found in their inherent grain boundary instability in the form of cellular coarsening. The fact

that grain boundary cellular coarsening is severe indicates a high driving force for coarsening

exists, apparently as a consequence of the high y-y' lattice mismatch. So, when a path for rapid

diffusion is present, in the form of a grain boundary, microstructural degradation or cellular

coarsening, occurs. However, single crystals exhibit excellent microstructural stability, despite

the high driving force for y' coarsening.

At temperatures below about 1600°F, another significant strengthening mechanism is found in the

NiAlMo-base alloys. At these temperatures, several types of Ni-Mo precipitate particles form in the y matrix phase, in addition to the y' precipitate phase. The precipitates have the general microstructural appearance shown in Figure 11. These precipitates are present as one or more of the

phases: metastable Ni2>fo (body-centered orthorhombic), metastable Ni3 Mo (D022) and equilibrium Ni^Mo

(Dig) . The strengthening effect associated with these precipitates, in combination with large negative lattice mismatch, is substantial at temperatures below about 871°C (1600°F) (as shown in the

Mechanical Properties Section)

To this point, consideration has been given only to microstructural features of alloys heated to tem- peratures above the y' solvus at some time during processing, as the nature of the DR process necessitates exposure to these high temperatures. As mentioned, a number of microstructural features of

the consolidated material play a part in DR response, however, the mechanical properties of the re- sulting material do not appear to depend, to any significant extent, on microstructural features

established directly by RSP. In contrast, processing at temperatures below the y' solvus allows

features influenced by RSP to be retained in the final microstructure to a much greater extent. Structures in this condition have not been considered during the DARPA/AFML/PWA blade alloy development program, consequently mechanical property data are not available. However, a range of microstructural features can be obtained and some of these are briefly presented below.

192
O.Sfim

Figure 11: NixMo precipitate structure in a NiAlMo-base alloy after solution

heat treatment for 2 hours at 1316°C (2400°F) followed by 100 hours

at 760°C (1400°F) . Bright field (a) and dark field (b) TEM images.

Compacted Forged 60% Reforged 60%

1900° F (1038° C) - 200 Hours

lOO^im

Figure 12: Various volume fractions of unrecrystallized particles in a recrystallized matrix produced by compaction and forging of

NiAlMo-base superalloy RSR powder.

For example unrecrystallized powder particles can be retained, with various geometries and volume

fractions, in a recrystallized matrix, as shown in Figure 12. The precipitate morphology in the

recrystallized matrix and unrecrystallized particles can also be controlled by RSP and TMP. Figure 13

shows unrecrystallized particles in a recrystallized matrix consisting of fine grains of Y and y'.

These unrecrystallized particles have a very fine-scale morphology as shown in Figures 13(b) and 14.

The unrecrystallized area shown in Figure 14 has a y' matrix with some isolated areas of Y and a-BCC,

Mo-rich precipitate particles. The morphology of these unrecrystallized powder particles is established by the RSP response of the material. In addition the areas of y phase will develop the

NixMO (x = 2, 3, 4) precipitates discussed above on heat treatment at temperatures below about 871°C

(1600°F)

193
50>im 10*xm

Figure 13: Optical (a) and TEM replica (b) photomicrographs of vacuum-hot compacted

NiAlMo-base superalloy powder showing unrecrystallized particles in a recrystallized matrix.

Figure 14: TEM bright field image showing detail of unrecrystallized NiAlMo alloy powder particle.

It is apparent that a variety of microstructures can be produced by RSP followed by suitable TMP. However, correlation of these microstructures with mechanical properties, in most instances even in

general terms, is not possible at present.

III. MECHANICAL PROPERTIES

The NiAlMo-base superalloys have an advantage of approximately 83°C (150°F) in stress-rupture capability over MAR-M200 at high temperatures, as shown in Figure 15. This curve represents data for a number of alloys, such as alloy 185, having the composition: Ni-14Mo-7Al-6W-0. 04C in weight percent.

The curve is essentially the same for material tested in the [112] or [111] crystallographic orientations. These orientations have nearly the same stress-rupture capability and are both anomolously

strong in these orientations compared to MAR-M200.

194
(704*0 (816*0 (927*0 (1038*C)

1950*F (1066*0

2000* F (1093*0

100 2050*F (1121*0

2100*F (1149*0

10

Larson-Miller Parameter. C = 20.0

Figure 15: Larson-Miller plot comparing stress-rupture life

behavior of NiAlMo-base alloys and Mar-M200.

The [112] orientation has a better combination of mechanical properties for turbine blade application

in that the elastic modulus is lower than for the [111] orientation (36 vs. 42 X 10^ psi) which results in better LCF capability. Directional recrystallization processing is essential for producing

the [112] orientation because, as mentioned in Section II, this orientation cannot be produced by

current directional solidification processing practice.

Another property of importance, which is strongly influenced by solidification rate, is environmental

resistance associated with low concentrations of reactive elements. During single crystal production

by DS the slow rates inherent in this process combine long mold-metal reaction times with large melt/

mold surf ace-to -volume ratios. The combination of such favorable conditions for mold/metal reaction(s)

with considerable dendritic segregation makes it difficult to maintain reactive element concentrations

at desired levels or to maintain uniform concentration throughout a casting. These difficulties are

not encountered in RSP-DR produced material and uniform reactive element concentrations can be obtained in crystals made by this process.

Little mechanical property data are available for alloys processed and tested at low temperatures

(less than 871°C (1600°F)), however a peak in hardness is observed on heat treatment at about 760°C

(1400°F) after a sub-solvus partial solution at 1246°C (2275°F) for two hours. The hardness profile

spanning temperatures from about 593°C (1100°F) to 982°C (1800°F) for a time of 20 hours is shown in

Figure 16. This peak is associated with formation of Ni-Mo precipitate phases in the y matrix phase,

as described in the preceding section.

The yield strength of this alloy at 760°C (1400°F) after heat treatment for 50 hours at 760°C (1400°F)

is 1240 MPa (178,000 psi). This compares favorably with INIOO, on a strength-to-weight basis, which

has a yield strength of 1040 MPa (150,000 psi) at 704°C (1300°F)

In summary, the high temperature strength capability of these alloys is excellent. There also appears

to be potential for application at temperatures below about 871°C (1600°F) , provided further development of these alloys is undertaken with this objective in mind.

195
Alloy X Extruded + 2275* F(1 246'’ C)/2/0a + 20 HR Gradient Bar

Figure 16: Rockwell C hardness profile of NiAlMo alloy powder

consolidated by extrusion and heat treated for 2 hours at 1246 C (2275 F) followed by 20 hours at the indicated temperatures.

IV. FUTURE DIRECTIONS

For high temperature structural applications of crystals grown in the solid state by DR processing,

it will be important to control, reduce or eliminate defects such as growth twins or retained grains

in the directionally recrystallized product. The frequency of these defects is known to depend on alloy composition and TMP conditions. Additionally, it may become possible to favorably modify the

grain boundary/ stacking fault energy ratio to reduce twin fault probability. Alternatively, it should

be possible to control the prior and final crystal orientations and growth velocities to minimize

formation of growth twins and to better define the TMP conditions and textures to avoid formation of

large numbers of twin-related grains during deformation. These improvements should lead to enhancement of creep resistance, stress-rupture life, stress-rupture ductility and fatigue capability.

Alloy development will focus more on undercooling and its relationship with microstructure, rather

than cooling rate per se. Melt thermal history, minor elements in the alloy and ceramic environment

(if any) will be important areas of effort. Clean melting will be a major thrust area. It will almost certainly become possible to specify and control the level and distribution of reactive elements

in the final part. This combined with uniform distribution of major alloying elements, should lead

to significant improvement in oxidation/corrosion resistance of RSP materials.

Considerable effort will be directed at preservation of RSP microstructural features (e.g., by

Laserglasing or dynamic compaction) through consolidation and TMP, in order to derive the maximum

RSP benefit. Production of finer particles or generally faster cooling rates will result in greater

undercooling and the opportunity to obtain a wider range of useful microstructures. This will also

provide the opportunity for exploitation of the superplastic deformation behavior of microcrystalline

structures produced by RSP. Finally, development of dispersion strengthened structures and extended

solid solutions remain essentially unexplored areas of considerable potential in nickel-base superalloys produced by RSP.

V. SUMMARY AND CONCLUSIONS

Current interest in RSP evolved from the expected benefit to be derived from rapid cooling rates and successes achieved by the combination of conventional P/M and the GATORIZING process. The RSR process,

in combination with TMP including DR processing, when applied to NiAlMo-base alloy compositions, re- sulted in exceptionally strong Ni-base superalloys for high temperature application. The basis for

this strength lies in the detailed nature of microstructural effects produced by Mo saturation including large values of y~y' lattice mismatch and low A1 diffusivity in the y matrix. Considerable

potential remains for the application of these alloys at temperatures up to 871°C (1600°F)

196
Future effort will be directed at a number of areas such as reduction of defect content in solid-state,

DR-processed crystals, optimization of reactive element distribution for improved environmental

resistance, clean melting and exploitation of the deformation behavior of microcrystalline structures.

Finally, preservation of RSP-produced microstructural features through consolidation and TMP, in order

to maximize the direct benefit of RSP, will be an area of considerable future interest.

REFERENCES

1. M.C. Flemings, "Solidification Processing," McGraw-Hill, New York, 1974.

2. J.B. Moore and R.L. Athey, Design News, January 1970, p. 17.

3. N.E. Paton and C.H. Hamilton, "Superplastic Forming of Structural Alloys," AIME, Warrendale, PA,

1982.

4. A.R. Cox, J.B. Moore and E.C. van Reuth, Proc. of Third International Symposium on Superalloys,

Claitor's, Baton Rouge, LA, 1976, p. 45.

5. E.M. Breinan, B.H. Kear and C.N. Banas, Physics Today, November 1976, p. 44.

6. B.H. Kear and E.R. Thompson, Science, Volume 208, May 23, 1980, pp. 847-856.

7. R. Field and H. Fraser, private communication.

8. W.B. Pearson, "Lattice Spacing and Structures of Metals and Alloys," Pergamon Press, New York,

1958.

9. I.L. Mirkin and O.D. Kancheev, Met. Science Heat Treat., nos. 1 and 2, 1967, pp. 10-13.

10. E. Aigeltinger and P. Reynolds, unpublished research.

ACKNOWLEDGEMENTS

The views and conclusions contained in this paper are those of the authors and should not be interpreted as necessarily representing the official policies, either expressed or implied, of the Advanced

Research Projects Agency of the U.S. Government.

197
THE MECHANICAL BEHAVIOR OF RSPD MATERIALS

M.R. Jackson

General Electric Corporate Research and Development

Schenectady, New York 12301

R.W. Smashey

General Electric Aircraft Engine Group

Evendale, Ohio 45215

L.G. Peterson General Electric Gas Turbine Division

Schenectady, New York 12345

ABSTRACT

A number of Ni, Fe and Co-base alloys have been evaluated in structures formed by rapidly

solidified plasma deposition processing. AUoys have included both conventional superalloy

compositions and experimental compositions. Tensile and rupture behavior have been evaluated

for some alloys in orientations within the plane of the deposit. Results of tests to evaluate

thermal fatigue resistance are described.

Introduction

Rapid solidification has received considerable attention in the past five years for metal processing (1,2). The technique offers the possibility of achieving unique mechanical properties through microstructural

control. Ultrafine microstructures free of macrosegregation are produced in compositions developed for more conventional metals processes. More useful, potentially, is the increased range of compositions that can be generated and the metastability that can be produced in terms of extended solubility and nonequilibrium phase formation.

Plasma spraying in a low-pressure environment has been used to produce structural alloys by rapid

solidification plasma deposition (RSPD) (3,4). Ultrafine microstructures free of microsegregation have been

produced in thick deposits of nearly full density. For the past year, a program has been directed at examining

the usefulness of the RSPD process for application of high-temperature superalloy materials (5). Areas being

addressed have included the technology associated with the process, detailed physical descriptions of the

plasma gun, plasma-powder interactions, and the materials technology.

The present discussion will describe some of the effort dealing with the materials technology.

Specifically, mechanical behavior and modulus of elasticity have been evaluated for a variety of compositions. Possible influence on behavior due to unmelted powder particles incorporated into the RSPD

structure has been investigated. Chemistry effects on intermediate temperature (750-950C) ductility have

been considered. Bond strengths have been measured for RSPD materials deposited on non-RSPD substrates

of the same composition. Finally, the resistance of RSPD materials to thermal fatigue exposure has been

determined. Structures produced using a variety of plasma facilities were sectioned so that thermal stresses would be maximized either parallel to or perpendicular to the plane of plasma deposition.

Experimental Procedure

Powders of alloys to be deposited were gas atomized by AUoy Metals, Inc., Troy, MI, or Special Metals

Corp., Ann Arbor, Ml, or were vacuum atomized by Homogeneous Metals, Inc., of Herkimer, NY. AU powders

used for spraying had been sieved to recover the -400 mesh fraction, that is, particles 3^m and smaller in diameter. Gas analyses were performed on aH the powder lots.

Depositions were performed in a low-pressure environment using plasma guns manufactured by two

manufacturers. Conditions used during deposition are listed in Table I. Substrates were, steel or copper

plates, .3 cm thick and approximately 7 cm by 15 cm. The transverse direction between gun and substrate was parallel to the long axis of the plate. Total powder feed of as much as 2.3 kg was performed after the

plate had been preheated to ~850C using the plasma as the heat source. Pre-heating and deposition generally were complete within 20 minutes.

198
TABLE I SELECTED RSPD PARAMETERS

Plasma Gun A B

Plasma Gas Parameters Ar 1.4 1/sec He .3 1/sec Ar .84 1/sec H2 .14 1/sec Gun Voltage 55V 52-53

Gun Amperage 1300A 900A Chamber Pressure 60 torr 40 torr Powder-Feed Rate 2.9 g/sec 1.25 g/sec Powder Injection Internal, 20“ angle External

Carrier Gas Ar .12 1/sec Ar .025 1/sec Gun-to-Substrate Distance 25-33 cm 39 cm

Traverse Rate 15-20 cm/sec 5 cm/sec

Test bars for tensile, stress rupture, elastic modulus, and oxidation resistance were normally produced

with the test bar axis in the plane parallel to the substrate surface. Slugs were heat treated prior to machining to test bar geometry. In general, tensile tests were performed in an inert atmosphere or vacuum

using a strain rate of .04/min. Tests were conducted over a range of temperatures from 20C to 1060C.

Rupture testing was done in air at temperatures of 610C to 1060C.

Thermal fatigue behavior was measured in a simulated engine thermal shock (SETS) facility. Specimens

"5.1 cm in length were machined as a wedge tapering at a 15° angle from .51 cm wide at the base over a distance of 1.52 cm. The narrow end of the wedge was machined to a .064 cm radius. Wedges were machined

from a cylindrical shell deposit (Figure 1) such that the sample was oriented in a longitudinal or short transverse relation to the shell deposit. Again, samples were cut as blanks, heat treated, and then machined.

The SETS test consisted of heating the wedges end-on in gas burners from 425C to 1040C within a time period

of 10 seconds, holding at 1040C for 40 seconds, cooling to 480C in 10 seconds, and further cooling to 425C in another 10 seconds. The initiation of cracking is monitored by periodic visual inspection. Crack growth rates can be measured after initiation.

Bond strength was measured using a test bar similar to that in the other mechancial test programs.

Material for test bar blsinks was electrodischarge machined from a thick deposit on a thick substrate so that

the center of the gage to one end of the machined bar was RSPD material. Tensile strength and rupture

resistance of the bond test bars were measured.

Results and Discussion

AUoy Selection

A study to determine what components within advanced jet engines would combine substantial

performance or cost benefits with rapid production implementation indicated that stationary components in the hot section of the engine were high priority candidates. Initial design studies predicted that part life might be increased two- to foui^fold by use of a composite structure in place of current monolithic

structures. The composite would use materials with low thermal expansion coefficient on the hot side, and

high thermal expansion coeffient on the cool side. This ordering of materials would effectively lower thermal

strains produced by thermal gradients. The desired imt^ovement in part life would be achieved with little change in mechanical properties of the individual layers of the composite relative to properties of the present

monolithic alloys.

Based on the composite analysis, seven compositions were chosen for RSPD which com positionally are

identical to commercially available cast or wrought materials. They were chosen on the basis of mechanical

properties and thermal expansion behavior of conventionally produced materials. The alloys included powder

compositions equivalent to the compositions of Rene' 80, Rene' 120, Udimet 700, Heist elloy-X, Mar M-509, A286, and FeCrAlYTi material. These compositions are listed in Table II.

In addition, six experimental compositions were chosen for RSPD. These compositions, designated Cl

through C6, are also listed in Table II. Several considerations entered into selection of these alloys. It had

been observed previously that many of the strong superaUoys, such as Rene' 80 or IN -73 8, showed a loss in ductility at about 900C when tested as RSPD material. An alloy simulating the composition of the '3' phase of Rene' 80 showed good tensile ductility, while alloys with 1' content equal to or greater than that of Rene' 80

showed poor ductility (see Figure 2). One consideration for the alloys chosen was to determine whether other

strong alloys would show the poor ductility, such as those strengthened by (Cl based on IN-718), or by a Tirich y (C6 based on A-286 or U-212). Another consideration related to the expectation that composites at

199
TABLE n NOMINAL COMPOSITION OF EXPERIMENTAL ALLOYS

Alloys

Designation Ni Fe Cr Mo fb Ti

Weight Percents

Al Si Mn C B Co W

Rene 80 59.7 .2 14.0 4.0 — 5.0 3.0 0.2 0.2 0.17 0.015 9.5 4.0 0.03 Zr Rene 120 59.2 .2 9.0 2.0 - 4.0 4.3 0.1 0.1 0.17 0015 10.0 7.0 3.8 Ta 0.07 Zr

Hastelloy-X 46.3 18.5 22.0 9.0 - - - 1.0 1.0 0.10 - 1.5 0.6 Mar-M509 10.0 1.5 22.5 - - - - 0.4 0.1 0.60 0.01 60.89 - 3.5 Ta 0.5 Zr A-286 26.0 52.8 15.0 1.3 - 2.1 0.3 0.7 1.5 0.04 0.005 - - 0.3V FeCrAlYTi - 72.8 21.6 - - .5 4.6 - - - - - - ,5Y Udimet 700 52.1 <1.0 15.0 5.0 - 3.5 4.5 0.2 0.15 <.06 0.03 18.5 -

Cl 61.91 9.5 12.38 3.26 9.48 1.63 1.15 0.29 0.28 0.08 0.04 - -

C2 54.95 18.06 12.39 4.90 7.10 1.22 0.69 0.29 0.28 0.08 0.04 - -

C3 51.98 18.24 16.98 6.60 3.99 0.82 0.70 0.29 0.28 0.08 0.04 - -

C4 48.94 18.43 21.67 8.33 0.81 0.42 0.70 0.29 0.29 0.08 0.04 - -

C5 39.88 33.50 17.43 5.07 0.81 2.11 0.48 0.30 0.29 0.09 0.04 - -

C6 30.52 49.05 13.05 1.72 0.83 3.86 0.24 0.30 0.30 0.09 0.04 - -

high temperature might show significant metallurgical instabilities at interfaces due to interdiffusion. For

this reason, concentrations of many elements in the alloy series were held nearly constant to minimize the

diffusional effects. AUoy C4, which is similar to Hastelloy-X, includes small amounts of Al, Cb, and Ti to minimize interdiffusion with the other alloys. Alloys Cl through C4 make up one series, with C2 and C3

having compositions intermediate between Cl and C4, while alloys C4 through C6 make up a second series. The end point compositions were based on commercial alloys and were chosen to be consistent with the

preliminary design analysis. A final consideration in the choice of experimental alloys was the intention of

minimizing strategic element content, including Co, Cr, and Ta.

Structure Characterization

Oxygen and nitogen analyses of the powders are shown in Table III. As can be seen, oxygen for the -400

mesh particles is in the range of 200-300 wppm for most of the alloys, regardless of manufacturer. These

values are typical of laboratory scale (20-80 kg) atomizer production. Nitrogen appears to be more dependent

on source, being below 60 wppm for Cl through C6. Whether these lower values are due to processing or melt

stock is not known.

Particle size was measured in a particle analyzer. Despite the large range of compositions represented

by the six alloys. Cl through C6, the particle size distributions of the -400 mesh fraction are nearly identical.

Figure 3 shows the value of the particle diameter for which 90%, 50%, or 10% of the particles are smaller.

The average values for the six aUoys are; 90%/43.2y<^m, 50%/25.9>-m, 10%/12.8yu.m. The analyses shown

here are also typical of analyses for the commercial alloy compositions. Figure 4 shows a typical -400 mesh

powder fraction. There is a contradiction between SEM observations of particle size and those indicated by

the particle analyzer. The analyzer indicates 10% of the particle volume is in particles larger than 43 Am,

% y IN ALLOY

Figure 1. RSPD SETS wedge configuration. Figure 2. Tensile ductUity in RSPD Ni alloys.

200
TABLE ni GAS ANALYSIS OF -400 MESH POWDER

Content (wppm) Content (wppm)

Alloy 0 N Alloy 0 N

Cl 184 38 A-286 790 142 C2 253 21 HasteUoy-X 587 231 C3 252 14 Rene' 80 306 182 C4 269 16 Rene' 120 180 76 C5 286 56 FeCrAlYTi 673 120 C6 310 12 Mar M-509 321 153 U-700 141 30

TABLE IV GAS ANALYSIS OF RSPD DEPOSITS

Content (wppm) Content (wppm)

AUoy 0 N AUoy 0 N

Cl 341 57 A-286 540 867 C2 380 43 Hastelloy-X 196 125 C3 401 26 Rene' 80 299 161 C4 395 30 Rene' 120 276 115 C5 345 67 FeCrAlYTi 1887 632 C6 317 25 Mar M-509 79 211 U-700 116 65

yet particles larger than 40^Rm are almost never seen. The analyzer results are useful on a relative basis,

however. If the measured "10% smaller" diameter moves below about lO^m, the powder tends to feed poorly

in the powder feeders and through the powder-injection ports.

Gas analyses for deposits shown in Table IV indicate oxygen pick-up following deposition ranges from

near zero to as much as ~200 wppm. For the alloys shown, the degree of pick-up appears to be related to A1

content of the alloy. Pick-up presumably occurs on the heated deposit surface while the plasma is on adjacent regions.

Densities of as-sprayed materials were generally above 97% of the theoretical full density values. This

is shown in Table V, together with the results of density measurements on heat treated samples. High

temperature heat treatments produce essentially full density structures, as shown in Figure 5.

Microstructures of as-deposited materials have been discussed elsewhere in the conference proceedings

(6,7). The heat treated microstructures depend strongly on alloy chemistry. AUoys with low!':'' solvus

temperatures, or noX' at aU, can undergo significant in situ grain growth following deposition while the

structure is still hot. Grain sizes of 50-100/Mn result in these materials after heat treatment at 1200-1250C.

The ^'strengthened alloys are fine grained (-y^m) after deposition, and grow to only 25-50yLAm grain size

during isothermal annealing.TABLE V EFFECT OF HEAT TREATMENT ON DEPOSIT DENSITY

0 Density (g/cm ) after 1 h dry H„ treatment at AUoy As-Deposited 960C 1060C 1160G 1260

Cl 8.10 8.14 8.10 8.32 8.33 C2 8.04 8.00 8.01 8.12 8.35 C3 8.09 8.08 8.06 8.10 8.20 C4 8.09 8.11 8.13 8.19 8.24 C5 7.90 7.97 7.96 8.02 8.15 C6 7.95 7.91 7.89 7.97 7.96

3 Density (g/cm ) after 4 h treat at As-Deposited 982C 1093C 1149C 1204C 1260C

Rene' 80 8.11 8.11 8.11 8.14 8.19 8.22 Rene' 120 8.39 8.39 8.41 8.41 8.44 8.44 HasteUoy-X 8.22 8.25 8.30 8.30 8.30 8.30 Mar M-509 8.80 8.83 8.83 8.86 8.86 8.83 A-286 7.75 7.75 7.83 7.86 7.86 7.89 FeCrAlYTi 6.92 6.95 7.0 7.03 7.03 7.06

Figure 3. Particle size analysis of -400 mesh

powders of alloys C1-C6.

Figure 4. Scanning electron micrograph of -400 mesh Mar M-509.

201
I

Figure 5. Transverse micrographs of RSPD U-700 after 1250 C/4h + 1080 C/4h + 843 C/24h + 760C/16h.

Mechanical Behavior

Tensile results for the RSPD conventional alloys are shown in Table VI. All alloys were solution treated

at 1250C/4h, except A-286 and Hastelloy-X, which were solution treated at 1205C/4h. Commercial aging

treatments were given for A-286, Rene' 80, Rene' 120, and U-700. To compeu^e these alloys to their

conventionally processed counterparts, the following observations can be made. RSPD versions of Rene' 80, Rene' 120, and Mar M-509 are stronger than conventionally processed materials from room temperature to 650-750C, they are nearly equal in strength over the range 750C to 950C, and are somewhat weaker at higher

temperatures. Both Hastelloy-X and A-286 are weaker in RSPD form than conventionally processed over the

entire test range.

Properties of hot isostatically pressed (HIP) powder processed versions of these alloys are shown in Table VII. These data provide an estimate of the effect of unmelted powders on the mechanical behavior of RSPD materials, since HIP powder provides an essentially unmelted particle structure. Tensile strength

variations between HIP and RSPD are not large, although RSPD versions of Rene' 80 and Rene' 120 are

greater than HIP at high temperature. More significant, however, is the trend for tensile ductility of the

RSPD material to be greater than for HIP material. A severe unmelted particle structure may be expected

to approach the HIP behavior.

TABLE VI HEAT-TREATED RSPD CONVENTIONAL

ALLOY TENSILE PROPERTIES

Alloy

Test Temp

("F) ("O

0.02% Yield Strength

(ksi) (MPa)

0.02% Yield Strength

(ksi) (MPa)

Ultimate

Tensile Strength

(ksi) (MPa)

Elong. (%)

Reduction

of Area

(%)

A-286 72 22 83.5 575.7 95.1 665.7 144.2 994.2 22.5 42 7 1200 649 63.2 435.7 79.2 546.0 105.9 730 4 20.3 36.6 1400 760 49.0 337.8 60.0 413.6 63.7 439.2 12.8 35.3 1600 871 12.7 87.5 14.6 100.6 16.1 II 1.0 41.3 71 6 1800 983 4.0 27.5 4.6 31.7 10 4 71.7 38 2 500

FeCrAlY 72 22 49.3 399.9 53.5 368.8 79.5 548.1 20.2 59.0 1200 649 31.3 214.4 36.1 248.9 37.0 255.1 50 7 86.3 1400 760 11.1 76.5 14.7 101.3 16.7 804.6 74 3 94 1 1600 871 5.7 39.3 6.7 46.2 8.5 58.6 73.2 91.2 1800 983 2.5 17.2 3.8 26.2 4,4 30.3 80.3 74 5 Mar-M509 72 22 56.3 388.1 68.9 475.0 162.6 1121.0 24.3 22.3 1200 649 69.2 477.1 72.7 501.2 149.5 1030.7 22 1 20.2 1400 760 26.0 179.2 36.5 251.6 77.4 533.6 30.8 27.9 1600 871 32.4 223.3 38.7 266.8 47.0 324.0 19.2 20.2 1800 983 19.4 133.7 23.8 164.0 28.3 195.1 6.2 6.7 Hastelloy-X 72 22 48 8 336.4 54.3 374.4 110.5 761.8 44.8 52,0 1200 649 18.3 126.2 22.8 157.2 71.3 491.6 48 6 41 6 1400 760 30.8 212.3 36.7 253.0 56.9 392.3 27.0 34.0 1600 871 23.0 158.5 31.2 215.1 37.2 256.4 17.3 18.0 1800 983 12.0 82.7 15.7 108.2 17.4 119.9 2 9 3.0 Rene 80 72 22 127.6 879.8 138.8 957.0 210.7 1452.7 12.7 16.3 1200 649 90.3 622.6 115.8 798 4 183.2 1263.1 9.8 15,3 1400 760 96.0 661.1 113.1 779.8 140 9 971.5 4.3 7.0 1600 871 57.0 392.9 71.6 493.6 88.9 612.9 3.4 2.5 1800 983 17.7 122.0 29.0 199.9 46.0 317.2 4.6 2.0 Rene 120 72 22 105.2 725.3 125.0 861.8 180.2 1242.4 5 4 5.9 1200 649 113.2 780.4 133.1 917.7 178 9 1233.4 9.6 13.7 1400 760 110.3 760.5 135.8 936.3 151.0 1041.1 6.6 9 8 1600 871 72.5 499 8 93.8 646.7 111.9 771.5 5.5 6.3 1800 983 30.0 206.8 44.5 306.8 53.5 368.8 5.4 3.0

202
TABLE VII HEAT-TREATED HIP POWDER TENSILE PROPERTIES

Alloy

Test Temp

rP) (“O

0.02% YS

(ksi) (MPa)

0.2%

(ksi)

YS

(MPa)

UTS

(ksi) (MPa)

Elong. (%)

R, of A. (%)

A-286 72 22 66.1 455.1 74,8 515,7 113.9 785.3 20.2 43.6 1600 871 12,7 87.5 17,6 121 3 19.1 131.6 28.9 39.2 1800 973 6,9 47.5 8.8 60.6 10.0 68.9 37.1 36 6 FeCrAlY 72 22 51.3 353,7 56.4 388.8 80.3 553.6 21.6 54,4 1600 871 5.1 35.1 6.8 46.8 11.6 79.9 88.8 94,9 1800 983 3.4 23.4 3.8 26,2 4.2 28.9 60,9 75.0 Mar-M509 72 22 85.6 590.2 107,9 743.9 170.8 1177.6 8,1 10 9 1600 871 21.2 146.1 34.5 237,8 39.5 272.3 10.5 12.0 1800 983 14.0 96.5 18.0 124.1 22.8 157.2 5.5 6.9 Hasielloy-X 72 22 50.3 346.8 59.6 410.9 115.7 797,7 28.4 35.3 1600 871 21,3 146.8 24.8 170.9 25.2 173.7 5.7 5.9 1800 983 7.5 51.7 10.6 73.1 1 1.5 79,2 5.5 4.9 Rene 80 72 22 117.8 812.2 126.9 874.9 161.5 1113.5 5.7 4.8 1600 871 --no lest*- 1800 983 14.8 102.0 22.3 153.7 26.5 182.7 2.2 1.9 Rene 120 72 22 120.6 831.5 131.4 906.0 197.1 1359.0 7.1 9.8 1600 871 32,2 222.0 52.4 361.3 70.3 484.7 3.9 2.0 1800 983 18.2 125.4 28.3 195.1 39.8 274.4 7.1 7.0

•Specimens repeaiedly fail in ihreaded section Tensile properties of the experimental RSPD alloys Cl through C4 are shown in Figure 6. AH alloys were heat treated 1220C/lh + 900C/4h + 760C/12h. AUoy C4 is chemically similar to HasteUoy-X and alloy

C5 to A-286, with C5 being intermediate in chemistry between C4 and C6. The figure compares the

experimental alloys to RSPD and conventional versions of A-286 and HasteUoy-X.

0. 2 %

YIELD STRENGTH

(MPa)

0 . 2 %

YIELD STRENGTH

(ksi)

ULTIMATE

TENSILE STRENGTH

(MPa)

200

180

160

140

120

100

80

60

40

20

0

ULTIMATE

TENSILE STRENGTH

(ksi)

TEMPERATURE (°C) TEMPERATURE (°C)

Figure 6. Yield and ultimate strength and

ductility of experimental alloys

C4-C6 compared to RSPD and

wrought A-286 and Hastelloy-X. REDUCTION

IN AREA

(%)

TEMPERATURE (“C)

203
On the basis of yield strength, both C5 and C6 are considerably stronger them conventional wrou^t A286 or RSPD A-286. At elevated temperatures, C5 is as strong or slightly stronger than C6, even though the

concentration of precipitation strengthening elements is much greater in C6. Presumably, this is due to the

greater Ni content of C5 and its enriched solid solution strengthening (Mo, W, and Cr). Alloy C4 is stronger

than conventional Hastelloy-X over the entire temperature range. It is stronger than RSPD Hastelloy-X

below ~850C, and equivalent to it above that temperature. The pattern of ultimate tensile strength (Figure

6b) behavior for these alloys, compared to conventional and RSPD versions of Hastelloy-X and A-286, is identical to that for yield strength. In both figures, it can be seen that above 900C, alloy C4 is stronger than

both C5 and C6.

Ductility, as represented by reduction in area in tensile tests, is shown in Figure 6c. The experimental

alloys show reasonable ductility over the entire temperature range. Both C4 and RSPD Hastelloy-X show

considerable reduction in ductility with increasing temperature, but the loss is less severe for C4. No strong

ductility minimum is observed for any of the experimental alloys.

Similar comparisons can be made for rupture testing of the experimental RSPD aUoys, as shown in Figure 7. All data for Cl through C4 are for the high temperature solution treatment. There is a smooth

progression in properties with composition for C2 through C4, but Cl is markedly stronger than the others.

AUoy Cl compares favorably with lN-718, the conventionally wrought alloy on which the Cl composition is based. Also shown for comparison is a curve describing conventional Hastelloy-X. At the highest test temperature, C4 is approximately 2.5 parameters (R,h) below wrought Hastelloy-X. The shortfall in rupture

life decreases with decreasing temperature or increasing stress. An extrapolation of the data suggests C4

may be more rupture resistant than wrought HasteUoy-X at stresses greater than about 70 MPa (10 ksi). A similar plot is shown in Figure 7 for alloys C4 through C6. AUoy C6 is nearly one parameter better

than wrought A-286, and aUoy C5 continues in that same trend, being stronger than A-286 over the entire

test range, and stronger than wrought Hastelloy-X above 60 MPa (9 ksi). The rupture results for C5 and C6

are consistent with tensile behavior. However, C4 is stronger in tension than C5, but is weaker in rupture.

These same three alloys, C4 through C6, are compared to the RSPD versions of A-286 and Hastelloy-X

in Figure 8. The experimental aUoys are 1.0 to 2.3 (R,h) parameters better than the RSPD conventional

compositions over the range of test conditions where comparisons can be made. This corresponds to a threeto ten-fold improvement in time-to-failure. Also shown in the same figure is a curve representing the

average properties reported for RSPD Rene' 80 and RSPD Rene' 120. The data for the experimental alloys

suggest that compatible composition series between aUoys like Rene' 80 or Rene' 120 and aiUoys like H£isteUoy-X or C4 may result in properties useful in high temperature composite components.

Stress rupture results for the RSPD version of conventional aUoys are shown in Table VIll. In stress rupture, the RSPD materials are consistently weaker than comparable cast or wrought materials. RSPD

Rene' 80 and Rene' 120, high /', average about 1.6 (R,h) parameters lower than cast in the lower temperature

test range, and this difference becomes much greater at the high temperature end of the test range. RSPD

HasteUoy-X (no /') is about the equivalent of wroi^ht HasteUoy-X at 760C (1400F) but degrades to about 4 (R,h) parameters less than wrought at 980C (1800F). RSPD Mar M-509 varies from five to seven (R,h)

parameters weaker than cast, the difference becoming greater with increasing temperature. At 760C (1400F)

RSPD A-286 approaches the stress rupture strength of wrought A-286, however, at 870C (1600F), it is much

STRESS

(MPa)

P (R,h)

34 36 38 40 42 44 46 48 SO S2

STRESS

(ksi) STRESS

(MPa)

P (R,h)

34 36 38 40 42 44 46 48 50 S2

STRESS

(ksi)

P(K,h) P = T (20 + log t)

Figure 7. Stress rupture behavior of RSPD and wrought alloys.

204
TABLE Vni

HEAT-TREATED RSPD CONVENTIONAL ALLOY

STRESS RUPTURE PROPERTIES

Alloy

Test Temp.

(“F) (°C)

Stress

(ksi) (MPa)

Life

(h)

P(C--

(K,h)

= 20) (R,h)

Eiong. (%)

R. of A. (%)

A-286 1400 760 25.0 172.3 31.3 22.2 40.0 9.1 10.7 1600 871 6.0 41.3 43.8 24.8 44.6 60.7 63.3 1800 983 2.0 14.1 1289.6 29.4 53.0 — 47.1 FeCrAlY 1400 760 5.0 34.4 58.7 22.4 40.4 89.1 89.2 1600 871 3.0 20.6 32.1 24.6 44.3 44.5 40.8 1800 983 1.5 10.6 0.1 23.8 42.9 79.8 89.2 Mar-M509 1400 760 20.0 140.6 2058.6 24.1 43.3 5.6 1.8 1600 871 16.0 110.3 28.5 24.6 44.2 3.5 3.6 1750 955 7.5 51.7 64.3 26.8 48.2 4.3 1.9 1760 960 8.5 58.6 13.5 26.0 46.8 - 2.9 1800 983 5.0 34.4 43.7 27.2 48.9 0.54 0.9 1800 983 9.5 65.5 2.2 23.0 41.4 3.5 1.0 Hastelloy-X 1400 760 20.0 140.6 69.9 22.6 40.6 13.6 13.6 1490 811 10.0 68.8 185.5 24.1 43.4 4.1 5.2 1600 872 5.0 34.4 208.6 25.5 45.9 1.2 5.1 1600 871 6.0 42.2 38.8 24.7 44.5 4.2 2.2 1800 983 3.0 21.1 7.2 26.2 47.1 - 3.4 Rene 80 1600 871 34.0 239.1 7.6 23.9 43.0 9.4 8.6 1750 955 7.5 51.7 197.8 27.3 49.2 39.1 34.1 1760 961 8.5 59.8 79.3 27.0 48.6 24.4 26.3 1800 983 5.0 35.2 737.2 28.7 51.6 • 45.7 1800 983 12.0 84.4 12.2 26.4 47.6 12.5 17.1 Rene 120 1400 760 70.0 492.3 64.8 22.6 40.6 7.4 5.5 1600 871 34.0 239.1 25.4 24.5 44.1 11.3 7.6 1750 955 7.5 51.7 78.4 28.1 50.6 61.9 62.5 1760 961 8.5 58.6 63.1 26.9 48.4 69.6 49.6 1800 983 5.0 34.4 70.2 27.4 49.4 134.7 60.6 1800 983 12.0 84.4 7.6 26.2 47.2 55.2 40.1

I ^Threads broke off on one end after test weaker than wrought Hastelloy-X. FeCrAlY also is much weaker in stress rupture than HasteUoy-X sheet.

The stress rupture strengths of all six RSPD conventional alloys is a reflection of their very fine grain size which has been shown by many others to be detrimental to good-high temperature stress rupture and creep

strength.

Of special interest is the very high rupture ductility of the highT ' alloys, Rene' 80 and Rene' 120, in the

960C (1760F) to 980C (1800F) test range. An RSPD Rene' 120 sample tested in 980C (1800F) tensile yielded

elongation and R of A ductilities of 5.4% and 3%, respectively. The same material, sample geometry, and in the same heat treated condition yielded elongation and R of A stress rupture ductilities of 134.7% and 60.6%,

respectively, an increase of about 25X. RSPD Rene' 80, although somewhat less ductile than Rene' 120,

showed a similar behavior. Visual examination of the failed stress rupture bars showed the entire gauge

length to be uniformly elongated, with the actual failure occurring without significant localized necking.

This behavior shows the RSPD hi^ 6

' alloys to be very deformation rate sensitive, a property which may be

utilized in subsequent forming operations. It also suggests that the ductility minimum observed to occur in this type of alloy during tensile tests may not be of significance in regard to actual component part

performance since most components in actual operation are more rupture or creep limited than tensile

P|R,h|

34 36 38 40 42 44 46 48 50 52

Figure 8. Comparison of stress rupture

behavior of RSPD alloys Rene' 80, Rene' 120, A-286, and Hastelloy-X

to RSPD experimental alloys C4-

C6.

STRESS

(MPa)

STRESS

(ksi)

P(K,h) P = T (20 + log 1)

205
TABLE IX

RSPD/NON-RSPD HASTELLOY-X, TENSILE PROPERTIES

Sample Test Temp.

(“ F) rc)

0.02% YS

(ksi) (MPa)

0.2% YS

(ksi) (MPa) UTS

(ksi) (MPa)

Elong. 1 (%) R of A. (%) RSPD Hasielloy-X 72 22 48.8 336 4 54.3 374 4 110.5 761 8 44 8 520

1200 649 18.3 126.2 22.8 157.2 71.3 491 6 48 6 41 6 1400 760 30.8 212.3 36.7 253.0 569 392.3 27.0 34.0 1600 871 23.0 158.5 312 215.1 37.2 256 4 17.3 18.0 1800 982 12.0 82.7 15.7 108.2 17.4 119.9 2.9 3.0 Bond A 72 22 37.3 257.2 — 75.8 522.6 11.7 10.7 1200 649 22.3 153.8 26.5 182.7 49.5 341.3 8.0 70

1400 760 22.9 157.9 28.8 198.6 49.5 341.3 8 4 10.0 1600 871 20.4 140.7 26.6 183.4 36.4 251.0 4.4 49

1800 982 13.8 95.2 18.3 126.2 20.0 137.9 1.3 2.9 Bond B 72 22 37.9 261.3 — — 95.3 657.1 25.8 29 1 1200 649 22.3 153.8 29.4 202.7 54.5 375.8 10.0 18.7 1400 760 22.6 155.8 28.3 195.1 58.2 401.3 194 18.0 1600 871 22.5 155.1 27.8 191.7 36.9 254 4 7.3 7.0 1800 982 14.5 100.0 18.5 127.6 20.1 138.6 4.5 10

Bond A — Grit Blast, 60 Grit, SiC + Vaporhone + Single Pass RTA Clean Bond B - Surface Ground + Vaporhone + Triple Pass RTA Clean

(All samples heat treated 1205°C (2200°F)/4 h/HeQ)

The bond strength between an RSPD material and the substrate on which it is deposited has been

evaluated for HasteUoy X, shown in Table IX. Reverse-transferred arc (RTA) cleaning, using single or triple pass cleaning was performed prior to deposition. Without such a cleaning process, bond strength is genersdly so low that test bars cannot be produced. The strengths observed at room temperature are equivalent to those seen in tests of RSPD to cast IN-738 (4).

Results of the comparisons indicate:

100% RSPD Hastelloy-X is stronger than RSPD Hastelloy-X to non-RSPD Hastelloy-X bonded

material below 870C (1600F); this difference becomes insignificant at and above 870C (1600F).

RTA cleaning apparently has a beneficial effect on RSPD Hastelloy-X bond strength. The strength

of the bond made after the triple RTA cleaning pass is superior to that after a single pass. RSPD Hastelloy-X to non-RSPD Hastelloy-X bond ductilities are much lower than that of 100%

RSPD Hastelloy-X up to 870C (1600F).

AU types of RSPD Hastelloy-X deposited from an external powder injection gun, bonded or not, have a ductility minimum beginning about 870C (1600F), which is not typical of conventional sheet.

Triple RTA cleaning enhances bond ductility significantly when compared to single pass cleaning.

The results of stress rupture testing of wrought HasteUoy-X/RSPD HasteUoy-X tanary composite

samfies are given in Table X. The 760C (1400F) and 810C (1490F) test points are about the equivalent of monolithic RSPD HasteUoy- X; the 870C (1600F) point is a particularly low value suggesting some interface

problem. Of special significance are the low rupture ductility values measured on all three samfies tested

under the same condition.

TABLE X

HEAT-TREATED WROUGHT HASTELLOY-X/RSPD HASTELLOY-X

(METCO DEPOSITS) STRESS RUPTURE PROPERTIES

Specimen

No.

Test

(°F) Temp

(°C)

Stress

(ksi) (MPa)

Life (h) PIm C=20

Elong. (%)

R. of A. (%) 578 1400 760 20 137.9 55 40.4 7.5 7.0 579 1490 810 10 68.9 126 43.7 1,9 2.0 580 1600 871 5 34.4 6.2 42.8 0.3 2,6

Thermal Fatigue Resistance

The behavior of RSPD materials when subjected to thermal fatigue conditions is of special interest

when selecting component applications, including the potential advantages of multi-metal RSPD fabricated

designs. The behavior of fine grain material typical of that produced by RSPD requires evaluation.

Accordin^y, it is the objective of this task to provide additional insight into the thermal fatigue behavior of RSPD materials.

A test matrix was selected to evaluate the thermal fatigue resistance of several RSPD monolithic and j multi-metal material combinations via Simulated Engine Thermal Shock (SETS) testing which included the

206
following comparisons:

RSPD Rene' 80 wedges produced using different vacuum plasma deposition system vs. conventionally cast Rene' 80. RSPD Rene' 80 wedges oriented in the longitudinal and short transverse directions relative to the

original spray direction.

Longitudinal (SETS 0 and short transverse (SETS II) oriented standard SETS wedge samples were cut from RSPD cylinders made using three different processes. The longitudinal RSPD samples were compared to equiaxed Rene' 80 in SETS I. The short transverse samples were compared to equiaxed Rene' 80 in SETS test

II. Both SETS I and U were run to a total of 10,000 cycles, respectively, rather than the 5,000 cycles

originally planned. This extended testing was necessitated due to the failure to produce cracking in any of

the RSPD samples during the first 5,000 cycles of either test. The thermal profile used for the first 5,000

cycles consisted of a 10 sec heating period from 41 5C to 1040C, a 40 sec hold at 1040C, cooling to 480C in 10 sec and further cooling to the original 415C in another 10 sec. With this cycle, cracking was induced in the

conventional Rene' 80 samples of SETS I before the completion of the first 250 cycles. In SETS II, one of the

conventional Rene' 80 samples cracked before 250 cycles and the second cracked before accumulating 500

cycles. At the end of the first 5,000 cycles in both SETS I and II, none of the RSPD samples had yet cracked.

Edge views of each wedge after 5,000 cycles are shown in Figure 9. This confirms previous thermal fatigue

results for RSPD Rene' 80 (4).

In order to induce cracking in aU the wedges so that some estimate of relative crack propagation rates might be obtained, the thermal cycle for the next 5,000 cycles for both SETS I and SETS n was modified. In the modified cycle, the temperatures at the end of each period were held constant; however, the hold time

for each period was reduced from 10 sec to 5 sec, thereby cutting the entire cycle time in half and increasing

the relative heating and cooling rates. With this more severe cycle, all the longitudinal SETS samples

developed cracking before 6,000 cycles (less than 1,000 cycles at the more severe conditions). Under the

selected conditions and parameters the short transverse samples of SETS II showed one of two wedges made

using the external injection systems to have cracks before 5,300 total cycles and the second wedge cracked

before 5,600 cycles. One of two wedges made by the internal injection Low-Mach fM’ocess also cracked

before 5,600 cycles, the second cracked before 6,000 cycles. The Hi-Mach internal injection wedges lasted

the longest, but were found to have cracked before they went through a total of 7,000 cycles.

LOM »1ACH. RSPD, R’ 80 HIGH HACH, RSPD R' 80

SHORT TRANSVERSE SETS 5000 CYCLES

Figure 9. Edge-cxi view of SETS wedges after 5000 cycles to 1040C.

207
The results of SETS I and II lead to the following conclusions:

There was little difference in the relative ability of longitudinal and short transverse orientated Rene' 80 to withstand thermal shock.

The various RSPD microstructures produced by different RSPD equipment and process controls had

little effect on thermal fatigue resistance.

Without exception, RSPD Rene' 80 had greater than lOX improvement over conventionally cast Rene' 80 with respect to cumulative cycles to crack initiation. SUMMARY

Rapid solidification by plasma deposition (RSPD) produces a simultaneous RS process and consolidation

to nearly full density. The process can be applied to superaUoys with Ni, Fe and Co bases to produce low

oxygen (generally 5 400 ppm), homogenous fine-grained structures. Heat treatment can result in full densification.

Excellent combinations of tensile strength and ductility result for the RSPD structure. For some

compositions, RSPD is significantly better than conventionally processed material. The tensile ductility

minimum at 750-950C previously observed for 6

' strengthened alloys Rene' 80 and IN-73 8 has been found to be present principally in higher strength materials.

The low ductility in tensile tests is not observed in stress rupture, suggesting that its importance in component applications may not be great. Rupture resistance in general is lower than for conventionally

produced material, primarily due to the fine grain size. However, for many applications alloys that cannot be

conventionsdly processed to particular geometries can be produced by RSPD.

Bond strength and ductility are improved by reverse-transferred arc cleaning. Ultimate strength at 20C

for RSPD to conventional HasteUoy^X is 655MPa (-^95 ksi), essentially equivalent to earlier results for IN738.

In thermal fatigue exposure, RSPD Rene' 80 was uncracked in 5,000 simulated engine thermal shock

exposures, while the cast Rene' 80 cracked before completion of the first 250 cycles. Increasing the severity

of the thermal shock exposure did lead to cracking of the RSPD alloys. Whether stresses were maximized

parallel or transverse to the deposition plane had no effect on the behavior. Also, RSPD structures produced

in a variety of gun types behaved nearly identically.

Acknowledgement

This program has been funded by DARPA, through the offices of Lt. Col. L.A. Jacobson, with H.L. Gegel of AFML serving as technical monitor. The authors wish to acknowledge the aid and suppxjrt of P.A. Siemers, M.G. Benz, A.M. Johnson, J.R. Rairden, R.L. Mehan, R.W. Smith, J.S. Smith, L. Stoffer and D. Rigney, all of General Electric Company, and H. Lipsitt of AFML.

References

(1) D.J. Looft and E.C. vanReuth, "Rapid Solidification Processing: An Overview," 1 Proc. of the Int. Ccnf. on Rapid Solidification Processing, Reston, Va., p. 1, 1977.

(2) R.J. Patterson, n, A.R. Cox and E.C. vanReuth, "Rapid Solidification Rate Processing and Application

to Turbine Engine Materials," J. of Metals, p. 34, Sept. 1980.

(3) M.R. Jackson and J.R. Rairden, "Thermal Fatigue Resistance of Plasma Sprayed Thick Sup>eralloy

Deposits," Fall Meeting Program, TMS-AIME, Pittsburgh, p. 74, Oct. 1980.

(4) M.R. Jackson, J.R. Rairden, J.S. Smith and R.W. Smith, "Production of Metallurgical Structures by

Rapid Solidification Plasma Deposition," J. of Metals, p. 23, Nov. 1981.

(5) P.A. Siemers, "Rapid Solidification (RSPD) for Fabrication of Advanced Aircraft Gas Turbine

Components," Contract F33615-81-C5156, Interim Report, Jan. 15, 1982, Apr. 15, 1982, July 15, 1982,

Air Force Systems Command, AFWAL, Wright-Patterson Air Force Base, Ohio 45433.

( 6) T.L. Cheeks, M.E. Glicksman, M.R. Jackson and E.L. Hall, "Temperature Effects on Plasma Deposition

of Ni-Base Materials," this conference proceedings.

(7) A.M. Ritter and M.R. Jackson, "Microstructural Characterization of RSPD Structures," this conference

proceedings.

208
NOTCHED TENSILE PROPERTIES OF RAPIDLY SOLIDIFIED AND

INGOT-BASED ALUMINUM ALLOYS

Shinhoo Kang and Nicholas J. Grant

Massachusetts Institute of Technology

Cambridge, MA 02139

ABSTRACT

In the development of high-strength aluminum alloys, the focus has been on alloy

toughness as well as on strength, stiffness, fatigue behavior, corrosion, etc. In this

paper, tension properties (oyj5

, a^f^, elongation) of various rapidly solidified (RS) and

ingot-based aluminum alloys were examined and correlated with their notched tensile properties. Compiled data have shown that a linear relationship exists between the yield

strength (ays) and the notch yield ratio (o[»^js/aY 5

) ; there is a poorer correlation between

the ductility and the notch yield ratio. On the basis of Linear Elastic Fracture Mechanics

(LEFM), the linear relationship is extended to relate ay^ with Ktq. Under a given set of

processing parameters, the extended equation, = fCoyc), could be utilized to indicate

whether the toughness of an alloy (taking into account the related processing steps) in any

way encourages or discourages further development.

Overall results indicate that RS materials are not better than I/M alloys in terms of

toughness, and raise the question of "why." One possible reason may be the presence of excess oxide in the form of continuous stringers in the inadequately worked RS-PM alloy.

Tests are planned to attempt to evaluate the oxide issue.

I . Introduction

In alloy development for aircraft structural components, fracture toughness of the alloys has

been getting more attention along with strength and other properties than before. In order to determine alloy toughness, various testing methods are being used. Of these methods, the plane strain compact test is indicated to be one of the most accurate ways to measure Ktq of lower ductility alloys

such as the high-strength aluminum alloys. Unfortunately, not only is the preparation and testing of

the compact specimens costly and time-consuming, but it is not easy to obtain valid Kjq results with

respect to test specifications.

Sharp notch tension testing has been proposed to measure the fracture behavior of thick sections

under plane strain conditions. This test is recognized as an inexpensive and quick method to determine alloy toughness, using a reasonable notch radius (p $ 0.0007"). A circumferential notch is not

the same as an infinitely small crack radius; however, a well-established stress concentration factor

concept minimizes the difference between the two approaches. Even though there are many remaining

problems in notch tension testing, including the alignment of the specimen, the notch radius, etc.,

this method seems very convenient to obtain a fast, inexpensive idea regarding a material's toughness,

with a reasonable accuracy.

1 1 . Results and Discussion

Compiled data of the tensile properties of many aluminum alloys [1-5] revealed interesting relationships with the notch yield ratio (NYR = Kaufman and Johnson [6] had shown that the

ratio of sharp-notch strength to yield strength is more indicative of an alloy's toughness than is the

ratio of notch strength to ultimate tensile strength in comparing a variety of aluminum alloys. Figures 1 through 3 show typical relationships between NYR and tensile properties of I/M Al-Li-Zr alloys

[4]. It should be noted that ay5 and oyjc show a linear behavior with respect to the NYR, whereas the

elongation values are not very sensitive to the change of the NYR, particularly when the ratio is less

than one. Figures 4 and 5 are summaries of the alloy systems studied in this program. The following

observations are warranted:

209
80

- 70

« 60

50

n

1—1—^—!1—: ^ ^ ! r

70

^

i ! 1

^ 1

i 1 1

i

- • 4*

. d“°.

# C 60 “ - o ~

• • •• •

(/) . <?

*

• t %

£ ° 4

OIt50

• _ o _

• •

(/) o

A

it ^ —

>

•• 40 _ _

A

A

^ 1 1 ^ i 1 1 1 1 ^ 10 12 14 16 18

Notch Yield Ratio

10 12 14 Notch Yield Rotio

Fig. 1. vs. NYR for I/M Al-Li-Zr Alloy. Fig. 2. Oyg vs. NYR for I/M Al-Li-Zr Alloy.

Ext. Ext.

Ext. Ext.

Ext,

+ S.T.

+ S.T.

+ S.T.

+ S.T.

No Ext. + S.T

+ W.Q. (R.T.) + 200°C (V.T.)

+ W.Q. (200°C) + 200°C (V.T.)

+ W.Q. (275°C) + 225°C (V.T.)

+ W.Q. (275°C) + 275°C (V.T.)

+ W.Q. (R.T.) + 200°C(V.T.)

I

S.T, V.T.

= Extrusion after casting

= S.T. at 552°C (0.5 h)

= Various Time

12 14 Notch Yield Ratio

Fig. 3. Elongation vs. NYR for I/M Al-Li-Zr Alloy.

Fig. 4. ayr vs. NYR for Various Aluminum Alloys. Fig. 5. Elongation vs. NYR for Various A1 Alloys

210
1) a linear relationship exists between (or and the NYR,

2) relatively low-strength alloys, with high NYR values, show a less steep slope than do the relatively higher strength alloys with low values of the NYR,

3) rapidly solidified aluminum alloys do not show any advantage relative to the Oyr-NYR relationship as compared to the I/M aluminum alloys, and

4) elongation values are not good indicators of NYR values.

It has been commonly observed that rapid solidification processes result in higher volume fractions of oxides than that found in ingot alloys. When the oxides are present in the form of relatively continuous stringers, after comparatively small amounts of hot reduction of small billets by

the extrusion process, grain boundary decohesion associated with the oxide stringers can be enhanced

greatly, resulting in lower fracture resistance. The frequently observed highly delaminated fracture

surfaces of some RS aluminum alloys are good examples of longitudinal grain boundary decohesion.

Changes in powder-size distribution, oxide content, extrusion ratio and subsequent hot and cold work

(based on the powder compact diameter) and other processing variables should result in improvement in

the fracture patterns and behavior.

On the basis of Linear Elastic Fracture Mechanics, the relationship between ays and NYR was ex- tended to explore the relationship between oyr and Kj^, using presently existing equations. A relationship between omjc and Ktq, using a 50% notched cross sectional area, was derived by Irwin [7] and

later modified by Tada et al [ 8 ] on the basis of the stress intensity factor; the relationship follows:

= 0.454 (1)

where D is the major diameter of the notched tensile specimen. This equation has given reasonable es- timates of Kjr for values of CT[\j7s/^YS ^ without a plastic zone correction. However, for an alloy which shows sufficient plasticity, the plastic zone correction, with an extended notch depth,

must be taken into account. According to the elastic stress field solution, the size of the plastic

zone around the crack is

rYS 2it

2

( 2 )

For the plane strain condition, the normal stress on the 0=0 plane can be as high as 3ay5

. If the

effective yield stress in plane strain is 3ay 5 , the plastic zone correction becomes

2tt(3c YS‘

1 Sira YS

(3)

Practically, plane strain does not exist at the specimen surface (crack

! average effective yield stress is much lower than 3ay

5

. Irwin [9] used

i tive yield stress, which modified equations 2 and 3 to:

I

I

!

tip). As a co nsequence, the

1 . 68ay 5 (

/

2 /2 ay

3

) as the effec-

(4)

•ff

I A first-order correction by Irwin [9] for ry^ (= r,^^/2/2) as the effective increase of notch depth

j leads to:

!

I

I where

I

'^NTS _ X /, ovc 0.233 ~ 2 J

( 5 )

211
X

/irU^-Oy^

Further corrections by Tata et al [8] result in:

^NTS

®YS

4x (1

( 6 )

where

On the other hand, a linear relationship between Qy^ and NYR can be expressed below with reasonable

accuracy:

Oy^ = a(NYR) + b = ^(aj^-j-^/ay^) + b (7)

where a is the slope and b is the intercept on y axis. By combining equations 1, 6 and 7, the following equations are obtained:

For an alloy with cf[^j 5 /oy

2

$ 1.1,

0.454/D

'IC {o YS

b ay3) ( 8 )

For an alloy with 1.1 $ ^njs^^YS ^

u 5 _ 6 . 5 . 4 ^, . 2 ,.3

^IC ^°YS ^'°°YS ° ^YS ^IC

® °YS

‘ ^IC (9)

where I

p j! c

= , d = 4 (ttD)^, and e = 4ttD. °

!

Figures 6 and 7 were obtained from equations 8 and 9 after computer computations. Hahn and Rosenfield [10] derived a relationship based on crack opening displacement:

^IC= KsE(!)’'^d]'^2 (10) I

where E is the elastic modulus, f is the total volume fraction of second-phase particles and D is

the diameter of cracked particles: As pointed out in their work, equation 10 cannot explain the fact

that the value decreases with an increase in Oy,. at constant volume fraction of secondary phases.

However, the cross-hatched areas in Figs. 6 and 7, which have practical applicability, are instructive of this fact. It is also to be, noted that the left side of the inflection points in Fig. 7 must

be out of LEFM (Linear Elastic Fracture Mechanics) range, which is the basis of compact and sharp , notch tension testing. |

I In order to understand the relationship between the elongation values and the NYR (or toughness),

j as a first approximation, the elongation is expressed as follows: ;

°YS

+ 6 e

( 11 )

o'

212
Fig. 6. Predicted Kjq vs. ays for Aluminur.i Alloys with NYR $1.1.

Fig. 7. Predicted vs. aye- for Aluminum Alloys with 1 .[ $ NYR $ 1.7.

where e is elongation and a and g are constants. Equations 7 and 11 lead (if one lets NYR = x) to:

X =

( 12 )

Therefore,

g _ g/(-a)

ax - (g-b) (g-b)/a - x

where a(the slope) < 0 and a > 0. Hence, the final form of equation 13 will be:

B A - NYR (14)

where A = (g-b)/a > 0 (b >> g in most cases) and B = a/(-a) > 0. The plot of equation 13 is shown in

Fig. 8; the hatched area of Fig. 8 corresponds to the elongation-nonsensitive-region with respect to

Fig. 8. Elongation Behavior with Respect to NYR.

213
NYR. It is to be noted from equations 7 and 11 that the elongation behavior, i.e., equation 13, is

directly related to equation 11. Therefore, the elongation behavior (i.e., elongation values are not indicative of alloy toughness: [NYR]) is consistent with the fact that a linear relationship exists

between and NYR, as far as the first approximation, equation 11, holds. Typical example of the

elongation behavior with respect to the NYR can also be found in reference 11.

The derived relationships between Oy^ and Kj^ (equations 8 and 9) can be used for the following

1) getting a fast idea about a material's toughness (Kjp) along with strength values, not by

compact testing, but by several notch tension tests^at relatively low cost,

2) knowing about a material's limits in terms of strength and toughness for further development

of an alloy system or for comparison between competing alloys.

[1] D. P. Voss, "Correlation of Processing Parameters, Microstructure, and Mechanical Properties of

High Strength Powder Metal 1 urgi cal ly Produced Aluminum Alloys," Eng. Doctor Thesis, RheinischWestfal ischen Technischen Hochschule, Aachen, Germany, Nov. 1979.

[2] S. Kang and N. J. Grant, Unpublished work.

[3] H. G. Paris, F. R. Billman, W. S. Cebulak and J. I. Petit, Proc. of the 2nd Int'l Conf. on Rapidly Solidified Processing and Tech ., Reston, Va., March 23-26, 1980, p. 331.

[4] T. H. Sanders, Jr., "Factors Influencing Fracture Toughness and Other Properties of Al-Li Alloys'

Final Report, Naval Air Development Center Contract No. N62269-76-C-0271 , June 14, 1979.

[5] 0. R. Singleton, "High Strength, P/M, A1 Mill Products," Reynolds Metals Co., 2nd Interim Report

4/1/80 - 6/30/80, Contract No. F 3361 5-79-C-51 61

[6] J. G. Kaufman and E. W. Johnson, Proceedings, ASTEA, Am. Soc. Testing Mats., Vol . 62, 1962,

pp. 778-793.

[7] G. R. Irwin, "Onset of Fast Crack Propagation in High Strength Steel and Aluminum Alloys," Re- port No. NRL 4763, Washington, D.C., Naval Research Laboratory, 1956.

[8] H. Tada, P. Paris and G. Irwin, The Stress Analysis of Cracks Handbook , pp. 2. 6-2. 7, Hellertown,

Pa., Del Research Corporation, 1973.

[9] G. R. Irwin, "Plastic Zone Near a Crack and Fracture Toughness," Proc. of the 7th Sagamore Mater-l ials Research Conf ., Syracuse, N.Y., Syracuse University Research Institute, 1960. ji

I;

[10]

G. T. Hahn and A. R. Rosenfield, Met. Trans., Vol. 6A, 1975, p. 653.

[11]

W. Wang, "Advanced Lithium-Containing Aluminum Alloys Made from U1 trasonical ly Gas Atomized Par- ticulates," Sc.D. Thesis, M.I.T., Cambridge, MA, 1982, p. 131.

purposes

III. References

214

I 3

'
RAPIDLY SOLIDIFIED O.D, IN-100: STRUCTURE AND PROPERTIES

Charles H. Smith

General Electric Company

Corporate Research and Development

Schenectady, NY 12301

and

Nicholas J. Grant

Massachusetts Institute of Technology

Cambridge, MA 02139

ABSTRACT

An oxide dispersion stabilized IN-100 with 2 vol . % of 10 to 40 nm Y

2 O

3 particles added

was processed from fully alloyed rapidly solidified powders. Excellent dispersion of the

oxides viz. AI 0 O

3

, Y

2 O

3 and YAIO 3 was obtained, resulting in a particle diameter of 21 to

79 nm and an IPS less than 0.2 ym. Full characterization revealed the as-extruded bar to have submicron equiaxed grains with a substructure that was highly worked and strained, con- taining a high dislocation density with extremely high room- temperature strength. The optimum structure was achieved by using a solid-state crystal growing process named ZAP (Zone

Aligned Polycrystals). Large columnar grains with (100) fiber texture and a high grain as- pect ratio (GAR) were produced. The optimum 100-hour rupture strength was 152 MPa (22 ksi) at 1093°C (2000°F) and 172 MPa (25 ksi) at 982°C (1800°F). Rupture elongations varied between 2 and 9%.

1 . Introduction

Oxide dispersion stabilized (ODS) superalloys offer great promise as advanced high-temperature

materials. The ability to combine the benefits of oxide dispersion stabilization for high temperatures and y‘ strengthening for low and intermediate temperatures was demonstrated using the mechanical alloying process [1]. Since the first reported successes in 1970 by INCO, the mechanical alloying

technique for preparing OD superalloys has been applied by a number of investigators [2-4]. The process is a high energy dry milling process in which a mixture of metal powders and premelted master alloy

powders or particulates and very fine oxide particles is subjected to continuous deformation agglomeration, fracturing and rewelding in an appropriate grinding mill and ball charge in an inert atmosphere. Usually the powders are then consolidated by hot extrusion. The consolidation is followed by

zone annealing (ZAP) above the y' solvus to accomplish oriented recrystallization by formation of

large elongated grains.

Although these alloys have exhibited the best mechanical properties and alloy stability at high

temperature to date, the high cost of processing restricts their use considerably.

The purpose of this work was to develop a process and to fully characterize selected oxide dispersion stabilized Ni-base alloys. Full characterization involved determining the oxide (and y')

volume fractions and their respective interparticle spacings; the grain size, shape and texture; and

the nature of the (intragranular) structure. The relationships between structure and properties were

investigated.

2. Experimental

Rapid solidification via powder metallurgy eliminates large scale segregation common in cast

structures; each powder particle is in effect a micro-ingot containing the full compliment of alloying

elements in the correct proportions. Processing [5] was accomplished by starting with atomized prealloyed powders; vacuum atomized powders and RS, inert powders (Osprey Metals) were examined in this

215
study. The chemical compositions are given in Table 1. The powders were manufactured from master re- melt cast ingots. Subsequent processing to incorporate the fine (10 to 40 nm) Y 2 O3 particles by inert

gas ball milling was accomplished as described elsewhere [5]. The powder conglomerates after such a process are irregularly shaped, Fig. 1. Table 2 lists the consolidation treatments which were applied.

Table 1 . Compositions (Weight Percent) of Alloys Studied

Alloy Ni Co Cr Mo Al Ti W Ta , C B Zr Fe Y 203

N-2 (1 ) bal 20 20 10 0.5 - - - - - - 1.18

N-3 (

1 ) bal 20 20 10 0.5 - - - - - - 2.76

N-4 ( 2 ) bal 14 10 3 5.5 4.6 - 0.16 0.012 0.04 0.1 -

N-5 ( 2 ) bal 14 10 3 5.5 4.6 - 0.16 0.012 0.04 0.1 1.2

IN853 (3) bal - 20 - 1.2 2.4 - - 0.007 0.07 - 1.3

MA6000 E (3) bal - 15 3.5 4.5 2.0 4.0 2.0 - 0.01 0.15 - 1.1

(1) Solid solution strengthened: Homogeneous Metals Co.

(2) IN-100 type; y-y' with about 60 v/o y': Osprey Ltd.

(3) Mechanically alloyed M-MO materials: Benjamin et al [4,6,12]

Table 2. Processing Parameters - Evacuation and Extrusion

Evacuation

Extrusion

Maximum Force

(tons)

Alloy Temp (°C) Time (h) Pressure (ym) Temp (°C) Upset Running

N-2 650 24 0.5 1066 290* 280

N-3 650 24 0.5 1149 290* 280 i N-4 650 4 30.0 1149 330 275

'

N-5 650 4 30.0 1149 300* 270

300-ton press

3. Results and Discussion

As-Extruded Structures

Table 3 lists important structural characteristics of the alloys. The grain sizes of the ex- truded bars were of submicron dimensions and essentially equiaxed. Some reports [6,7,12] have claimed

that mechanically alloyed ODS alloys recrystallize after extrusion to submicron, equiaxed grains while

retaining a high dislocation density. It is difficult to conceptualize a recrystallized structure

with a high dislocation content. The TEM micrographs of Figs. 2 to 4 show highly deformed structures

with dislocation substructures. The dislocation content was estimated to be greater than lo” cm/cm^.

These alloys did not recrystal 1 ize during extrusion. It is the authors' opinion that recrystallization could have occurred during the time when the canned alloy is at the extrusion temperature (1 hour)

prior to extrusion. Grain sizes, as measured by the linear intercept method, yielded about 0.4 ym for

N-2, N-3, and N-5, and 0.6 ym for N-4.

Transverse to the high angle grain boundaries in alloys N-2, N-3 and N-5 are bands of dislocation

tangles which define a cellular type of structure. Cell sizes vary as follows: N-2 and N-3, 0.013

to 0.042 ym; and N-5, 0.009 to 0.039 ym. Figure 4 demonstrates that the cell walls along a grain's

high angle boundary are thin and well defined, while the cell walls across the grain width are thick

and irregular. Like most ODS alloys, the processing of the alloys in this work had a major effect on the oxide particle distribution and on the structure of the matrix. The range of particle sizes is shown in Fig. 5. A typical particle size distribution measurement is shown graphically in Fig. 6 .

216
Fig. 1. Appearance of ball mill blended

alloy N-5 with Y

2 O2

.

Fig. 2. TEM micrograph of alloy N-2; as extruded, transverse view.

Fig. 3. TEM micrograph of alloy N-5; as extruded, longitudinal view.

Fig. 4. TEM micrograph of as-extruded

N-5 showing cell structure.

Fig. 6. Typical oxide particle

size distribution for

alloy N-3.

Fig. 5. TEM micrograph of as-extruded

alloy N-2 showing fine Y

2 O

2 oxide dispersion.

217
Table 3. Structural Characteristics of Nickel-Base Alloys (As Extruded)

N-2 N-3 N-4 N-5

Ext. Temp. (°C) 1066 1149 1149 1149

Ext. Ratio (Area) 20:1 16:1 16:1 16:1

Theoretical Density (g/cc) 8.42 8.33 7.75 7.75

Measured Density (g/cc) 8.51 7.82 7.92 7.99

% Theoretical Density 101 94 102 103

Volume % Y 2 O

3 (added) 2.0 5.0 -

I 2.0

Avg. Grain Size (pm) 0.41 0.31 0.62 0.43 i Avg. Oxide Particle Size (pm) 0.079 0.034 0 . 112

* 0.021

Grain Aspect Ratio (GAR)

, , ** 1-2 1-2 1-2 1-2

IPS (p) 0.319 0.121 - 0.105

Oxide particles from surface oxidation, primarily in the form of Al^O^.

IPS = [(A - n ap)/n] ' , where 3

p = [(ird)/4] , d = mean oxide particle size,

A = area on micrograph, and n = number of oxide particles in A,

'k'k'k Density of INIOO (cast).

Phase Extraction and Analysis

Following methods outlined in the literature [ 8 ], separations of the oxides and y' from the alloys were made. For alloy N-5, y' precipitates were extracted from a sample heat treated as follows:

(1260°C/2h/AC + 1232°C/0.5hMC + 954°C/2h/AC + 843°C/24h/AC) . The analyzed w/o oxide values for N-3

and N-5 are 2.96 and 2.65, respecti vely . The y' content for N-5 was determined to be about 59%. TEM

studies revealed coarser precipitates varying in size from 0.1 to 0.4 pm dispersed throughout the matrix; they in turn are surrounded by hyperfine 200A y' particles. Both types of y' appear to be spher-i

ical and are readily seen in Fig. 7. Both fine and hyperfine y' precipitates are coherent; these LI 2 ordered structures introduce superlattice reflections in the selected area diffraction patterns. This

is shown by the diffraction pattern insets in Fig. 7 for a (100) zone pattern. The precipitates have

also been identified in dark field using a ( 100 ) y' superlattice reflection.

Fig. 7. TEM micrograph (A) of y' precipitates in N-5 after ZAP treatment; corresponding

(100) zone SAD pattern, inset, showing y' superlattice reflections. (B) Dark

field using a ( 100 ) y' superlattice reflection.

Stability Studies

The structural stabilities of the processed, as-extruded alloys were assessed by Rockwell "C" ' hardness (at 20°C) versus temperature profiles; see Fig. 8 . Alloys N-2, N-3 and N-5 maintain hardness

218
Fig. 8. Rockwell C hardness (at 20°C) versus annealing temperature

for alloys N-2 to N-5.

stability to 1093°C (2000°F) and above but at different hardness levels, based on dislocation, cell

boundary and grain boundary pinning by the fine oxide dispersions. Recovery, followed by partial

recrystallization, are responsible for the drop in hardness at the higher temperatures.

Room-Temperature Tensile Tests

The room- temperature tension properties for the several alloys are listed in Table 4, along with

data on the several mechanically alloyed OD materials. It was determined that the strength of ODS

alloys increased generally with increasing volume fraction of oxide [9,10]. Hence the increases in

strength of alloys N-2 and N-3 with 2 and 5% additions of Y20^, respectively, were expected. The ex- ceedingly high strength value of as-extruded N-5 was the result of its combined volume fraction of y'

(59%) and Y 203 ( 2 %) added. The tensile behavior of these alloys is typical of that of the more complex

ODS Ni-base alloys in which the yield strength approaches the ultimate strength, indicating less plastic yielding. Of importance in understanding this behavior is the state of the final material strucTable 4. Room-Temperature Tensile Data

Alloy Condition Ref.

Yield Strength

ksi (MPa)

UTS

ksi (MPa)

Elong.

(%)

R.A.

(%)

N-2 As Ext. 207 (1427) 209 (1441) 2

I 4

N-3 As Ext. 239 (1648) 246 (1697) 2 3

N-3 As Ext. + ZAP 3"/h at 1300°C 114 (789) 135 (926) 3 3 N-4 As Ext. 223 (1539) 223 (1539) 2 4

N-5 As Ext. 326 (2253) 326 (2253) 1 2

N-5 As Ext. + ZAP 3"/h at 1260°C + HT* 189 (1302) 201 (1389) 2

N-5 As Ext. + 1243°C/2h/AC + HT 156 (1076) 174 (1198) 3 4 i

1

1 IN853 As Ext. at 1066°C 295 (2037)

1

j IN853 As Ext. + lOh at 1177°C 6 176 (1212) 1

1 IN853 As Ext. + lOh at 1260°C 6 100 (689) 146 (1007) 12 16 i

! MA6000E As Ext. at 1038°C 13 186 (1284) 188 (1294) 2.5 3.0 j MA6000E As Ext. at 1010°C then hot rolled

at 1010°C

13 185 (1275) 186 (1284) 2.0 6.0

1

I *HT = 1232° C/0.5h/AC + 954°C/2h/AC + 843°C/24h/AC

1

1

1

219
ture which results from the effects of such variables as the temperature for hot extrusion, the deformation strain and the strain rates which are responsible for the level and distribution of the

stored energy of deformation. The severe cold working which results from the high strain rate extru- sion at the extrusion temperature leads to a submicron grain size and a high level of residual strain ! energy stored within the substructure. Tensile data after ZAP and recovery heat treatments are also

included in Table 4 for N-3 and N-5. Annealing at temperatures greater than 1093°C caused a drop in the yield/tensile ratio and an increased work hardening response. Clearly any strength increase de- | rived through particle dislocation interactions was lost due to recovery, recrystallization and grain

growth processes.

Recrystallization and El evated Temperature Stress-Rupture Strength

A large grain size is required for optimum high-temperature strength. The recrystallization tem- perature of the alloys was the most important variable of the starting materials. The recrystallization response for alloys N-3 and N-5 was studied and found to be different. For N-5, the response

was delayed and optimized at 1243°C. For N-3 the best response was at 1316°C. Alloy N-3 exhibited

a response that was entirely dependent upon the material's prior thermomechanical history. In con- trast, the delayed response of N-5 is attributed to the higher level of stored energy present and the

presence of a large volume content of precipitates which must first be solutionized before there can

be any response. A fine, uniform distribution of oxides proved essential for the desired storage of

energy. H|

Using a solid-state crystal growing technique called Zone Aligned Polycrystal (ZAP), large col^[

umnar grains were developed for alloys N-3 and N-5. Wi*

The maximum ZAP rate is thought to be controlled by the mobility of the interface between the

recrystallized and the as-worked material. For N-3 the interface temperature was >1300°C and speci- men movement varied from 7.6 cm (3.0 inch)/hour to 15.2 cm (6.0 inch)/hour; for N-5, the interface

temperature was >1260°C and motion rates varied from 3.8 cm (1.5 inch)/hour to 7.6 cm (3.0 inch)/hour.

Evaluation of ZAP treated material showed that the dispersoid size was not affected by the treat- ment. Dislocations were virtually absent and strong (100) recrystallization textures were developed.

The mechanism by which a very coarse-grained, elongated textured structure is produced from

another textured, highly cold-worked structure is unknown, but two factors are postulated as influencing grain-boundary mobility, namely, 1) the orientation of the grain boundary with respect to the

crystallographic orientation of the growing grain, and 2) the directional effects of the dispersoids.

The combination of these two effects yields not only irregular shaped grains but also a preferred

recrystallization texture (Fig. 9).

Stress-rupture behavior of N-5 is shown in Figs. 10 and 11. The stress for a 100-hour life at 1093°C is 22 ksi (152 MPa); the curves show the classical flat slopes of OD alloys. Figure 12 demonstrates that this alloy has comparable strength properties to the best previously reported alloys

(IN853 and MA6000E). Alloy N-5 has 59% y' while MA6000E is reported to have 39% y' [13], with a 100-

hour rupture stress of 24 ksi at 1093°C.

With a stress for 100-hour rupture life at 1093°C of 22 ksi, alloy N-5 has exhibited hightemperature strength values second only to MA6000E. Its strength exceeds that of some of the best

superalloys reported to date, namely DS MAR-M 200, DS eutectics (y/y'- 6) and single crystal alloys

(PWA 1409). See Table 5 and Fig. 12. The very simple method of processing [5] employed for N-5 makes

the method worthy of serious consideration.

Table 5. Comparison of Stress for 100-Hour Rupture Life for Several Superalloys

Alloy

100-Hour Rupture Strength, ksi (MPa)

Ref. 982°C (1800°F) 1038°C (1900°F) 1093°C (2000°F)

Mar-M 200 (DS) 14 29 (200) 20 (138) - MAR-M 200 (Cast) 14 26 (179) 16 (110) 9 (62)

INIOO (Cast) 14 25 (172) 18 (124)

1 MAR-M 509 (Co-base) 14 15 (103) 12 (83) 8 (55) i PWA 1409 (Mono crystal) 15 30 (207) - i

i

220
Fig. 9. The grain structure of N-5 test

bar after ZAP treatment (7.6 cm/h

at 1260°C) plus heat treatment.

Fig. 11. Log stress versus log rupture

time for alloy N-5 (ZAP treatment) at 1093°C.

Fig. 10. Log stress versus log rupture

time for Alloy N-5 at 982°C.

Fig. 12. Comparison of stress values for

lOOh life versus temperature for

a variety of Ni-base alloys.

4.

References

1. Benjamin, J. S., "Dispersion Strengthened Superalloys by Mechanical Alloying," Met. Trans., 1,

2943 (1970).

2. Gessinger, G. H., "Mechanical Alloying of IN-738," Met. Trans., 7A, 1203 (1976).

3. Kramer, K, H., "Dispersion Strengthened Superalloys," Powder Metallurgy International, 9, 105

(1977).

4. Merrick, H. F., Burwick, L. R. and Kim, Y. G., "Development of Oxide Dispersion Strengthened Turbine Blade Alloy by Mechanical Alloying," NASA CR-135150, January, 1977.

5. Smith, C. H., "Process Development and Characterization of an ODS Nickel Base Superalloy," Ph.D.

Thesis, MIT (1982).

6. Cairns, R. L., "Effect of Annealing on Structure and Properties of a Dispersion Strengthened

Superalloy, IN-853," Met. Trans., 5, 1677 (1974).

7. Sellars, C. M. and Petkovic-Luton , R. A., "Creep of Dispersion-Strengthened Alloys," Materials

Science and Engineering, 46, 74 (1980).

8. Kriege, 0. H. and Sullivan, C. P., "The Separation of Gamma Prime from Udimet 700," ASM Trans.

Quarterly, 61, 278 (1968).

221
9. Zwilsky, K. M. and Grant, N. J., "Dispersion Strengthening in Copper-Alumina System," Trans. AIME, 22, 371 (1961).

10. Zwilsky, K. M. and Grant, N. J., "Copper-Silica and Copper-Alumina Alloys of High Temperature

Interest," Trans. AIME, 209, 1197 (1957).

11. Smith, C. H. and Grant, N. J. , to be published. ^

12. Cairns, R. L., Curwick, L. R., and Benjamin, J. S., "Grain Growth in Dispersion Strengthened

Superalloys by Moving Heat Treatment," Met. Trans., 6A, 179 (1975).

|

I * 13. Merrick, H. F., and Kim, Y. G., "Characterization of an Oxide Dispersion Strengthened Superal loy,l

MA6000E, for Turbine Blade Applications," NASA CR-159493, May, 1979.

14. The International Nickel Co., Inc., "High Temperature, High Strength Nickel Base Alloys," Third « Ed., N. Y. (1977).

15. Ver Snyder, F. L. and Shank, M. E., "The Development of Columnar Grain and Single Crystal High \ Temperature Materials Through Directional Solidification," Material Science and Engineering, 6, i|

213-247 (1970).

5. Acknowledgements

The authors are pleased to acknowledge support by DOE

Howmet for a fellowship which permitted completion of this

under Contract No. research study.

EG-77-S-02-4466 and by
THE MICROSTRUCTURE OF TOOL STEELS PROCESSED FROM RAPIDLY SOLIDIFIED POWDERS

S. Kumar, A. Lawley and M. T. Koczak

Drexel Uni vers i^"y, Philadelphia, PA

ABSTRACT

A detailed microstructural examination has been made of T15 tool steel processed from

rapidly solidified powders. Gas atomized powders were screened into four size fractions

(^840ym; 250-840ym; 44-lOOym; S44ym) , hot isostatical ly pressed to full density and heattreated. Two HIPing temperatures, two austenitizing temperatures and three tempering

temperatures were included in the study. Atomized and consolidated powders were characterized by means of optical and electron microscopy, x-ray diffraction and hardness. Cubic

and hexagonal MC carbides are present in each size fraction of the atomized powder in the

form of interconnected carbide networks. Carbide size decreases and particle hardness

increases with decreasing particle size. MC, MgC and M2 aCe carbides are present after

HIPing in the form of individual particles in a matrix of a-ferrite. The volume fractions

of MC carbides ('\/16%) and MgC carbides ('\'17%) are essentially independent of prior particle

size and HIPing temperature. For both types of carbide and for each powder size fraction,

the carbide size distribution is skewed to higher values with an increase in HIPing temperature. At a given HIPing temperature, carbide size distribution is broader for the coarser

size fractions. M23 C 6 carbides can go completely into solution during austenitizing and do

not reprecipitate on tempering. No new carbide-types form on tempering. The final volume

fraction of MeC is a function of austenitizing and tempering temperatures.

INTRODUCTION

Powder metallurgy (P/M) high speed tool steels have been a commercial reality for over a decade.

Compared to their ingot metallurgy (I/M) counterparts, the P/M tool steels exhibit superior alloying

homogeneity coupled with fine grain size and a uniformity and fineness of carbide dispersions that cannot

be achieved by conventional casting and working processes. The microstructural characteristics of the

P/M material are reflected in improved grindabil ity, bend strength, tool life, and consistent machinability. In addition, a smaller out-of-roundness occurs during heat- treatment compared to the I/M

product; this is a consequence of the relatively isotropic microstructure of the P/M tool steels.

Commercial P/M tool steels are consolidated from gas atomized powder covering a broad particle

size range, typically <1200ym. Hence, a range of cooling rates is inherent in the as-atomized powder

prior to consolidation, with the finer particle sizes falling in the rapid solidification regime. A primary purpose of the present study is to examine the possible role of powder particle cooling rate

on the microstructure and properties of high speed tool steels. To this end, gas atomized powders

have been screened into specific size fractions and evaluated in powder form and after consolidation

to full density. Microstructural observations on T15 are presented here.

EXPERIMENTAL PROCEDURE

T15 tool steel powder of weight composition 1.52%C, 4.67%Cr, 5.22%Co, 12.19%W, 4.85%V, 0.31%Mn,

0.46%Si, 0.0063%02, 0 . 1 %N 2 , balance Fe was atomized in nitrogen. The four powder size fractions

studied were 2840ym, 250-840ym, 44-lOOym and ^44ym. Bars of each size fraction were HIPed at 1195°C

or 1130°C. The higher temperature is that used by Crucible, Inc. for its commercial P/M high speed

tool steels. The lower temperature was used as a possible means for controlling grain size and carbide size during densification. Heat-treatment studies of pieces cut from the HIPed bars included

two austenitizing temperatures (1176°C and 1225°C) and three tempering temperatures (538°C, 552°C and

565°C). The HIPing, austenitizing and tempering temperatures are coded in Table I.

An initial appreciation for overall powder shape and surface condition was obtained utilizing SEM.

High resolution SEM of polished and etched powder cross-sections was used for carbide characterization.

Two-stage replicas of the etched powder cross-sections were also examined by means of TEM. Carbide

223
Table I: Coding System for Powder Consolidation and Heat Treatment

Code

HIH2

A1A2

T1T2

T3

Interpretation t

Hot Isostatically Pressed at 1130°C V Hot Isostatically Pressed at 1195°C ( Austenitized at 1176 C for 4 mins, and oil-quenched L Austenitized at 1226 C for 4 mins, and oil-quenched

J Tempered for (2+2+2) hrs. at 538°C I Tempered for (2+2+2) hrs. at 552°C

\ Tempered for (2+2+2) hrs. at 565°C f

phases in the powder were identified by means of x-ray diffraction. For purposes o^ identification carbides were chemically extracted from the powder or consolidated material (1). Carbide volume fracand

s 1 ze d 1 s t r 1 bu 1 1 On s were determined by means of a combination of selective etching (1) quantitative metallography. ctv,iniiy

RESULTS AND OBSERVATIONS

1 . Atomized Powders

The powder particles are generally spherical in shape with small satellites present in each size fraction. Debye-Scherrer diffraction patterns of extracted carbides from each siL fraction are con5 the presence of cubic and hexagonal MC carbides. Powder cross-sections corresponding to three size fractions are shown in Figures 1, 2 and 3. A transition from cellular to a dendritic solidification mode is evident with decreasing particle size. In the cellular mode, carbides exist as interconnecting networks. Martensite needles are visible in the coarse powder. Figure 1

Figure 1: Atomized powder particle cross-section (SEM) from 550-

700um screen fraction. Etch:

Kalling's reagent + HCl/Picral.

Figure 2. Atomized powder particle cross-section (SEM) from 110-

130um screen fraction. Etch:

Kalling's reagent + HCl/Picral.

Figure 3. Atomized powder particle cross-section (SEM) from <37um screen fraction. Etch: Kalling's

reagent + HCl/Picral

2. HIPed Material

Micrographs (SEM) of polished and etched samples corresponding to the coarsest and finest size fractions are shown in Figures 4 through 7. The carbides are coarser following HIPing at the higher

224
Figure 4: HIPed material (HI) from 250-840ym

screen fraction. Etch: Kal ling's reagent. SEM.

Figure 5: HIPed material (H2) from 250-840pm

screen fraction. Etch: Kal ling's reagent. SEM.

Figure 6: HIPed material (HI) from -44um screen Figure 7: HIPed material (H2) from -44ym screen

fraction. Etch: Kalling's reagent. SEM. fraction. Etch: Kalling's reagent. SEM.

Figure 8: Size distribution for M0 C in HIPed (HI) Figure 9: Size distribution for MgC in HIPed (H2)

material; 250-840ym screen fraction. material; 250-840ym screen fraction.

225
temperature (H2). For a given HIPing temperature (HI or H2), the carbide size is essentially indepen

dent of prior particle size fraction. The matrix is a-ferrite and the carbides of the type MC, MgC,

and M23 C 6 . 1n all cases. Carbide volume fractions are summarized in Table II, as a function of particle size fraction and HIPing temperature. For MC and MgC carbides, the volume fraction is not a sensitive function of size fraction and HIPing temperature. Table II. However, the size distribution

is skewed to higher carbide sizes with an increase in HIPing temperature. Figure 8 and 9. For a given

HIPing temperature, the carbide size distribution is dependent on the prior particle size fraction;

it is broader for coarser size fractions. No selective ecthing technique exists for the unambiguous

identification of M23 C6

.

Table II: Carbide Volume Fraction as a Function of Powder Size Fraction and HIPing Temperature

Size Fraction HI H2

MC 250-840wm 0.162 0.179

f44iim 0.144 0.141

MeC 250-840vim 0.167 0.166

544pm 0.175 0.152

3. HIP + Heat-Treated Material

Representative micrographs (SEM) of the T15 tool steel after austenitizing (quenching) and

austenitizing + tempering ere compared in Figures 10 and 11, respectively. Martensite is visible in

the quenched condition. MC and MgC carbides are present at each stage of heat- treatment; in addition,

M23 Ce is retained after austenitizing at Al. After austenitizing at A2, all the M23 Ce carbides are dissolved and they do not reprecipitate during tempering. No new carbide-types form on heat-treatment

Volume fraction data for MgC carbides are given in Table III. A significant fraction of MgC goes

into solution during austenitizing and then reprecipitates on tempering. As expected, more MgC goes

into solution at A2 than at Al. The volume fraction of MgC increases with increasing tempering

temperature, and the size distribution is skewed to coarser particle sizes, irrespective of the prior

particle size fraction. A similar evaluation of the effect of heat- treatment on the volume fraction

of MC is currently in progress.

Table III: Volume Fraction of MeC as a Function of Heat Treatment

Size Fraction Al Al+Tl A1+T2 A2 A2+T1 A2+T2

250-840pm(H2) 0.0667 0.0721 0.0865 0.0543 0.0611 0.0752

544pm (H2) 0.0688 -- 0.1032 — 0.045 0.072

DISCUSSION

1. Atomized Powders

In gas atomization, a stream of liquid metal is broken-up by means of a high pressure gas stream.

The resulting liquid droplets spheroidize in free fall before solidification due to surface tension.

Particle collisions can occur before and/or after solidification (2) and this is responsible for

satellite formation. Cubic MC and hexagonal y-MC carbides have been identified in all the powder size

fractions examined. y-MC is not present in the HIPed material. This implies that y-MC is not an equilibrium phase but is present in the atomized powder due to the high cooling rate and that y-MC

carbide goes back into solution when thermal energy is supplied to the powders. A transition from a cellular to a dendritic mode of solidification with decreasing particle size can be explained in terms of the cooling rate T, the thermal gradient G, and the solid/liquid interface velocity R. Similar observations have been reported by Smugeresky (3). Levi and Mehrabian (4) have pointed out

that as the droplet size decreases, the time for the start arid completion of solidification decreases,

the interface velocity at equivalent fractions solidified increases and the G to R ratio decreases.

Consequently, as the droplet size and the G to R ratio decreases, a change from a cellular to a dendri

tic structure can be expected.

226
Figure 10: HIPed (H2) + austenitized (A2) mate- Figure 11: HIPed (H2) + heat-treated (A2+T3)

rial from 250-840ym size fraction. Etch: material from 250-840ym size fraction. Etch:

Kalling's reagent. SEM. Kalling's reagent. SEM.

Needles of martensite are visible in the coarsest size fraction but are not evident in the finer

powder particles. This can be explained in terms of the high degree of supersaturation of the

austenite associated with increasing cooling rates ( 2 ).

2. HIPed Material

After HIPing, MC, MgC and M23 Ce carbides are present. This is to be expected since the slow

cooling inherent in the HIP cycle leaves the consolidated material in the annealed condition. These

carbides are known to be present after annealing ( 1 ).

No appreciable dependence of volume fraction of MC and MgC with HIPing temperature and prior

powder particle size exists. This implies that the equilibrium level of solubility is achieved. A comparison of Figures 4, 5, 6 and 7 shows that for a particular HIPing temperature, the carbides are approximately the same size. Further, the carbides are coarser following HIPing at the higher temperature.

A close examination of the size distribution for MgC carbides shows that for a given HIPing

temperature (HI or H2), the 250-840ym powder size fraction exhibits a wider spread of carbide sizes

than does the ^44ym powder screen. For similar volume fractions, these observations imply that the

rate of carbide nucleation was higher in the finer size fraction. This is probably related to (i) the

higher surface area per unit volume, (ii) the larger number of particle-particle contact points per

unit volume, (iii) the finer distribution of the carbide network and (iv) the higher degree of supersaturation with decreasing particle size. These four structural differences may affect nucleation

rate separately or in combination. A similar comparison of the MC carbide distribution shows that

the 544ym powder has a narrower size spread and a smaller average carbide size than the 250-840ym

screen.

Only MC carbides are present in the atomized powder. Assuming that the volume fraction of MC

present in the powder decreases with increasing cooling rates, then during HIPing, a higher nucleation

rate should exist in the finer size fraction. This follows from the higher degree of supersaturaticn

of the matrix and the surface area enhancement in the fine powders. In the coarser powder particles,

existing MC carbides can serve as sites for precipitation, thereby enhancing growth.

3. HIP + Heat-Treated Material

The presence of M23 C 6 in the consolidated tool steel after austenitizing at A1 confirms that this

temperature is not high enough to dissolve all this type of carbide which was present after HIPing.

Therefore maximum strengthening of the matrix requires a higher austenitizing temperature, e.g. A2, at which all the M23 CG is dissolved. At A2, about 10% of the MgC present after HIPing is dissolved in

the matrix, cf. Tables II and III. MgC starts to reprecipitate out on tempering, and this accounts

for the increase in volume fraction. This carbide coarsens with increasing tempering temperature,

with the resulting skew in the carbide size distribution curve to larger carbide sizes.

227
CONCLUSIONS

1. Cubic HC and hexagonal y-MC carbides are present in the four size fractions of atomized powder ,) examined. y-MC is shown to be metastable.

2. HIPing temperature is important in dictating the size and size distribution of carbides; it has

a stronger influence than the prior powder particle size. 1 3. The volume fractions of MC carbides (-vieX) and MgC carbides (''-17%) are essentially independent /

of prior particle size and HIPing temperature.

4. A powder processing cycle, namely HI + A2 + T2, has been selected for tool steel powder consolida^'

tion and heat-treatment as a precursor to evaluation of tool steel properties and performance. j

ACKNOWLEDGMENT

This research is supported by the Office of Naval Research. The authors wish to thank Dr. B.A. MacDonald and Dr. D.E. Polk for their interest and involvement in the program.

'j REFERENCES

1. Blickwede, D.J., Cohen, M. and Roberts, G.A., "The Effect of Vanadium and Carbon on the Constitu-J

tion of High Speed Steels", ASM Trans., p. 1161, Vol . 42, 1950. j

2. Fischmeister, H.F., Ozerskii, A.D. and Olsson, L., "Solidification Structure of Gas-Atomized

High-Speed Steel Powders", Powder Metallurgy, p. 1, Vol. 25, No. 1, 1982.

3. Smugeresky, J.E., "Characterization of a Rapidly Solidified Iron-Based Superalloy", Metall. Trans.

A., p. 1535, Vol. 13A, Sept. 1982.

I 4. Levi, C.J. and Mehrabian, R., "Heat Flow in Atomized Metal Droplets", Metall. Trans B., p. 21,

Vol. IIB, March 1980.
STRUCTURE AND PROPERTIES OF A POWDER-PROCESSED HIGH-CHROMIUM IRON

S. Mohan*, A. Lawley*, L. R. Woodyatt** and K. S. Brosius***

*Drexel University, Philadelphia, PA **Bethlehem Steel Corporation, Bethlehem, PA ***Crucible Research Center, Pittsburgh, PA

ABSTRACT

Microstructure-property relationships have been developed for a powder-processed high

chromium iron alloy (Fe-20%Cr-l .5%Mo-2.5%C) with particular interest vested in the possible

effects of rapid solidification of the powder on final properties. The alloy was gas

atomized, screened into three size fractions (119 to 213 ym; 74 to 107 ym; < 44 ym), hot

isostatical ly pressed (HIPed) to full density, and heat-treated (HT) to give a structure

consisting of M

7C3 carbides in tempered martensite. Microstructures were characterized in the powder form, after HIPing and HIP + HT. Properties evaluated were hot hardness, hot

compressive strength, fracture toughness and abrasive and adhesive wear resistance. Comparisons were made with a chill cast alloy of similar composition. Each of the above

properties was relatively insensitive to powder particle size. The strength and temperature dependence of strength of the powder-processed and cast forms were similar but there

were significant differences in fracture toughness and wear resistance between the powderprocessed and cast forms of the alloy. The level of fracture toughness in the powder

metallurgy material was significantly lower than that of the cast material. In terms of

abrasive wear, the powder-processed material was inferior to the cast material at low

stress levels but was comparable at high stress levels. Resistance to adhesive wear was marginally superior in the powder-processed material compared to the cast material. These

property differences have been analyzed and interpreted in terms of the constituents,

scale and morphology of the associated microstructures.

INTRODUCTION

High chromium white cast irons are used extensively in the mining industry due to their high

level of abrasion resistance coupled with adequate toughness [1]. These alloys are also used as the

roll material in rolling mill r'^lls, particularly in continuous hot plate and strip mills. Conventional casting of these highly ct I loyed materials results in both microscopic and macroscopic segregation of alloying elements, with attendant problems in heat-treatment, with property levels and in the

uniformity of properties. Since alloying flexibility, chemical homogeneity and fine-scale microstructures are inherent in powder-processed alloys, this approach could offer a technically and

economically viable alternative.

In this study, detailed relationships between microstructure and properties have been established

for a powder-processed high-chromium iron. Particular interest was vested in the possible role of

powder particle cooling rate, including the rapid solidification regime, on hot strength, fracture

toughness, abrasive and adhesive wear. Microstructure and property comparisons were made on a cast

alloy of similar composition.

EXPERIMENTAL PROCEDURE

Powder of weight composition Fe-20.6%Cr-l .48%Mo-2.44%C-0.5%Mn-0.41%Si-0.007%02-0.12%N2 was nitrogen gas atomized and screened into three size fractions, 119 to 213 pm, 74 to 107 pm, and ^ 44pm.

These size fractions are designated A, B and C, respectively and reflect approximate cooling rates of

10^°C/s, 105°C/s and 5 x 10^°C/s respectively. Powder characteristics, microstructure, and constitution were identified utilizing SEM and x-ray diffraction.

Blended and screened powders from each size fraction were HIPed to full density at 1121°C.

HIPed material was soaked at 1010°C for 3 hours, air cooled, and tempered at 575°C for 3 hours.

229
I ^

i,

Figure 1(a). Atomized powder; 119 to 213 urn Figure 1(b). Atomized powder; ^ 44 urn size

size screen. Arrows point to austenite screen. SEM. dendrites. SEM.

Consolidated material was characterized by means of high resolution SEM and x-ray diffraction. '

Hot hardness, hot compressive strength and static fracture toughness (ASTM E399-78a) were determined for the HIPed + heat-treated material. Abrasive wear resistance was measured by means of < the dry sand-rubber wheel test (ASTM G65-80) at two stress levels [2]. Resistance to adhesive wear

was evaluated by means of the crossed-cyl inder test.

A chill cast alloy (code D) of similar composition to that of the P/M material was heat-treated ,} to the same hardness level as that of the P/M material. The microstructure and properties of this ; material were evaluated as described above.

RESULTS AND OBSERVATIONS

1 ) Powders

jtii Representative micrographs (SEM) of powder from the coarsest and finest size fractions examined

are compared in Figure 1. Particles are generally spherical in shape with relatively smooth sur- faces. At each size fraction, small satellites are attached to the larger particles. Dendritic-type

solidification structures are evident. Dendrites were identified (EDAX) as being austenitic.

M/Cs-type carbides are present in each particle size range.

2) HIPed and Heat-Treated Material

Representative micrographs (SEM) of two of the size fractions are illustrated in Figure 2. The

scale of the microstructure is similar in Figures 2(a) and 2(b) and consists of eutectic carbides

Figure 2(a). HIPed and heat-treated material; Figure 2(b). HIPed and heat-treated material;

119 to 213pm size screen. A = eutectic carbides, < 44 pm size screen. SEM.

B = area showing secondary carbides in tempered

martensite. SEM.

230
Figure 3. Chill-cast material. A = eutectic Figure 4. Hot compressive strength as a carbides, B = secondary carbides. SEM. function of temperature for P/M (A,B,C)

and chill-cast (D) material.

and secondary carbides in a matrix of tempered martensite. For comparison, a representative micrograph (SEM) of the chill cast and heat-treated material is shown in Figure 3; coarse eutectic

carbides and smaller secondary carbides are present in a matrix of tempered martensite.

3)

Hot Strength

Results of the hot compressive tests are shown in Figure 4. The temperature dependence of

strength, and strength level at a given temperature, are similar in each of the three powder size

fractions and the cast materials. Hot hardness shows a similar dependence on temperature.

4)

Fracture Toughness

Fracture toughness levels for the P/M and chill cast material are compared in Figure 5. For the

HIPed + heat-treated P/M material, fracture toughness is largely independent of prior powder particle

size. The fracture toughness of the cast alloy is about twice that of the P/M-processed material.

For tool steels, when toughness is measured by means of the Charpy test, P/M processed material is generally superior to the cast and wrought product [3].

5)

Wear Resistance

Low stress abrasive wear resistance, measured in terms of the wear factor, for the P/M and cast

materials is compared in Figure 6. Linder these conditions, the P/M material is less resistant to wear than the cast form. In high stress tests, the level of abrasive wear resistance is similar in the two forms of material. No effect of prior powder particle size was detected.

Under conditions of adhesive wear, the P/M-processed material exhibits a wear resistance

slightly higher than that of the cast alloy. No significant dependence of adhesive wear resistance

on prior powder particle size was observed.

DISCUSSION

1 ) Microstructure

The matrix of the chill cast alloy consists of highly supersaturated (retained) austenite which

is destabilized during austenitizing. Thus, fine-scale secondary carbides are precipitated. In comparison, after HIPing, the P/M material has a ferritic matrix low in alloying elements so that

secondary carbides do not precipitate when the material is austenitized. This is substantiated by

231

COMPRESSIVE

STRENGTH

(GPA)
Figure 5. Fracture toughness of P/M (A,B,C) and Figure 6. Low stress wear data for P/M (A,B,C)

chill-cast (D) material. and chill-cast (D) material.

the observation that the volume fraction of carbides ('vO.23) is similar in the HIPed P/M material

before and after heat-treatment. The spherical secondary carbides in the P/M material, formed during

HIPing, remain after heat-treatment. In the chill-cast material, the carbide volume fraction in- creases from 0.19 to 0.29 after heat-treatment. The P/M-processed material has a much finer microstructure than the cast form, for each of the prior particle screen sizes.

2) Properties

It is the size and shape of the eutectic carbides that differ most in the P/M and cast materials

This difference is primarily responsible for the observed differences in properties and wear response

between the two forms.

A. Hot Strength

The level of hardness and compressive strength and their dependence on temperature are similar in both the P/M processed and cast material. Since the size and shape of the eutectic

carbides are significantly different in the two materials, strength and hardness must be controlled

primarily by the matrix phase. The loss in compressive strength and hardness at elevated temperature

is associated with matrix softening. Depletion of carbon from the martensitic matrix leads to soften

ing of the matrix phase. Softening also results from growth and/or coalescence of the carbide

particles.

B. Fracture Toughness

Qualitatively, the superior fracture toughness of the cast alloy compared to the P/Mprocessed material is understood in terms of the corresponding fracture surface morphologies.

Figure 7. Extensive fine-scale microvoid coalescence and fine cleavage are evident on the fracture

surfaces of the P/M-processed material. Microvoids develop around the small carbides with little

232
Figure 7(a). Overload fracture surface of HIP + heat-treated P/M material (119-213 pm size

screen). Arrows point to microvoids. SEM.

Figure 7(b). Overload fracture surface of chillcast + heat treated material. A = cleaved

carbides; B = dimpled rupture.

associated shear. Other fine features appear to be the flat facets which occur in clusters adjacent

to the dimples; these are related to the fine-scale structure of the matrix. Crack initiation takes

place at carbide-matrix interfaces and this is followed by microvoid coalescence. In the cast

material, the relative absence of fine carbides eliminates sites for microvoid initiation. The

crack propagates primarily through the matrix between the coarse network of eutectic carbides, with

some associated cleavage of the coarse (eutectic) carbides.

C. Wear

In the low stress abrasion testing of the P/M material, silica particles (size 212-300 pm;

hardness 900 KHN) can readily gouge out the small carbides (size '^ 7. pm; hardness 1700 KHN) from the

martensitic matrix (hardness 700 KHN). Thus these small carbides make only a small contribution to abrasive wear resistance. In comparison, the larger elongated clusters of carbides in the cast form

are more resistant to gouging. Removal requires cutting and/or chipping of the carbides.

In the high stress test, the small eutectic carbides in the P/M processed material are readily

removed from the wear surface by spalling under the influence of the AI 2 O 3 abrasive particles (size

212 to 300 pm; hardness 2000 KHN). Thus, the carbides are only partially effective in reducing

material loss. The AI 2O3 particles are able to cut through both the matrix and the carbide clusters

in the cast alloy with relative ease. Optical metallography of wear surface cross-sections and SEM

of wear surfaces confirm the above differences in the mechanisms of abrasive particle - wear specimen interactions.

In terms of adhesive wear the P/M material is marginally superior to the cast material. This is due primarily to the higher surface area density of the carbides in the P/M material compared to that

of the cast material. Thus, the mean free path of the intervening matrix is smaller in the P/M alloy

than in the cast material. This results in an increased contact area between the tungsten carbide

cylinder and the M^C^ carbides in the P/M material.

ACKNOWLEDGMENT

This work was supported by Bethlehem Steel Corporation. The authors are indebted to Dr. John

M. Chilton for his advice and interest in this program.

REFERENCES

[1] F. Maratray, " Choice of Appropriate Compositions for Chromium-Molybdenum White Irons ", Climax

Molybdenum S.A., Paris, France (1971). ~~

[2] F. Borik, Rubber Wheel Abrasion Test, SAE Preprint 700687. Society of Automotive Engineers,

1970.

[3] A. Kasak and E. J. Dulis, Powder Met., Vol . 21 (2), p. 116, 1978.

233
THE EFFECT OF POWDER COOLING RATE AND THERMAL EXPOSURE

ON THE TENSILE PROPERTIES OF A RAPIDLY SOLIDIFIED

NICKEL BASE SUPERALLOY

D. B. Miracle, K. A, Williams, and H. A. Lipsitt

AFWAL Materials Laboratory

Metals and Ceramics Division

Wright-Patterson AFB OH 45433

ABSTRACT

Although much work has been conducted in the area of rapid solidification, little has

been done to correlate the structure/property relationships of a material with differences

in the cooling rates of the starting powders. This study was designed to investigate the

influence of pov/der cooling rate, within the range of 10^ to 10^°C/sec, on stock extruded

from rapidly solidified superalloy powder. Powder particle sizes ranged from less than 53

microns to approximately 250 microns, representing a change in cooling rate of approximately

one order of magnitude. The effect of long time exposure at high temperatures on mechanical

properties of the consolidated alloy was also determined. The results showed no effects on properties due to differences in cooling rate. Long time exposure at high temperature was found to degrade the properties of equiaxed material tested at room temperature and 1400°F

mainly because of phase changes, Ostwald ripening of the y', and grain boundary cellular

coarsening.

Introduction

In the mid-1970's, a program was begun to develop a new family of turbine blade materials with

the increased temperature capabilities necessary for advanced high-performance engines. The program

was conducted by the Government Products Division of Pratt and Whitney Aircraft, under the sponsorship

of the Defense Advanced Research Projects Agency and the AFWAL Materials Laboratory. The Ni-Al-Mo

system was chosen as a base, and alloys were produced by the Rapid Solidification Rate (RSR) powder

metallurgy process, utilizing cooling rates between 10“* and 10^°C/sec [1]. These high quench rates

have been shown to be beneficial in other alloy systems [2], but little was known about the effects

of rapid cooling in nickel -base superalloys. A project was thus initiated at the AFWAL Materials

Laboratory to assess the role that this range of cooling rates played in the tensile properties of one of the alloys produced--Al loy 185, a y' strengthened Ni-base superalloy which contained 73.8% Ni,

15% A1 , 9% Mo, 2% W and 0.2% C (atomic percent).

Conventional powder-making techniques, with cooling rates in the range of 102-10‘*°C/sec, have

essentially eliminated the problem of macrosegregation associated with conventional casting techniques.

Microsegregation persists, however, on the order of the secondary dendrite arm spacing. As the

cooling rate of the powders increases, the secondary dendrite arm spacing (DAS) becomes smaller, and

the material becomes more homogeneous. In the limiting case, at very high quench rates, homogeneous

nucleation may occur and microsegregation may effectively be eliminated. This has been shown previously

with conventional superalloy powders produced by the RSR technique, where the volume fraction of

homogeneously nucleated, or "micro-crystalline," powders increases as the quench rate increases

(particle size decreases) [3]. It has also been shown that at very high quench rates, an amount of

solute in excess of the equilibrium quantity may be quenched in solution during the solidification

process. This excess quantity may precipitate during subsequent thermal or thermomechanical treatment

with such a size and distribution as to be effective as another precipitation strengthening phase.

These were possible effects of cooling rate on tensile properties to be investigated by this study.

Any candidate material for high temperature application must exhibit a microstructure that is stable at its use temperature for extended periods of time. Some of the microstructural instabilities

upon long-term exposure at 1700°F for 250 hours have been documented for Alloy 185 [4], but the

effect of these instabilities on mechanical properties had not been established. The documentation

of the effect of these instabilities upon the tensile properties was also an objective of this

program.

234
Table I

Mesh Fractions

MESH PARTICLE QUENCH

SIZE SIZE (y) RANGE

- 60/+ 80 250-177 ~ 10‘* °C/sec

- 80/+100 177-149

- 80 < 177

-140/+170 105-88

-140 < 105

> -270 < 53 ~ 105 °C/sec

Procedure

The effect of cooling rate was isolated as follows: powder runs of Alloy 185 were precision-sieved

to several powder mesh size distributions. In general, a different average cooling rate was associated

with each of the sieved mesh sizes. Since the surface-to-vol ume ratio increases as particle size

decreases, the smaller particles experience a higher quench rate. This has been verified experimentally

for other superalloys and is shown in Figure 1 for MAR-M 200 by plotting secondary DAS as a function

of particle size [5]. Six mesh size distributions were used and are given in Table 1, along with the

predicted range of cooling rates associated with each distribution. The powder was placed in stainless

steel cans, hot outgassed, and then extruded at 2000-2300°F with reductions in area of 26:1. The

extrusions were then de-canned and machined into tensile specimens. The specimen geometry has been

used by this Laboratory before and appears in the literature [6]. Half of each of the extrusions was heat-treated for 250 hours at 1700°F prior to machining. The exposures were done in a Lindberg box

furnace in a flowing Ar atmosphere.

Figure 1. Variation of Secondary Dendrite Arm Spacing with Particle Size.

The tensile tests were conducted on an Instron tensile machine using a strain rate of 0.016

cm/cm per min. Strain was measured with a linear variable differential transformer (LVDT). A majority of the tests were done , at room temperature, 1400°F, and 1800°F, although a small number of

DR'ed specimens were tested. at 1600°F. Temperature was controlled to ±10°F in all cases. All tests were run in air.

Microstructural characterization was done with the aid of an ETEC SEM and a Philips 300 TEM. Specimens were polished using standard metallographic techniques and etched with a glyceregia etch.

235
TEM thin foils were produced with a Fischione polishing unit using a methanol-butanol-perchloric acid

solution as the electrolyte and utilizing voltages of approximately 20-25 volts at temperatures from

-20 to -50°C.

Results and Discussion

A representative photomicrograph of the as-extruded material is shown in Figure 2. The material

appears to be fully recrystallized and exhibits a grain size in the range of 16-20 ym, A coarse

incoherent y' surrounds grains of nominally semi-coherent y' precipitates in a y matrix. Other

phases (apparent at this scale of magnification) have been identified by TEM studies and electron-probe

microanalysis as an a-Mo precipitate and carbides. Electron diffraction in the TEM revealed the

presence of two other phases residing between the y'--a phase that has been identified as a Pt 2 Mo isotype (nominally NijMo) and a phase with the DO22 structure (nominally Ni 3 Mo). No microstructural

variations were apparent for different mesh sizes.

Figure 2. Microstructure of As-Extruded Material.

Tensile properties are shown in Figure 3. Approximately ten to twenty tensile tests were made

at each temperature, including specimens from all powder mesh sizes. Differences in powder mesh size

produced no systematic variation in either the stress or strain levels achieved. Thus, it appears

that any cooling rate effects that may have been present in the powder (i.e., supersaturation, microcrystallinity, etc.) have been lost after conventional processing, that the effects were too small to measure, or that the effects manifest themselves under different circumstances (i.e., low cycle

fatigue or oxidation, for instance).

Post-consolidation processing does have an effect on the microstructure and properties. Material

which was directionally recrystallized (DR-ed) and heat treated subsequent to extrusion is slightly

weaker than the as-extruded equiaxed material. The properties are shown in Figure 4. The equiaxed

material has a mild <111> (high elastic modulus) texture with a mean y' size of 380-400 nm. The

grains in the DR'ed material are completely aligned in the high modulus orientation, the mean y' size

is much larger (500-600 nm), and it has no Pt 2 Mo or DO22 type precipitates. In a precipitationstrengthened material, the increase in flow stress is a function of the size and distribution of the

strengthening phase. In the DR'ed and the equiaxed materials, the effects of the variations in the

modulus and the size and distribution of the strengthening precipitates almost offset one another

such that the flow strength of the DR'ed material is only slightly weaker than that of the equiaxed

material

Equiaxed 185 exhibits a mid-range ductility minimum. This has been seen in other superalloys

and has been attributed to a variety of mechanisms [7,8]. In this case, one additional possible

explanation is strain localization and subsequent premature fracture. The strength of the alloy is seen to decrease only slightly from room temperature to 1400°F. However, y' experiences a dramatic

increase in strength in this range with a maximum around 1400°F [9]. Because the strength differential between the y and y' is greatest around 1400°F, the potential for strain localization is also

the greatest. Indeed, strain localization at precipitate-free regions at the grain boundaries of

precipitation-strengthened alloys has been documented as the mechanism for premature fracture in A1
300 -

200 400 600 800 1000 1200 1400 1600 1800 TEMPERATURE (»F)

Figure 3. Tensile Properties of As-Extruded Material.

Figure 4. Tensile Properties of Directionally Recrystallized Material.

alloys and steel 110]. It should be noted that a ductility maximum is seen in the DR'ed material at 1400°F.

The mi crostructural response of the equiaxed alloy to an exposure of 1700 F for 250 hrs is shown

in Figure 5. The most dramatic change is the occurrence of a discontinuous precipitation reaction at the grain boundaries, resulting in a drastic coarsening of both the y and y phases. This coarsened

region is up to 10 ym in width and represents a significant loss in material strengthened by fine y

237
precipitates. The matrix y' is seen to coarsen somewhat from an average y' size of 380-400 nm to 670-700 nm--often by "linking up" with adjacent particles of y'- This is possible because the very

fine Pt 2Mo and DO22 type precipitates previously residing in the regions between the y' have gone

into solution. The solution temperatures for these phases is between 1550-1650°F [11]. Therefore,

at 1700“F, they solution rapidly, thus no longer providing a barrier to the y' coarsening. The

solutioning of the Pt 2 Mo and D022 type precipitates provides an excess of Mo in the matrix which

precipitates as large, randomly distributed particles with various morphologies. These precipitates

have been previously identified as a-Mo [12]. None of the high aspect ratio phases previously seen

and tentatively identified as NiMo were seen in significant quantities. This has been attributed to variations in actual 185 composition [13].

Figure 5. Microstructure After 1700°F/250 Hours.

The effect of the elevated temperature exposure on the tensile properties of Alloy 185 is shown

in Figure 6. A drastic drop in yield, ultimate, and fracture stress levels was noted at room temperature and 1400°F, with little change in the material's behavior at 1800°F. A slight decrease in room temperature ductility was also observed. The strength degradation appears to result from many

factors including, but not limited to, Ostwald ripening of the y' and cellular coarsening at the

grain boundaries. The overriding factor, however, may be the loss of precipitate strengthening by

the very fine Pt 2Mo and DO22 type precipitates.

At 1800°F, the alloy exhibited the same level of strength, irrespective of its previous microstructural condition. It was observed, after testing at 1800°F, that the microstructures were now equivalent. Thus, during heat up, all material experienced the following: the loss of the very fine

ordered precipitates and the subsequent precipitation of the excess Mo as a-Mo; the coarsening of

phases at the grain boundaries due to a discontinuous reaction; and partial solutioning of the y'

particles, resulting in a mean size of approximately 430 nm.

Concl usions

From this investigation, no effect of cooling rate on the tensile properties of Alloy 185 was noted (after consolidation) for the range of powder particle sizes studied. The alloy exhibited a slight decrease in strength from room temperature to 1400°F and a sharp decrease between 1400°F and

1800°F. A mid-range ductility minimum was seen at 1400°F in the equiaxed material. A mid-range

ductility maximum was seen at 1400°F in the directionally recrystallized material. A long-term

exposure of 1700°F for 250 hours resulted in a significant degradation of the equiaxed tensile

properties.

Acknowledgments

The authors are grateful to the Rapid Solidification Processing Group at the Government Products

Division, Pratt & Whitney Aircraft Company, for supplying the material for this project and for

directionally recrystallizing the needed bar stock. Mr. Attwell M. Adair, of our laboratory, is the

Program Manager for the rapid solidification rate (RSR) research and development program.

238
300 -

Figure 6. Tensile Properties After 1700°F/250 Hours.

References

[1] M. R. Glickstein, R. J. Patterson II, and N. E. Schockey, Rapid Solidification Processing , p. 46, edited by R. Mehrabian, B. H. Kear, and M. Cohen, Claitor's, 1978.

[2] C. M. Adam, R. G. Bourdeau, and J. W. Broch, Application of Rapidly Solidified Alloys,

AFWAL-TR-81-4188, February 1982.

[3] P. R. Holiday, A. R. Cox, and R. J. Patterson II, Rapid Solidification Processing , p. 246,

edited by R. Mehrabian, B. H. Kear, and M. Cohen, Claitor's, 1978.

[4] K. A. Williams, D. B. Miracle, and H. A. Lipsitt, Rapid Solidification Processing:

Principles and Technologies, II , p. 129, edited by R. Mehrabian, B. H. Kear, and M. Cohen,

Claitor's, 1980.

[5J P. R. Holiday, A. R. Cox, and R. J. Patterson II, Rapid Solidification Processing , p. 246,

edited by R. Mehrabian, B. H. Kear, and M. Cohen, Claitor's, 1978.

[6] H. A. Lipsitt, D. Shechtman, and R. E. Schafrik, Met. Trans. A, 6A, p. 1991, 1975.

[7] D. A. Woodford, Met. Trans A, 1 2A , p. 299, 1981.

[8J D. A. Woodford and R. H. Bricknell, Superalloys 1980 , p. 633, edited by J. K. Tien, S. T. Wlodek, H. Morrow III, M. Gell, and G. E. Maurer, ASM, 1980.

[9J P. H. Thornton, R. G. Davies, and T. L. Johnston, Met. Trans, J_, p. 207, 1970.

[10] E. Hornbogen and H. Kreye, J. Mater. Sci., ]7_, p. 979, 1982.

[llj P. L. Martin and J. C. Williams, private communication.

[12] P. L. Martin and J. C. Williams, private communication.

[13] Unpublished research by High Temperature Materials Research Group, AFWAL Materials Laboratory.

239
TENSILE BEHAVIOR OF RSR IRON-ALUMINIDES

M. G. Mendiratta, S. K. Ehlers, and D. K. Chatterjee

Systems Research Laboratories, Inc. 2800 Indian Ripple Road

Dayton, OH 45440-3696

H. A. Lipsitt

AFWAL Materials Laboratory

Wright-Patterson Air Force Base, OH 45433

ABSTRACT

Ordered binary alloys in the composition range Fe-25 to -50 at% A1 have been fabricated by blending and hot extruding prealloyed RSR powders. The tensile properties were determined from room temperature to 700°C. Alloys containing up to 40 A1 were found to possess a good combination of room-temperature tensile strength and ductility. The stoichiometric FeAl underwent a brittle-to-ductile transition at ~450°C. In the DO3 alloys,

the strength decreased significantly with temperature, presumably due to the approaching

DO 3

-> B2 transformation in the vicinity of 540°C; however, the single-phase B2 alloys

exhibited good strength retention at high temperature. The slip systems, dislocation con- figurations, and fracture modes have been determined. Deformation mechanisms governing

the tensile flow and fracture behavior of these alloys are presented.

Introduction

Two ordered structures, one having DO 3 crystal lattice (based upon Fe 3 Al composition) and the

other the B2 crystal lattice (based upon FeAl), exist on the iron-rich side of the Fe-Al phase diagram. At room temperature, the DO 3 and B2 orders are maintained in the composition range ~ 23.8 to 34.8 A1 and ~ 35 to 51 A1 , respectively. These ordered alloys are considered to be highly oxidationresistant [1,2]; and, therefore, have potential for high-temperature structural applications. A number of studies have been conducted to determine the slip systems in these materials [3-6]; however, very limited effort has been expended on the tensile behavior [7-9]. Although the operative

slip of { 110 } < 111 > provides more than a sufficient number of independent systems for polycrystalline

ductility, these materials are generally thought to be extremely brittle and difficult to fabricate.

Previous studies on these ordered alloys have been carried out on materials prepared by conventional

ingot-metallurgy techniques. In the present study, the tensile behavior of a number of powderproduced (P/M) alloys, covering a wide range of compositions from Fe-25 to -50 A1 , has been investigated. The tensile flow and fracture properties of these P/M alloys were determined from room temperature to 700°C. In addition, extensive SEM fractography and thin-foil TEM were carried out in an attempt to understand the deformation mechanisms.

Experimental Procedures

Two lots consisting of spherical, prealloyed -140 mesh RSR powders (Pratt and Whitney) having

compositions of Fe-25 and -50 A1 were utilized to prepare alloys having nominal compositions of Fe-25,

-35, -40, and -50 A1 . The compositions Fe-35 and -40 A1 were prepared by mechanical blending of

appropriate quantities of the two powder lots. The Fe-25 A1 powders were consolidated by hot extrusion at 900°C with an extrusion ratio of 16:1, while the Fe-35, -40 and -50 A1 compositions were con- solidated by hot-extrusion at 1000°C with the same extrusion ratio. The chemical analysis of the

powder lots showed them to contain 0.03, 0.03, and 0.006 wt% carbon, oxygen, and nitrogen, respectively, and very small amounts of metallic impurities. The machined tensile specimens of all four

alloys were subjected to a homogenization anneal of 900°/24 hr. The Fe-25 A1 alloy was further heat

treated at 500°C for 100 hr to produce a high degree of DO 3 order. The tensile tests were carried out

in an Instron machine at a crosshead speed of 0.025 cm/min in air from room temperature to 700°C. For TEM, thin discs (- 0.25-mm thick) were cut by precision low-speed diamond saws, mechanically ground

to a thickness of ~ 0.125 mm, and electrolytically thinned utilizing the double-jet technique. The

240
electrolyte consisted of one part HNO 3 and two parts methyl alcohol by volume. The polishing was

carried out at -40OC and 10 V. Fracture modes were examined by scanning electron fractography.

Results

I n i ti al Microstructure

Metal lographic examination of as-extruded-plus-heat-treated specimens of all the alloys revealed

an equiaxed recrystallized grain structure with an average grain size of ~50 ym for Fe-25 Al and ~30 ym

for the remainder of the alloys. The prior powder-particle boundaries decorated by submicron-size impurity particles (presumably oxides, nitrides, or carbides) were also observed metal lographical ly;

however, in general, the recrystallized grain boundaries did not coincide with the prior powder boundaries. Electron-probe microanalysis carried out on several grains indicated that the blended alloys

were homogeneous with regard to Al and Fe content. As observed by TEM, the 500°C/100 hr heat treatment produced in the Fe-25 Al alloy thermal D03 -order domains (size ~ 0.3 ym) having highly irregular

boundaries. The size of the domains and the sharpness of the superlattice reflections in the

electron-diffraction patterns indicated that a high degree of DO 3 order had been achieved. In contrast, electron-diffraction patterns in the Fe-35 Al alloy showed sharp and intense superlattice

reflections corresponding to the B2 order and extremely weak DO 3 reflections, indicating that this

composition is in the B2 phase field very near the D03 -to-B 2 boundary. TEM examination of Fe-40 and

-50 Al revealed the presence of the single-phase B2 order. In the three B2 alloys, no thermal antiphase domain boundaries (APB) were observed, presumably due to the long-time homogenization at high

temperature (900OC/24 hr).

Tensile Properties

Fe-25 Al . Figure 1 shows the stress-strain (a - e) curves for different test temperatures. At

room temperature, this alloy exhibited a high yield strength, oo. 2 > of ~ 103 ksi and an elongation to fracture, e^, of 8%. It is apparent from the a - e diagrams that work-hardening rate increases with

temperature up to 300°C, above which it decreases appreciably. The maximum ductility was exhibited at

' 54QOC, above which it decreased slightly.

% EGR. STRAIN

Fig. 1. Engineering stress-strain curves Fig. 2. Engineering stress-strain curves for different

for different test temperatures test temperatures for Fe-35 Al

for Fe-25 Al

Fe-35 Al and -40 Al . For Fe-35 Al the a - e diagrams are given in Fig. 2. The a - e diagrams

for Fe-40 Al exhibited a similar trend to that for Fe-35 Al and, hence, are not included here. At

room temperature, ao ,2 values of 70 and 92 ksi and sf values of 6 and 2.5% were obtained for Fe-35

and -40 Al , respectively. Both alloys exhibited very high work-hardening rate at room temperature.

Figure 3 shows the variation of ao. 2 > £f» percent reduction in area, R.A., with test temperature

for Fe-35 Al . Up to 6 OOOC, ao .2 decreases slowly with temperature; however, above this temperature it

drops rapidly. A ductility maximum is observed at 500°C, above which both R. A. and ef decrease. A similar trend in the temperature dependence of ao. 2 > ^nd R.A. was also observed for Fe-40 Al

241

EGR.

STRESS

(ksi)
Fe-50 A1 . Figure 4 shows the a - 0 curves for the stoichiometric FeAl . No ductility was observed up to 450°Ci at this temperature the alloy exhibited some degree of plasticity. As in the case

of Fe-35 and -40 A1 , a maximum in ductility was exhibited at SOQOC, above which the ductility decreased with temperature.

Figure 5 shows the variation of ao ,2 with temperature for the four alloys investigated. For the

Fe-25 A1 alloy, a rapid drop in strength'occurred above 200°C. However, in comparison, the Fe-35 and

-40 A1 alloys exhibited a significantly lower decrease in strength up to 6 OOOC; at this temperture

the yield strength was ~60 ksi, and above this temperature the strength decreased rapidly.

(MPo) (ksl)

Fig. 3. Variation of yield strength. Fig. 4. Engineering stress-strain curves for different test

percent elongation to fracture, temperatures for Fe-50 A1

and percent reduction in area with temperature for Fe-35 A1

Fracture Modes

Fe-25, -35, and -40 A1 alloys exhibiting ductility at room temperature fractured by the transgranular cleavage mode. The Fe-50 A1 alloy with no ductility exhibited intergranular fracture. With

increasing test temperature, increasing proportions of tearing and dimples were observed; and at 500°C

where a maximum in ductility occurred, the fracture surfaces consisted of a mixture of large and small

transgranular dimples. With increasing test temperatures, a decrease in ductility was observed, and

the fracture mode consisted of an increasing proportion of grain-boundary facets containing small

cavities. Examples of this transition in fracture mode are given in Figs. 6 (a) and 6 (b) which show

the room- temperature cleavage and high-temperature intergranular-cavity fracture for Fe-35 A1

Fig. 5. Variation of yield strength

with temperature for Fe-25,

-35, -40 and -50 A1 alloys.

Deformation Structures

Fig. 6 . Fracture-mode transition in Fe-35 A1 a) transgranular cleavage at room temperature, b) intergranular cavities at 600°C.

Previous studies on single-crystal and polycrystal DO 3 and B2 Fe-Al alloys [3-6,10] have shown

that at room temperature, the slip direction is <111>. Slip occurs in the <111> direction either by

242

EGR.

STRESS

(ksi)
the movement of perfect superlattice dislocations or by their dissociation to produce APB trails. In the DO 3 lattice, the superlattice dislocations consist of four ordinary a'/4 <111> dislocations separated by nearest-neighbor (NN) and next-nearest-neighbor (NNN) APBs. a' is the lattice parameter of

the DO 3 unit cell and is twice the length of that for the disordered BCC lattice. In the B2 lattice,

the superlattice dislocation consists of two ordinary a/2 <111> dislocations separated by NNAPB. The

ease or difficulty of glide of these superlattice dislocations and their imperfect variants governs

the flow and fracture behavior in these alloys. Specific TEM contrast and diffraction conditions

required to characterize these features have been reviewed by Marcinkowski [11]; and, in the present

study, the same procedures have been followed.

Fe-25 A1 . The room-temperature deformation in the Fe-25 A1 alloy was accomplished by the movement of ordinary dislocations and the formation of NN and NNNAPBs; no coupled dislocations were

observed in any grains. Figure 7(a) is a bright-field electron micrograph showing the a'/4 <11 1> ordinary screw dislocations; the general motion of these dislocations through the grain is indicated

by the slip traces left behind. As can be seen, the slip traces are highly curved with a significant

angular spread, indicating that the glide dislocations have cross-slipped extensively during deformation. The movement of the ordinary dislocations produces APB bands is shown in Fig. 7(b), which is a dark-field micrograph taken with the superlattice reflection (200). At 200 and 300°C, where higher

work hardening was observed, TEM examination revealed the presence of coupled superlattice dislocations in a number of grains.

Fe-35 and -40 A1 . For these alloys, room- temperature deformation occurred by movement of superlattice dislocations of the a/2 <111> Burgers vector, as can be seen in the electron micrographs of

Figs. 8 (a) and 8 (b). Numerous paired dislocations can be observed; that these were superlattice

(b) (b)

Fig. 7. TEM micrographs showing room- temperature

deformation substructures in Fe-25 A1 a) bright-field micrograph showing crossslip of ordinary dislocations, bj darkfield micrograph showing slip-induced NNAPBs.

Fig. 8 . TEM micrographs showing superlattice

dislocations in Fe-35 A1 deformed at room temperature a) bright field, b)

weak-beam dark field (g/3g).

243
dislocations was deduced from the constancy of the separation distance of the component dislocations

under the g+ and g' diffraction conditions. Figure 8(b) is a weak-beam dark-field micrograph from

which the separation distance of the dislocations in the screw orientation for the Fe-35 A1 alloy was ; measured to be ~ 50 8, which is in agreement with the value obtained by Ray, ^ [12]. Dark-field ‘ examination utilizing superlattice reflections {100} did not reveal unambiguously the formation of

slip-produced APBs. At all temperatures up to 700°C, the Burgers vector was found to be <111>.

Fe-50 A1 . At room temperature, _sl ip occurred predominantly in uniformly distributed planar * bands. Fig. 9(a). The standard g

• B = 0 analysis revealed that most of the dislocations had a Burgers vector <111> and were in screw orientation. It was not possible to resolve the components of

the superlattice dislocations by weak-beam dark-field observation, most probably due to the high APB energy of this alloy. Although the predominant slip direction was <111>, a few dislocations of the

<100> type were also observed. In contrast, at 500°C the predominant slip direction was found to be

<100>, with a few dislocations being of the <111> type. Figure 9(b) shows the 500°C dislocation

structure, with <100> dislocations mainly in the edge orientation. Numerous loops and dipoles, * characteristic of thermally assisted deformation, were also observed.

(a) (b)

Fig. 9. Bright-field electron micrographs showing dislocation 1

structures in Fe-50 A1 a) specimen deformed at room «

temperature, b) specimen deformed at 500°C. f

Discussion

|

Room-Temperature Flow and Ductility

The room- temperature ductility of the iron-aluminides from Fe-25 to -40 A1 is due to the ability

j of the <111> dislocations to move freely. For the Fe-25 A1 alloy, the values of NNAPB and NNNAPB i energies (77 and 85 ergs/cm^, respectively) are relatively low [12] and the nucleation of superlattice ' dislocations is difficult [13-15]; therefore, in this alloy the ordinary dislocations which create

APB faults move and cross-slip freely. Such dislocation character is responsible for the observed i parabolic low work hardening and a reasonably high ductility. Also, since the applied stress must I overcome APB energies, the observed yield strength is quite high. i

With increasing A1 content in the B2 alloys, the NNAPB energy increases [12], and the nucleation

j of superlattice dislocations is energetically favored [14-15]. Thus, in Fe-35 and -40 A1 alloys | numerous paired dislocations and no slip-produced APBs were observed. The motion of superlattice dis-

' locations produces very high work hardening presumably due to the Taylor-type locking of superlattice ! dislocations on neighboring parallel slip planes [16] and to the production of APB tubes at jogs ! [3,17]. Also, the equilibrium separation distance between the superlattice dislocations is suffi- ; ciently large (~ 50 8

)

to prevent easy cross-slip. These factors combine to decrease the ductility ( of these alloys with increasing A1 content, although an increase in yield strength was observed.

Extrapolation of the ao .2 curve for the Fe-50 A1 (Fig. 5) to room- temperature indicates that the | alloy would have yielded at room temperature at a stress of ~ 180 ksi. This value should be considered in the context of the reported value of critical resolved shear stress (CRSS) of ~ 68 ksi for

dominant <111> {110} slip [4]. This CRSS value is extremely high. Thus, the observed lack of ductility'

in the Fe-50 A1 alloy at room temperature may be due to the yield strength exceeding the fracture | strength, thereby causing premature fracture. , High-Temperature Properties

The yield strength drops rapidly in the Fe-25 A1 alloy with temperature; while in the Fe-35 and

244
-40 A1 alloys, the strength decrease is significantly more gradual (Fig. 5). The Fe-25 A1 alloy

ij undergoes a second-order DO 3

-> B2 phase transformation at a low temperature of 540°C; the approach to this transformation is thought to be responsible for the rapid drop in high- temperature strength. In contrast, the Fe-35 and -40 A1 alloys are thought to be fully ordered up to their melting temperaj tures, thus exhibiting a less severe strength drop with temperature.

{ All iron-aluminides investigated exhibited a loss of ductility with increasing temperature above

; 600°C (e.g.. Fig. 3). This loss of ductility was accompanied by a transition in fracture mode from

transgranular-dimple formation to intergranular-cavity formation. Further experiments involving

varying strain-rate tensile tests [18] have shown that this loss of ductility is both stress and

j diffusion enhanced. Generally the intergranular cavities form in engineering alloys which are sub-

; jected to creep deformation; during the creep process the impurities segregate to the grain boundaries

and aid in cavity nucleation. In the present case the strain rate was significantly higher than that

I generally occurring during creep deformation; however, diffusion of some impurity species may be

j sufficiently rapid. The observed loss of ductility in the iron-aluminides indicates the presence of

I an instability which may have important engineering implications.

Conclusions

1. The RSR Fe-25 to -40 A1 alloys are ductile at room temperature. Maximum ductility and yield

strength are exhibited by Fe-25 A1 . Ductility decreases with increasing A1 content.

2. In Fe-25 A1 , the room- temperature plastic flow is governed by the movement and cross-slip of

j ordinary dislocations creating APB bands. In contrast, with increasing A1 content, the plastic flow

I occurs by the movement of superlattice dislocations.

j 3. The Fe-50 A1 alloy is totally brittle at room temperature and exhibits a brittle-to-ductile

transition at ~ 450°C.

4. The yield strength of the Fe-25 A1 alloy drops rapidly with temperature, presumably due to the DO 3 B2 transformation at a low temperature of 540OC. In contrast, the Fe-35 and -40 A1 alloys,

showing no transformation, are reasonably strong up to 600°C ( 09.2 = 60 ksi).

5. All the alloys studied exhibited a high- temperature ductility loss and associated intergranular cavity fracture mode.

References

[1] C. Sykes and J. W. Bampfylde, J. Iron Steel Inst. 130 (II), 389 (1934).

[2] E. R. Morgan and V. F. Zackay, Met. Prog. 129 (1955).

[3] R. C. Crawford, Phil. Mag. 3^, 529 (1976).

[4] T. Yamagate and H. Yoshida, Mater. Sci. Eng. 1_2, 95 (1973).

[5] Y. Umakoshi and M. Yamaguchi, Phil. Mag. A 537 (1980).

[ 6 ] Y. Umakoshi and M. Yamaguchi, Phil. Mag. A 711 (1981).

[7] G. P. Rauscher, Jr., F. C. Perkins, and J. F. Nachman, "Study of Factors Influencing Ductility

of Iron-Aluminum Alloys," Technical Report under Contract AT(11-1)742 (U.S. Atomic Energy

Commission, University of Denver, Denver, CO, 1961).

[ 8 ]

[9]

[ 10 ]

[ 11 ]

[ 12 ]

[13]

[14]

[15]

[16]

[17]

[18]

M. J. Marcinkowski , M. E. Taylor and F. X. Kayser, J. Mater. Sci. 1_0, 406 (1975).

G. Sainfort, P. Mouturat, P. Pepin, J. Petit, G. Cabane, and M. Salesse, Memoires Scientifiques

Rev. Metallurg., LX 125 (1963).

M. J. Marcinkowski and N. Brown, Acta Met. % 764 (1961).

M. J. Marcinkowski "Theory and Direct Observation of Antiphase Boundaries and Dislocations in

Superlattices," in Electron Microscopy and Strength of Crystals (G. Thomas and J. Washburn, Eds.)

(Interscience Publishers, N.Y., 1963), p. 333.

I. L. Ray, R. C. Crawford, and D. J. H. Cockayne, Phil. Mag. 1027 (1970).

H. J. Leamy and F. X. Kayser, Phys. Stat. Sol. 34, 765 (1969).

H. J. Leamy, F. X. Kayser, and M. J. Marcinkowski, Phil. Mag. 763 (1969).

H. J. Leamy, F. X. Kayser, and M. J. Marcinkowski, Phil Mag 779 (1969).

M. J. Marcinkowski and G. Lakso, J. Appl . Phys. 2124 (1967).

C. T. Chou, and P. B. Hirsch, Iss. Electron Microsc. Anal., 459 (1981).

M. G. Mendiratta, Systems Research Laboratories, Inc., Dayton, OH, work in progress under Air

Force Contract F33615-81-C-5059.
MICROSTRUCTURE AND MECHANICAL PROPERTIES

OF RAPIDLY-SOLIDIFIED WHITE-CAST IRON POWDERS

k

Lawrence E. Eiselstein , Oscar A. Ruano

Oleg D. Sherby***, and Jeffrey Wadsworth

Metallurgy Group, Poulter Laboratories, S.R.I.

International, Menlo Park, California 94025

Centro Nacional de Investigaciones-Metalurgica

,

Av. de Gregoria del Amo, S/N, Madrid-3, Spain

Department of Materials Science and Engineering

Stanford University, Stanford, California 94305

Metallurgy Department, Lockheed Research Laboratory

Palo Alto, California 94304

ABSTRACT

Rapidly solidified white cast iron powders containing 2.4 and 3.0%C, as well as a 3.0%C

alloy containing 1.5%Cr, were investigated. The as-quenched microstructures of these white

cast iron powders consist of retained austenite, cementite, and a small fraction of martensite. Upon heating to 600°C the microstructure changes to a mixture of fine cementite

(~1 ym diam.) and ferrite of grain size 1 to 3 pm.

It is shown that densi fi cati on of comoacts of the white cast iron powders can be en- hanced by multiple phase transformations through thermal cycling. This transformation warm

pressing at intermediate temperatures leads to dense compacts at low pressures and short

times. Because the transformation temperature is low, ultrafine structures are found in the

densified compacts.

The influence of these fine microstructures on the mechanical properties of the white

cast irons was determined at intermediate and room temperatures by tension and compression

tests. The white cast irons were found to be superplastic in the temperature range of

630—725°C. Maximum tension elongations of 1410% were found for the 3.0%C + 1.5%Cr material,

940% for the 3.0%C material, and 480% for the 2.4%C material (in all cases tested at 700°C at

a strain rate of 1% per minute). Negligible cavitation was noted during superplastic deformation of these white cast irons. Under superplastic conditions, the white cast irons exhibited

a high strain rate sensitivity exponent m = 0.5 and activation energies for plastic flow nearly

equal to the activation energy for grain boundary diffusion. These values are in agreement

with models for superplasticity based on grain boundary sliding accommodated by slip mechanisms

adjacent to the grain boundary. Room temperature compression tests of the white cast irons

revealed large strain-to-fai 1 ure values in the order of e

== 0.25. In tension tests, a ductility of over 7% was obtained in rapidly solidified 3.0%C white cast iron at room temperature

in contrast to nil ductility for the as-cast material.

INTRODUCTION

Powders with a very fine microstructure can be obtained by the rapid solidification rate (RSR)

process which makes use of advanced jet technology for their production [1]. Such fine microstructures are very attractive because of the potentially beneficial properties they can show in industrial applications. For such applications, however, the densified compacts must retain the fine

microstructure of the individual powders. For this reason, the compacting methods should avoid

coarsening of the microstructure.

In the present investigation, the microstructures and the annealing behavior of ferrous powders

246
DAS

(u

Figure 2 . DAS vs. particle size. CA

particles have smaller dendrites than AA

particles, but both are in same order of

magnitude.

Figure 3. Relative cooling rate vs. particle size for reported extremes

pf the coefficient n used to calculate

T.

Figure 4. Approximate cooling rates vs. particle

size. AA and CA R-95 are contrasted with AA and

RSR C-101.(5)

263
RAPID SOLIDIFICATION STUDIES OF NisAl ALLOYS*

C. C. Koch, J. A. Horton, C. T. Liu,

0. B. Cavin, and J. 0. Scarbrough

Metals and Ceramics Division

Oak Ridge National Laboratory

Oak Ridge, Tennessee 37830

ABSTRACT

Rapidly solidified foils of Ni3Al (24 at. % A1 ) with and without additions of boron

(500 and 1000 ppm by weight) were prepared in an arc-hammer apparatus. The bend ductility,

microhardness, and microstructure of the foils were studied as a function of solidification

rate, which was defined by foil thickness. Boron-free Ni3Al samples exhibited brittle

behavior over the entire range of solidification rates used (foil thicknesses 20 to 100 ym).

However, boron-containing samples, which were ductile for foil thicknesses greater than

45 ym showed a transition-to-brittl e behavior in the range of thickness of 30 to 45 ym. The

high hardness of the rapidly quenched foils (450 DPN) remained approximately constant in this transition range. Electron microscopy revealed a complex anti-phase boundary (APB)

structure that could not be correlated to the mechanical behavior. The suppression of the

beneficial boron segregation to grain boundaries is believed to be responsible for the

ducti 1 e-to-brittle transition on increasing solidification rate.

Introduction

Intermetal 1 ic compounds, such as the iron or nickel aluminides (e.g., Ni3Al, NiAl, Fe3Al , FeAl )

,

have unique properties that make them potential structural materials for elevated temperature applications. The ordered crystal structures in these materials reduce the atomic mobility at elevated

temperatures resulting in good strength, stability, and creep resistance. Unlike conventional

alloys, the static strength of many aluminides shows an increase rather than a decrease with

increasing test temperatures [1,2]. The presence of aluminum provides an oxidation resistant AI2O3

film. However, the major problem with using aluminides as structural materials is their low ductility

and brittle fracture [1,3], particularly at low temperatures. Because of poor fabricabil ity and lowfracture toughness, the aluminides have not been considered as serious candidates for structural

appl ications.

The nickel aluminide, Ni3Al {y'), is an important strengthening constituent of nickel-base

superalloys. While single crystals of Ni3Al (which has the LI2 structure) exhibit extensive ductility [4], polycrystalline forms are extremely brittle, suggesting grain boundary embrittlement. The

intergranular fracture observed supports the idea that weak grain boundaries — either due to their

structure or impurity segregation — cause brittle fracture with little plastic deformation in the

grains. While there has been research aimed at improving the ductility of intermetal 1 ic compounds

like Ni3Al for over 20 years, only recently has there been significant success in this area. Two approaches to this problem have provided solutions: (1) rapid solidification processing, and

(2) microalloying. Rapid solidification processing of Fe3Al (30 at. % A1 ) powders has provided finegrained homogeneous alloys which exhibit useful ductility (5-8% elongation at room temperature) [5].

Recently, Inoue et al . [6] have reported improvements in ductility and strength of LI2 compounds in Ni-Al-X (X = Cr, Mn, Fe, Co, or Si) alloys by rapid solidification using the melt-spinning technique.

Elongations of 4 to 27% were achieved. It was suggested that this improvement in ductility was due to rapid solidification producing a lower degree of long-range- order, fine-grain-structure, and

suppressing-grain -boundary segregation. Microalloying — the addition of small quantities of an element (<1%) for controlling the chemistry and cohension of grain boundaries — has been found to dramatically ductilize Ni3Al. Aoki and Izumi [7] first discovered that boron additions of 500 and

1000 ppm by weight imparted room temperature ductilities of 35%. Liu and Koch [8] later found that

boron concentrations of 250 to 1000 ppm along with control of the thermal -mechanical treatment pro- duced ductilities of greater than 50% elongation.

*Research sponsored by the Division of Materials Sciences, U.S. Department of Energy under

contract W-7405-eng-26 with the Union Carbide Corporation.

264
As part of the goal of our present research program to understand the effects of boron on the

mechanical behavior of Ni3Al, we have prepared a series of Ni3Al samples, with and without boron additions, by rapid solidification in an arc-hammer apparatus. We report here preliminary results of the

ductility, hardness, and microstructure of these Ni3Al alloys as a function of quench rate. Since

boron diffusion to the grain boundaries must be necessary to provide its beneficial effects, this

experiment was designed to see whether suppressing the segregation of boron would prevent its ductilizing effect and thus provide information on the mechanism responsible.

Experimental

Samples from the previous study [8] of Ni—24 at. % A1 (referred to as Ni3Al in this paper) con- taining 0, 500, or 1000 ppm of boron were cut into small pieces for rapid solidification in the arc- hammer apparatus. The solidification rate in the arc hammer is a function of hammer velocity, mass of

the liquid alloy, as well as melt superheat. A model of the arc hammer [9] has provided estimates of

solidification rates of 10^ to 10^ degrees per second for metallic glass foils typically 20 to 40 ym

thick. In this paper we shall use foil thickness as a measure of solidification/quench rate. The

foils produced in the arc hammer for different hammer speeds and alloy masses were typically 1.5 to 3.0 cm in diameter by 20 to 100 ym thick. The structure of the foils was determined by x-ray diffraction on a Philips automated diffractometer using copper Ka x-radiation. The integrated intensity of

pairs of fundamental and superlattice lines (200:100 and 220:110) were measured in an attempt to extract the long-range-order parameter, S.

The bend ductilities of rectangular strips cut from the foils were obtained by bending around

progressively smaller cylinders until cracking and then fracture occurred. The fracture strain, e^, was calculated by

= t/(2r + t) ,

where r = radius of the cylinder and t = thickness of the foil. The foil thicknesses were measured

between the conical points of a dial gauge equipped with a slip clutch and accurate to about 1 to 2 ym.

If a foil could be bent 180° on itself so that r = 0, e"*" = 1, the material is considered to be quite

ductile.

Microhardness measurements were made on the foils in a Wol pert-Griess hardness tester with a Vickers diamond indentor and a load of 100 g.

Approximately 3-mm-diam disks were cut from the foils and thinned el ectrolytical ly for examination by transmission electron microscopy. The thinned foils were examined in an analytical electron

microscope. Philips EM 400T equipped with a field-emission gun.

Experimental Results

The crystal structure of all the rapidly solidified foils was found to be the ordered Ll2-type by

both x-ray and electron diffraction. Well defined superlattice reflections were observed. X-ray

diffraction and electron diffraction revealed a high degree of preferred orientation in the quenched

foils with a (100) type texture. This preferred orientation prevented a straightforward determination

of the long-range-order parameter, S, by the conventional technique [10]. However, Starke et al . [11]

have developed a procedure for extracting S in heavily textured CU3AU (LI2 structure). We are presently employing this procedure in our work and the results will be reported elsewhere. In any case,

it was observed that even the fastest quench rates used in this study could not suppress the orderedcrystal structure which contains a substantial degree of long-range order.

The results of the bend tests are presented in Fig. 1 as a plot of fracture strain versus sample

thickness; sample thickness being inversely related to quenching rate. Data for Ni3Al samples which

do not contain boron (represented by open circles) all exhibit brittle behavior with values of fracture strain less than 0.01. Conversely, Ni3Al samples containing 500 and 1000 ppm of boron show

completely ductile behavior for sample thicknesses greater than approximately 50 ym. The foils can be

bent 180° on themselves giving values for the defined fracture strain of one. However, for sample

thicknesses between about 30 to 45 ym, that is for faster quenching rates, a marked change in ductility is observed. Over this range of sample thickness (quenching rate) the fracture strain drops

dramatically from = 1 to e'^ < 0.01. Within this transition range, ductile and brittle behavior

was often observed within a given specimen. Cracks were seen to form from the edge of the bent

specimen at very low strains (<0.01) but then stop. The specimen could then continue to be bent back

upon itself. Parts of several samples which exhibited brittle fracture behavior could be bent 180° on themselves. The brittle fracture behavior appeared to be intergranular in nature. A scanning

electron micrograph of the fractured surface of a 26-ym-thick specimen containing 1000 ppm boron.

265
fO

0.5

0.2

0.1

0.05

0.02

0.01

0.005

0.002

0.001 20 30 40 50 60 70 80 90 100 SAMPLE THICKNESS (/im) —

•^QUENCH RATE

=—1—

V—r

• \

*‘ '*1—I—I—

- 0 0% BORON - ~ ^ 500 ppm BORON ~

1000 ppm BORON _

Fig. 1. Bend fracture strain, versus sample

thickness. Boron-doped specimens showed

a ductile-to-brittle transition around a thickness of 45 ym. All the specimens

without boron additions exhibited brittle

behavior.

presented in Fig. 2, indicates the grain boundary separation. Auger electron spectroscopy studies are in progress to define more precisely the nature of the grain boundary fracture.

Microhardness is plotted against specimen thickness (inverse quench rate) in Fig. 3. The bars on the data points represent plus or minus twice the standard deviation of at least ten measurements.

The hardness of rapidly quenched specimens (thickness <60 ym) is approximately 450 DPN. This magnitude has been obtained for heavily cold-rolled single crystals of Ni 3

Al [12] and cold-rolled

polycrystalline Ni 3

Al containing boron additions [13]. The hardness decreases with specimen

thickness, that is, with slower cooling rates. A 45-ym-thick foil was subsequently annealed at 1000°C

for 15 min. The hardness dropped from 440 DPN to a value of about 240 DPN, comparable to the values

for annealed, bulk Ni 3Al. It should be noted that the transition range of specimen thickness from

Fig. 2. Scanning electron micrograph of a brittle

fracture surface of Ni 3Al containing

1000 ppm boron.

500 -

400 -

ITRANSITIONI

brittleI range 300 -

ductile

}}

o 0% BORON

A 500 ppm BORON

1000 ppm BORON

200

annealed

1000 °C FOR 15 min _J 1 20 30 40 50 60 70 80 SAMPLE THICKNESS (ym) 1 QUENCH RATE

90 100 110

266

Fig. 3. Hardness versus sample thickness.
ductile-to-brittle behavior lies within the region where the hardness is essentially constant.

Likewise, no obvious differences in hardness can be observed as a function of boron concentration.

Thus, it would appear that the hardness of the rapidly quenched Ni 3Al samples is mainly affected by

the bulk structure, over which boron concentration has little influence. However, the bend ductility

is expected to be controlled by grain boundary properties, which are dramatically altered by the

presence of boron.

The microstructures of the rapidly solidified Ni 3

Al samples as revealed by transmission electron

microscopy have the following features. The average grain diameters of the specimens varied from

about 1.0 to 3.5 pm as measured by the linear intercept method. The more rapidly quenched specimens

have the smaller grain size. For example, for Ni 3

Al samples containing 1000 ppm B, a specimen

50 pm thick had an average grain size of 0.70 pm compared to a specimen 90 pm thick with an average

grain size of 3.4 pm. There was no obvious dependence of grain size on boron concentration. The

above grain sizes are for the center of the specimen cross sections. The grain size on the specimen

surface of the 50 pm thick, 1000-ppm B specimen ranged between 0.25 and 0.50 pm which was about

one-fourth to one-half the size of the grains in the center. This morphology of larger, textured

grains in the bulk, and smaller grains on the surface of alloys solidified in the arc hammer has been

observed previously [14] for (Fe,Co) 3 V alloys with the LI 2 structure. The main features within the

grains are dislocations and APBs. The dislocation density is in general very low, some grains being

completely free of dislocations. The APBs are readily observed in rapidly quenched Ni 3Al, as already

noted by Inoue et al . [6] but in contrast to conventionally prepared material [15] in which APBs are rarely observed. In Fig. 4 a bright-field micrograph typical of the APB structure observed in the

smaller grains is given. Curved APBs surrounding small domains are seen in the center of the grain,

while straighten APBs, approximately normal to the grain boundary, separate larger domains. In the

larger grains the APB structure illustrated in Fig. 5 is often observed. In this case there are several regions within the grain of the small curved APBs surrounded by larger domains, which are now

not confined to the edge of the grains. The domain size was measured by the linear- intercept method

with the number of intercepts increased by one third to account for the out-of-contrast APBs for

micrographs made with <200> and <220> diffraction vectors. The domain size (small domains) for the

boron-containing specimens is about 50 nm in diameter, essentially independent of solidification rate,

while for the boron-free samples domain diameters of about 80 nm are observed. The domain size at the

surface where the grains were smaller was approximately the same as in the center of the specimen.

The tentative conclusion is that boron has reduced the diameter of the small domains by 30 to 40%.

Annealing at 1000°C for 15 min increased the average grain diameter to 6.5 ym and removed all the

Fig. 4. Bright-field electron micrograph of a Ni 3

Al sample containing 1000-ppm B; the APB structure in small grain. Foil

normal and electron-beam direction near [114].

Fig. 5. Dark-field electron micrograph on the

[100] superlattice reflection of a Ni 3Al sample containing 1000-ppm B; the APB structure in large grain. Foil normal

and electron-beam direction near [001].

267
Di scussion

While more research is required to clarify several of the experimental observations made in this ' preliminary study, some tentative explanations of the results can be given.

A simple rationalization of the solidification rate dependence of the ductile-brittle transition

in boron-containing samples is that at sufficiently fast cooling rates the beneficial segregation of

boron is suppressed leaving the grain boundaries brittle. If this explanation should hold, it implies

that (1) harmful impurity elements (such as sulfur) can segregate to the grain boundaries faster than can boron, or (2) impurity segregation is not responsible for the weakness of grain boundaries in

Ni 3Al but the weakness is an inherent property of the grain boundary structure. Clearly, experimental

verification of this suggestion by, for example. Auger spectroscopy is needed.

Rapid solidification by the arc-hammer process has not increased the ductility of boron-free

Ni^Al in contrast to the melt-spinning study of Inoue et al . [6] on Ni-Al-X (X = Cr, Mn, Fe, Co, or

Si) LI 2 materials. Their explanation for the ductilizing effects of rapid solidification is based on prevention of segregation to the grain boundaries and small domain size effects. They reported their

grain sizes to range from 2 to 5 urn. Their domain sizes as measured from Fig. 1 [6] by the same technique as used here were 30 to 75 nm. The domain sizes reported here ranged from 50 to 80 nm which is similar to their domain sizes. The quench rates used here were probably faster as judged by specimen

thickness (their 100 um thick ribbons and wires as compared to 20 to 100 ym thick produced here) and as judged by grain size, their 2 to 5 ym compared to 0.7 to 3.4 ym produced here. In this study, rapid ' solidification by itself did not ductilize the pure Ni 3Al even though the solidification rate was faster.

Of the several possible mechanisms that might be responsible for the high hardness of our most rapidly solidified samples, dislocation density and anti-phase domain size must be dismissed. The

dislocation density is low and similar for all specimens and the anti-phase domain size is also essen- tially independent of quenching rate over the range of samples studied. In addition, the domains

observed in our Ni 3Al samples appear to be too coarse to produce any significant hardening [16]. The

differences in grain size and the degree of long-range order remain as the likely variables to explain

this hardness difference, but confirmation of this suggestion must await the results of our quantitative x-ray study.

The complex APB structure in rapidly solidified Ni 3Al revealed by transmission electron

microscopy promises to provide additional information on the APB energy in this highly ordered alloy

and thus on its mechanical behavior. A detailed structural study is in progress, utilizing the

analytical and microdiffraction capabilities of the Philips EM 400T/FEG.

References

[1] N. S. Stoloff and R. G. Davis, Pvog. Mater. Soi, _12(^)> 1~34 (1966).

[2] B. H. Kear, C. T. Sims, N. S. Stoloff, and J. H. Westbrook, eds., "Ordered Al 1 oys — Structural

Applications and Physical Metallurgy," in Proa. 3rd Bolton Landing Conf., Claitor's Publishing

Division, Baton Rouge, Louis., 1970.

[3] H. A. Lipsitt, D. Shechtman, and R. E. Schafrik, Metall. Trans. A llA, 1369 (1980).

[4] S. M. Copley and B. H. Kear, Trans. AIMS 239 , 977 (1967).

[5] S. K. Ehlers and M. G. Mendiratta, paper presented at the Annual TMS-AIME Meeting, Dallas, Texas,

Feb. 14-18, 1982.

[6] A. Inoue, H. Tomioka, and T. Masumoto, J. Mater. Sai. Lett. _1, 377-80 (1982).

[7] K. Aoki and 0. Izumi, Nippon Kimoku Gakkaishi 43 , 1190 (1979).

[8] C. T. Liu and C. C. Koch, to be published in Proceedings of Conference on Trends in Critical

Materials Requirements for Steels of the Future, Conservation and Substitution Technology for

Chromium, Vanderbilt University, Oct. 4—7, 1982.

[9] D. M. Kroeger, W. A. Coghlan, D. S. Easton, C. C. Koch, and J. 0. Scarbrough, J. Appl. Phys. _W,

1445-53 (1982).

[10] B. D. Cullity, Elements of X-Ray Diffraction, Addi son-Wesl ey Publishing Co., 1956, pp. 363—78.

268
[11] E. A. Starke, Jr., J. C. Ogle, and C. J. Sparks, Jr., pp. 372-^0 in Advances in X-Ray Analysis,

12 ed. C. S. Barrett, J. B. Newkirk, and G. R. Mallett, Plenum Press, New York, 1969.

[12] E. M. Schulson, The Rote of Grain Size on the Brittle to Ductile Transition of the Strongly

Ordered Alloy NigAl, Progress Report (DE-AC02-81ER10907) Thayer School of Engineering, Dartmouth

College, Hanover, New Hampshire, March 1982.

[13] C. T. Liu, unpublished research. Oak Ridge National Laboratory, 1982.

[14] E. H. Lee, C. C. Koch, and C. T. Liu, "Rapidly Solidified Amorphous and Crystalline Alloys,"

pp. 375—79 in Materials Research Society Symposia Proc. 8, ed. B. H. Kear, B. C. Giessen, and

M. Cohen, Elsevier Science Publishing Co., New York, 1982.

[15] Y. Calvayrac and M. Fayard, Acta Metall. ^4, 783 (1966).

[16] R. G. Davis and N. S. Stoloff, Trans. Metall. Soc. AIMS 230 , 390 (1964).

269
MICROSTRUCTURAL CHARACTERIZATION OF RSPD STRUCTURES

A.M. Ritter and M.R. Jackson

General Electric Corporate Research and Development

Schenectady, New York 12301

ABSTRACT

When plasma deposition of melted particles is performed in a low-pressure environment,

solidification occurs rapidly. For high-temperature Ni-base superaUoy compositions, particles

smaller than 37;/m are injected into the plasma, accelerated and heated simultaneously, and

flattened either against a substrate or against rapidly solidified plasma deposited (RSPD) material

just deposited. Thickest splats seen in thin foils by transmission electron microscopy are

generally Assuming these are discs created by flattening the largest particles, volume

considerations indicate the discs are on the order of 100 yUm in diameter. Over a considerable

range in deposit temperature, little variation in splat dimension is observed, but substantial

change in the structures of grains and within grains is observed. For substrates heated to approximately 900C before deposition, the solidification structures indicate a quench rate

equivalent to that attained in melt-spun ribbon.

Introduction

Application of rapid solidification technology to Ni-base high-temperature superalloys has been in progress for a number of years (1,2). Powder atomization, self-quenching (laser or electron beam processing),

and melt-spinning (3-5) have received the most attention because of the high cooling rate achievable and the

relative ease of experimentation. Plasma deposition has also received some study, but most of this work was

done at atmos[rfieric pressure, either in air or an inert gas shroud (6,7). The structures formed tended to be

low in density (~'90% of the theoretical density) and to have considerable oxide content (oxygen concentration

of -'1 w/o). For low-pressure plasma spraying (LPPS), low oxide content and nearly full density were achieved

in MCrAlY coatings. Deposition of structural aUoys as free-standing bodies has been termed rapid

solidification plasma deposition (RSPD) (8,9). Such deposits have been found to have excellent structural

integrity and the process is being developed for potential high-temperature components (10).

Because of the interesting mechanical properties displayed by RSPD alloys, a study of as-deposited

microstructure and control of that microstructure was initiated. A characterization of the as-deposited

microstructure will be presented here.

Experimental Procedure

Alloys with chemistries near or equal to the following nominal composition were chosen: .14Cr, .095Co,

.05Ti, .04Mo, .04W, .03A1, .0017C, .0003Zr, .00015B, and balance Ni, by weight fraction. Powders produced

by both gas atomization and vacuum atomization were purchased sieved to pass through a 400 mesh screen, which results in a maximum particle size of 37/Am for spherical powder. These powders were fed at approximately 10-20 kg/h into an Ar-He plasma (Electro-Plasma, Inc. gun) at 52V, 1300A, 60 torr tank

pressure, or into an Ar-H

2 plasma (Metco gun) at 56V, 700 A, 60 torr tank pressure. A deposit was formed on

substrates of steel or copper located at a distance of 25 to 45 cm from the gun. Deposits approximately 1 cm

thick were made either with or without relative translation between gun and substrate.

Slices were cut with a diamond saw from deposits both parallel and perpendicular to the plane of the

deposit. These were for light optical metallography or were mechanically and electrochemically thinned to perforation for transmission electron microscopy. The electrochemical thinning was a jet milling technique

using a 20% perchloric acid/80% methanol electrolyte. Light optical metallography of polished mounted

samples was sometimes complemented with platinum shadowed replica or extraction replica electron

microscopy. Extraction was performed using a 10% Br/90% methanol solution.

Transmission electron microscopy was performed using an Hitachi H-600 or a JEOL JSEM-200

microscope. Bright field, dark field, and electron diffraction imaging were performed for phase analysis and

crystal orientation observations.

270
Results and Discussion

Low-Temperature Deposition

Typical powder particles eire shown in Figure 1. Although the largest particle passing through a 400

mesh screen is 37/' m in diameter, a very wide range of sizes actually is included in the powder fed to the

gun. For both gas atomized and vacuum atomized powders, the particle size distribution includes 0.2 below

10-12/-'m, 0.5 below 25/'m and a mean size of 28/^m, on a volume basis. The degree of segregation within

these particles will depend on particle size, with dendritic segregation being most severe in the coarsest powders.

In the deposits from EPI plasma guns, particles are observed (Figure 2) which have a structure very

similar to the dendritic structure of the as-atomized powder. Some appearance of shock deformation is seen

in the large partide. The question arises as to whether these partides did not melt in the plasma, or whether

they mdted and re-solidified in transit. Experiments have been suggested, such as homogenizing particles

before spraying so that no dendritic structure would exist (11). Any dendritic structure in spherical partides

within the deposit would then be an indication of melting and re-solidification in transit. However, a

consideration of segregation kinetics suggests such experiments will not answer the question.

Figure 1. Micrograph of -400 mesh powder. Figure 2. Spherical partides in RSPD alloy.

Homogenization in the liquid wiU occur over a distance approximated by /Dt, wheje D ^s the diffusivity

and t is the diffusion time. The diffusivity in the liquid will be on the order of 10 ^ cm'^/sec. Since the

distance over which powder travels from gun-to-substrate is ~.3 m and the particle vdocity is greater than

300 m/sec, the transit time is on the order of a millisecond. The homogenization region from /St will be on

the order of l^m. For a particle heated just above the liquidus and then re-solidified, its segregation pattern

will be essentially identical to that which would exist if the partide never had melted. A similar argument

could be made for the lack of segregation in an homogenized particle that was mdted and re-solidified.

This simple analysis suggests such experiments will not darify whether these partides ever melted. It

is of interest, however, that the analysis also suggests partides arriving fuHy molten may be segregated

liquids with compositional distributions essentially the same as in the atomized powder.

Similar questions arise for microstructures such as seen in Figure 3. The structure appears to consist of

dendritic cores sheared on what was formerly interdendritic material. Most likely interpretations include a partide heated prior to impact to a temperature intermediate between the high melting point of the cores and the lower melting point of the interstices; or alternatively, a partide heated on impact by conversion of

kinetic to thermal energy, again to just above the melting point of the interstices. Of course, resolidification of a fully melted partide could have progressed partially prior to impact. A lower

magnification view of an alloy more prone to the shearing behavior is seen in Figure 4. No such question is raised for partides that were fully mdted at impact. Transmission dectron

microscopy (TEM) is shown in Figure 5. A fofl. taken paraUd to the plane of the deposit. Figure 5a, shows an

equiaxed morphology with average grain size well below 1/A.m in diameter. The structure, deposited at approximately 900C, shows a very homogeneous grain size, ho evidence of dendritic or cellular features, and

infrequent occurrence of carbide partides. There is no tS-'V nodular eutectic at grain triple points, such as is observed in cast material. For the aUoy carbon content of 0.15 w/o, a carbide volume fraction of about 0.015

is expected, if all the carbon were present as carbide. The apparent carbide fraction is considerably less them

is seen in conventionaUv cast structures of the same composition. Most of the carbides observed in the RSPD

271
material {ffoduced at 900C are fine, on the order of .01-.2/-m, and exhibit an electron diffraction pattern

consistent with an crystal lattice. Some of the very fine carbides may be MC. Occasionally a much

larger carbide is observed, on the order of 5^m. These exhibit an MC crystal lattice diffraction pattern,

close to that of TiC. It is believed that these carbide particles are ones that were present in the atomized

powder, and that the powder particles did not experience a sufficient time-temperature pulse to solution

these carbides into the liquid. As noted earlier, transit times for powder particles from injection into the

plasma to deposition on the substrate is of the order of a millisecond.

Figure 5. Tranmission electron micrograph.

a) Parallel and b) tranverse-to- plane of low-temperature deposit,

c) Dark field view of b.

272
The distribution of the M

2 oCg carbide particles is displayed in Figure 5b, for a foil taken perpendicular

to the plane of the deposit. In<lvidual splatted discs show a demarcation of their interfaces with adjacent

splats by the presence of the fine carbides. The splats vary in thickness, from a maximum of about 4/^ rn to very thin discs. It is expected that the 4yi/m splats are produced by spreading of the largest powder particles

(37^ m), and that the thinner discs resiit from the entire range of molten droplets that reach the substrate.

Of course, the very finest particles probably are vaporized by the plasma gases, while coarser fractions (still very fine) are not vaporized but are too small in mass to ta’eak free of the plasma jet as it is diverted around

the substrate. The vaporized particles and the low-mass particles are lost to the deposition process.

Within the individual splat layers, a columnar grain morphology is observed, with the grain width

corresponding to the equiaxed grain diameter seen in Figure 5a. Dark field microscopy indicates that for the

'-^OOC deposition temperature, grains grow from one splat layer through the next, sometimes extending over

three splat thicknesses (Figure 5c). The structure within the splats again indicates a lack of segregation, with

carbides only marking inter-splat interfaces.

The well-defined splat interfaces can be interpreted as evidence of a lack of general re-melting by a molten droplet spreading over a just-solidified splat. For particle sizes, powder-feed rates, and particle

densities in the plasma, particles spreading across the surface cover any single location on that surface

approximately 350 times per second. If arrival times from particle-to-particle were identical, this would

require a solidification time 1/350 sec, or .0029 seconds. Because of the statistical nature of the powder

feeding and injection, the lack of re-melting suggests a much shorter solidification time.

Oxygen contents of -400 mesh powders for conventional Ni-base superalloys vary typically from 120-

300 wppm. The lower value is for material in production lots where atomizing chambers are only infrequently

opened to atmosphere, while the hi^er value is for material producd by a batch process, with an atomizer

cleaning sequence immediately before production. If all the oxygen were on powder surfaces, these levels

would correspond to the equivalent of 5-15 atom layers, depending on particle size distribution and the degree

of roughness of the powder surfaces. During RSPD processing, another 50-150 wppm of oxygen is incorporated into the final deposit. This total corresponds to approximately 0.1 a/o oxygen, which would be

expected to produce 0. 2-0.3 v/o oxide in the microstructure. A first consideration might predict that the

oxide would be at inter-splat interfaces, because these might correspond roughly to the outer surface of the

original powder particles. However, droplet spreading on the rough RSPD surface (Figure 6) may involve

considerable deformation and mixing of the droplet-oxygen layer boundary. This may produce a more general

distribution of oxide particles. Observations of oxides in the fuUy melted and splatted structures is extremely infrequent. It is possible that some fraction of the oxygen is [x-esent in the carbide particles as oxy-car tides.

Figure 6. Schematic view of the roughness of

the exposed RSPD surface.

Higji-Temperature Deposition

Figure 7. Transmission micrograph trans- verse-to-plane of high-temperature

deposit.

When a greater energy density is concentrated on the deposition surface, such as by decreasing the gunsubstrate separation or by eliminating relative motion between gun and substrate, changes in microstructure can occur. In situ annealing can occur if temperatures are very great (12), but even at temperatures which do

not significantly alter microstructure after deposition, considerable variations from those shown in Figure 5 can be produced. A deposit made using a Metco plasma gun with no translation of gun relative to substrate is shown in Figure 7. This is from a fott perpendicular to the plane of the deposit for a structure produced at ^1150C. The thickness of individual splats is similar to the lower temperature deposits, but the columnar

273
grain size and the corresponding equiaxed grain size on sections parallel to the deposition fiane is more than an order of magnitude greater than for the lower temperature deposits. Dark field microscopy indicated

growth of individual columnar grains through eis many as 10 Sfdat interfaces. Again, cartHdes marked the

splat interfaces, but the carbides were MC-type carbides, with unit cell dimensions close to TiC. Within the

columnar grains was a considerable substructure shown in Figure 8. The structure appears to be cellular in nature, and this is also seen on the sections parallel to the deposition plane. Cell boundaries have a pronounced demarcation from small carbides as well. The Y' produced in the sample as the deposit cooled

following termination of deposition at a temperature above the t' solvus is evident in Figure 8.

Comparison to Other RS Processes

Powder atomization processes are generally rapid solidification methods for powders less than ^-lOO A m

in diameter. Even for much finer powders, cooling rates are such that dendritic structures frequently result. Cooling rates can be estimated from the spacing between secondary dendritic arms found in the powder

particles. However, in the splat-quenched RSPD structures no evidence is seen of dendritic solidification in the low-temperature deposits, while a cellular solidification microstructure is observed fot higher temperature deposits. This comparison to powder structures indicates a cooling rate greater than lO^K/sec.

A similar comparison can be made with structures from melt spun ribbon. A sample of an RSPD Nibase superaUoy was re-melted and rapidly solidified by a melt-spinning operation, and the resulting ribbon was examined by TEM. A fine, equiaxed structure was observed in foils prepared parallel to the plane of the

ribbon (Figure 9). Replica electron microscopy of sections polished transverse to that plane showed that the

majority of the ribbon thickness consisted of columnar grains. The equiaxed structure seen in TEM was

essentially identical to that observed in the RSPD material, with a grain size for the equiaxed structure of

less than 0.5/'.m. The microstructural similarity for the same composition produced by RSPD and by meltspinning suggests^ th£g; cooling rates for the two processes are also comparable. A cooling rate of

approximately 10 -10 K/sec can be estimated for the RSPD structures from such a comparison. This is a reasonable rate compared to estimates made based on splat thickness, the maximum solidification interval

and the thermal gradients imposed by the process.

Figure 8. Transmission micrograph transvers e-to-s plat boundary and splat

interior.

Figure 9. Transmission micrograph through

the thickness of a melt-spun rib- bon. »

274
Summary

Microstructures of as-deposited RSPD Ni-base superalloys are characterized by a layered assembly of

discs generally less than 3A^m in thickness, which formed as liquid droplets impacted the surface and spread

prior to re-solidifying. Within the discs are columnar grains, with the long axis of the grains perpendicular to the plane of the disc. In the plane of the discs, the columnar grains are equiaxed. Grain growth occurs from

splat-to-splat, with nucleation of new grains eliminating further growth in about three splat thicknesses at 900C deposition, and in about 10 splat thicknesses at 1150C deposition. The lower temperature depositions

appear to have solidified in a planar mode, with splat boundaries marked by fine carbides. The higher

temperature deposition shows a cellular microstructure, with both splat boundaries and cell boundaries

marked by carbides. A fraction of the microstructure consists of spherical particles that appear to have

arrived at the substrate without melting, and other

^ particles that appear to be melted only along

interdendritic interstices. Cooling rates in excess of lO^K/sec are expected.

Acknowledgement

The authors wish to note the technical discussions with J.R. Rairden, R.W. Smith and J.S. Smith (all of

General Electric Company). Depositions were made by S.F. Rutkowski, W.A. Seaman and J.A. Resue.

Metallographic preparation was performed by J.A. Resue.

References

(1) D.J. Looft and E.C. vanReuth, "Rapid Solidification Processing: An Overview," Proc. of the Int. Coif, on Rapid Solidification Processing, Reston, Va., p. 1, 1977.

(2) R.J. Patterson, H, A.R. Cox and E.C. vanReuth, "Rapid Solidification Rate Processing and Application

to Turbine Engine Materials," J. of Metals, p. 34, Sept. 1980.

(3) P.N. Ross and B.H. Rear, "Surface Segregation in Rapidly Solidified Powders," Proc. Int. Coif, on Rapid

Solidification Processing, Reston, Va., p. 278, 1977.

(4) P. Duwez and R. Willens, "Rapid Quenching of Liquid Alloys," Trans. AIME, 227, p. 262, 1963.

(5) E. Breinan, B. Rear, C. Banas and C. Greenwald, "Surface Treatment of Superalloys by Laser Skin

Melting," Proceedings of the 3rd Int. Sym. SuperaUoys: Metallurgy and Manufacture, Edited by Rear, et

al, Claitors Publishing division. Baton Rouge, LA (1976) p. 435.

(6) D.R. Mash and I. MacP. Brown, "Structure and Properties of Plasma Cast Materials," Matls. Engr.

Quart., p. 18, Feb. 1964.

(7) R.W. Cahn, "Rapid Solidification by Plasma Spraying," Proc. Int. Conf. on Rapid Solidification

Processing, Reston, Va., p. 129, 1977.

(8) J.R. Rairden, M.R. Jackson, M.F.X. Gigliotti, M.F. Henry, J.R. Ross, W.A. Seaman, D.A. Woodford and

S.W. Yang, "Coatings for Balanced Environmental/Mechanical Behavior," Proc. Int. Conf. on HighTemperature Corrosion, NACE, San Diego, p. 72, 1981.

(9) M.R. Jackson, J.R. Rairden, J.S. Smith and R.W. Smith, "Production of Metallurgical Structures by

Rapid Solidification Plasma Deposition," J. of Metals, p. 23, Nov. 1981.

(10) P.A. Siemers, "Rapid Solidification (RSPD) for Fabrication of Advanced Aircraft Gas Turbine

Components," Contract F33615-81-C-5156, Interim Reports, Jan. 15, 1982, Apr. 15, 1982, July 15, 1982,

Air Force Systems Command, AFWAL, Wrigth-Patterson Air Force Base, Ohio 45433.

(11) H. Lipsitt, Materials Laboratory, AFWAL, Wri^t-Patterson Air Force Base, Ohio, private communicatiOTi, 1982.

(12) T.L. Cheeks, M.E. Glicksman, M.R. Jackson and E.L. Hall, "Temperature Effects on Plasma Deposition

of Ni-Base Materials," this conference proceedings.

275
PRECIPITATION SEQUENCES IN A RAPIDLY SOLIDIFIED Ni-Al-Mo-W ALLOY

Dennis M. Dimiduk

AFWAL Materials Laboratory

Metals and Ceramics Division

Wright-Patterson AFB OH 45433

ABSTRACT

Rapid solidification P/M techniques have been used in developing nickel -base superalloys with high molybdenum content. Previous studies on these consolidated powders have

shown that a sequence of precipitation reactions occurs which results in the formation of

Ni Mo compounds in many of the Ni-Al-Mo-W alloys. The present study centers around the

thermal stability of these Ni Mo phases and their effects on the stability of the y'

strengthening phase. Long-te^m aging at elevated temperatures has been performed in order

to establish the dissolution temperatures of the D1 , DO 22 . Pt 2 Mo and BCC stereotypical

phases. The results of TEM observations on the heat-treated alloys, as well as the implications that these precipitate phases have for the design of Ni-Al-Mo base alloys, will be

presented.

Introduction

One of the many fascinating areas of application for rapid solidification technology (RST) over ’ the last five years has involved nickel-base superalloy development. In particular, P/M techniques | have allowed the Ni-Al-Mo ternary base system to be explored over a much broader range of composition^

than that previously explored in the y/y' directionally solidified eutectics. RST has allowed

homogeneous, equiaxed grain morphologies to be studied in addition to aligned grain structures. | Recent studies of this alloy base have included a wide range of quaternary and higher-order elemental#

additions [1]. Cox, et al., in reporting mechanical test results of aligned grain forms of these ^ alloys, demonstrated that a W-containing alloy, RSR 185, had very attractive high-temperature creep |

strength U>2]. Martin, et al . , in performing microstructural studies, indicated that after high- J temperature extrusion and air cooling, the matrix of this alloy was not simply a solid solution of i f.c.c. Y> but rather that a number of ordered precipitates of the class Ni Mo existed between the | cuboidal particles [3]. In another study Pearson, et al . , demonstrated th^t the presence of DO 22 - i type (Ni 3 (MoTa)) coherent precipitates had a favorable effect upon the creep properties of single

crystals [4]. The composition which they studied was a Ta-containing version of a Ni-Al-Mo base . alloy also investigated by Martin, et al . [3]. Kersker, et al . , have recently shown in the same j composition that the coherent D 022 -type precipitate can serve to stabilize the y'

,

LI 2 -type phase up

j

to the disordering temperature of the DO 22 [5]. ]

I In light of these observations, the present study was initiated to determine the critical order- I ing temperatures (T ) of the various Ni Mo phases which form in the RSR 185 alloy. Additionally, the ^ precipitation of the b.c.c. a-Mo phase ^as studied through the temperature of dissolution. The j intent of these studies was to identify the temperature range over which the reported microstructural 1 effects exist and consequently determine the maximum temperature at which the mechanical properties ; could be affected by the presence of these phases.

Experimental Procedure

The material investigated was prepared as powder by the rotary atomization process at Pratt and

Whitney Aircraft and has a nominal composition of 72.6Ni-6.8Al-14.4Mo-6.2W-0.4C (wt. %). Powder of

-140 mesh size was canned, out-gassed, and extruded with a 26:1 reduction ratio at 2300°F. Prior to heat treatment, the stainless steel can was removed from the extruded bars by machining. Slices (3 cm long) of the bars were prepared for heat treatment in air. Initial microstructural evaluation

showed the extrusions to be consistent with those of previous studies [3]. Uniform grains having an ordered "y matrix" of both the DO22 and Pt 2 Mo (Ni 2 Mo) isomorphous phases existed, with a uniform

distribution of y' and an MgC-type carbide throughout the grains. All samples were given a y'
solution heat treatment at 2390°F for 1 hr,, followed by an ice-water quench. This treatment resulted

in a very uniform distribution of y' of about 0.1 ym edge length throughout the grains, with the y'

necklace entirely solutioned. Examination by TEM showed the matrix to be disordered,

A sequence of up-quenching experiments was chosen for the determination of T for each of the

phases. A 3 cm bar of the alloy was placed in the furnace at the desired temperature for initiating

the aging treatment. After 50 hr. the bar was water quenched, and a 5 mm slice was removed for preparation of TEM samples. The remainder of the bar was then inserted into a furnace at the next temperature of interest and aged for an additional 50 hr. This procedure was chosen so that if the

ordering reactions for the Ni Mo phases were very rapid and occurred upon cooling (as does the y'

reaction), the T could still^be determined by the change in precipitate morphology which is expected

to occur after s8lutioning and re-precipitation. The T of the Ni Mo phases was determined by

examining the temperature range 1350-1700°F in 50-deg. ^increments' Temperatures as high as 2300°F

were required in studying the dissolution of the a-Mo phase. TEM examination was carried out in

order to observe the ordered "matrix" phases, y' morphology, and grain-boundary phases and morphology.

• FUNDAMENTAL f.c.c. • Ni^Mo ORDERING

X LI2, T' SUPERLATTICE — DO22 ORDERING

( 002 )

( 022 ) ^

* ' ^

%

\ •

. \ • ( 000)

( 020 )

. lOOOA

Figure 1 TEM micrographs and selected area diffraction patterns showing Ni£Mo phase:

(a) B. F. , g = (020) (b) D.F. using 1/3 (022) reflection (c) [100] zone-axis pattern with

DO22 and Ni 2 Mo (d) indexed trace of (c)

Results/Discussion

Structural identification of the Ni Mo-type phases by electron diffraction is most readily

accomplished by examination of zone-axi s^patterns which contain <220> and, more importantly, <420>

type reciprocal lattice vectors. Individual variants of these phases cause intensity peaks to occur

at one-third, one-fourth, one-fifth of the unit <420> distances, depending upon the particular

phase(s) present. These occurrences have been discussed in some detail in the literature and will

not be discussed here [6-8]; however, it is these features of electron-diffraction patterns which

were used to detect the Ni Mo phases during these studies.

X

277
In this W-containing alloy, the DO 22 - and Pt 2 Mo-type phases develop during aging in the temperature range 1350-1500°F. The D1 -type (Ni 4 Mo) structure was not observed in this alloy at any point

during aging of the solution-triated and water-quenched samples. Separate studies of RSR-185 aged in the as-extruded condition did reveal the D1 phase when the alloy was aged below 1500°F. It is

expected that this discrepancy can be accounted for by the non-equilibrium constitution of the matrix

which exists both in the as-extruded and in the solution-treated-and-quenched samples. Figs 1(a) and

(b) are bright- and dark-field micrographs, respectively, of the 1400°F-aged samples. The diffractior

pattern and schematic trace in Figs 1(c) and (d) show that reflections from the DO22 phase exist

along with those from the Pt 2 Mo-type phase.

Further aging of the RSR-185 samples indicated that the DO 22 a solution temperature between

1500 and 1550°F, while the dissolution of the Ni 2 Mo phase did not occur until 1600°F was reached.

These results appeared to be consistent with those recently reported by Martin and Williams [9];

however, the non-coherent Ni 2 Mo reported in their work was not observed here. A phase of similar

morphology was observed above 1650°F; however, convergent-beam electron-diffraction studies indicate

that this phase has a structure type and lattice parameter consistent with an M 12 C-type carbide and

that the orientation relationship and lattice parameter are such that some of the <220> and <420>

reciprocal lattice vectors are divided into thirds, just as they are by the Ni 2 Mo phase [10] .

Figure 2 (a) One quadrant of [lOO] zone with (Oil) satellites (b) D.F. using (Oil) and

satellite spots (c) D.F. using 1/3 (022) (d) D.F. using 1/3 (022)

Of potentially greater interest in this study are the observations made on the samples aged

between 1600 and 1700°F. In this temperature range the Ni 2 Mo is going into solution and, simultaneously, the b.c.c. a-Mo phase is precipitating throughout the matrix. Fig 2 shows an <001> diffraction pattern along with bright- and dark-field images of a microstructural feature which has been

shown to be unique to the early stages of a-Mo precipitation. The high-aspect-ratio regions shown in

Fig 2(b) give rise to the pairs of satellite reflections associated with the (Oil) reflections in Fig

2(a). These high-aspect-ratio features (referred to here as a'-Mo for discussion purposes) are

278
characterized by the satellite pairs on the {110}-type reflections on all zones containing {110}.

The exact nature of the a'-Mo is not yet understood and cannot be readily explained by those features

which are known to cause satellite reflections in electron-diffraction patterns [11,12]. Particularly

puzzling in this case is the fact that the satellite pairs occur only on the superlattice reflections

and not on the fundamental reflections. Inspection of Figs 2(c) and (d) indicates that regions

depleted in one variant of the Ni^Mo phase exist adjacent to a siiigle corresponding variant of the

d'-Mo. This precipitation reaction remains the subject of further study.

Figure 3 TEM B.F. micrographs of water-quenched RSR-185 after: (a) long-term aging at 1600°F

(b) same sample after 50 hr. at 1650°F

Continued aging of the RSR-185- al loy at 1700°F results in complete precipitation of a-Mo in two different orientation relationships. These relationships were part of a separate study and were

determined to be the Nishiyama-Wasserman and Pitsch orientation relationships for f .c.c. :b.c.c.

systems [13]. It was determined that the a-Mo phase did not completely solution until the solvus

temperature for the matrix y' had been exceeded (about 2300°F).

Throughout this study the size and morphology of y' was observed after each aging treatment. It

is important to note that very little coarsening of the y‘ occurs during aging through 1600°F or while the Ni Mo phases exist as a "matrix". Fig 3(a) is a bright-field micrograph of a sample aged

in 50-deg. il^crements throughout the range 1350 to 1600°F for a total time at temperature of 300 hr.

This image is compared with the same sample after an additional 50 hr at 1650°F. A marked change

occurs in both the y' size and morphology after the additional time at higher temperature. Although

279
the mechanistic evidence for Ni Mo phases restricting y' coarsening has not yet been obtained, the

evidence in Fig 3 indicates that^coherent Ni 2 Mo is an effective stabilizer of the y' phase. In light

of this evidence and the evidence put forth by Pearson, et al . [4], one would expect that alloys

designed for use below 1600°F could incorporate such precipitates for microstructural and, thereby,

property control. In the event that the Ni-Al-Mo system is explored as a polycrystalline alloy with

Ni Mo precipitates, efforts would have to be devoted to finding ways of detering the discontinuous

pricipitation reaction known to occur in the grain boundaries of these alloys. Several studies have

shown that such reactions result in a drastic decrease in tensile and rupture ductility [1,14].

Summary

The critical ordering temperature for the Ni 3 Mo phase in RSR-185 was found to be between 1500

and 1550°F, while the Ni 2 Mo phase was found to be stable for about another 100 deg. or to 1600°F.

A microstructural feature (a'-Mo) which appears to be a precursor phase to the b.c.c. a-Mo was

shown to precipitate after long times at 1600°F. This feature gives rise to paired diffraction

intensity on the {110} superlattice reflections but not on the fundamental (220}-type reflections.

The y' precipitate phase appears to be stabilized by the presence of coherent Ni 2 Mo in the "y"

regions of the alloy. This stabilization appears to be effective to a maximum temperature equal to the T^ of Ni 2 Mo (1600°F).

Acknowledgments

The author wishes to thank Dr P. L. Martin for many helpful discussions regarding this study.

The assistance offered by Mr M. J. Kaufman and Professor H. L. Fraser, in performing convergent-beam

electron-diffraction analysis, is also acknowledged.

References

[1] A. R. Cox, R. G. Bourdeau, R. J. Patterson, E. H. Aigeltinger, R. E. Anderson, J. W. Broch, and

T. D. Tillman, Application of Rapidly Solidified Superalloys , AFWAL-TR-80-4051 , May (1980).

[2] R. E. Anderson, A. R. Cox, T. D. Tillman, and E. C. van Reuth, Rapid Solidification Processing

II , p. 416, edited by R. Mehrabian, B. H. Kear, and M. Cohen, Cl ai tor's (1980).

[3] P. L. Martin, H. A. Lipsitt, and J. C. Williams, Rapid Solidification Processing II , p. 123,

edited by R. Mehrabian, B. H. Kear, and M. Cohen, Cl ai tor's (1980).

[4] D. D. Pearson, B. H. Kear, and F. D. Lemkey, Creep and Fracture of Engineering Materials and

Structures , p. 213, edited by B. Wilshire and D. R. J. Owen, Pineridge Press (1981).

[5] M. M. Kersker, E. A. Aigeltinger, and J. J. Hren, Proceedings from Thirty-Ninth Annual EMSA

Meeting , p. 46, edited by G. W. Bailey, Claitor's (198iy!

[6] S. K. Das and G. Thomas, Phys. Stat. Sol. (a), vol . p. 177 (1974).

[7] M. Yamamoto, F. Shohno, and S. Nenno, Trans. JIM, vol. 21, p. 475 (1978).

[8] E. Ruedl , P. Delavignette, and S. Amelinckx, Phys. Stat. Sol., vol. 28, p. 305 (1968).

[9] P. L. Martin and J. C. Williams, "Microstructural Development During Aging of Ni-Mo-Al-X

Alloys," International Conference on Solid-Solid Phase Transformations , ASM, (1982).

[10] D. M. Dimiduk, Private Research (1982).

[11] P. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley, and M. J. Whelan, Electron Microscopy of

Thin Crystals , Butterworth and Co., London (1967).

[12] G. Thomas and M. J. Goringe, Transmission Electron Microscopy of Materials , John Wiley and Sons,

Inc. , New York ( 1979)

[13] D. M. Dimiduk, "Orientation Relationships between BCC Mo and FCC y in a Ni-Al-Mo-W Superalloy,"

to be published in Met. Trans. A.

[14] K. A. Williams, D. B. Miracle, and H. A. Lipsitt, to be published.

280
EFFECT OF THE RAPID-SOLIDIFICATION METHOD

ON MICROSTROCTURES AND PROPERTIES OF Al-Li ALLOYS

P. J. Meschter, J. E. O'Neal, and R. J. Lederich

McDonnell Douglas Research Laboratories

St. Louis, Missouri 63166

ABSTRACT

Extrusions of the alloy Al-3Li-l.5Cu-0.5Co-0.2Zr have been prepared from vacuumatomized powder, ultrasonically atomized powder, and roller-quenched flakes. The two

powder-production methods yield extrusions of comparable microstructures and mechanical

properties and are thus equally suitable for scale-up to produce large batches of Al-Li

alloy powders. The roller-quenched flakes were difficult to consolidate to 100% density,

and the ductility of the flake extrusion was adversely affected by interflake debonding.

The microstructure and properties of Al-3Li-l.5Cu-0.5Co-0.2Zr were shown to be more sensitive to variations in alloy composition than to the particulate production method

or rapid-solidification rate.

Introduction

Most of the numerous previous investigations of rapid-solidification processing of Al-Li alloys

have concentrated on the influence of alloy chemistry, processing variables, and heat treatment on microstructure and properties. Rapidly solidified particulates have generally been produced by only

one method in any one study, and few properties data exist on the same alloy prepared by different

rapid-solidification (R/S) methods. The object of this investigation was to isolate the effect of the

R/S method on the microstructure and properties of the low-density, high-modulus, high-strength alloy

Al-3Li-l.5Cu-0.5Co-0.2Zr, and to develop criteria for deciding which R/S method is most suitable for

production of large volumes of Al-Li alloy particulates.

Experimental Procedure

Rapidly solidified powder and flakes for this study were prepared by vacuum atomization

(Homogeneous Metals, Inc., Herkimer, NY), ultrasonic atomization, and twin roller quenching (both

performed at McDonnell Douglas Research Laboratories). Salient characteristics of the particulates

produced by these three methods are summarized in Table 1. Average solidification rates were calculated using the correlation with dendrite arm spacing developed by Matyja et al.[l] Powders produced

by vacuum atomization and by ultrasonic atomization are comparable in average particle size and dendrite arm spacing and hence have similar average solidification rates. The average solidification

rate in the roller-quenched flakes is more uniform than in the powders and is approximately an orderof-magnitude larger.

Powders and chopped flakes were canned, cold consolidated, vacuum degassed at 400°C, and extruded

into 1.25 cm diam axisymmetric extrusions at 400°C and an extrusion ratio of 16:1 by Nuclear Metals,

Inc., Concord, MA. Chemical compositions of the extrusions are listed in Table 2. The two powder

extrusions had comparable chemistries, while the flake extrusion was low in Li and Zr and thus had

significantly lower strength levels. Samples from the extrusions were solution annealed at 530°C for

1 h. 190°C aging curves of all samples were similar and showed peak hardness after 4 h aging time.

281
TABLE 1 COMPARISON OF RAPID SOLIDIFICATION (R/S) METHODS

Characteristic of

R/S Method

Vacuum

Atomization

Ultrasonic

Atomization

Twin Roller 4

Quenching J

Quench medium Vacuum (expansion of

pressurizing gas)

Argon at 8.3 MPa flowing at Mach 2, pulsed at ~ 80 000 HzStainless steel ^

rolls at 1725 rpm 3

Particulate shape Spherical powder Spherical powder Flakes

Particulate dimensions Avg. diam = 46 pm 80% of powder between

32 pm and 58 pm diam

Avg. diam = 59 pm

80% of powder between

38 pm and 89 pm diam

5 cm X 0.5 cm

j X 125 pm thick ^

-

Dendrite arm spacing (D.A.S.) 0.7-4 pm

D.A.S. at average

particle diam = 2.4 pm

0.7-5. 5 pm

D.A.S. at average

particle diam =2.9 pm

0 . 6- 1.0 pm a

Average solidification rate - 10^ K/s - 10^ K/s - 10^ K/s 1

ANALYZED

TABLE 2 COMPOSITIONS OF Al-3Li-l .5Cu--0.5Co-0.2Zr EXTRUSIONS

Analyzed Composition (wt%)

Rapid Solidif.ication Method Li Cu Co Zr

Vacuum Atomization 3.14 1.63 0.44 0.21

Ultrasonic Atomization 3.06 1.40 0.46 0.18

Roller Quenching 2.68 1.47 0.32 0.10

Microstructure and Properties

MicrosCructures of solutioa-annealed and of peak-aged Al-3Li-l.5Cu-0.5Co-0.2Zr extrusions produced from the three types of particulates are shown in Figures 1 and 2. The sizes and distributions

of constituent particles, 5

' precipitates, incoherent Co-containing dispersoids, and subgrains are similar for the two powder extrusions (Figures la, lb, 2a, 2b, 2d, and 2e). The partial recrystallization in the flake extrusion (Figure Ic) is attributable to the smaller Zr concentration in this

extrusion, and the smaller density of constituent particles in the flake extrusion is attributable to the lower Li concentration. Thus the rapid-solidification method and solidification rate (> 10^ K/s)

have little effect on the solution-annealed or peak-aged microstructures of the alloy.

I

Mechanical properties of solution-annealed and peak-aged extrusions are summarized in Table 3. The ductility of the alloy is low regardless of particulate type or heat treatment. Peak-aged

strengths of extrusions made from vacuum atomized and ultrasonically atomized powders are virtually ' identical, in keeping with the similarity of the microstructures. The flake extrusion does not have ai

higher ductility than the powder extrusions even at a considerably lower yield strength level. 8

Fractographs of the peak-aged extrusions shown in Figure 3 reveal fracture largely along subgrain ^

boundaries in all three types of extrusions. In addition, the vacuum-atomized powder extrusion shows

relatively large cleaved regions (Figure 3a), which are not present in the fractographs of the other

two extrusions. No convincing explanation for this effect exists at this time. The ultrasonically j atomized powder extrusion shows some evidence of interparticle debonding (Figure 2b). The flake

extrusion shows extensive interflake debonding, aggravated by the existence of occasional pre-existing

voids on flake boundaries as a result of incomplete compaction in the original extrusion. Flake

debonding adversely affected strength and ductility in the peak-aged alloy. Mechanical testing of a flake extrusion of a composition comparable to those of the powder extrusions is required to quantita- ' tively deteriQine the extent of mechanical properties degradation caused by flake debonding. *

282
Figure 1. Optical microstructures of solution-annealed Al-3Li-l.5Cu-0.5Co-0.2Zr extrusions produced

from a) vacuum atomized powder, b) ultrasonically atomized powder, and c) roller quenched

flakes

TABLE 3 MECHANICAL PROPERTIES OF Al-3Li-l.5Cu-0.5Co-0.2Zr EXTRUSIONS

Solution-Annealed Solution-Annealed and Peak-Aged

Process

0.2% Yield

Stress, MPa (ksi)

Elongation

% 0.2% Yield

Stress, MPa (ksi)

Elongation

%

Vacuum Atomization 270 (39.2) 3.6 515 (74.6) 0.7

Ultrasonic Atomization 211 (30.6) 4.4 509 (73.5) 1.6

Roller Quenching 151 (21.8) 6.1 399 (57.9) 1.1

283
(d) i 1 0.5 ^un

Figure 2. Microstructures of Al-3Li-l.5Cu-0.5Co-0.2Zr (a) - (c) solution annealed 530°C/1 h: (a) vacuum atomized powder, (b) ultrasonically atomized powder, and (c) roller quenched flakes,

(d) - (f) 530°C/1 h plus peak-aged 190°C/4 h. Dark field, 6^ reflection (d) vacuum

atomized powder, (e) ultrasonically atomized powder, and (f) roller quenched flakes.

284
Figure 3. Microstructures of tensile fracture surfaces of peak-aged Al-3Li-l.5Cu-0.5Co-0.2Zr

extrusions prepared from (a) vacutim atomized powder, (b) ultrasonically atomized powder,

and (c) roller quenched flakes.

Conclusions

1) Vacuum atomization and ultrasonic atomization are methods of powder production of Al-Li

alloys which yield powders of similar morphology, size range, and average solidification rate. Consolidated forms produced from the two types of powders have similar microstructures and properties.

These two methods are thus equally suitable for scale-up to produce large quantities of powder.

2) Rapidly solidified powders are preferable to flakes or splats because they are more readily

consolidated to 100% density and show superior interparticle bonding under comparable consolidation

processing conditions.

3) Microstructures of solution-annealed and peak-aged Al-3Li-l.5Cu-0.5Co-0.2Zr extrusions are more sensitive to variations in chemistry, particularly Li and Zr levels, than to original particulate

type or rapid solidification rate, provided that this rate is > 10^ K/s.

4) The low ductility in peak-aged Al-3Li-l.5Cu-0.5Co-0.2Zr is not directly attributable to particulate type, except that interflake debonding adversely affects ductility. Current research on Al-Li alloys at McDonnell Douglas Research Laboratories centers on the specific effect of the

precipitate-former Cu, high hydrogen levels, and the interaction between incoherent dispersoids and

substructure, as alternative sources of low ductility in this alloy.

Acknowledgment

This research was sponsored by McDonnell Douglas Corporation Independent Research and Development

(IRAD) program.

Reference

[1] H. Matyja, B. C. Giessen, and N. J. Grant. J. Inst. Metals 96, 30 (1968).

285
!

AGING BEHAVIOR OF DEVITRIFIED Ti RICH

ALLOYS CONTAINING METALLOIDS

Sung H. Whang and Y.Z. Lu

NORTHEASTERN UNIVERSITY

Boston, MA 02115

ABSTRACT

Rapidly quenched Ti rich glassy alloys containing B and Si were made and devitrified

as functions of temperature (550 - 950°C) and time. Different aging responses for Ti,--

Zr.^Nii^Bi^ and Ti^^Zr.^Ni .^Si were observed from the annealing experiments. For the

former .initial moaest^aging response was followed by a strong second response after a long-hour annealing at 700°C while for the latter the first response only was detected.

Also, the second aging in the Ti-Zr-Ni-B alloy develops very high hardness as well as a brittle phase.

Microstructure of heat treated TiggZrgAl .Ni ,gB .

glass at 950°C for 2 hours shows

that grain size remains within sub-micron range Tn diameter while precipitate grows into a half-micron size of acicular shape.

Introduction

Recently, a unique microstructure exhibiting high hot and cold strength has been synthesized

via devitrification of glassy alloys containing metalloids [1][2][3]. Particularly, this group of

alloys, in which sub-micron equiaxed metalloid dispersoid and metalloid-insoluble metal matrix form

a network structure, gives rise to a microstructure stable at high temperature with unique mechanical properties. Similarly, a large quantity of metalloid elements ('\^up to 20 at %) can be dissolved

in Ti alloys via glass formation and later precipitated out by annealing. The following is a preliminary study of devitrified microstructure and mechanical properties of Ti rich glasses.

Experimental

An alloy button prepared in an arc furnace under inert gas atmosphere was splat-quenched using

hammer and anvil technique. The sample foils were examined by x-ray diffractometer to determine

amorphousness. Subsequently, the crystallization of the glassy alloy was studied by DSC-2 calorimetry with a heating rate of 80 K/min and a temperature range up to 1000 K. To study aging behavior

the samples were wrapped in thin tantalum foil and sealed in quartz capsules under vacuum (10 torr)

the capsules samples were heat treated in a furnace as functions of temperature and time. Hardness

of the samples was measured by Vickers diamond pyramid tester using 50 and 100 g loads.

Results

a) Glassy State

Table 1 shows measured hardness (Hv) and crystallization temperatures (Tx) for some Ti-Zr-M-B

(Si) alloys and as-quenched Ti-6A1-4V alloy, where M is Cr, Mo, Ni and A1 . Young's modulus was estimated from a correlation between hardness and Young's modulus in glassy alloy [4]. Each alloy has more than one exotherm. The temperature at which the strongest peak occurs was indicated with an

286
Table 1 Properties of Ti-Rich Glassy Alloys Before and After Crystallization

Alloy As-Quenched

Structure Microhardness

Hv (GPa) EstimatedYoung's Modulus

(E) (GPa)

Crystal 1 i zation Temperature (°C) Aged; 750°C

2h, Hv (GPa)

Ti^O^riof^ilO^lO Glass 6.31 (G) .115 357,452 9.7

^''68^’"6*^4^''l0®12

G1 ass 6.27 (G) ->.114 407,477 9.6

Ti^qZi^io^'IiO^^O Glass 4.94 (G) -.87 413,544* 8.8

Glass 5.83 (G) '.106 419, 551*, 567 6.2

T1-6A1-4V Cryst. 4.28 (C)

*Major exotherm

asterisk in the table. It is interesting to note

that TiyQZr.^Ni, does not exhibit any strong

exotheniial peak as Shown in Fig, 1(b).

The crystallization temperatures of Ti^ Zr^nNi^^-

Si.Q are similar to those of JVasP

539°C) [5]. The Ti all6?s &8nt^?ning B show higher hardness in both glassy and crystalline states than the Ti alloys with Si. Overall,

hardness values at aged state are as much as 10 - 80% higher than those at glassy state.

b) Heat Treatment

As-quenched glassy alloys were crystallized

isochronal ly at a temperature range of 550 - 950°C

for 2 hours. Measured hardness for Ti.,^Zr,„Ni,„-

^^10 ZO^'^IO^^' 1 0^1 D ®^^oys was plottea against

anrtealing tfemperatuf'e as shown in Fig. 2. For both alloys lower temperature annealing gives rise

to higher strength than higher temperature annealing, which is expected from the growth kinetics of

grain and precipitate. The fact that in Ti.,.Zr..-

Niio^lQ ha<"dness does not change in the range ot^

756 - 950°C may be indicative of insignificant

microstructural change above 750°C. Isothermal

annealing was conducted on Ti.,„Zr. „Ni ,„B, „ and

"'"^'ZO^’^IO*'^^' lO^Uo 700°C below 3'^transus^shown in

Fig. 3. Initial aging response presumably comes from a-Ti and titanium-metalloid compound. However, the second aging response for the Ti-Zr-NiB alloy occurs after ten hour annealing. Also, it

is interesting to point out that similar double

aging was observed in Ti„„ ^Zr^^B^ but not in

Ti

90 -xZrioSix, where x [6^^

Figs. 1(a) and 1(b) One of two exotherms for the

Ti-Zr-Ni-Si (a) is very strong in contrast with

two weak peaks in the Ti-Zr-Ni-B (b).

287
79

11“

11 .

10.

9.0

8. 0

“

* 10^10

1 I M E = 2 h

X = B,Si

7.0-

6. 0 ^^-^-

5.0^ A.Q.

' ^0 ~650 750 850

Annealing temperature, “c 9^

Fig. 2 Isochronal annealing (2h) at a temperature

range of 550“C - 950°C for

10^10 ''^Vo'

Fig. 3 Isothermal annealing at 700°C at a range of

0.5 - 40 h for Ti^piZr,,-

Ni,„B,^ and Ti.^ZfVoNn^-

Si 10^10 '70 10'” 10 10-

Microstructure

Bright field TEM micrograph of TiggZr^Al ,Ni

,

qB,alloy aged at 550°C for 4 hours shown in Fig. 4(a)

exhibits very fine early stage crystallization ana grain growth. The corresponding diffraction pat- tern shown in Fig. 4(b) also confirms the fact that grain size is indeed very small from the continuous

first debye ring. A number of bright spots inside the first ring indicate the existence of precipitates having a larger cell size, probably borides.

In Figs. 5(a) and 5(b), microstructure of the Ti-Zr-Al-Ni-B alloy annealed at 950°C for 2 hours

exhibits sub-micron grains and randomly scattered needle shape precipitates. From these micrographs

grain growth in the devitrified Ti alloy is not serious even at the elevated temperature and the grain

size remains within a sub-micron range while precipitate grows large enough to alter mechanical properties of the alloy.

288
Fig. 4(a) Electron Micrograph of Ti^cZr^-Al .Ni ,p,- Fig. 4(b) Diffraction pattern of 4(a).

B ^2 3ged at 550°C for 4 hours.

Fig. 5(a) Micrograph of the Ti-Zr-Al-Ni-B alloy Fig. 5(b) Higher magnification micrograph of 5(a).

annealed at 950°C for 2 hours.

289
Discussions

Glass forming ability of Ti alloy containing Si is superior to that of Ti alloy containing B. * For example, glass forming was observed in binary Ti-Si but not in binary Ti-B. Also, Ti,.Zr,p,Ni,Si

alloy forms glass upon quenching while TigQ_ Zr,QNi B^q alloys require full 10% Ni to go to giassy

state. "

)

I It is speculated that precipitate reaction corresponding to the second aging response shown in | the isothermal annealing of the Ti-Zr-Ni-B alloy may not be shown in the DSC-pattern as well as in

the isochronal plot probably because of its high temperature and sluggish kinetics. The reason for

the double aging in the Ti-Zr-Ni-B alloy may be as follows. For the initial age hardening, orthorhombic TiB or unknown (Ti, Zr)B type may be formed (based on phase diagram information). At the same

time, the formation of hexagonal ZrB^ or perhaps (Zr, Ti)B

2 can be possible thermodynamically although ! nucleation and growth of this compound may be sluggish. Therefore, the second aging is believed to be j associated with the precipitation of this compound. On the other hand, single aging response for the ' Ti-Zr-Ni-Si alloy is consistent with that for Ti-5Zr-lSi [7], in which (Ti, Zr)(-Si^ was identified as

'

a G.P. zone precipitate instead of TirSi.,. The devitrified Ti alloys in this experiment are brittle { except the Ti-Zr-Mo-Ni-Si . Furthermore, "^extensive brittleness develops in these alloys after long

time annealing at a higher temperature. This is coincident with the microstructure of the annealed

Ti-Zr-Al-Ni-B alloy in which acicular shape precipitates were found. However, for the Ti-Zr-Mo-Cr-Ni- | Si alloy, a significant portion of ductility remains after devitrification and annealing at higher

temperature (700 - 1000°C).

Acknowledgements

The authors wish to thank Mr. G. Mouchawar and K. Carlson for their assistance in carrying out

this experiment. This work has been supported by the Office of Naval Research, Arlington, VA, which

is gratefully acknowledged.

References

[1]

Ranjan Ray, J. Mat. Sci., 16 (1981) 2924.

[2] Ranjan Ray, Proc. Mat. Res. Conf., Rapidly Solidified Amorphous and Crystalline Alloys, B.H. Kear,

B.C. Giessen and M. Cohen, eds., 8 (1981) 435.

[3] C.C. Wan, Proc. Mat. Res. Conf., Rapidly Solidified Amorphous and Crystalline Alloys, B.H. Kear,

B.C. Giessen and M. Cohen, eds., 8 (1981) 441.

[4]

[5]

[ 6 ]

[7]

S.H. Whang, D.E. Polk and B.C. Giessen, Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai

Japan (1981) 1365.

D.E. Polk, A. Calka and B.C. Giessen, Acta Met., 26 (1978) 1097.

Y.Z. Lu and Sung H. Whang, TMS-AIME Fall 1982, St. Louis, MO, to be published.

H.M. Flower, P.R. Swann and D.R.F. West, Titanium Science and Technology, eds., Jaffee & Burte,

2 (1972) 1143.

290
DIRECTED ENERGY SURFACE PROCESSING OF METALS

T. R. Tucker and A. H. Clauer

Battelle's Columbus Laboratories

Columbus, OH 43201

Abstract

Lasers and particle beam generators have been employed to develop a variety of pro- cesses for improving the surface properties of metals. These processes include laser or electron beam transformation hardening, surface melting, cladding, alloying, ion implantation, and pulsed laser shock hardening. Consistent with the theme of this conference,

only those processes involving melting and solidification are considered here.

All modifications involving fusion of either the base material or a surface alloying

agent rely on the energy beam as a localized and directed source of heat input. The

relative merits of lasers or electron beams depend on the degree of beam control and the

efficiency with which the energy is coupled to the workpiece. For example, the electron

beam can be scanned by pattern generators and efficiently absorbed by most metals to produce controlled surface melts. The introduction of a powder alloy to the process, however,

presents problems for an electron beam but increases the coupling efficiency of a laser

beam. Such trade offs are examined with reference to recently reported results.

The process speeds and melt depths involved in surface melting, cladding, and alloy

mixing, insure that rapid solidification rates are achieved. Initially, much development

was directed towards the production of glassy metallic surface layers by these techniques.

In the last several years, progress has been reported in that regard, but problems asso- ciated with residual stress, thermal stability of microstructures, and thin sections persist. Of more engineering signip'cance now are those processes which involve the more

modest solidification rates (<10^ °K/S). In general, metastable phases and/or refined

microstructures can be expected. However, other important factors influencing properties

are melt convection, epitaxial growth, oxidation tendency, and thermal shock resistance.

Introduction

The improvement of metallic surface properties by directed energy processing methods depend on

the ability of the energy beams to rapidly heat a thin metallic surface layer in a controlled manner.

The rapid process rate allows refined microstructures, enhanced alloy solid solubilities, and re- duced or dispersed inclusions to be produced. Control of the applied energy is necessary to localize

the heat pulse to as thin a surface layer as is practical and thereby reduce bulk property changes.

Thermal input control is also necessary to maintain process consistency.

The two most prominent directed energy sources used in surface materials processing are the

high-power continuous-wave laser and the electron beam. Both sources can be scanned, focused, or

shaped to perform a variety of process functions, although the methods to accomplish these ends differ

in the two cases. To the extent that the two energy sources can be made to provide an equivalent

thermal input, they are interchangeable.

While lasers and electron beam's can perform similar functions, each has advantages and disadvantages in particular applications. The e-beam is a more energy efficient source. A high percentage

of the operating electrical power is produced as e-beam power. Moreover, the electrons are efficiently absorbed ( 80-90%) by most metals. The e-beam is easily focused and scanned over short distances

by electromagnetic lenses and deflectors. On the negative side, e-beams generally require vacuum

operation and cannot be effectively propagated long distances. In addition, e-beams are not suitable

for processes involving powder streams because of the charging effects of isolated particulates in

the beam.

291
TABLE I. LASER AND ELECTRON BEAM SURFACE PROCESSING GEOMETRIES

BEAM MODE PART MOTION E-BEAM LASER APPLICATIONS

Point Fast Rotation

and Linear

X Skin melting, glazing

Defocused

Poi nt 2D Raster X X Cladding, surface melting

Effective Line

Source

Linear X X Melting, cladding, consolidation

Square Patch Stationary or LinearX X Cladding, consolidation, transformation hardening

Lissajous Scan Stationary or LinearX Cladding, consolidation, transformation hardening

By comparison, the laser is not as energy efficient. Operating electrical efficiencies of 5-10%

are common. CO2 laser wavelength absorption efficiencies are typically 10% for most metals. This

disadvantage can be mitigated by use of absorption enhancers or in processes involving particulates.

Offsetting the energy disadvantage for the laser are the following positive factors: operation in

gaseous atmospheres, insensitivity to particles or magnetic fields or electrically insulating sample

regions, long propagation distances, and the ability to be shaped to a uniform power distribution.

Process Geometries

A variety of process schemes for covering surface areas are possible depending on the desired

depth of treatment, substrate geometry, beam characteristics and material additions required. Some of

the more general approaches are outlined in Table 1 and illustrated in Figures 1 through 6. A com- bination of beam shaping and part motion are usually employed. The processes in the table go from

the highest solidification rates at the top to the slower rate processes at the bottom.

The beam can be used at focus with a high speed relative to the sample surface to yield extremely high heating and cooling rates. This method is described in the next section. By applying the

beam at or near focus and scanning the part in a raster fashion (Figure 1), area coverage can be

achieved with overlapping beam passes. The raster coverage involves linear motion in one direction

perpendicular to the beam axis and a slower translation (or stepping motion) in a mutually perpendicular direction. The two motions must be coordinated to result in the desired overlapping factor.

Rotational motion can be substituted for linear translation to cover cylinders, shafts, or spiral

surfaces

A focused beam can be rapidly oscillated to generate an effective line source of energy. The

e-beam is particularly suited to this type of treatment (Figure 2). By sweeping the beam with a highfrequency triangular waveform and combining motion perpendicular to the oscillation direction, surface areas can be covered in a manner analogous to the raster coverage pattern. Rapid oscillations in both

the X and Y directions produce a rectangular patch of energy (Figure 3). This uniform energy pattern

can be scanned across a sample or pulsed onto a specimen for a selected time interval.

Laser beams can also be rapidly scanned to produce line or uniform rectangular thermal inputs.

However, the scanning mechanisms for laser beams are different from that of e-beams. Rotating or vibrating reflective optics are required for laser beams. Orthogonal rotating polygon mirrors

(Figure 4) can be coordinated to yield a small rectangular raster scan pattern. The same coverage can

be obtained with oscillating mirrors, but in this case motions are restricted to harmonic oscillations

The sinusoidal oscillations give rise to Lissajous patterns rather than uniform raster coverages.

292
Figure 1. Area coverage of a flat sample by

raster scanning

Figure 2. Oscillation of a focused e-beam to produce effectively a line source

of energy input

Figure 3. Rapidly scanned e-beam in two dimensions to input a uniform rectangular

energy input

Figure 4. Formation of a uniform rectangular

laser beam profile by use of rotating polygonal mirrors

Focusing Mirror

Oscillating

Flat Mirror

Lissajous Pattern

Orthogonal Oscillation Technique

e»*m InlafMiion Tschniqu*

Figure 5. Formation of a rectangular laser

beam profile by use of oscillating

mi rrors

Figure 6. Formation of spatially averaged

square beam profile by use of

faceted optics

293
The laser beam can also be shaped to produce approximately uniform rectangular energy distributions without beam scanning. The use of segmented optics (Figure 6) provides for overlaying smaller

segments of a single larger high-power beam. The integration of the original beam distribution

averages its non-uniformities and generates a rectangular rather than round shaped cross section.

These differences are important for some of the slower processes of interest such as homogenizing

surface melts, coatings consolidation, and transformation hardening. Moreover, the integrated beam

can be reimaged and manipulated to tailor the input beam distribution to the part dimension and

geometry.

Amorphous Surface Layers

Early development of directed energy surface processing technology was concerned with the production of glassy metallic layers on crystalline substrates. The goal of this development was to derive the strength, ductility, and electrochemical advantages of amorphous metals on the surfaces

of bulk parts to improve their resistance to corrosion, fatigue, and/or wear.

The first use of laser radiation to rapidly surface melt an alloy producing glassy material was reported by Breinan and coworkers in 1976 (1,2). The method, which they termed 'laser glazing',

employed a stationary laser beam that was brought to focus in the plane of the specimen's top surface.

The specimen was fastened to a turntable which translated under the laser beam as the rotary stage

turned. The sample thereby received a spiral melt treatment analogous to the grooves in an audio

recording.

Figure 7 shows the cross section of an individual melt pass in a glass forming alloy (NiCrSiBFe)

glazed by the method described. The fusion zone exhibits the featureless behavior after etching,

which is characteristic of glassy phases. This particular alloy along with the PdCuSi alloy reported

in reference 1 are known to be relatively easy metallic glass formers.

Although the rapid surface melting of separated beam passes readily produces glassy stripes on surfaces of these alloy types, the production of an extended amorphous surface layer by this method

is more difficult. The source of the processing difficulty lies in the overlapping pass geometry as

in Figure 8. The fusion zone layer of this M2 tool steel sample, which was melted by a scanning

e-beam, has no glassy phase present, but the edge of the melt zone of each overlapping beam pass is

clearly visible. When glassy metal formers are skin melted, crystallization occurs along and just

below these overlap zones.

To date, there have been two reports of continuous beam processed surfaces which have retained an amorphous microstructure - the first employing focused laser radiation and a specially prepared microcrystalline alloy coating, and the second using a scanned electron beam 'line' source.

Laser Processed Glassy Layer

Tucker and Ayers reported glassy overlapped laser melt passes on alloy AMS 4777 which had been

prepared as a coating on a mild steel coupon (3). The alloy coating process was chosen so as to preserve alloy chemical composition and produce a fine-grained microstructure (4,5).

For the skin melting process, the applied laser power was high (9 kW) with a rapid rotational

speed (600-900 cm/s). After 30 overlapping laser passes, an uncracked featureless surface layer was produced with an average melt depth of 25 pm. The totally glassy layer appeared to be about 10 pm

deep.

The retention of glassy phase was attributed to the starting coating microstructure, the extremely

rapid process rate, and the properties of the alloy employed. AMS 4777 has a relatively low quench

rate threshold for glassy phase formation, high thermal stability, smooth flow characteristics, and

good ductility in the amorphous state. It is likely that this combination of properties is necessary

for the laser skin melt process to be successful.

E-beam Processed Glassy Surfaces

Bergmann and Mordike (6) modified the scanned line source electron beam surface treatment to produce noncrystalline surface layers. The effective line source beam, which was produced by oscillating

a focused e-beam with a 10 kHz triangular wave form, was scanned at speeds of 20-40 cm/s across samples of FeCrBC and NiNb alloys (7). A beam size of 60 pm was reported.
Figure 7. Cross-section of glassy-metallic

laser surface melt zone in NiCrSiBFe

alloy. (From Tucker & Ayers, Ref. 3)

Figure 8. Overlapping e-beam melt passes on M2 tool steel (from Kurup et.al.,

Ref. 14)

Figure 9. Surface smoothness of a laser clad

layer of NiCr on mild steel substrate

by concurrent vibration of specimen

(from Powell & Steen, Ref. 15)

Figure 11. Microstructural non-uniformity in

laser surface melted MarMOOZ processed at 91 cm/s (from Lipscomb

et.al.. Ref. 17)

Figure 10. Epitaxial solidification on a laser surface melted track in

white cast iron

Figure 12. Overlapping heat-affected zones

in laser melted FeCQ

g alloy

295
By employing the described process conditions with the appropriate beam power, a condition was reached whereby a region could remain molten over the time span of several line oscillations. Thus,

a melt front would proceed along the line at the process speed, and no overlapping solidification

layers need be produced. This method resulted in a beam interaction time over several line oscillations that was nearly as short as that of the laser skin melt process above but without the solidification between passes.

Practical Considerations

While progress has been made in demonstrating the technical feasibility of using directed energy

processing to produce glassy metal surfaces, there are several practical considerations inhibiting

development of applications using this technology. The principal limitation imposed on the process

is alloy suitability. The number of alloys that can be beam treated to a glassy state appears to be

smaller than the set that can be produced by other rapid solidification methods.

Another limitation of glassy coatings is the thickness of the amorphous layer. It is unlikely

that glassy zone thicknesses can exceed 20 ym in practice. This limit is more restrictive than that

of melt spun ribbon for example. It is not possible to build up glassy layer thickness by repetitive

treatments because subsequent tempering effects immediately below each rapid melt pass would crystallize the underlying layer. It is theoretically feasible to build up alternating layers of glassy and

crystalline structures, but not bulk glassy material

Another factor influencing performance of skin melted layers is the residual tensile stress

distribution imposed by the processes. The shrinkage stresses arising from the cooling of the

solidified layer on a cooler substrate can be quite large. Cracking is often a problem with either

crystalline or glass-forming alloys. Thus, the ductility of the alloy in the glassy state is an important property as far as beam processing is concerned. An increased sensitivity to fatigue

damage can also be assumed due to the residual stresses.

Taking these considerations together, it is clear that the materials combinations and useful

applications to result from the beam processing approach would be highly specialized and restrictive.

This is not to say that the development in this area to date has been unsuccessful, but rather that

it will probably not have the general applicability that was initially expected.

Crystalline Surface Layers

Surface melting and surface alloying using directed energy beams have the advantages of providing

control of chemistry, microstructure, and properties at the specimen surface different from that of

the bulk material. Although surface melting alone does not usually change the alloy chemistry, the

microstructural changes can benefit wear, erosion and corrosion properties. These benefits derive

from refinement of the microstructure leading to higher hardness and smaller potential failure sites,

and from the resolutioning and refinement of phases which are deleterious to corrosion or wear

resistance. Many of the results of numerous studies on microstructure and processing are contained in

various conference proceedings (8-11).

The principal advantage of the electron beam over the focused laser for glassy surface melting is

its small spot size. CO2 laser beams have sufficient power, time stability and line scan capability

through the use of rotating polygon optics. Their diffraction-limited focused beam sizes, however,

are simply too large to apply the line scan approach for this application. A laser beam of shorter

wavelength would be more focusable, but such a beam is not practically available in high-power con- tinuous mode at present.

In any of the directed energy beam surface melting and alloying processes there are several major

concerns. These are surface cracking, surface smoothness, porosity, and microstructural uniformity.

The cracking is a result of residual surface tensile stresses created by solidification and thermal

shrinkage on cooling after melting. Crack alignment in general occurs parallel, perpendicular and,

at times, at 45 degrees to the beam traverse direction. The occurrence of cracking will depend on the solidification temperature range and hot strength of the alloy, the temperature differences

developed between the substrate and the melt, the rate of cooling, the degree of plastic strain

developed in the substrate layer, and the presence of any solid state transformations which may occur

in either the cooling melt zone or the underlying heat affected zone. Minimizing the surface cracking

tendency would considerably enhance the viability of surface melting processes as a material improvement process. This problem has received little attention from a basic research standpoint because its

complex nature does not lend itself to simple modeling assumptions.

296
ll Surface smoothness is affected by the ripple effects (12,13) in surface melt tracks similar to

i the ripples observed in weldbeads. In addition to the surface ripples, there may be a trough formed

i{ in the center of the track or an undercut or overflow along the sides of the track. A minimum amount

j of as-treated surface roughness is desired to avoid or minimize any secondary machining of the surface,

i This is particularly important if more rapidly solidified, and therefore thinner, surface layers are

j

desired for applications such as wear resistance which require smooth surfaces.

I The degree of roughness can be addressed by modifying the laser or e-beam power, beam size,

j scanning rate and cover gas within the constraints of accomplishing the desired depth of melt and

! microstructural or property modification. Lower beam powers and more rapid beam scanning rates tend

to enhance the surface smoothness but also develop thinner melt layers. E-beam melting also provides

relatively smooth surfaces under certain conditions (14). Powell and Steen found that vibrating the

substrate during laser surface alloying improved the surface smoothness, reduced porosity in fused

j powder coatings and reduced the cracking tendency (15). An example of a smoother surface profile

I from a surface cladding operation is shown in Figure 9. Similar effects should also apply to simple

j surface melting.

Pores in the melted layer are of less concern than are cracks, but nevertheless are best avoided

if possible. The porosity is usually nominally spherical and therefore can be attributed to entrapped

gas in the melt layer. One cause is residual gas from that originally present in the surface layer,

e.g., in preplaced powder, plasma-sprayed coatings or in solution in the alloy (16). Another cause

is the vaporization of a constituent of the alloy, or reaction of an alloy constituent with the

atmosphere. Porosity attributed to oxidation of the carbon has been observed in laser surface melted

cast irons with inadequately shielded melt regions. Also an increased tendency for porosity with increasing phosphorous content has been reported in laser surface melts of gray cast iron (16).

Uniformity of the microstructure and chemistry is of major importance to obtaining treated

surfaces with acceptable, consistent properties. Variations in microstructure and chemistry are

caused by variations in the temperature gradient, solidification rate, cooling rate, and the convective

and mixing (e-beam) currents in and through the melt layer. There may be significant variations in

the size and morphology of the microstructure and in distribution of the phases through the melt layer. Because of the variation in the temperature gradient and rate of solidification, the solidification

structure often varies through some range of cellular to dendritic characteristics and shows a coarsening of the microstructure and increasing segregation from the bottom to the top of the melt

zone. Another effect is that the presence of the substrate in intimate contact with the melt layer

promotes epitaxial solidification relationships between the resolidified melt layer and the substrate.

This is particularly evident in the laser surface melted white cast iron shown in Figure 10, where

there is a clear correlation between the iron carbide dendrites in the melted layer and the carbide

orientation in the underlying grains.

Another source of non-uniformity of microstructure is caused by the presence of convective

currents in laser melted pools and this plus electromagnetic stirring in electron beam melted pools.

These melt currents create local differences in temperature and therefore solidification conditions,

which often show up as "swirl-like" patterns in metal lographic cross sections of the melt zone. These

effects have been observed in a number of alloys, but primarily in nickel-base superalloys (17,18).

These swirl patterns originate from local variations in growth direction, size of dendrites or cells,

phase distribution and segregation. A laser melt showing very noticeable variations of this type is

shown in Figure 11. Alternating copper- and silver-rich bands have been observed in copper-silver

alloys by Beck et al (19). The magnitude of these variations will depend on the alloy's solidification

characteristics, the complexity of the microstructure, and the melting conditions such as melt depth,

preheat, and super heat. Modification of these conditions could afford some degree of control over

this convective or mixing affect.

Another cause of non-uniformity is the heat affected zone surrounding the melt layer. This not

only affects the material below the melt layer, but with the overlapping passes used to obtain area

coverage, it affects the surface material along the edge of each succeeding pass. A representative

example is shown in Figure 12. The darker lines delineating each succeeding pass define the heat

affected zones intersecting the surface. Although heat affected zones cannot be avoided, their

effects can be minimized by modifying the process parameters, selection of alloy composition, and

post surface melting heat treatment. The extent of the heat affected zone is decreased by decreasing

beam power and interaction times; these conditions will also decrease the melt depth. Alloy compositions less sensitive to aging or tempering during rapid excursions to higher temperatures will

also minimize this effect. Post-processing heat treatment may tend to make the melt layer more

uniform in structure. The presence of a hardened heat affected zone may be desirable where improved

support is useful to support a thin surface melted or alloyed layer against wear or erosion.

297
In surface alloyed melts, in addition to the above effects, there is also the need for uniform

mixing of the surface deposited alloy addition into the melt layer. This mixing occurs by both

diffusion and the convective and electromagnetic (e-beam) currents. It is also influenced by the

melting conditions used. In surface alloying of mild steel with molybdenum an indication of these

effects was observed (20). The surface alloying was accomplished by melting a molybdenum powder

layer into the steel surface. Linder rapid melting conditions, the beam power was high enough that

the molybdenum layer melted first and was mixed with the melted steel below primarily by convective

mixing. This resulted in swirl patterns as shown in Figure 13a which appear to consist of iron rich

melt regions having been carried up into the melt layers. A similar effect was observed for chromium

surface alloyed in steel via an EDAX elemental map of the chromium distribution in the melt layer (21)

The mixing is not homogeneous, but in general appears to be satisfactory for most applications.

A different situation can be developed if the alloy coating has a higher melting temperature than

the substrate. For example at a low enough laser beam intensity a plasma sprayed molybdenum layer on

a steel substrate was heated above the melting temperature of the underlying steel before the

molybdenum melted. Melting then began at the molybdenum-steel interface, favoring diffusion over con- vective mixing of both constituents from the initiation of melting. This resulted in a much more

uniform appearing melt zone. Figure 13b, indicating a more uniform distribution of molybdenum. There

was still not uniformity of composition; point EDAX analysis of molybdenum in the melt layer in

Figure 13b gave top center and bottom analyses of 33.2, 27.2, and 28.7 weight percent molybdenum.

Alloyed surface layers having the highest additive concentrations at the top of the layer have often

been observed (22).

An example of the relative influence of convection and diffusion of alloying elements in a laser

melt is shown in Figure 14. In this case, the beam interaction time, x-j, was 5 X 10"5 seconds, and

the surface melt time was im (2 to 5) xi 2 X lO"^ sec. Using these parameters, the convective

flow distance, Sq can be estimated from

Sc (t^ n/p)^^^ (1

)

where p is the material density and n is the melt viscosity, and the diffusion distance from

(x D)^/^ (2) d ^ m

-3 -4 where D is the liquid diffusion coefficient. These relations give Sc 1 X 10 cm and Sj 1 X 10

cm, or Sc % 10 S^. These distances are consistent with the mixing effects represented by the flow

lines from the interdendritic zones in the left half of Figure 14. The convective flow lines of the

material from the interdendritic regions do extend about 10 to 15 ym. In addition the amount of

diffusional smearing of the existing compositional lamellae near the surface of the melt is consistent

with a nominal 1 ym distance.

Since the effects of directed energy beam processing relate to the in-depth thermal effects in

the material being processed, it is necessary for a thorough understanding of the observed affects

and basic studies of microstructure and property modification to have an analytical or modeling

method to calculate local temperatures and heating and cooling rates. A number of models have been

developed (23-25) and used (26) for determining these quantities, but the agreements with the

observed effects is not yet sufficiently accurate to be of use as an independent analytical technique.

Better methods for including the amount of absorbed energy from the beam, convective flow in the melt,

heats of fusion and temperature dependent thermal properties are required to improve the predictive

capabilities of the modeling approach.

Properties

As more directed beam materials processing work is being accomplished, increasing attention is

being paid to the resulting property changes. The properties of primary interest are. wear, erosion

and corrosion resistance - a direct result of the surface modification aspects of the process.

Several studies have provided encouraging results. Laser surface melting has produced enhanced

corrosion resistance in Mg-Zn (27) Al-Fe (28) alloys. In the material not surface melted, large

intermetal 1 ic particles created galvanic corrosion cells on the material surface. Surface melting

and rapid solidification caused these phases to be dissolved and subsequently maintained in solution

or highly refined in the resolidified surface layer, thereby reducing corrosive attack.

298
Figure 13. Appearance of laser melted steel coupons alloyed with molybdenum and carbon;

(a) conditions favoring convective mixing, (b) conditions favoring diffusive mixing.

Inserts show higher magnification views of alloyed zones.

Figure 15. Improvement in wear resistance of

laser surface treated gray cast

iron (from Bell, Ref. 30)

Figure 14. Mixing patterns in a rapidly lasersurface melted layer of FeMoC

alloy

299
Cavitation erosion resistance of Al-Fe bronze was also increased by laser surface melting (29).

The wear resistance of gray cast iron has been reduced by surface melting (30). In this case the

surface melting develops a more rapidly solidified hard, white cast iron surface on a gray cast iron

substrate. The improvements in wear are shown in Figure 15.

The use of surface melting or alloying in fatigue sensitive applications is limited by the

residual tensile stresses developed during solidification of the surface layer. Laser surface melting

of Ti-6A1-4V (31) significantly decreased its fatigue life. While the rapidly solidified microstructures contributed to this loss of fatigue strength, there is little doubt that the residual

surface tensile stresses did also. The surface melted specimens contained numerous parallel surface

cracks after fatigue testing whereas the unmelted specimens had single fracture origins.

For applications where residual tensile surface stresses will degrade the useful life, as in

cyclic fatigue, fretting fatigue, stress corrosion cracking, hydrogen embrittlement, there may be

benefits to be gained by application of a properly controlled alloy chemistry and microstructure.

In addition to these effects, consideration of how to minimize or avoid the effects of residual

tensile stresses must also be included in the process design and development.

Summary

Progress in the field of laser and e-beam surface processing of metals during the last 2 years

can be summarized in the following conclusions.

In the case of glassy metallic surface layers by rapid surface melting, several methods have

demonstrated success on small extended surface areas. However, the limited number of suitable

alloys, thin amorphous case depths, and high residual tensile stresses limit the applicability of

such glassy layers. These difficulties have decreased the interest in skin-melting processes in

spite of the technical progress.

With respect to the microstructure encountered in rapidly solidified beam processed materials,

metallurgical phases are comparable to those found in specimens prepared by other rapid quenching

processes. That is to say, extended alloy solid solubilities, metastable phases, retention of hightemperature phases to room temperature, and glassy material can be produced. It is significant to note that no phases unique to beam processing have been reported to date.

While the crystallographic phases produced by beam processing are similar to other rapidly

solidified materials, the microstructural arrangement of these phases can be different. For the most

part these differences can be attributed to epitaxial solidification effects and melt convection

inherent in the rapid heating rates employed.

In surface alloying, cladding and coatings consolidation, significant progress has been

demonstrated in the areas of microstructural refinement, alloy mixing and metallurgical bonding of

treated layers. Process development has extended coverage to somewhat larger areas and various

sample geometries. Many of the effects inhibiting this development relate to surface smoothness,

consistency over larger process areas, and thermal stresses.

Finally, the engineering, property data base for beam processed materials needs to be greatly

improved. Data on wear and corrosion of surface melted and alloyed samples have recently appeared in

the literature, but this data is insufficient to draw any general trends or make meaningful com- parisons with competing technologies. Property evaluation should receive the most attention in this

field in near term studies.

Acknowledgement

Portions of this work which were performed by the authors was supported by the Army Research

Office, Research Triangle Park, North Carolina.

300
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22. Gnanamuthu, D. S., Applications of Lasers in Materials Processing, American Society for Metals,

Metals Park (1979), p. 177.

23. Cline, H. E. and Anthony, T. R., J. Appl. Phys., p. 3895 (1977).

24. Sekhar, J. A., Mehrabian, R., and Fraser, H. L., Lasers in Metallurgy , The Metallurgical Society

of AIME, Warrendale, Pennsylvania (1981), p. 207.

301
25. Chande, T. and Mazumder, J., Lasers in Metallurgy , The Metallurgical Society of AIME, Warrendale,

Pennsylvania (1981), p. 165.

26. Tucker, T. R. , Clauer, A. H., Walters, C. T., and Ream, S. L., Lasers in Metallurgy , The

Metallurgical Society of AIME, Warrendale, Pennsylvania (1 981 ) , p. 53.

27. Kattanis, T. Z., Lasers in Metallurgy, The Metallurgical Society of AIME, Warrendale, Pennsylvani

(1981), p. 1.

28. Draper, C. W. , Vandenbury, Preece, C. M., and Clayton, C. R., Rapidly Solidified Amorphous and

Crystalline Alloys , Elsevier Science Publishing Co. (1982), p. 529.

29. Draper, C. W. , Lasers in Metallurgy, The Metallurgical Society of AIME, Warrendale, Pennsylvania

(1981), p. 21.

30. Bell, T., Metallurgia, ^

(3), p. 103 (1982).

31. Bayles, R. A., Meyn, D. A., and Moore, P. G., Lasers in Metallurgy , The Metallurgical Society of

AIME, Warrendale, Pennsylvania (1981), p. 127.

302
INTERFACE RESPONSE IN LASER-MELTED ALLOYS

S. M. Copley, D. Beck and M. Bass

University of Southern California

Los Angeles, CA 90089-0241

ABSTRACT

By scanning the surface of a metal or alloy with a focused CW-CO2 laser beam,

it is possible to melt and solidify it at high rates. In Ag-Cu alloys, we have

shown that sufficient kinetic undercooling can be attained by this method to form

the metastable extended solid solution. For a Ag-38.3 at pet. Cu alloy scanned at 34 cm s“^ to produce the metastable solid solution, thermodynamic considerations

indicate that an undercooling of at least 52°C was attained. Metallographic examination of melt trails revealed the presence of regularly spaced bands appproximately

marking sequential positions of a planar melt-solid interface. These bands correspond to fluctuations in concentration of at least 1.1 at pet. Their presence suggests

that the interface may have an oscillating steady state motion. Implications of the

banding phenomenon will be discussed with respect to mass at thermal transport at the melt solid interface and characteristics of the interface response function.

Introduction

The name rapid solidification processing (RSP) has been applied to solidification processes in

which the cooling rate is greater than 100°Ks“' [i] . In such processes, it has long been recognized

that equilibrium may not exist at the melt-solid interface. Deviations from equilibrium may occur

for two reasons; (1) suppression of nucleation of the equilibrium phase(s) due to rapid cooling;

and (2) kinetic undercooling due to high solidification rates.

We summarize here the results of an investigation of the rapid solidification of Ag-Cu alloys.

Rapid solidification was attained by scanning the surface of Ag-Cu alloy specimens with a continuous

wave, carbon dioxide laser beam. Two aspects of our research will be described: (1) estimation of

melt-solid interface temperature; and (2) observations of variations in solute concentration along

the direction of local interface motion (banding) and speculation as to its origin.

Laser Melting and Solidification

In our experiments, the specimen was mounted to a turntable and rotated under the laser beam.

Single trails resulted from a single rotation. Multiple overlapping trails were produced by multiple

rotations accompanied by a continuous translation of the turntable perpendicular to the laser beam.

The laser beam had a Gaussian spatial distribution of intensity and diameter (l/e^) of 196 ym. A flow of inert gas was maintained over the specimen surface. Absorption of the incident radiation

was enhanced by a graphitic coating. Conduction of heat from the substrate to the turntable was increased by use of a conductive cement.

Viewed in a frame of reference that is stationary with respect to the laser beam, laser melting

and solidification may be considered to be a quasi-steady-state process. Accordingly, the shape of

the melt pool, which is shown schematically in Fig. 1, does not change with time and the local interface velocity is given by the equation

V = Vq cose (1

)

where Vq is the velocity of the specimen with respect to the laser beam and 0 is the angle between

the normal to the melt-solid interface and the specimen velocity. For the melt pool shown in Fig. 2,

the interface velocity in the center plane varies from 0 at the bottom of the pool to 0.95 Vq at the

top.

303
Figure 1 Interface velocity varies with angle e along the back edge of the melt pool.

Figure 2 Melt pool sectioned perpendicular to the direction

of specimen translation. (Marker indicates 7.5 ym)

304
A photomicrograph of a melt pool sectioned perpendicular to the direction of motion is shown in

Fig. 2. The specimen velocity was 34 cm s~' and the incident beam power was 1090W. Both the Cu-rich

phase (dark, dispersed) and the Ag-rich phase (light, continuous) can be seen in the unmelted region.

It was not possible to dissolve the Cu-rich phase at a beam velocity of 34 cm s"‘. After laser melting and solidification, the composition in the melt trail was inhomogeneous as evidenced by decoration

of the convection swirls on the left hand side of the melt pool. In order to attain a uniform composition in melt trails formed at 34 cm s" , a laser premelting step involving scanning the surface

10 cm s"' using the same beam power was employed. This resulted in a uniform composition as can be seen at the right hand side of the melt pool, which overlapped a trail premelted at 10 cm s" ‘

It is clear by examination of Fig. 2 that the melt is in contact with both the Ag-rich and Cu-rich

equilibrium phases during solidification. Thus, in laser melting and solidification, deviations from

equilibrium must be the result of kinetic undercooling due to the high interface velocity rather than

the suppression of nucleation due to the high cooling rate.

Melt-Solid Interface Temperature

The Ag-Cu equilibrium diagram shows two equilibrium phases, ct (Ag-rich) and 3 (Cu-rich), which

form from the melt by a eutectic reaction at 1052 °k [2]. Both Ag and Cu are face centered cubic and

years ago Duwez, Willens and Klement showed that metastable extended solid solutions could be formed

across the a-3 phase field by "splat cooling" [3]. In laser melting and solidification experiments,

it is very difficult to measure directly the melt solid interface temperature. We have taken an indirect approach to determining the interface temperature which yields an upper bound. In our approach we determine Tq(C), the maximum temperature at which it is thermodynamically possible to form the metastable extended solid solution from a melt of the same composition (C). We examine melt

trails produced at different powers and velocities employing x-ray diffraction. If the metastable

extended solid solution is detected, the melt-solid interface temperature must have been less than

To(C).

In order to determine Tq(C), we must know the Gibbs free energy of mixing (AG^ix) a function

of composition for both the solid and liquid phases. We have calculated aG^^^ for the solid phase

employing the Lumsden solution model [4,5], and aGL.| for the liquid phase, extrapolating thermochemcal measurements made at higher temperatures. Details of the calculations are published elsewhere [6].

Figure 3 is a plot of aG^^x for the solid and liquid phases calculated at 1003°K. Two points

on the Tp(C) curve (see Fig. 4; corresponding to 1003°K are established by noting the compositions

where aGL^-^ = aG^- . For C < 0.381 a/o Cu or C > 0.676 a/o Cu, the Gibbs free energy decreases if

liquid is transformed to solid of the same composition and thus such a transformation is possible.

Points on the metastable solidi and liquid! curves at 1003°K are determined by the common tangent

rule.

Figure 4 shows a phase diagram for the Ag-Cu system. The equilibrium diagram is based on that

of Hansen and Anderko [2]. The Tq curve and metastable solidi and liquid! curves were calculated

previously by us [6]. Triangular and diamond shaped points on the diagram are used to indicate

compositions where laser melting and solidification experiments have been carried out with the objective of producing and detecting the metastable phase.

Initial experiments involved scanning electron microscopic examination of trails laser melted

at a scan speed of 0.7 cm s”^ in Ag electroplated with Cu. Primary Ag-rich dendrites and a very

fine plate-like eutectic microstructure were observed throughout the melt zone [7]. Further experiments were conducted at a scan speed of 10 cm s~^ on alloy compositions indicated by the triangles

in Fig. 4. The metastable extended solid solution was detected by x-ray in single trails for all

three alloys [8]. Thus, the maximum possible interface temperatures were 795, 700 and 779°K for the

Ag-25 at pet Cu, Ag-50 at pet Cu and Ag-75 at pet Cu alloys, respectively.

Still further experiments were carried out employing a scan speed of 34 cm s~^ with alloy compositions indicated by the diamonds in Fig. 4. In these experiments, the surface was covered by multiple overlapping trails. The metastable extended solid solution was detected by x-ray in the Ag-18

at pet Cu, Ag-28.2 at pet Cu and Ag-38.3 at pet Cu alloys. [6] Thus the maximum possible interface

temperatures were 838,771 and 727°C for these alloys, respectively. No metastable solid solution was detected at the compositions indicated by the open diamonds. It is believed that the metastable

extended solid solution was probably present initially in the trails but decomposed to a and 3 due to additional heating associated with the multiple overlapping procedure.

Figure 5 shows the {111} reflections from an x-ray scan of the laser melted and solidified layer

for the Ag-38.3 at pet Cu alloy (eutectic composition). Four peaks were detected in the as-irradiated sample (0.0 sec) corresponding to the Ag-rich equilibrium solid solution (a), a Ag-rich meta305
Fig. 3 AG[|,-jx versus composition curves for solid and liquid

phases at 1003°K illustrating determination of points

for the iQ-curve, metastable solidi curves and metastable

liquid! curves.

Ag ATOMIC PERCENT Cu Cu

Fig. 4 Ag-Cu phase diagram showing Tp-curve, metastable solidi curves and metastable curves. Diamonds and triangles

indicate compositions investigated.

306
Intensity

Fig. 5 The {111} diffracted intensity for an x-ray scan of a laser melted and solidified layer on a Ag-38.3

at pet Cu alloy.

Fig. 6 Schematic drawing showing proposed structure of

laser melted trail in Ag-38.3 at pet Cu alloy.

307
stable solid solution designated a', the metastable extended solid solution having the same composition as the alloy (y) and the Cu-rich equilibrium solid solution (b). Also shown in Fig. 5 is the

effect of annealing at 420°C for various periods of time on diffracted intensity. As would be expected,

the intensity of the Y-peak decreased while that of the a-peak increased with increasing annealing

time.

The x-ray response of the Ag-38.3 Cu alloy can be understood if the upper region of the trails

solidified as the metastable extended solid solution while the lower region solidified with the two- phase eutectic microstructure as shown schematically in Fig. 6. This arrangement phases is likely

because the lower region solidified with a lower interface velocity and hence a higher interface

temperature than the upper region. Such an arrangement should produce a strong Y-peak and broad

a'-peak for the Ag-rich eutectic phase corresponding roughly to the metastable solidus. The a and

B peaks were diffracted from regions where the metastable extended solid solution decomposed to form

the equilibrium phases. No separate B' peak corresponding to the Cu-rich eutectic phase was detected

because it was masked by the B peak owing to the steepness of the metastable B solidus.

The observed lattice parameters for a, a', B and y phases in the Ag-38.3 at pet Cu alloy are shown in parentheses in Fig. 7. The arrows labeled %ax indicate the lattice parameters of S terminal solid solutions corresponding to the maximum solid solubility limit. The values of the lat-

| tice parameters observed for the a and B phase indicate that they were formed by solid state precipi- ; tation. The arrow for the y phase is based on the work of Duwez et al . [3]. The lattice parameters

of the Y phases in the Ag-38.3 at pet Cu alloy and the other alloys investigated are in good agreement i with those predicted from Duwez et al.'s work. The lattice parameter of the a' phase is significantly

smaller than that predicted from the metastable solidus indicating a more Cu-rich composition. This

observation suggests that solute trapping may occur at the melt-a' interface during solidification.

0<

(TU

h—

UJ<

(T

<CL

UJU

H

I-<

Fig. 7 Comparison of observed lattice parameters to those of Duwez et al
Solute Banding

We have previously reported the occurrence of variations in solute concentration along the direc

tion of local interface motion in laser melted and solidified trails [ 9 ]. These variations produce

solute bands that approximately mark sequential positions of the melt-solid interface during solidifi

cation. Figure 8 shows such bands revealed by etching in a Ag-50 at pet Cu alloy. Statistical

analysis of the results of an energy dispersive x-ray analysis of the compositions of the light and

dark bands is given in Table 1 and showed that to the 99.9 confidence level they were different.

The average atomic fraction of Ag was found to be at least 1.1 at pet greater in the light bands than

the dark bands. Actually the difference was probably greater because the sampling volume was approxi

mately 1 urn in diameter due to the depth of penetration, while the bands being studied are less than

1 ym in width.

Fig. 8 Heavily etched bands - Cu 50 at pet Ag.

Table 1. Concentrations Found For Points in Figure 8 Aq Concentration

Pair Number

(Points) Light Band Dark Band

1 ( 1 & 2 ) 80.3 78.6

2 (3&4) 71.0 70.1

3 (5&6) 71.8 70.1

4 (7&8) 71.6 70.2

5 (9&10) 69.7 69.3

6 ( 11 & 12 ) 70.0 70.0

7 (13&14) 71.2 69.9

8 (15&16) 71 .7 69.9

9 (17&18) 69.9 69.2

10 (19&20)

(off picture)

79.4 78.1

We have proposed that the solute bands observed in laser melted and solidified trails are evidence for planar, oscillating, steady-state interface motion. Further, we have suggested that such

motion might be explained on the basis of heat and mass transport at an undercooled planar interface,

if sufficient information regarding the interface response functions were known [ 6 ]. Although,

neither kinetic theories nor experimental data are available which define these functions, thermodynamics places some limitations on them. These limitations can be conveniently described by a diagram developed by Baker and Cahn [lo]

Figure 9 shows a Baker-Cahn type diagram, we have calculated for the Ag-Cu system. The range

of solid compositions, C^, that can form with a net decrease in free energy from a liquid of composition, Cl , at a specific interface temperature lie within or on the closed loop corresponding to that

temperature. For steady-state solidification, C

3 must equal the alloy composition, Cq.

309
Fig. 9 The range of compositions that can form with a net decrease in free energy from a liquid of composition

Cl . For steady-state solidification must equal

the alloy composition Cq.

Consider, for example, steady state solidification of an Ag-38.3 at pet Cu alloy (eutectic composition). The diagram. Fig. 9, indicates three cases:

(i) For interface temperatures greater than 1003°K, e.g. 1052°K, it is thermodynamically impossible

for a metastable solid solution with the eutectic composition to form from a melt with any

composition.

(ii) For an interface temperature equal to 1003°K, a metastable solid solution with the eutectic

composition can form but only from a melt with the same composition.

(iii) For interface temperatures less than 1003°K, e.g. 983°K, it is possible for a metastable solid

solution with the eutectic composition to form from melts with compositions ranging from Ag-24

at pet Cu to Ag-61 at pet Cu.

In case (iii), kinetics dictates which value of C|_ lying within the allowed limits will occur.

With the appropriate interface response functions, it may be possible to find oscillating steadystate solutions for the coupled heat and mass transfer equations at the moving interface. We speculate that if such solutions exist, it might be possible to determine experimentally interface response

functions through measurement of the spacing and concentration variation of solute bands. An oscillatory solution, which can be shown as a closed curve, see arrows in Fig. 9, may under certain

circumstances lead to cellular interface breakdown through a diffusional instability described by

Baker and Cahn [lO] . For a melt composition approaching the limiting Cu-rich composition, a downward

fluctuation in solid composition may cause the melt composition to move outside the range where steady

state solidification is possible. In this case, the plane front interface would break down and lag

back to a lower temperature. Evidence for such breakdown has been observed in laser melted and

solidified eutectic melt trails (see Ref. 6, Fig. 12).

Acknowl edgment

This work was supported in part by the National Science Foundation, Metallurgy Program, Grant

No. DMR 78-07532.

310
References

[1] M. Cohen, B.H. Kear and R. Mehrabian, "Rapid Solidification Processing - An Outlook" in Rapid

Solidification Processing , edited by R. Mehrabian, B.H. Kear and M. Cohen. Claitor's Publishing Division, Baton Rouge, LA (1980).

[2] M. Hansen and K. Anderko, Constitution of Binary Alloys, McGraw Hill Book Co., New York, NY,

1958, p. 18.

[3] P. Duwez, R. H. Willens, and W. Klement, Jr., J. Appl . Phys., 1960, vol . 31, pp. 1136-37.

[4] John Lumsden, Thermodynamics of Alloys , The Institute of Metals, London, 1952, p. 335.

[5] B. E. Sundquist, Trans. TMS-AIME, 1966, vol. 236, pp. 1111-22.

[6] D. G. Beck, S. M. Copley and M. Bass, Met. Trans. 1 3A , 1879-89 (1982).

[7] S. M. Copley, M. Bass, E. W. Van Stryland, D. G. Beck, and 0. Esquivel, "Microstructures of

Surface Alloyed Ag-Cu Films Produced by Laser Melt Quenching," Proc. Ill Int. Conf. Rapidly

Quenched Metals , vol. 1, pp. 147-50, The Metals Society, London, 1978.

[8] D. G. Beck, S. M. Copley, and M. Bass, "Constitution and Microstructure of Ag-Cu Alloys Produced

by Continuous Laser Melt Quenching," Laser and Electron Beam Processing of Materials , C. W. White

and P. S. Peercy, Eds., pp. 734-39, Academic Press, NY, 1980.

[9] D. G. Beck, S. M. Copley and M. Bass, Met. Trans. 1 2A , 1687-91 (1981).

[10]

James C. Baker and John W. Cahn, Sol idification , ASM, Metals Park, OH, 1971, pp. 23-58.

311
INFLUENCE OF ELECTRON BEAM SCANNING PARAMETERS ON SURFACE

MELTING AND HEAT TREATMENT CHARACTERISTICS

P. R. Strutt, M. Kurup and A. Taquir

Metallurgy Department, University of Connecticut,

Storrs, CT, 06268

]. ^Introduction

The unique feature of directed energy beam materials processing is the ability to form fully

consolidated and perfectly bonded rapidly solidified layers on a surface. The actual energy source

may produce electrons, ions, or photons, for instance, and in each case the particles may be focussed

into a concentrated beam. This beam may then be rapidly deflected across a substrate surface by

optical or electromagnetic means, thereby forming a rapidly solidified layer. Other rapid solidification processing methods, in contrast, cannot directly form such a layer, since the as-formed

material is either produced as a powder or ribbon.

The particular advantages of using an electron beam as an energy source are (i) the high efficiency of energy coupling between the beam and material [1], and (ii) the facility with which the

beam may be deflected electromagnetically [2]. Recent studies of "electron beam glazed" hard

iron-base materials shows there to be a considerable enhancement in:

(i) microhardness [3]

(ii) wear resistance [4 and 5]

(iii) resistance to chemical attack [5].

These specific features show the technological significance of the technique in increasing the

performance of components used in cutting and machining and in applications requiring high wear

resistance. A particular application would be in the development of mainshaft bearings used in

high performance military engines capable of operation at speeds up to 3 million DN [6]; D is the diameter in millimeters and N the number of revolutions per minute. Presently there is a need for an inner-race with a surface that will withstand excessive bearing and wear forces,

larger Hertzian stresses, and corrosion pitting [7]. However, the body of the race must possess

good toughness and notch fatigue properties. Obviously, electron beam surface melting provides a method for producing the required dual properties. Using the technique it would thus be possible

to develop an inner race with a surface possessing good rolling contact fatigue [6] and good

corrosion resistance and a body combining moderate strength with good toughness and notch fatigue

properties [6].

Naturally, the sophisticated application of "electron beam glazing" just discussed requires a precise knowledge of factors determining the homogeneity, morphological characteristics, tempering

characteristics, and surface topography of wide surface melted traces. The present paper summaries

recent studies of these aspects and considers specific aspects in detail, such as, the circumstances

in which a scanning beam either acts as a "linear heat source" or "moving point source".

2. Surface Area Melting

An electron beam may rapidly scan across a surface in a variety of ways either by two-axis

electromagnetic deflection [2] or by single-axis electromagnetic deflection combined with a mechanical stage translation [8-11]. In essence however, all types of scanning may be characterized as being either of the (i) moving point source or (ii) linear heat source type, these two modes are now considered in some detail. A separate section on two-axis electromagnetic deflection is included, this, in essence, is an alternative form of point source melting.

312
2.1. Moving Point Source Melting

We envisage in Fig. 1 that a strip of rapid solidified material is formed by an oscillating

electron incident upon a moving substrate surface. Constant velocity of the beam as it moves to and fro requires that the field in the electromagnetic deflection yoke responds linearly to an applied triangular waveform excitation voltage. This is achieved by employing a low inductance

coil (in the microhenny range) driven by ultra-high rate response current amplifiers. Presently,

a high degree of linearity is attainable at frequencies < 15 kH using a triangular waveform excitation signal. Under these conditions a uniform depth of melting is obtained with no detectable

dwelling of the beam at the amplitude maxima.

As the situation is depicted in Fig. 1 the movement of the specimen translation stage per cycle

of beam oscillation is nearly equal to the beam diameter. Under this condition the entire strip

of glazed material is formed with minimal overlap of successfully melted traces. Generally,

however, any degree of overlap may be obtained simply by selecting appropriate values of oscillation frequency (f) and specimen stage velocity (V).

Figure 2 shows an actual processed surface formed by a scanning point heat source where the

familiar surface ripple structure is well evident along the individual traces. The material in

Fig. 1 is a high carbon (1.5 wt.%), low alloy (1.5 wt.% Cr) steel whose composition is given in

Table I. In this particular instance the beam oscillation frequency is low (62.5 hertz) and beam

has a power of 170 watts and a focussed diameter of 0.035 cm. The value of w, the width of the

rapidly solidified strip, is '^-2 mm; V is 1.7 cm.s"^.

The successively overlapping zones of the strip of material seen in Fig. 2 are also shown in

the corresponding longitudinal section in Fig. 3. The striking feature in Fig. 3 is the homogeneous appearance of the "as-glazed" martensitic structure throughout the entire surface melted

region with no evidence of microstructural change where successive zones overlap. The homogeneous

nature of the martensitic structure is further confirmed by careful examination at high magnification in the scanning electron microstructure, a typical field of view is shown in Fig. 4. In- terestingly, this microstructure is similar in type to that formed in the same material by thermomechanical processing followed by cryogenic quenching [12]. The fact that the microstructure can

be produced by a short duration-single step process is thus rather intriguing both from a metallurgical viewpoint and in terms of electron beam processing technology.

The absence of any change of the martensitic structure in the zone overlap regions indicates

that the ambient temperature in a zone is above that for the y ^ a transition when the adjacent

zone solidifies. It has been demonstrated that a significant delay time between the formation

of adjacent zones results in appreciable microstructural modification where the zones overlap.

These findings are serendipitous in that they show that an undesirable microstructural feature

may readily be avoided when forming rapidly solidified surface layers.

2.2. Two-axis Electromagnetic Deflections

Surface scanning studies have shown that particularly smooth surface may be formed using a scanning pattern obtained by applying two triangular waveform signals of nearly equal frequency

to the X and y axes of the beam deflection yoke. This mode has been used previously for generalized surface heating [2]. The succession of rectangular Lissajous figures traced out in this

scanning mode are shown in Fig. 5 where x and y axis deflections of the spot on the substrate have

maximum values of a and b respectively. The rectangular path results from the progressively

increasing amount of phase lag between the two frequencies. As the beam sweeps diagonally across

from point A (where the phase lag is zero) the path intercepts the line BC when the x axis

deflection amplitude attains it maximum value: the intersection point is at a distance

(bAf/f)/(l + Af/f) below point C. After this the path next intercepts line CD when the y axis

amplitude attains it maximum value. The distance of this intercept point to the left-hand side

of point C is a Af/f. In proceeding in this way it is readily apparent how the successive rectangular type of scanning pattern is obtained. A simple consideration shows that the angle 0 in Fig. 5 is given by

tan 0 b 1 a 1 + Af/f ( 1 )

Also by geometrical construction the spacing d between two adjacent diagonal traces is given by

313
2abAf/f

{a^O + Af/f)^ +

( 2 )

It should be noted at this juncture that no mention has been made of the actual beam spot diameter

w. When considering a beam of finite width the diagonal lines in the pattern will represent the

path traced out by the center of the spot.

In considering the nth cycle, the beam path intercept on line BC is a distance y away from

point C:

^ ^ (2n - l)(Af/f)b

^n 1 + Af/f (3)

Equation (3) can be used to find n, the number of cycle required to sweep out the entire area ABCD,

simply by noting that y will then be equal (or almost equal) to b. The precise condition is given

by

"

n>-^'^>n-l (4)

It should be noted that, although the time to complete a frame is n/f, the line spacing is frequency independent and depends on the ratio Af/f.

Typical micrographs showing sequential changes in topography in two-axis electromagnetical ly

scanned surfaces are shown in Figs. 6 and 7, in this case the substrate material is M2 high speed

steel, see Table I for the composition. Low magnification micrographs of regions glazed for a frequency of 60 Hz are shown in Fig. 6 where Fig. 6(a) shows part of successive glazed traces

obtained for Af/f = 1/150. With this value of Af/f the number of cycles per frame is 75 and the

surface finish obtained after a complete frame is shown in Fig. 6(b). As is evident from eqn. (2),

the effect of decreasing the ratio Af/f is to reduce the spacing distance d and as a result to produce overlapping of melting zones when w > d where w is the beam diameter. An intriguing finding is the remarkable increase in surface uniformity as Af/f is progressively decreased. This

became strikingly evident when Figs. 6(b), 6(c) and 6(d) were compared, for which the corresponding

values of Af/f are 1/150, 1/450 and 1/600; in each case the value of f is 60 Hz. The experimental

and calculated values of d are given in Table I for results obtained both at 30 Hz and at 60 Hz. Figures 7(a), 7(b) and 7(c) show the surface topographies of the specimens corresponding to Figs. 6(b), 6(c) and 6(d) at a higher magnification.

The overlapping of glazed zones is clearly seen in cross-sectional view, as shown in Fig. 6. It seems reasonable that the rather dramatic increase in surface uniformity achieved by repetitive

melting and solidification arises from the dispersion or evaporation of surface contaminants.

This cleansing promotes progressi vely better melting of the solid surface by molten liquid during

each successive cycle. By examination of the micrographs it is evident that it is possible to eliminate the "ripple structure" and to obtain a surface with minimal periodic topographical

variation.

2.3 Line Source Surface Melting

As previously discussed microstructural change due to transient heat flow from adjacent passes

may be avoided by ensuring that the temperature in a zone exceeds that for the y ^ a transition

when the neighboring zone solidifies. Another approach, however, is to wave a linear heat source

across the entire width of the specimen, see Fig. 9. Such a source is conveniently produced by

rapidly oscillating an electron beam and by slowly moving the substrate material. The analysis

to be presented examines the range of parametric values required to area glaze in this particular

mode; the parameters include power density, oscillation frequency, and translation velocity.

The requirement for a rapidly oscillating beam to act as a linear melting source is that in

any one cycle the material at one extremity remains molten until the beam returns on completion

of the cycle. In order to express this condition analytically we use the Greenwald heat transfer

analysis [12]. On the basis of this particularly simple analysis the peak temperature attained

at point on the surface, such as the reference extremum previously mentioned, is given by

equation 5.

314
T_ (5) = abT^/^ + T,

where Tj^ is the residual temperature due to the undissipated heat from previous cycles. Here we neglect^the ambient chamber temperature Tq since it is relatively small. The interaction time

is T and the parameters a and b are given by

a = b = q /, ( 6 )

In this equation a. and k are the thermal diffusivity and thermal conductivity and qo the incident

beam power density (in watts per cm^). If we assume that the actual spot oscillates in a perfectly

linear manner across the surface at a frequency of f then x is equal to d/2wf, where d is the spot

diameter and w the width of oscillation as measured on the specimen surface. In substituting

d/2wf for T in equation 5 we obtain equation 7 for the temperature difference (Tp -

(Tp - J^) = ab(d/2w)^/2 . (l/f)^/2 (7)

The residual temperature Tp at the extremum may be determined by applying N number of iterations

to an expression in ref. [13]. The result of this is given by equation 8.

T^ = z'J'ab{(n/f)^/^ - (n/f - d/2wf)^/^}; (8)

Equations 7 and 8 express the irradiation temperature (Tp - Tp) and the residual temperature Tp

in terms of parameters, which include the number of cycles N; their ratio is given by equation 9.

V(T„ - T„)

E-|n 1/2 d/2w) 1/2

(d/2w) 1/2 (9)

Now, for material at the reference position to remain molten when the beam returns there after a cyclic excursion the value of Tp must equal the melting point temperature It is now convenient to define the ratio Tp/(Tp - Tp) by the parameter C, where Tp is equal to T^ ('\^1500°C for

iron). The condition that the vaporization temperature not be exceeded (Ty '^^2900°C) requires that

5

> 1 , an alternatively that the irradiation temperature (Tp - Tp) be less^than T^. This now gives

equation 10.

(n - d/2w)^'^^ = C-(d/2w)^^^ (10)

The solution to the problem, therefore, is to determine the value of N satisfying equation 10.

This value is, in effect, the number of cycles required to maintain a molten state on the substrate

surface along the line of the oscillating incident beam. We may thus regard the quantity N/f the

interaction time to produce a molten strip of material of width w. In envisaging the progressive

movement of this strip across the entire substrate as a quasi-continuous process the translation

velocity (V) is therefore equal to (fd)/N, see equation 11.

V = (fd)/N (11)

The requisite parameters to produce a continuous rapidly solified strip across a substrate

surface of given width (w) with a beam diameter d are found by first iteratively solving equation

6 to find the value of N. The inverse of this is (from equation 11) the normalized translation

velocity (V/fd) and Fig. 10 shows this parameter as a function of normalized width (w/d) for

values of Tp of 1250, 1500, and 1750°C; these temperatures correspond to values of 5 of 1.40,

1.00, and 0.71.

Once values of frequency (f) and beam diameter (d) have been selected to find the velocity

V we then determine the value of the parameter b and hence q^, the power density from equations

5 and 6. For convenience this equation may be rewritten to give beam power P as a function of

beam diameter, oscillated beam width and applied frequency, see equation 12,

P = (d/2)3/2(wf)’/2; (12)

315
For an example we consider the required beam power and velocity V to produce a 1.25 cm wide

track, it is assumed that d is 0.035 cm. In this particular case (w/d) is 35.71, calculation and

graphical solution (Fig. 10) shows V/fd) to be 0.01190, with f = 1,000 hertz the value of V is 0.42 cm.s-T. Solution of equation 12 shows that P is 554 W.

Actual examples of wide zones formed in M2 high speed steel are shown in Fig. 11 where the

width is 5-7 mm and the frequency 1000 hertz; in each case the beam diameter and power is 0.035 cm and 563 watts respectively. The shallow zone (Fig. 11a) is formed with a translation velocity V of 4.35 cm.s"^ and is seen to have a reasonable constant depth with a mean value of 62 urn, see Table II. The effect of reducing the translation velocity V to 1.75 then to 0.44 cm.s"^ is to significantly increase the melt depth to 510 ym & 1,250 ym. The deep zones in Figs, lib and c, nevertheless, are uniform and symmetrical in shape. The dimensions of the five width zones discussed are given in Table II, these include those shown in Fig. 11; the process parameters are also given. The parameter P* is the threshold power to initiate the formation of a molten zone

as calculated by the preceding analysis. For each value of P* there is a corresponding calculated

value of the peak temperature. It would seem reasonable to attribute a value exceeding 2900°C

to penetration of the beam (or cavity formation).

In the case of the particularly wide zone (3.2 cm) shown in Table II a threshold power of

2292 watts is required to initiate wide molten zone formation at a beam oscillation frequency of

1000 hertz. Interestingly, a significant decrease in the peak temperature is obtained by increasing the frequency by an order of magnitude without any significant power change, i.e. 2106

as compared with 2290 watts. The physical consequence of the value of Tp below that for vaporization shows that "cavitation" can be avoided by applying a sufficiently high beam oscillation

frequency.

3. Fluid Flow Effects

3.1 Compositional Homogenization

Previous studies of laser [14-17] and electron beam glazing [4] have shown that significant

segregation of alloying elements may result from the connective flow within the melt zone. The case of the high speed steel M2 is particularly interesting, since the flow induced segregation

is frequently sufficient to produce regions of differing phase upon subsequent solidification [17].

As a consequence a pattern of light etching austenite and dark etching ferrite within a crosssectional view is clearly indicative of the familiar connective flow field [14]. This type of effect

is not observed in similar alloys with carbon contents exceeding that corresponding to the

peritectic composition. Nevertheless, the effect of connective flow induced segregation is still

clearly discernible in these materials.

In the studies of singly melted zones in molybdenum-base tool steels the degree of compositional variation increased with increasing beam velocity. Interestingly, a similar effect is

discernible in the wide zones in M7 high speed steel in Fig. 8. The "mottled" appearance of the zone in Fig. 8b as revealed upon etching shows a more complicated compositional variation than in

the case of a single-pass "point source" melted zone. Nevertheless, the dark etching "swirls" are clearly indicative of a connective flow pattern. Such an effect is suggestive of the pattern of

"connective cells" considered originally by Lord Rayleigh and later by Landau in the "rotons" in

superfluidity theory, see Normand, Pomeau, and Velarde [18].

3.2 Surface Topography

In contrast to the familiar "ripple" pattern in Fig. 2, and the smooth surface effect produced

by successive remelting [2], yet another topography in produced by linear heat source melting.

These surfaces appear to be particularly smooth and have a "matt" appearance. Microscopic studies

of such surfaces on M2 tool steel reveal a particularly intriguing topography consisting of a fairly uniformly distributed small "hillocks". In such an as-melted in Fig. 12 the "hillocks"

are smooth in appearance and within the "valleys" there appears to be a fine filamentary eutectictype structure. The effect of progressively polishing and etching the surface is revealed in

Figs. 13a-d. In these micrographs the typical cellular structure is initially evident within the

"hillocks" and eventually throughout as the "hillocks" are polished away. The filamentary structure

quite clearly exists at, and very near to the surface. Systematic experiments now in progress

reveal that the spacing between the "hillocks" decreases with increasing translation stage velocity

V. An interesting explanation for the modulated "hillock" type appearance is that it results from

fluid flow created by "connective cells" of the type just discussed [18]. A rigorous mathematical

316
physics analysis of this type of phenomenon reveals a result which is contrary to what might be

intuitively anticipated [18]. Namely, that the valleys correspond to fluid "up-flow" from within

the interior of the liquid body and the "hillocks" to fluid "down-flow". This phenomenon is part

of a continuing study which is systematically considering the effects of the various processing

parameters.

Acknowl edqement

Appreciation is expressed to the Office of Naval Research for support of the current project

under Contract N00014-78-C-0580 and to Dr. Bruce MacDonald for his interest and encouragement in

the program.

References

[ 1 ]

[ 2 ]

[3]

[4]

[5]

[ 6 ]

[7]

[ 8 ]

[9]

[ 10 ]

[ 11 ]

[ 12 ]

[13]

[14]

[15]

[16]

[17]

[18]

P. R. Strutt, Mat. Sci. Eng., ^ (1980), 239.

P. R. Strutt, Mat. Sci. Eng., ^ (1981), 87.

Y. W. Kim, P. R. Strutt and H. Nowotny, Met. Trans., lOA (1979), 881.

B. G. Lewis and P. R. Strutt, Journal of Metals, Nov., 1982, p. 37.

A. W. Ruff and L. K. Ives, Wear, 7^ (1982), 285.

B. L. Averbach, Metal Progress, Dec., 1980, p. 19.

T. V. Philip, Metal Progress, Feb., 1980, p. 52.

W. Gruhl , B. Grzemba, G. Ibe and W. Hiller, Aluminum, ^ (1977), 1 77.

R. Mayer, W. Dietrich and D. Sundermeyer, Weld. J. (Miami), ^ (1977), 35.

J. D. Ayers, T. R. Tucker and R. J. Schaefer, Proc. Second Int. Conf. on Rapid

Solidification Processing, Claitor's Publishing Division, Baton Rouge, LA, 1980, p. 212.

T. R. Tucker, J. D. Ayers and R. J. Schaefer, "Laser and Electron Beam Processing of

Materials", Academic Press, New York, 1980, p. 760.

H. Sunada, J. Wadsworth, J. Lin and 0. D. Sherby, J. Mat. Sci. Eng., ^ (1979), 35.

L. E. Greenwald, E. M. Breinan and B. H. Kear, Laser-Solid Interactions and Laser

Processing, Am. Inst, of Physics, NY, 1978, p. 189.

S. M. Copley, D. Beck, 0. Esquivel and M. Bass, Laser-Solid Interactions and Laser

Processing, Am. Inst. Physics, NY, 1979, p. 161.

R. Mehrabian, S. Kou, S. C. Hsu and A. Munitz, Laser- Solid Interactions and Laser

Processing, Am. Inst, of Physics, NY, 1978, p. 129.

R. H. Kear, J. W. Mayer, J. M. Poate and P. R. Strutt, Metallurgical Treatises, AIME,

1981, p. 321.

P. R. Strutt, B. G. Lewis and M. Kurup, Electron-Beam Laser Glazing of Iron-Base

Materials, Annual Process Report 1981, ONR Contract N00014-78-C-0580.

C. Normand, Y. Pomeau and M. G. Velarde, Rev. of Modern Phys., ^(3) (1977), 581.

317
TABLE I

Major Alloyins Elements (wt%) in Hard Iron-Base Materials

C Cr W V Mo Co Ti Ni Mn Si Fe

Ferro-TiC(SK) 5.0 3.5 - - 3.0 3.0 21.0 0.5 - - 64.0

M2 0.84 4.5 6.3 2.0 5.0 0.12 _ _ 0.30 0.30 Bal M42 1.09 3.7 1.4 1.2 9.6 7.9 0.24 0.40 Bal M7 1.01 3.6 1.7 1.8 8.3 - 0.29 0.29 Bal

52150 1.48 1.34 _ 0.01 0.02 _ 0.02 0.56 0.07 Bal

P P* f

TABLE II

w ^

1 ( cm . s )

d Tp

(watts) (chO) (hertz) (mm.

)

(pm) i°h

563 310 1 ,000 6.8 0.44 1 ,240 2636

563 405 1 ,000 6.1 0.88 510 3069

563 611 1 ,000 6.2 1.75 220 3848

563 848 1 ,000 5.0 4.35 62 5143

3000 2292 1 ,000 32.0 1 .00 250 5379

2106 10,000 32.0 1.00 2627

318
Fig. 2 . Actual view of the surface of 52150

bearing steel formed by an oscillat- —L- -jpg beam (f=62.5 hertz).

Schematic diagram showing a wide

strip of rapidly sol idified material

being formed by an oscillating beam.

p-jq, 4 . Detailed microstructure in a section

Fig. 3 . Longitudinal section of the surface shown in Fig. 3. layer melted by an oscillating point

source in 52150 bearing steel, note

the absence of an "overlapping

effect", frequency is 62.5 hertz.

319
Fig. 5 . Schematic diagram showing the x-y

deflection pattern formed by two nearly equal frequency triangular

waveforms

Fig. 6 . Low magnification of "glazed"

surfaces on M2 steel using the

scanning mode illustrated in Fig. 5; conditions are given in text.

320
I

I

Fig. 7 .

Higher magnification micrographs of

surface shorn in Fig. 6.

Fig. 9 . Schematic diagram of linear heat source melting.

I

Fig. 10. Graph of Normalized Velocity as a

^ function of Normalized Width for

linear source scanning.

Normalized Width (w/d)

321
Fig. 11 . Cross-sectional views of wide zones

in M7 steel formed by linear heat source melting. For conditions

see text and Table II.

1mm.

J

Fig. 12 .

"As glazed" surface of material

glazed using a linear heat source, material - M7 steel.

Fig. 13 .

Surface topography and microstructure of similar areas as shown in Fig. 12 following

successive amounts of light polishing and subsequent etching.

322
IMPROVED IRON BASE ALLOYS

MADE USING A RAPID

SOLIDIFICATION PROCESS

Ranjan Ray and Viswanathan Panchanathan

Marko Materials, Inc., North Billerica, MA

and

Saul Isserow

Army Materials and Mechanics Research Center

Watertown, MA

ABSTRACT

Initial results of investigation on selected iron base alloys using a Rapid Solidification Powder Metal Technology currently being developed at Marko Materials under

a DARPA/AMMRC program are reported. The various alloys studied include commercial precipitation hardenable (PH) stainless steels, iron aluminide base alloys and high speed

steels. These alloys were modified with specific amounts of boron and processed by

melt spinning as rapidly solidified powders followed by hot extrusion. The consolidated alloys were subjected to various heat treatment procedures for optimum mechanical

properties.

A 30-35% improvement in tensile strength combined with good ductility has been

achieved in boron modified PH stainless steels compared to the conventional alloys.

Rapidly solidified Fe-Al-B alloys show high specific strength and modulus of elasticity in the temperature range 1000-1400°F. The other noteworthy attributes of the

iron aluminide base alloys are outstanding oxidation resistance, good low cycle fatigue

strength, and high thermal stability at elevated temperatures. Very high elongations,

typically up to 280%, were achieved in these alloys at high temperatures(1600°F) indicating the possibility of superplastically hot deforming these materials into shaped

objects.

A rapidly solidified boron modified tool steel showed significant improvement in

tool life in cutting (turning) tests, compared to its commercial counterpart.

323
INTRODUCTION

During the past three years, Marko Materials has been engaged in the development of a rapid solidification powder metal technology to produce improved iron-based alloys containing certain specific

amounts of boron. Rapid solidification processing of such boron-containing ferrous alloys produces

generally a solidified alloy having a chemically homogeneous, metastable microcrystalline phase(s).

Upon subsequent heat processing during hot consolidation, the metastable phase decomposes into an aggregate of stable phases including a fine dispersion of borides. The fine boride particles enable re- tention of an ultrafine recrystallized (matrix) grain structure and contribute to the microstructural

stability and enhanced mechanical strength of the alloys at elevated temperatures. Marko practices

its rapid solidification powder production process as follows: The alloys are produced by the method

of melt spinning as chill cast thin ribbons which are brittle as a consequence of metastable phase formation; the chill cast (brittle) ribbons are readily pulverized into powders of desirable size ranges

by a rotating hammer mill. Powders are consolidated into bulk forms by standard powder metallurgical

techniques e.g. hot extrusion.

In August 1981, Marko began a DARPA/AMMRC program aimed at evaluating the suitability of the aforesaid iron base alloys of different classes for potential defense hardware application (e.g. gas turbine

engine parts). One major thrust of this program was to demonstrate the feasibility of scaling up Marko 's powder production process to a pilot-plant size, thus providing a basis for a new manufacturing

technology. This paper highlights some of the important results accomplished on three classes of alloys during the first year of the program.

PROCESS SCALE U P

Equipment has been designed, fabricated and installed to batch produce large quantities (150-300

pounds) of rapidly solidified powder. This equipment (Figure 1) is based on the principles of fabrication of multiple ribbons chill cast by melt spinning followed by in-line pulverization of the ribbons

directly off the rotating substrate using hammer mill into powders. At the time of writing, the

testing and debugging of the scaled-up powder production system have been initiated. The chief attributes of this process are anticipated to be (a) ability to manufacture all powder particles with uniform quench history, and (b) ability to produce ecomonically a high yield (95%) of good quality powder

quenched at a very high rate under the effect of conductive cooling.

ALLOY STUDIES

Commercial iron-base alloys combining high strength with good oxidation/corrosion resistance at high temperatures are primarily the hardenable stainless steels or modifications based on them. The

various martensitic stainless steels of the straight chromium grades (e.g. AISI 420, 440, etc. ) as well

as the semiaustenitic and martensitic pricipitation hardenable (PH) stainless steels (e.g. 17-7 PH,

AM 350, 17-4 PH, etc.) are useful to 900°F. Strengthening due to martensite formation or precipitation

hardening in these steels decreases rapidly above 900°F due to tempering and overaging effects.

Austenitic precipitation hardenable stainless steels are an outgrowth of extensive compositional

modification of austenitic stainless steels for retention of high strength above 900°F. These alloys

are better known as iron-base superalloys and are considered to be forerunners of the more widely used

nickel and cobalt base superalloys. These highest temperature commercial iron-base superalloys, with

capabilities up to 1200°F, include Discaloy, A-286 and V-57. Most of these alloys in this system con- tain about 15% chromium to provide corrosion resistance, 25% nickel to stabilize the austenite structure,

1-2% molybdenum for solid solution strengthening and 2-3% aluminum plus titanium for precipitation

hardening. Higher temperature conventional wrought alloys of this type appear unlikely to be accepted

by industry because of a tendency towards phase instability as alloy content is increased. With increasing nickel content in these alloys (replacing iron), improved stability is achieved, but at a higher cost. The major goal of the DARPA/AMMRC program is to optimize and evaluate selected iron base alloys

of Marko via rapid solidification technology (RST) for potential high temperature structural applications in the intermediate temperature range of 800-1200°F. These new RST alloys will be specifically

developed to evolve as potential replacements, on the basis of cost-effectiveness, superior properties

and/or low contents of strategic metals (e.g. Cr, Co etc.) for conventional precipitation hardenable

(PH) stainless steels, titanium alloys and iron base superalloys.
Figure 1. RST powder production

facility at Marko

Materials

MODIFIED PH STAINLESS STEELS

Alloys of one category which were undertaken for the present investigation are based on conventional PH stainless steels modified to contain 1.6-2. 4 weight percent boron. Preliminary exploratory

screening of a number of different PH stainless steels with different boron concentrations led to the

selection of the two alloys for detailed studies as listed in Table 1.

TABLE 1 Nominal Chemistry of Boron Modified Precipitation

Hardenable Stainless Steels

Alloy No. Composition (weight percent)

RST 1.1 PH 15-7 Mo containing 2.1% Boron [Bal Fe, 15.1 Cr, 7.1 Ni, 1.2 A1 , 2.2 Mo,

.07 C and 2.1 B]

RST 1.27 Custom 450 containing 2.2% Boron [Bal Fe, 15 Cr, 6 Ni , 0.8 Mo, 1.5 Cu,

0.27 Cb, 0.05 C, and 2.2 B]

325
The RST 1.1 alloy is a modification of a semiaustenitic PH stainless steel. Conventionally pro- II

cessed semiaustenitic PH steel (annealed or solution treated) exhibits microstructure consisting of tt

austenitic matrix with an average grain size of 20 microns[l]. Transmission electron microscopic examin?

ation of thinned, electropol ished sections of as solidified ribbons of RST 1.1 alloy parallel to the i ribbon surface revealed microcrystalline structure consisting of equiaxed grains (Figure 2). There is no evidence of second phase precipitation at the grain boundaries. The size of the micro grains vary, but is typically 0.25 microns. Furthermore, the grains were in all cases subdivided into a cell structure. The cell diameter is very small, typically 300-500 A. Boron has negligible solid solubility

in steels. Pricipitation of MpB occurs above 58 p.p.m boron in commercial Fe-Cr-Ni base stainless

steels [2]. Rapid quenching inhibits the diffusional processes which cause second phase precipitation

and generated the present extremely homogeneous non equilibrium supersaturated solid solution phase

with microcrystalline morphology. Wood et al obtained by rapid solidification of melt up to 1.0%

boron in supersaturated metastable Fe-Cr-Ni base solid solution which decomposed upon aging at high

temperatures forming complex M23 B 6 borides at the grain boundaries with no obvious crystallographic

relationship with the matrix [3]. Figure 3 shows a transmission electron micrograph of RST 1.1 alloy

ribbon heat treated at 1750°F for 2 hours followed by air cooling to room temperature. The resultant

microstructure which grew from the original rapid solidification structure was still exremely fine.

The primary matrix grains remained small, about 0.3 micron, perhaps due to the pinning effect induced

by 0. 1-0.3 micron sized boride crystals nucleated at grain boundaries (indicated by arrows in the

micrograph). Additionally a very fine scale precipitate phase was observed within the grains. Although the as solidified ribbons were brittle, the heat treatment at 1750°F rendered them fully ductile

to 180° bending.

PREPARATION OF BAR STOCK AND MECHANICAL PROPERTIES

Alloy powders of the two boron modified PH steels were prepared by the RST melt spinning-pulverization process. About 20 pounds of -80 +400 mesh sized powders of each alloy were hot dynamically out

gassed and sealed off in 4 inch dia. mild steel cans for consolidation by extrusion. Extrusion was ac- complished at 1975°F at a high reduction ratio (20:1) to ensure complete densification of powders.

The commercial PH 15-7 Mo alloy is a semiaustenitic stainless steel generally produced and used in flat rolled sheet and strip products. After fabrication, through a series of relatively simple low

temperature heat treatments, it can be heat treated to a variety of strength levels depending on the

design requirements. The heat treatment consists of three basic steps: (1) austenitic conditioning,

(2) martensitic transformation, and (3) age hardening by formation of fine precipitates of Ni-Al intermetallic compound. The optimum heat treatment cycles of semiaustenitic PH steels are dependent on careful control of the interrelationship of austenite and ferrite forming constituents and their relationship to the total chemical balance of the system. It should be noted that the RST 1.1 alloy has a microcrystalline grain structure containing a fine dispersion of boride phase(s). It is very likely

that the alloy ctiemistry of the matrix was disrupted as a result of addition of boron and hence the

optimum heat treatment condition for RST 1.1 might be different than for commercial PH 15-7 Mo steel.

The heat treatments for the RST alloy were initially chosen based on the basic principles of physical metallurgy associated with the commercial PH steels. A reasonable wide latitude in the times and

temperatures of the various heat treatment sequences is possible to obtain usable mechanical properties.

However, the heat treatments employed in the present investigation were restricted to (a) austenite

conditioning around 1600°F, (b) subzero treatment around -100°F for 8 hours and (c) precipitation hardening between 840°F and 1020°F. The precipitation hardening effect is a complex interplay of three

separate and distinct reactions e.g. tempering of martensite, precipitation of an intermetallic com- pound and reversion of martensite to austenite. The degrees to which these reactions proceed depend

on the times and temperatures. A careful analysis of the complex aging phenomenon will be required to develop the optimum mechanical properties of the present alloy. The effect of selected precipitation

hardening treatments on the mechanical properties of the RST 1.1 alloy is shown in Table 2.

As illustrated in Figures 4 and 5, the RST 1.1 alloy in suitably heat treated condition possesses

tensile properties (0.2% offset YS and UTS) about 25-30% higher than a fully heat treated commercial

semiaustenitic (17-7 PH) stainless steel in the range from room temperature to 1000°F.

326
Figure 3. TEM picture of RST 1.1 alloy ribbon heat treated at 1750°F for 2 hours

327
TABLE 2

Room and Elevated Temperature

Properties of RST 1.1 Alloy

Aging condition Test 0.2% Offset Ultimate % % RA following sub- Yield Tensile Elongation

zero treatment Temperature Strength Strength

°F KSI MPa KSI MPa

24 hours at 75 255 1760 277 1911 0.4 0.6

840°F

800 173 1184 180 1242 0.6 1.0

72 hours at 840° F 75 230 1586 250 1722 3 6

800 173 1192 181 1248 3.6 4.5

24 hours at 75 135 928 192 1325 11.4 12.1

1024°F

1000 100 692 122 842 15.3 18.1

0.2% YIELD STRENGTH

230

190 KSI

150

110

200 400 600 800 1000 TEST TEMPERATURE (°F)

ULTIMATE TENSILE STRENGTH

280

240

200 KSI

160

120

200 400 600 800 1000 TEST TEMPERATURE (°F)

Figure 4. 0.2% yield strength of RST 1.1

alloy and commercial 17-7 PH steel as a function of test

temperature.

Figure 5. Ultimate tensile strength of RST

1.1 alloy and commercial 17-7 PH steel as a function of test

temperature.

328
An of the commercial semiaustenitic stainless steels such as 17-7 PH, PH 15-7 Mo, Ph 14-8 Mo etc. are generally capable of developing approximately the same tensile strength [1] ranging from a low

175,000 psi (1206 MPa) to a high 240,000 psi (1655 MPa). However, fracture toughness and ductility

of the semiaustenitic stainless steels all heat treated to the same tensile strength level are not the

same. It is also known that fracture toughness is an excellent indicator of the stress corrosion re- sistance of these steels at room as well as elevated temperatures. Stress corrosion resistance is generally enhanced with increasing fracture toughness. Work has been planned, therefore, to determine the

fracture toughness values of the present RST 1.1 alloy and compare them with the semiaustenitic stainless steels at the same tensile strength levels.

The RST 1.27 alloy of the present work is based on commercial Custom 450 alloy which is a martensitic precipitation hardenable steel. After high temperature solution treatment, the structure of this

steel is martensitic. Additional strengthening is achieved by an aging treatment at 900-1100°F. The

RST 1.27 alloy was prepared as bars similarly as the RST 1.1 alloy. Fine microstructural features were

observed in the optical photomicrographs of the as extruded bars (Figure 6). The extruded bars were

subsequently aged between 850 and 1020°F for different hours. Selected tensile data of the RST 1.27

alloy are given in Table 3 as a function of different age hardening treatments. The RST 1.27 alloy

was found to exhibit greater than 25% improvement in tensile properties (0.2% offset YS and UTS) than

Custom 450.

TABLE 3 Room and Elevated Temperature

Tensile Properties of RST 1.27

1 Condition

Test

Temp.

°F

0.2% Offset

Yield

Strength

Ultimate

Tensile

Strength

%

Elong.

%RA

KSI MPa KSI MPa

<\ged at RT 228 1571 251 1730 4.3 4.8

340°F for

hours 600 187 1288 208 1433 7 10

800 163 1123 188 1295 7.7 13

Aged at RT 212 1461 233 1605 7.2 11.5

1020°F for 600 184 1268 198 1364 3.1 9.4

4 hours

800 164 1130 179 1233 12 12 1050 111 765 122 841

k

1 11.3 24 Figure 6. Optical photomicrographs (lOOOx)

of as extruded RST 1.27 alloy bar (a) longitudinal section (b) transverse section.

Stress corrosion tests were carried out with tensile specimens of the RST 1.27 alloy in the sge

hardened condition in accordance with NACE TM-01-77. The environment used was distilled water solution

of 0 1 percent NaCl , 0.5 percent acetic acid, saturated with H 2 S gas. Time to failure was measured as

a function of applied stress. The RST 1.27 alloy was found to have significantly superior sulphide

stress corrosion resistance than Custom 450 (see Table 4).

TABLE 4

Stress Corrosion of Tensile Specimens in

0.1% NaCl + 0.5% Acetic Acid + H 2

S

Alloy Hardness

Rockwell C Applied Stress

KSI MPa

Time to Failure

Custom 450 41 60 414 10 hours

RST 1.27 42 60 414 56 hours

329
IRON-ALUMINUM-BORON ALLOYS

Alloys based primarily on iron and aluminum, relatively inexpensive elements with abundant and H

more secure reserves, offer excellent alternative possibilities to the current heat resistant alloys

containing chromium, nickel and cobalt. Iron-aluminum alloys containing 13 wt.% (24 at%) on higher B

contents of aluminum exhibit outstanding oxidation resistance and were considered as potential candi- H

dates for a wide variety of heat resistant applications as engineering components. Binary Iron-alumin-j

urn alloys have a strong susceptibility to microcracking during the ingot casting operation due to low

thermal conductivity and high thermal expansivity. This characteristic feature, coupled with large

grain growth at high temperatures, renders hot working iron-aluminum alloys containing high aluminum H

contents difficult. Poor workability and room temperature brittleness of these alloys have retarded H

their applications. Furthermore, iron-aluminum alloys with about 13 wt.% or higher contents of aluminu^

have stress rupture properties less than adequate for high temperature applications. Some of these

problems encountered in the iron-aluminum alloys can be circumvented by processing the alloys from prealloyed rapidly solidified powders of high purity. Recent efforts at Pratt and Whitney Aircraft generated improved strength and ductility combined with good oxidation resistance in Fe-Al alloys containing

13 wt.% aluminum, 1.5 wt.% titanium and 0.5 to 0.7 wt% boron which were compacted from rapidly solidified powders [4]. A new family of Fe-Al -B alloys was examined via RST at Marko under the present investigation. The

primary rationale for addition of boron in certain critical concentrations to Fe-Al alloys was to render

the alloy responsive to Marko's RST process (melt spinning-pulverization). Furthermore, boron was anti

cipated to generate microcrystalline structures stabilized by finely dispersed borides in the consolidated alloy. Strength and ductility are likely to benefit from such structures. Preliminary search for alloys which conformed to the RST processing constraint and displayed excellent high temperature oxidation resistance led to the formulation of the following optimum composition:

Bal Fe, 17-22 wt.% A1 and 3-4 wt.% B Additions of alloying elements such as Mo, W and V to the above alloys were explored for beneficial effects on creep strength and thermal stability via possible formation of stable borides. Table 5 lists

the three alloys which were studied in greater details in the present program.

The Fe-Al base alloys were prepared as bars by hot extrusion of RST powders (-80 +400 mesh) at 1750°F using a reduction ratio of 14:1. The extruded rods were further given an annealing treatment at 1740°F for 3 hours prior to testing of various mechanical properties (e.g. tensile, stress rupture and

low cycle fatigue). The alloy compositions of the present Fe-Al-B system are novel with little or no knowledge existent in the prior art about the various phase relationships in this ternary system. Hence

forth, a detailed understanding of the microstructural features of these new RST alloys in the consolidated state was deferred pending full determination of their mechanical properties in the desired tem- perature range. An immediate goal in the current program was to find out the potential usefulness as well as the drawbacks of the present alloys.

Elevated temperature tensile properties of the Fe-Al-B alloys are listed in Table 6.

Figure 6 shows the SEM micrograph of RST 4.2 alloy ribbon heat treated at 1750°F for 1 hour. The

microstructure consists of a fine dispersion of a pricipitate phase on a fine grained general matrix.

The size of the multi-faceted precipitates is about 0.5 micron. Assuming boron to be partitioned en- tirely in the precipitate phase, EDX (Energy Dispersive x-ray) analysis revealed the precipitate to be

an iron rich boride phase with a small amount of aluminum. The matrix composition, as a consequence of

the formation of the iron rich boride phase, shifted close to the intermetal 1 ic FeAL . compound. Fine

microstructural features were revealed in the optical photomicrographs of the RST 4.2 alloy in as ex- truded condition (Figure 8).

TABLE 5

Iron-Aluminum-Boron Alloys

A1 loy Composition, weight percent

Fe A1 B Other

RST 4.2 Bal 17.9 3.1

RST 4.4 Bal 20.8 4.1

RST 4.24 Bal 21.3 4.0 4.8 Mo 2.5 V

330
Figure 7. SEM photograph of RST 4.2 alloy

ribbon heat treated at 1750°F

for 1 hour.

(b)

Figure 8. Optical photomicrographs (lOOOx) of

extruded bars of RST 4.2 alloy.

(a) longitudinal section

(b) transverse section

Elevated Temperature Tensile

Properties of RST Fe-Al-B Alloys

Alloy Test

Temperature

°F

—

0.2% Offset

Y.S

Ultimate

Tensile

Strength

%

Elongation

% RA

KSI MPa KSI MPa

RST 4.2 800 157 1082 177 1223 0.2 0.2

[Fe-17.9 Al-3.1 B]

1000 118 813 140 967 7.6 8.2

1200 72 496 93 643 24.0 31.2

1400 26 181 58 397 56 41.5

RST 4.4 1000 118 813 144 992 5.4 7.9

[Fe-20.8 Al-4.1 B] 1200 79 547 101 693 18.2 24.2

1400 23 157 60 411 49.4 37.2

1600 9 60 14 97 280.8 71.7

RST 4.24 1000 137 947 162 1118 2.5 4.1

[Fe-21.3 Al-4.8 1200 78 538 115 793 8.1 11.6

Mo-2.5 V-4 B] 1300 51 348 94 646 5.1 8.9

331
SPECIFIC TENSILE STRENGTH

Figure 9. Specific tensile strength of RST

Fe-Al-B alloys and A-286 as a function of test temperatures.

Figure 10. Specific tensile strength at elevated temperatures of RST

4.24 alloy and commercial

titanium alloys.

THERMAL STABILITY

UTS

(KSI)

CYCLIC OXIDATION BEHAVIOR

Time (Hours)

Figure 11. Effect of isothermal annealing

at 1650°F on tensile strength of

RST 4.4 alloy in the range 1000-

1400°F.

Figure 12. Cyclic oxidation behavior at 1800°F of RST Fe-Al-B alloys as shown compared to 304 stainless

steel

332
An the alloys were found to exhibit low ductility i.e. less than 1% elongation at temperatures

below 800°F. At 1000°F or higher temperatures, the alloys exhibited reasonably good to excellent ductility combined with attractive tensile strength values.

The density values of the Fe-Al-B alloys range from 0.228-0.235 lb/inch3 (6.32-6.53 gm/cc) which

are generally 20% lower than iron base alloys. In the temperature range of interest, i.e. 1000-1400°F,

the present alloys exhibited significantly higher specific tensile strength values than iron base superalloy, A 286 as illustrated in Figure 9. The RST 4.24 alloy was found to have higher specific strength

than two commercial titanium base alloys at high temperatures (Figure 10). Low temperature tensile

properties are not satisfactory criteria for selecting a high temperature alloy. Such tests, made after

an alloy has been exposed to higher temperatures, are indicative of the change in properties that may be

expected due to high temperature exposure. Thermal stability of structures of alloys intended for long

term uses at high temperatures is considered an important characteristic. The RST 4.4 alloy of this

work showed excellent thermal stability as evidenced by little or no decrease in the ultimate tensile

strength values in the range 1000-1400°F (Figurell) after annealing the alloy at 1600°F for 100 hours.

Low room temperature ductility of the Fe-Al-B alloys might pose a problem during fabrication of

parts by conventional machining. However, perhaps due to equiaxed grained structure stabilized by

a large volume fraction of dispersoids e.g. borides, the RST Fe-Al-B alloys can be elongated at elevated

temperatures to exceptionally large strains without risk of rupture. High elongations (to 280%) at 1600°F have been achieved with the RST 4.4 alloy (Table 6) indicating superplastic behavior. Superplasticity might allow forging of these alloys into complex shapes at relatively low stress.

A potential problem that needs to be addressed involves inadequate high temperature stress rupture

strength values of the microcrystalline alloys. For example, at 1000°F, 100-hour specific creep

strength of the RST 4.4 alloy is comparable to A 286 (see Table 7). However, with increasing temperatures, the stress rupture strength of the fine grained RST alloy deteriorates at faster rate than the

conventional coarse grained A 286 alloy. The grain-boundary related creep rupture failure mechanism

becomes a serious concern for fine grained materials. Control of grain size in the present RST alloys

seems deemed important to enhance creep and rupture strength without excessively trading off tensile

and yield strength properties. Judicious alloying with proper elements to improve creep resistance

via decreased boundary diffusion offers an alternative option.

TABLE 7

Stress Rupture Properties

Alloy 100 Hour Specific Rupture Strength

[inch X 10"3]

1000°F 1100°F 1200°F

RST 4.4 320 197 105

A 286 325 265 205

Cyclic oxidation between 70 and 1800°F was studied with two RST alloys, 4.4 and 4.2. The specimens

were heated in static air environment to 1800°F for 50 hours and cooled to room temperature, and then

reheated. The changes of weight as a function of exposure time are shown in Figure \Z. When exposed

in air at high temperatures, the Fe-Al-B alloys form thin, continuous and tightly adherent oxide skin

over the entire specimen surface. The oxidation resistance of the present alloys was found to be ex- cellent as compared to 304 stainless steel.

Low cycle fatigue testing at 1200°F with smooth bar specimens indicated reasonably good fatigue

life of the RST 4.4 alloy. Some test data are given in Table 8.

333
TABLE 8

Low Cycle Fatigue Data of

RST 4.4 Alloy at 1200 °F

Test Conditions

Mode: Longitudinal Strain Control

A Ratio: 1; Frequency: 0.50 Hz Wave Form: Triangle

Maximum Engineering Stress, Max(KSI) 45.8

1 71.6 ' 74.4 101.5

Minimum Engineering Stress, Min(KSI) -50.0 1 -76.6 i -79.8

' -105.3

Measured Total Strain Range, (%) 0.37 0.70 ! 0.80 2.00

Measured Plastic Strain Range, (%) .02 j 0.11 0.24 0.81

Cycles to Failure: 55,587 ! 3,635 2,549 118

The present Fe-Al-B alloys have very high modulus of elasticity over a wide range of temperatur^

(Figure 15). For example, the RST 4.4 alloy exhibits 50% higher specific modulus than A 286 alloy in the range from room temperature to 1200°F.

Figure 15. Modulus of elasticity of RST 4.4

alloy and A-286 as a function of

test temperature. p

i:

p

BORON MODIFIED HIGH SPEED STEELS

Powder metallurgical technologies have recently been applied to the production of tool steels [5,6]r,

Because of their finer grain size, more uniform dispersion of fine carbides and improved alloy homogeneity, high speed steels processed by such powder metallurgical techniques exhibit, compared to ingot

cast materials, superior cutting performance, a better response to hardening heat treatments, improved

334
dimensional stability and improved grindability of cutting edges. In the present program, the princi-

' pies of RST have been applied to the design of improved high speed steels. The alloys were formulated

'i with compositions similar to conventional high speed steels modified to contain 0.45 to 1.35% boron,

j The addition of boron was intended to provide a fine dispersion of refractory metal borides in the

i martensitic matrix of the RST alloy after suitable heat treatments. These borides do not dissolve or

' deleteriously coarsen at elevated operating temperatures, giving enhanced hot hardness and high temperature strength.

ii One such alloy is RST 3.11 having composition similar to commercial M-4 high speed steel containing

0.65% boron. Its nominal composition is as follows: Bal Fe, 4.25 Cr, 4.0 V, 4.5 Mo, 5.75 W, 1.35 C and 0.65 B. The alloy was prepared as bars by hot extrusion of RST powders at 1975°F using a reduction

ratio of 20:1. Its microstructure is very homogeneous and fine grained dispersed with fine boride partj ides (Figure 13). The consolidated alloy was austenitized at 2175°F followed by oil quenching to room

: temperature. RST 3.11 was fully heat treated to a hardness value of Rockwell C 66 and then tested as

! a single point lathe tool bit in a cutting (turning) operation using 4340 steel with a hardness of

Rockwell C 35 as work material. The tool life versus cutting speed curve is shown in Figure 13. This

figure also shows, for the sake of comparison, the cutting life of a widely used conventional M-2

high speed steel. M-2 and M-4 high speed steels are generally used for turning operations with a cutting speed of 50 to 75 ft/mi n. From Figure 14, it is estimated at a cutting speed of 60 ft/min., the

RST 3.11 shows 70% improvement in tool life compared to M-2. Also, at 75 ft/min. cutting speed, the

improvement in tool life of the RST alloy is 30% over M-2.

Figure 13. Optical photomicrograph of RST 3.11

alloy extruded bar.

Magnification: lOOOx

(a) Longitudinal section

(b) transverse section

TURNING 4340 STEEL, ROCKWELL C 35

Cutting Speed- Feet/Minute

Figure 14. Tool life as a function of cutting

speed of RST 3.11 alloy and a commercial M-2 steel.

335
CONCLUSIONS

It is believed from the work done in the initial phase of the DARPA/AMMRC program, the feasibility i of applying Marko's rapid solidification powder metal alloy technology to synthesis of new and improved

iron base alloys has been demonstrated. Although the first generation alloys investigated under this

program cannot be considered as optimum, the results of the study are sufficiently encouraging to warrant further intensive development work. j- j

ACKNOWLEDGEMENTS *

The work is being sponsored by the U.S. Defense Advanced Research Projects Agency and monitored

by the U.S. Army Materials and Mechanics Research Center, Watertown, Massachusetts under the contract

No. DAAG 46-81-C-0062. The TEM and SEM photographic work was conducted by Dr. S. Sastri , Gillette

Company, Boston, Massachusetts.

REFERENCES H- 1. D. Peckner and I.M. Bernstein, Eds., Handbook of Stainless Steels, McGraw-Hill, Inc. (1977)

2. H.J. Goldschmidt, J. Iron Steel Inst., 209, p. 900 (1971)

3. J.V. Wood and R.W.K. Honeycombe, Proc. Second International Conference on Rapidly Quenched

Metals, N.J. Grant and B.C. Giessen, Eds., Section 11, M.I.T. Press, Cambridge, Mass., 107 (1976)

4. S. Das and E. Slaughter, Proc. Second International Conference on Rapid Solidification Proces-TM,

sing, Eds., R. Mehrabian, B.H. Kear and M. Cohen, Claitors p.354 (1980) :

5. T. Levin and R.P. Harvey, Metals Progress, vol 115, No. 6, p. 31 (1979)

6. A. Kasak, G. Steven and T.A

Congress, Detroit (1972)

. Neumeyer Society of Automotive Engineers, Automotive Engineering
BORIDE AND CARBIDE DISPERSION IN MELT SPUN SUPERALLOYS

S. C. Huang, K. M. Chang and E. L. Hall

General Electric Corporate Research and Development

P.O. Box 8 Schenectady, NY 12301

ABSTRACT

Melt spinning of superalloys has produced crystalline and in some compositions amorphous ribbons, which, upon heat treatment, yield unique microstructures. This paper

reports our preliminary investigation of the opportunity for developing new superalloys

through the melt-spinning, rapid solidification process. Two superalloys were chosen as the basis of our studies. One was a nickel-based superalloy into which various amounts of

boron were added to change the solidification and precipitation behavior. When the boron

addition was less than 1.0 wt.%, a crystalline cell structure developed with boride precipitation at cell boundaries. When the boron addition was greater than 1.0 wt.%, the ascast ribbon contained small particles (<200X) of Ni-base solid solution embedded in an amorphous matrix. In the other studied alloy which contains 0.43 wt.% carbon, only onethird of the ribbon thickness (near the ribbon surface contacted with the casting wheel)

developed a homogeneous solidification structure containing monocarbide precipitates less

than 200A in size. The high-temperature stability of borides and carbides in the above

systems will also be presented. These results have helped us to define guidelines for our future superalloy research using the rapid solidification technique.

Introduction

Alloys produced by rapid solidification usually develop microstructures unobtainable by conventional casting. These unique microstructural features include submicron grain size, extended solid

solutioning, ultrafine precipitation, metastable phases, and in some cases metallic glasses can be

produced [1]. Commercial alloys, originally developed using conventional processing, may exhibit

improved properties when prepared by rapid solidification and consolidation. With few exceptions,

however, alloy systems designed explicitly to take the full advantage of rapid solidification have

yet to be explored [2-4]

Fine carbide and boride particles strengthen grain boundaries as well as the bulk material.

Carbon and boron, however, tend to segregate during solidification, causing blocky precipitation at grain boundaries and thus give rise to suboptimal properties. Rapid solidification reduces solute

segregation and is therefore capable of extending the domain of alloying. Studies of carbides and

borides in iron-base alloys processed by rapid solidification techniques have been carried out [5-

7]. The present paper reports on the boride and carbide dispersion produced by the melt-spinning

process in nickel-base superalloys. In particular, the effects of boron and carbon on the ribbon

castability, microstructure, phase stability and ductility have been investigated. A specific aim

of the present studies is to demonstrate the capability as well as the limitation of rapid solidification processing in nickel-base superalloy development.

Experimental

One of the superalloys studied (alloy A) has a nominal composition as in Table 1. The gamma

prime fraction is about 30% and the concentrations of the prime carbide and boride-forming elements

are typical of most superalloys. Ribbons of alloy A and its variations with boron levels of 0.75,

1.45 and 3.0 weight percent were fabricated by the melt spinning technique in ambient air. The

other superalloy studied (alloy B, Table 1) was originally designed for the unidirectional growth of

y/y' “ TaC eutectics. Ribbons of alloy B were melt-spun in a vacuum. The ribbon cross-sections

were studied by optical metallography, the ribbon bottom-surface topography by microscopy using the

Normarsky-interference contrast, and the ribbon top surface by scanning electron microscopy (SEM).

337
TABLE 1. Nominal Compositions of Alloys Studied (w/o)

Alloy Ni Co Cr A1 Ti Nb Ta Mo W Re C B

A balance 14.0 17.0 1.75 3.5 0.65 _ _ 8.0 2.0 0.02 0.01 B balance 3.80 3.84 5.34 - - 11.37 3.01 4.33 6.44 0.43 -

These yielded information about the solidification structure and the ribbon-wheel contact characteristics. Furthermore, transmission electron microscopy (TEM) examinations of the ribbon microstructures were made on samples electrolytically thinned in the solution of 20% perchloric acid ini! methanol. For alloy B ribbons, microstructures at various positions through the ribbon thickness were studied by TEM on samples thinned from either one or both of the ribbon surfaces. Differential

scanning calorimetry (DSC) and X-ray diffraction were performed on the ribbons based on alloy A to determine the ribbon structure and phase stability as influenced by boron content. The ribbon ductility, as cast and annealed at 900°C for 1 hour, was tested by bending.

Results and Discussion

A. Effects of Boron on Castabilitv

Alloy A is in general difficult to cast into ribbon in air. The quench rate is too slow to prevent unacceptable oxidation of the ribbon surface, particularly at the beginning of each casting

run. Observations are that the sticking distance of alloy A ribbons on the casting surface are too short to provide adequate secondary cooling. Additions of boron greater than 1.0 wt.% to alloy A drastically improve air castability, giving fast enough solidification and secondary cooling that little oxidation of the ribbons occurs. The ribbon-to-wheel sticking distance increases with boron

additions greater than 1.0 wt.%, but the contact patterns [8] on the wheel-side surfaces of the ribbons fail to show any evidence of improved metal-wheel contact. Fig. 1. In fact, the alloy A ribbon

shows near-perfect contact with the wheel in regions between extended air pockets. Fig. 1(a), while

the ribbon of alloy A containing 3.0 wt.% boron shows uniform but poor contact with the wheel. Fig.

Kb). The beneficial effects of boron on the ribbon castability thus cannot be totally attributed

to its modification of the liquid-metal fluidity or surface tension, which may influence the meltwheel contact characteristics. Instead, it should be largely attributed to boron's effect on the thermal properties of the melt and the thermodynamics and kinetics of the solidification

process [9]

Fig. 1 The mel t-to-wheel contact patterns observed on the wheel-side surfaces of the air-cast

ribbons of (a) alloy A and (b) alloy A plus 3.0 w/ o boron. The castings were in the horizontal direction. I

B. Effects of Boron on Solidification Microstructures

The solidification structure developed in ribbons is most conveniently studied on the ribbon

top surface where liquid metal is 'decanted'. Figure 2 compares the ribbon top surface structure as influenced by the boron content. For alloy A, which contains only a trace amount of boron, the sur- face dendrites are about 1 pm in length, and only a few ternary branches are visible. Fig. 2(a).

Although the ribbon surface appears to have the [100] texture judging from the branching symmetry,

the [100] dendritic crystals tend to have rotation mismatches with neighbors. A detailed study [10]

has shown that the dendrites are, in fact, formed by independent nucleations and that every dendrite

constitutes a separate grain. When the boron content is increased to 0.75 wt.%, the [100] dendritic

structure is still maintained. Fig. 2(b). The dendrites are about 2 pm in length, and they tend to
be more intricately branched, having a ternary branch spacing of ~0 .2 pm, compared to ~0.5 pm in the

low-boron ribbons. Furthermore, some neighboring dendrites of similar orientation seem to group

together, being surrounded by the heavily segregated grain boundaries. The grains as conceivable in

Fig. 2(b) are about 5 pm in diameter. When boron is further increased to beyond 1.0 wt.%, the dendritic structure no longer exists. Fig. 2(c). Instead, the ribbon becomes essentially structureless, decorated only by some unidentified surface striations.

Fig. 2 The 'decanted' solidification structure on the top surfaces of superalloy ribbons having

boron concentrations of (a) 0.01 w/o, (b) 0.75 w/o, and (c) 3.0 w/o.

A comparison of the ribbon cross-sectional structure dramatizes the effect of boron on the ribbon solidification. Fig. 3. As shown in Fig. 3(a), a segregated crystalline structure develops when

B < 1.0 wt.%. Here, the solidification crystals nucleated at the wheel-side surface and grow unidirectionally toward the ribbon top surface, in the cellular morphology initially before a transiFig. 3 The cross section micrographs of superalloy ribbons containing (a) 0.01 w/o and (b) 3.0

w/o of boron. The wheel-side surface is shown at bottom.

339
tion to the branched dendritic morphology beyond the center of the ribbon thickness. Small additions of boron refine the structural size-scale. When the boron content is increased to B > 1.0

wt.%, a planar-front solidification must have occurred, since the resultant ribbon structure is uniform without obvious segregation. Fig. 3(b). As discussed previously, this swift change in the

solidification mode at the 1.0 wt.% boron level is not due to improvement in solidification heat

transfer, but due to the reduced atom mobility at the solid-liquid interface. Boron has been known

for its promotion of the formation amorphous phases [9]

Fig. 4 TEM microstructure in superalloy ribbons containing (a) 0.01 w/o, (b) 0.75 w/o, and (c)

1.45 w/o of boron.

The microstructures of the superalloy ribbons containing various amounts of boron were further

investigated at high magnifications by TEM. For B < 1.0 wt.%, boride particles precipitate essen- tially in the interdendritic regions. Fig. 4(a) and (b) . Apparently, boron has segregated during

the solidification process. The size, as well as the quantity, of the borides increases with

increasing boron content, with the boride diameter being ~0.01 pm in alloy A ribbons. Fig. 4(a), and

~0.1 pm in the ribbon containing 0.75 wt.% boron. Fig. 4(b). The specific boride phase and composition are yet to be determined. Fig. 4(a) and (b) also show the refinement of the dendritic structure as observed in Fig. 3(a) and (b) . Gamma prime precipitates cannot be resolved in the magnifications used, even though the superlattice diffraction spots can be detected. For B > 1.0 wt.%, the

as-cast ribbons show a homogeneous structure containing small precipitates of less than 200&, Fig.

4(c). Electron diffraction reveals that the precipitates are nickel-base solid solution particles

embedded in an amorphous matrix. This is in corroboration to the X-ray diffraction result of Fig. 5 that there are diffraction peaks of crystalline nickel phase overlapping with the broadened diffraction background due to the presence of an amorphous phase. Boron is believed to be entrapped during

the nonequilibrium plane-front solidification, resulting in the amorphous phase without boride

precipitation.

340
Z9 (DEGREES)

Fig. 5 X-ray diffraction pattern showing the coexistence of a crystalline and an amorphous phase

in the as-cast superalloy ribbon containing 1.45 w/ o boron.

C. Effects of Boron on the Annealed Microstructure

The ribbon of alloy A with 0.75 wt.% boron addition was annealed at 760°C for 24 hours. The

resultant structure as in Fig. 6 shows a uniform distribution of coherent gamma prime precipitation

about 0.03 pm in diameter. Compared to Fig. 4(b), the intercellular boride particles remain about

0.1 pm in diameter without significant coarsening. Neither the cell size nor the morphology is

affected by this low temperature anneal. When the ribbon is annealed at 900°C for 1 hour, the

resulting gamma prime precipitates are about 0.5 pm in diameter, and the boride particles coarsen to about 0.2 pm in diameter agglomerating to the grain boundary triple points. Coarsening of the grain

structure seems to just begin, perhaps made easy by the redistribution of the grain boundary boride

particles

Fig. 6 Microstructure in the superalloy ribbon containing 0.75 w/o boron, annealed at 760°C for

24 hours.

341
The differential scanning calorimetry (DSC) study reveals two low temperature reactions, at ~460°C and ~600**C, in those ribbons containing B > 1.0 wt.%. Fig. 7. The specific reaction temperatures vary slightly with boron concentration, but both reactions are absent in ribbons containing B

< 1.0 wt.%. After annealing at 460 C for 1 hour, the ribbon with 3.0 wt.% boron exhibits a microstructure similar to Fig. 4(c). However, it generates electron diffraction rings additional to those from the nickel particles, indicating the possible precipitation of borides from the as-cast

1 1 I I 1 I I I 300 400 500 600 700

TEMPERATURE (“0

Fig. 7 DSC curves for superalloy ribbons of various boron concentrations, based on alloy A.

Fig. 8 Microstructure in the superalloy ribbon containing 3.0 w/o boron, annealed at 900”c for 1 hour

342
[

[l

i

j

j

I

amorphous matrix. Annealing at 600°C for 1 hour precipitates no additional new phase. The nickel

and boride particles enlarge slightly to about 200X, but the matrix phase remains homogeneous with

no development of grain structure. Annealing at 900**C for 1 hour produces a unique microcrystalline

structure having uniform grains of approximately 0.2 |im in diameter [3,4], Fig. 8. There is no

obvious precipitation of gamma prime phase within grains and boride particles of about 0.1 pm precipitate at the grain boundary triple points. Preliminary analyses by electron diffraction suggests

the boride is MB,.

I D. Effects of Boron on_Jlibbon_ Ductility

!

li The ribbon ductility can be conveniently tested by bending to fracture. The fracture strain,

i e, is given by e = h/(D+h), where h is the ribbon thickness and D is the bend diameter at which the

' ribbon fractures. The test results. Table 2, show that the ductility of the as-cast ribbons

j decreases with increasing boron content up to B = 1.45 wt.%. Further increase in boron improves the

I ductility, presumably due to the increase in the degree of amorphous state. After annealing at

' 900*C for 1 hour, the ribbon of 0,75 wt.% boron becomes ductile. Ribbons of higher boron concentra-

! tions, however, become further embrittled after annealing. This result is contrary to that of ref.

I [3] which showed an improved ductility when amorphous structure containing less than 2 wt.% boron

I was transformed to microcrystalline structure.

TABLE 2. Ribbon Bend-Fracture Strain

Composition

Strain to bend-fracture

(as-cast)

Strain to bend-fracture

(900“C/1 hr)

alloy A

A + 0.75 w/o boron

A + 1.45 w/o boron

A + 3.0 w/o boron

1

* 1* 0.03 1*

0.01 0,004

0.04 0.006

Bent 180 degrees without fracture.

E. Microstructure in Ribbons Containing Carbon

The microstructure of a longitudinal section (parallel to the casting direction) in an alloy B ribbon is shown in Fig. 9. As shown in this figure, the microstructure varies through the ribbon

thickness, being homogeneous near the wheel-side surface, cellular at the ribbon center, and dendritic near the top surface. These structural variations are much more pronounced than in the boronmodified ribbons. Obviously, the microstructure developed in the carbon-modified alloys are more

sensitive to the cooling rate.

Fig. 9 Optical micrograph of a longitudinal section in an alloy B ribbon. The wheel-side surface

is shown at bottom.

343
Fig. 10 TEM overall cell structure near the wheel-side surface of alloy B ribbon.

Fig. 11 The subcell structure observed near the wheel-side surface of alloy B ribbon: (a) bright i field image, (b) the electron diffraction pattern, (c) dark field image using an MC car- bide reflection, and (d) dark field image using a gamma prime reflection.

344
The overall microstmctare at the wheel-side surface of the ribbon is shown in the TEM micrograph of Fig, 10. The structure consists of cells of varying sizes, generally 0.5 to 2.0 |im in

diameter. The boundaries between these cells are decorated with particles less than 20 nm in diameter, Fig. 11(a). In addition, a lage number of particles of similar size are seen inside the cellular grains, often forming a regular subcell structure. Both the boundary and sub-boundary particles

can be shown by electron diffraction to have the MC—carbide structure. Fig. 11(b). X-ray spectroscopy further reveals that the particles are rich in tantalum. The lattice parameter measured for

the carbide particles is approximately 0.426 nm (in comparison to 0^455 nm reported for the

stoichiometric TaC) , suggesting a carbide composition of Ta^ jjCq 45*

' Substitution of Ta for

other carbide formers may also contribute to the reduced lattice parameter. It is also shown by

dark field studies that within a single grain, all the TaC particles are of a single orientation,

and have fully coherent relationships with the matrix. Fig. 11(c). The electron diffraction pattern

also reveals the presence of gamma prime in the gamma matrix. The dark field image formed using the

gamma prime reflection shows that the gamma prime is also fully coherent, with a cuboidal morphology

and an approximate dimension of 30 nm. Fig. 11(d).

In the center of the ribbon, the size of the cellular grains has increased slightly, the mean

gamma prime size is approximately 50 nm, and the carbide particles at both the grain boundaries and

inside the grains are much larger, 50 to 100 nm in diameter. Furthermore, the amount of the

intragranular carbide precipitation is notably decreased. Near the ribbon top surface, the centers

of the primary cells are totally depleted of carbide. Fig. 12. The carbide particles which preferentially precipitate in the intercellular regions generally take an irregular shape. Some continuous precipitation along cell boundaries can also be observed.

Fig. 12 TEM microstructure near the top surface of alloy B ribbon.

The alloy B ribbons have been annealed at 1000*C for 2 hours. The annealed structure near the

center of the ribbon is shown in Fig. 13, The gamma prime precipitation is coarsened to about 0,1

|un in diameter, but the cell size remains unchanged, being stabilized by the cell boundary precipitation. It should be noted that the irregular carbide particles are stable under the present

annealing condition, showing no signs of changes in morphology, distribution or size. There is,

however, a second set of precipitation particles, which are rounded in shape, 0.02 to 0.1 (xm in

size, and primarily located at the cell boundaries. Preliminary analyses indicate that the latter

particles are rich in rhenium and are precipitated and coarsened during the anneal.

Conclusions

Boron additions to a superalloy have demonstrated drastic modifications in the melt spinning

castability, solidification structure, and heat treatment precipitation. More specifically, boron

affects the liquid-metal thermal properties, and when B > 1 wt.%, a continuous ribbon can be readily

345
Fig. 13 Microstructure near the center of alloy B ribbon, annealed at 1000°C for 2 hours.

cast in air without excess surface oxidation. Boron also refines the dendritic structure and

changes the solidification kinetics, and when B > 1 wt.%, a mixture of microcrystalline and amorphous phases can be obtained as is evident from DSC and X-ray diffraction. TEM studies further show I that the high boron ribbons contain small particles <20 nm in diameter of Ni-base solid-solution

embedded in an amorphous matrix. Annealing the ribbons at 600“c for 1 hour precipitates uniform

borides smaller than 10 nm, and annealing at 900**C for 1 hour produces 0.2 pm microcrystall ine

cells, which are decorated with cell boundary precipitation of 0.1 pm borides. This unique microcrystalline structure, however, embrittled the high boron superalloy ribbons, giving a bend fracture

strain less than 0.01.

A uniform microstructure in a su^ralloy containing 0.43 wt.% carbon can be produced only when

the cooling rate is greater than 10^ K/s. In the bottom third of the ribbon thickness near the | wheel-side surface, the microstructure consists of uniform subcells of about 0.1 pm in diameter,

having MC-carbide precipitation smaller than 20 nm at the subcell boundaries. With the reduced

cooling rate toward the ribbon top surface, the subcell structure gradually diminishes and MCcarbides ranging from 50 to 100 nm in diameter precipitate only in the intercellular regions. The

MC-carbides remain stable after annealing at lOOO^C for 2 hours, but particles of a second phase

precipitate and coarsen at the cell boundaries during the anneal. This unstable, rhenium-rich precipitation causes reduction in ductility and strength of the annealed ribbon.

Acknowledgments

j We thank R. J. Zabala, R. P. Laforce, L. Perocchi, and the late A. C. Rockwood for their technical assistance. We also thank S. W. Yang, M. R. Jackson, B. Cantor, and L. A. Johnson for stimulating discussions.

References

[1] P. Duwez, R.H. Willens, and W. Element, J. Appl . Phys., 31 (1960) 1136.

[2] T.H. Sanders, Jr., J.W. Johnson, and E.E. Underwood, Proc. 2nd Int. Conf. Rapid Solidification Processing, Reston, Virginia, March 1980, Ed. R. Mehrabian et al., Claitor ' s ,Publ ishing,

p. 141.

[3]

[4]

R. Ray, Proc. MRS Meeting on Rapidly Solidified Amorphous and Crystalline Alloys,

November 1981, Ed. B.E. Kear et al . , North-Holi and, p. 435.

C.C. Wan, as in Ref. 3, p. 441.

Boston,

!’

t.

i

346
J.V. Wood and R.W.E. Honeycombe, Mat. Sci. Eng., 38 (1979) 217.

E.R. Slaughter and S.E. Das, as in Ref. 2, p. 354.

J.J. Rayment and B. Cantor, Proc. 3rd Int. Conf. Rapidly Quenched Metals, Univ. of Sussex,

July 1978, Ed. B. Cantor, The Metals Society, P. 110.

S.C. Huang and H.C. Fiedler, Met. Trans., 12A (1981) 1107.

I.W. Donald and H.A. Davies, Phil. Mag., A42 (1980) 277.

S.C. Huang and A.M. Ritter, to be published in the Proc. AIME Symposium on the Chemistry and

Physics of Rapidly Solidified Materials, St. Louis, Missouri, September 1982.

E.K. Storms, The Refractory Carbides, Academic Press, 1967.

347
HIGH TEMPERATURE OXIDATION OF RAPIDLY SOLIDIFIED Ni-Al-Mo-W ALLOYS

G. H. Meier and F, S. Pettit

University of Pittsburgh

Pittsburgh, PA 15261

A. S. Khan

Pratt and Whitney Aircraft

West Palm Beach, FL 33402

ABSTRACT

The effects of Mo and W additions on the isothermal and cyclic oxidation behavior

of Ni-Al alloys in the temperature range of 760 to 1200°C are reported. The presence

of Mo and W was found to affect the selective oxidation of A1 in two ways. At higher

temperatures and in some concentrations these elements, particularly Mo, had the

beneficial effect of promoting the formation of an external layer of protective

AlgO^ in a similar manner to the well-known effect of Cr additions to Ni-Al alloys.

Under other conditions these elements, particularly W, had the detrimental effect

of promoting internal oxidation of the A1 , apparently by decreasing the alloy

interdiffusion coefficient. In the latter case the accumulation of MoO^ and W0.:j in the reaction product produced catastrophic oxidation rates, particularly at the

higher temperatures. Extensive spalling and large weight losses in cyclic oxidation

were observed wherever significant amounts of Mo- and W- oxides formed. The

modification of the alloys with additions of several percent Cr and Ta and a few tenths of a percent Y and Hf promoted the formation of an adherent, protective

AI

2 O

2 scale over the temperature range studied with corresponding excellent

resistance to both isothermal and cyclic oxidation.

Introduction

The major emphasis of this study is on the effect of Mo and W on the selective oxidation of

A1 from Ni-Al alloys to form a continuous external AI

2 O2 scale and the extent to which Ho and

W can promote catastrophic oxidation. Therefore, a brief discussion of these phenomena will be

presented.

The formation of a continuous external film of an oxide such as AI

2 O

2 as opposed to its

formation as internal oxide particles is essentially determined by a balance between the outward

flux of A1 through the alloy and the inward migration of the internal oxidation front. The

formation of an external scale is favored by increased A1 concentration, a higher alloy interdiffusion coefficient and lower oxygen permeability which can result from decreased Pq in the

atmosphere and the presence of other elements in the alloy which form oxides of high ^2 stability

(e.g. Cr in Ni-Cr-Al alloys). An additional factor favoring formation of a continuous AI

2 O2 film is a slow growth rate of the other surface oxides (transient oxides).

The critical A1 content in binary Ni-Al alloys for external AI

2 O

2 formation has been

determined (1) and found to decrease with increasing temperature. For example, Ni-6A1 is a NiO-former at 900°C but forms an external AI

2 O

2 film at 1200°C. The addition of Cr to Ni-Al

alloys is known to enhance external AI

2 O 0 formation at smaller A1 concentrations but there are

no studies of the effect of Mo or W on^AT202 formation. However, several studies have been

reported for the effect of W on Cr20o formation on Ni-Cr alloys. El-Dahshan, et al., (2)

found that additions of 40 wt.% W to'^Ni-lOCr promoted external Cr202 formation at temperatures

between 900 and 1000°C, presumably due to W acting as an "oxygen getter" in the same manner as Cr in Ni-Al. It was observed, however, that the W additions resulted in ultimate breakdown of

the Cr202 due to W causing a marked decrease in the alloy interdiffusion coefficient so that

the flux'^of Cr to the scale/alloy interface was not sufficient to maintain the Cr202 layer.

Similar effects were reported by Espevik, et al . , (3) for W additions to Ni-22Cr.

348
The oxidation literature contains several examples of catastrophic attack induced by high

refractory metal contents in Fe- and Ni- base alloys. Leslie and Fontana(4) found that Mo- containing Fe-25Ni-16Cr alloys underwent catastrophic oxidation at 900°C in static air even though

the oxidation resistance was good in flowing air. Brenner (5) found that binary Ni-Mo and Fe-Mo

alloys containing as much as 30^ Mo did not oxidize catastrophical ly at 1000° C. Pizzini, et al .

,

( 6 ) obtained similar results for Ni-Mo alloys. Brenner, (7) however, found that a number of

Fe-Ni-Mo and Fe-Cr-Mo alloys did undergo catastrophic attack. This attack was ascribed to the

accumulation of liquid MoO, at the scale/alloy interface. Peters, et al., ( 8 ) found that

Ni-15Cr-Mo alloys had good'^oxidation resistance in flowing atmospheres but Mo contents in excess

of 3% caused catastrophic attack in static atmospheres. A Mo-rich oxide was observed at the

scale/alloy interface suggesting the importance of MoO^ accumulation. Relatively little work

has been done on the effect of Mo or W on the oxidation of a-Al

2 CH -forming alloys, although

catastrophic hot corrosion of such alloys is well documented (9> 'It appears that the Mo and W play an important role in causing this type of degradation.

This paper presents the results of a study of the effects of Mo and W on the oxidation of

Ni-Al alloys prepared both by conventional casting and consolidation of rapidly solidified

powders. Results for the oxidation behavior of similar alloys modified by small additions of

Cr, Ta, Y, and Hf are also presented.

Experimental

The compositions of the alloys studied are listed in Table 1. These include alloys prepared

by conventional drop casting and by consolidation of rapidly solidified powders. Specimens were cut.

Table 1 Nominal Alloy Compositions

Concentrations (Weight Percent)

Preparation

Alloy m_ KL Mo W Cr la Y Hf Technique

Ni-Al Bal6

- - - - - Drop Cast

A Bal. 7 11 - - - - Drop Cast

B Bal. 7 14 - - - - Drop Cast

C Bal

7 17.5 - - - - Drop Cast

JG-185 Bal. 6.8 14.4 6.1 - - - Drop Cast

RSR-185 Bal. 6.8 14.4 6.1 - - - Rapid Sol

,

RSR-467 Bal

7.1 11.4 6.3 - - - Rapid Sol

Modif ied-467 Bal. 6.7 11 6 3 to 4 1.5 0.1 0.1 to 0.2

polished through 600 grit SiC, ul trasonical ly cleaned in ethanol, and dried. The oxidation

experiments were carried out in flowing air at temperatures between 760 and 1200°C in either a continuously-weighing microbalance or an automatic thermal cycling apparatus. The oxidation

morphologies were examined by optical and scanning electron metallography and X-ray diffraction.

Results and Discussion

The composition of the alloys was found to influence their oxidation behavior to a greater

extent than the preparation technique. Figure 1 compares the rates of oxidation at 1150°C

in flowing air. These curves are characteristic of the behavior of the alloys in the highest

temperature range studied (1100-1200°C) . The rates for RSR-185 and JG-185 are initially rapid

and then level off. However, this plateau is partially due to the evaporation of MoO^ from the

specimens as evidenced by condensation of M0 O

3 crystals in cooler parts of the reaction tube.

After times on the order of fifty hours the curves bend up sharply giving oxidation rates

faster than those for an NiO forming alloy such as >li- 6Al. The corrosion morphology typical

of RSR-185 and JG-185 is shown in Figure 2 for JG-185 exposed for 30 minutes at 1200°C.

The scale is composed of a dense outer NiO layer beneath which is a 5 ym thick layer

containing only Ni and Mo and which is indicated by X-ray diffraction to be NiMoO^.

Beneath this layer are the fragmented remnants of a layer which is an intimate mixture

349
Figure 1 Oxidation rates for several cast and RSR alloys in air at 1150°C.

of Al^O^, NiAl^O^, NiMoO^, NiWO^, NiO and, perhaps, MoO^. The dark internal oxides are AI

2 O

2 and tne'^bright^layer at the scaTe/metal interface is rich in Mo and contains significant amounts

of A1 . After 1 hour (Fig. 3) it is seen that the growth of the external scale has consumed

the internal oxidation zone leaving a ragged scale/metal interface. The X-ray maps of Figure 3 indicate the formation of an AI

2 O

2 film over some portions of the scale/metal interface. This

layer is not able, however, to completely seal the surface and the corrosion sequence repeats

itself. Observation of specimens exposed for longer times indicates that this repetition

continues to occur.

The preceding metal lographic results suggest the mechanism presented in Figure 4 for the

degradation of RSR-185 and JG-185 in the temperature range above 1100°C. Initially, (t,) an external NiO scale forms with internal AlgO^ precipitates in the alloy. As indicated by the

approximate values of the equilibrium oxygen partial pressures, the dioxides of Mo and W can form below the NiO but the trioxides are unstable. With continued reaction (t

2

) the dioxides of

Mo and W react with the NiO to form NiMoO^ and NiWO.. At this stage the external scale is analogous to that found on binary Ni-Mo alloys. (5,6) However, for the A1 -containing alloys

the trioxides of Mo and W form (t,) producing a liquid oxide. This, of course, requires the

Ppi at the scale/alloy interface to exceed the dissociation pressures of NiO. The evaporation of

2 MoO^ was observed in the present experiments indicating that the outer NiO scale cracks since

the NiO scales formed on binary Ni-Mo alloys are apparently impermeable to MoO^. (5,6) This

cracking could be the cause of the p^, becoming high enough to form the 1 iquid"^tri,oxides . On the

other hand, if the Ni activity below^2 the scale decreases much below 0.1 the trioxides could

become stable. In this case the liquid formation could be the cause of the scale disruptions.

Once the liquid forms the subsequent corrosion is similar to that observed in some hot corrosion

reactions. The liquid preferentially attacks the AI

2 O 0 or the interface between the Ni and AI

2 O 0 This leads to the incorporation of the internal Al20^ particles and islands of Ni into the liquid

corrosion product (t^). At the same time some areas'^of the alloy are forming continuous AI

2 O

2 films and the outer ^NiMoO^ layer is consumed by MoO^ evaporation.

350
Figure 2 Cross-section of cast Ni-7A1 -14Mo-6W exposed for 30 minutes in air at 1200°C.

The never able to become complete and are continally stripped and incorporated

into the liquid corrosion product (tc). At this stage the outer scale is NiO and below this is

a layer of NiO which contains Al^O^ 3 nd NiAl 20^. This layer forms by the evaporation of MoO^

from the liquid corrosion product. Below the ^NiO layers is a layer containing NiO, AI

2 O

3

, NiMoO^, and NiWO^, and NiO being formed by the oxidation of the Ni islands. At this stage the

internal oxides of A1 again form in the alloy substrate and the above mechanism goes through

another cycle. The more rapid oxidation rate for the cast alloy compared to the RSR alloy is

believed due to a less homogenous distribution of Mo and W in the former.

351
Figure 3 Cross-section of cast Ni-7Al-14Mo-6W exposed for 1 hour in air at 1200°C

Several additional observations support the above mechanism. One specimen of RSR-185 was^

exposed at 1150°C in argon in which the residual oxygen partial pressure was approximately 10~°

atm. The reduced oxygen pressure caused the NiO to grow much slower than in air and allowed a protective layer of AI

2O 0 to seal the surface and prevent the formation of significant amounts

of Mo-oxides. A similar'^effect was observed when 6 weight percent Cr was added to the alloy, even for exposures in air. In this case the Cr acted as a "secondary getter" and allowed a protective AI

2 O

2 scale to seal the surface.

Another supporting observation is the slow oxidation rate of the lower-Mo RSR-467 (Fig. 1).

This alloy develops a thin Al-rich scale and exhibits an oxidation rate comparable to Ni-18Cr-6Al

Apparently, the Mo and W behave as oxygen getters in a manner similar to Cr to promote an AI

2 O2 scale while the lower Mo concentration prevents sufficient accumulation of Mo-oxides to initiate

352
Air

ti Mo02(s) WO2

AI2O3.

NiO

Alloy

c:

- PO2 * 0-2 atm

--PO2

’®' 10’®atm

'PO2 10'^^ atm

- Pq2 J=10‘^®atm

PoMoO2/MoO3^,0'®atm

Ai r *2 n7o

TJiMot 4 \ \ V \

W02‘s)

A I2O3—

nr

. ^A^S!f°NV'o~ o'V'o T"o

^ NIMo04 _\ X, O\0 \t^ 0

+AI2O3+ NiO

AI203-^ /y UW [f

Figure 4 Model for the rapid oxidaton of Ni-Al-Mo-W alloys

the rapid attack observed for RSR-185. Once the Al^O^ layer becomes continuous the Pq below the

scale remains too low for the oxidation of Mo or w/ " The maximum in the weight ^2 change

for RSR-467 after 1 hour is due to the evaporation of transient Mo-oxides after the AI2O0 layer

has covered the alloy surface.

The exposure of ternary Ni-Al-Mo alloys has indicated a difference in the behavior of Mo and W. For example, the Ni-7Al-17.5Mo alloy has the same total refractory metal content (in terms of mole

fraction) as the Ni-7Al-14Mo-6W alloys but shows much slower oxidation rates (Figure 1). This

alloy was also observed to form a continuous (or nearly continuous) AI2O2 scale. Thus it appears

that Mo can effectively act as a secondary getter to promote external

^ AI2O3 formation since it

forms oxides with stabilities lying between those of NiO and AI2O3. The stabilities of the

oxides of W are similar to those of Mo which indicates W could also promote external AI2O3 formation.

The reason that it does not is believed to be a reduction in the alloy interdiffusion coefficient.

353
Diffusion data are not available for Ni-Al-W system but the interdiffusion coefficient in binary

Ni-W alloys decreases with increasing W concentration. (10,11) The diffusion coefficient of A1 in Ni-Mo alloys has been observed to decrease slightly with increasing Mo concentration (12).

However, the effect of Mo on the diffusion coefficient is apparently less than that of W.

An additional feature regarding the oxidation mechanism is that although considerable M0 O

3 evaporation was observed this evaporation does not seem to have either beneficial or detrimental

effects on the rapid oxidation process as the corrosion appears to be flow-rate independent.

Experiments run with inverted AI

2 O

2 crucibles covering the specimens produced the same amount of

attack as those run in flowing gas; This observation is different from that for Ni-Cr-Mo ( 8 ) where severe corrosion was only observed in stagnant atmospheres.

The oxidation behavior of the Mo-containing alloys was somewhat different at temperatures

below 1000°C. The rates of oxidation at 850°C are compared in Figure 5. All of the Mo-containing

alloys are NiO-formers and oxidize more rapidly than the Al 202-former Ni-18Cr-6Al, included in

Figure 5 for comparison. RSR-467 and Ni-6A1 oxidize at comparable rates and, exhibit similar

oxidation morphologies consisting of an external NiO scale and internal AI

2 O2 precipitates which

form continuous films in some locations.

Figure 5 Oxidation rates for several cast and RSR alloys in air at 850° C

The cast alloy JG-185 and RSR-185 oxidized significantly faster than Ni-6A1. This is due

primarily to profuse internal oxidation of these alloys as indicated in Figure 6 . A significant

difference from the high temperature behavior of the same alloys is that the internal oxides are

rich in Mo and W as well as A1 at low temperatures. These oxides form continuous networks allowing

rapid ingress of oxygen into the alloy. The oxides of Mo eventually begin to concentrate at the scale/alloy interface during lower temperature exposures, as indicated in Figure 6 , and once NiMoO^ begins to form significant spalling of the scale occurs on cooling. This is a significant

observation since the scales remain adherent in essentially all cases where NiMoO^ does not form.

354
The reason that this phase promotes spalling is the occurrence of a phase transformation at about

650°C with an accompanying large volume change (5,6) which promotes shattering even in bulk

NiMoO^. (This effect was also observed in the current work when NiO and MoO^ were heated together

and formed NiMoO^). Therefore, it is clear that alloy compositions or coatings which preclude

the formation of^NiMoO^ must be used for this class of alloy to have resistance to cyclic oxidation.

Figure 7 shows the cyclic oxidation behavior of RSR-185 over a range of temperatures. The

alloy shows significant weight losses at the higher temperatures whereas at 760°C it has continued

to gain weight after 100 hrs. However, even at this temperature, extending the time of oxidation

or shortening the cycle length will produce weight losses. Figure 8 is a typical microstructure

of a cycled specimen showing significant internal oxidation and Al-depletion.

Thermal cycling also has a significant effect on the oxidation behavior of the more resistant

alloys such as RSR-467 as indicated in Figure 9. After about 150 hours weight losses are observed which increase with increasing exposure time.

Modification of RSR-467 by small additions of Cr, Ta, Y and Hf results in an alloy which

forms an adherent AI

2 O 0 layer over the entire temperature range studied. This is illustrated

in Figure 9 for 850°^C'^and in Figure 10 for 1150°C.

355
6

“ IW i I I I I 0 20 40 60 80 100

TIME (hours)

Figure 7 Cyclic oxidation of RSR-185 in the temperature range 760 to 1150°C. Cycles

consist of 25 hours heating, 1 hour cooling.

Figure 8 Cross-section of RSR-185 exposed in cyclic oxidation for 200 hours at 870 °C.

356
Figure 9 Cyclic oxidation of RSR-467 and modified -467 alloys at 850°C. Cycles

consist of 50 minutes heating, 10 minutes cooling.

357
2-1

Cr, Ta, Y, and Hf Modified Alloya

100 200

TIME (hours)

Figure 10 Cyclic oxidation of modified 467 alloys and RSR-185 at 1150^ C.

Summary

300

The oxidation behavior of several Mo- and W-containing Ni-Al alloys has been studied over a range of temperatures. It was found that for temperatures above 1100°C the presence of Mo and W can promote the formation of a protective AI

2 O

2 scale in a manner similar to Cr additions. However

if the Mo concentration exceeds a critical value the formation of liquid oxides results in rapid

corrosion. This critical concentration lies between 11 and 14 wt. percent for rapid solidified

alloys containing about 7 percent A1 and 6 percent W. However, this value is probably lower for

less homogeneous, cast alloys and is lower if the alloys are exposed under cyclic conditions. The,

mechanism for the rapid corrosion has been described. The oxidation of W-free alloys has indicated

that Mo concentrations in excess of 17 wt. percent can be tolerated under isothermal conditions

without producing severe degradation. The detrimental influence of W is believed to be due to a reduction in the alloy interdiffusion coefficient.

I

"3

5

6

.

I

At temperatures less that 900°C all of the unmodified Ni-Al-Mo-W alloys were NiO formers.

Apparently, the decreased alloy interdiffusion coefficient due to the W in solution negates ; the oxygen-gettering effect of W and Mo and prevents the formation of a continuous Al^O^ scale.

The alloys containing the higher Mo concentration oxidize faster than Ni-6A1 due to tne'^profuse

formation of internal oxides of A1 , Mo, and W. The scales remained adherent on cooling until j NiMoO^ was formed. When this occurred extensive weight losses were observed during cyclic oxidation

The addition of Cr to Ni-Al-Mo-W alloys was observed to promote the formation of protective I AI

2 O 0 films. The modification of the Ni-Al-Mo-W alloys with additions of Cr, Ta, Hf, and Y 1 resulted in the formation of adherent, protective AI

2 O

2 scales which provided excellent cyclic| oxidation resistance over the entire temperature range"^studied. 3

358
References

I 1. F. S. Pettit, Trans. AIME, 1296 (1967).

2. M. E. El-Dahshan, D. P. Whittle, and J. Stringer, Corr. Sci . , 1^, 83 (1976).

3. S. Espevik, R. A. Rapp, P. L. Daniel, and J. P. Hirth, Oxid. Met., 14, 85 (1980).

4. W. C. Leslie and M. F. Fontana, Trans. A. S. M. , H, 1213 (1949).

5. S. S. Brenner, J. Electrochem. Soc., 102 , 7 (1955).

6. S. Pizzini, E. Ruedl , and L. Merletti, "Oxidation of Nickel -Molybdenum Alloys:

Thermogravimetric and Structural Analysis" in Metal-Slag-Gas Reactions

and Processes , Z. A. Foroulis and W. W. Smeltzer, eds., Electrochem. Soc.,

Princeton, N. J., 1975, p. 448.

7. S. S. Brenner, J. Electrochem. Soc., 102 , 16 (1955).

8. K. R. Peters, D. P. Whittle, and J. Stringer, Corr. Sci., 16, 791 (1976).

9. J. A. Goebel, F. S. Pettit, and G. W. Goward, Met. Trans., 1, 261 (1973).

10. J. M. Walsh and M. J. Donachie, Met. Sci. J. 1, 68 (1969).

11. V. N. Pimenov and V. E. Dgaste, Fiz. Metal. Metalloved., 590 (1973).

12. J. E. Bridge and G. N. Maniar, Met. Trans., 1, 1006 (1972).

ACKNOWLEDGEMENTS

The authors grateful ly acknowledge the contributions of D. Meikle, K. Andrew, E. Hewitt and C. Martin

in performing the experiments done at the University of Pittsburgh. The authors also wish to thank

R. J. Hecht and J. A. Miller for their constructive discussions and support.

This work is jointly sponsored and monitored by the U.S. Defense Advanced Research Projects Agency

and the Air Force Wright Aeronautical Laboratories

The views and conclusions contained in this paper are those of the authors and should not be interpreted as necessarily representing the official policies, either expressed or implied, of the U.S.

Defense Advanced Research Projects Agency, U.S. Air Force, or the U.S. Government.

359
RAPIDLY SOLIDIFIED CARBIDE POWDERS PRODUCED BY SPARK MACHINING OF Ti AND Zr

C. S. Pande and J. D. Ayers

Physical Metallurgy Branch, Naval Research Laboratory, Washington, DC 20375 USA

ABSTRACT | Fine rapidly solidified TiC and ZrC powders were obtained from commercial purity Ti and i iodide processed Zr by spark machining in kerosene, the carbide being formed because of a I reaction between the metal and carbon from kerosene. Powders produced by this process were

j studied by x-ray diffraction, and by scanning (SEM) and transmission electron microscopy l (TEM). Results from these studies show that depending on the processing condition in spark j machining, the resultant powders are either single phase carbides or contain, in addition,

unreacted metal or free graphite. Examination of the spark machined surfaces by x-ray I diffraction and by SEM revealed that they were covered by a carbide layer and that the ,i degree of carburization correlated with the carbon concentration of the spark generated 'i powders. ;

Single phase TlC-ZrC alloy powders were also produced by spark machining of the

metallic alloys. Heat treatment of these single phase powders produced a phase separation

into TiC and ZrC rich phases both with NaCl structure but with different lattice parameters.

p

r This study has thus demonstrated that equilibrium and metastable structure carbide ^ powders ranging in size from about 40 ym down to submicron sizes can be produced by this r technique. The rate of powder production was only a few grams per hour, but the unique | microstructures produced may make these powders technologically interesting.

^

I

I Introduction

|

I

It is well known that rapid solidification processing may result in materials with microstructure | and chemical homogeneity very different from that obtained by conventional means. The novel | properties are determined largely by the transport processes occurring at the liquid/solid interface.

However, a detailed study of the basic mechanism of solidification and its correlation with resultant • microstructure is needed to fully utilize the potential benefits.

j It was realized as early as 1924 [1] that spark machining generates carbide, oxide or nitride

^ layer resulting in a surface hardening effect. This effect has also been employed commercially. ^ However, the nature of the materials removed by the spark machining has received much less attention,

j If the dielectric is a hydrocarbon, the reaction products are metal carbides. These carbides are j produced not only on the metal surface, but are ejected from the surface into the dielectric. They are usually in the form of fine roughly spherical powders, of the size ranging from less than a micron in

diameter to up to 40 ym. These carbide powders are obviously produced by a solidification mechanism

similar to that in other rapid solidification processing. The detailed technique of the formation of , such metal powder carbides have been published elsewhere [2]

The present paper is concerned mostly t with the microstructure obtained by this technique.

i

I Formation of Carbide Powders

I The conditions under which these powders were produced are given in Table 1. After machining

with a graphite tool with kerosene as the dielectric, carbide powders were collected from the

dielectric, and washed to remove organic residue. The powders so obtained were examined by x-ray >

360
diffraction and by scanning (SEM) and transmission electron microscopy (TEM). Energy dispersive x-ray

analysis and electron energy loss analysis has been used to analyze Ti (and Zr) and carbon content,

respectively.

Table 1

Metals used in this study: Commercial purity Ti; Iodide processed Zr

Electrical discharge instrument: AGIE Type AB 30H which generates a square wave of variable frequency

Frequency of electric discharge: 400 to 25000 hertz using a 50% duty cycle

Dielectric

Kerosene

Cutting tool: Graphite (rotating)

Polarity of the discharge: Workpiece both positive (conventional) and

negative

Voltage of the discharge: 220 Volts

Further details are given in reference 2

Results

Titanium Carbide Powders

A low magnification micrograph obtained by scanning electron microscope (SEM) is shown in Fig. 1. The metal powders were seen to be nearly spherical, though a few particles did show odd geometries.

The size ranged from submicron to up to about 40 ym. X-ray examination of these powders showed strong

Tie peaks (with NaCl structure) as expected, but it also contained weak lines which were identified as due to a small amount of a phase Ti. How is Ti distributed in a given TiC particle? This was investigated by etching the powders for several minutes or longer in Kroll's reagent (5 parts HF, 10 parts HNO3 rest water) which dissolves Ti and Zr but not their carbides. The etchant attack is very

slow, due to the fact that the outer layer of the particles are more likely to be TiC, or coated with

TiC. However, where the attack was deep enough, the Ti metal was etched away revealing a TiC

dendritic structure (see Fig. 2). It was indicated that individual powder particles have probably

different volume fraction of TiC.

These powders were then examined in TEM. For this purpose the powders had first to be thinned.

The following technique of thinning gave consistently good results. The metal powders were sandwiched

between two thin sheets of vanadium (~5 mil thick) and cold-rolled. 3 mm disc were then punched out and thinned in the middle using ion thinning (argon ions at 5 KV) till perforation. Carbide particles

were found to be electron transparent near the perforation. Fig. 3 is the TEM micrograph TiC particle

showing the various grains inside a single TiC powder. Many micrographs indicated that the outer edge

of the particle is titanium carbide; inside the particle, however, metallic Ti is present. The SEM

and TEM, and x-ray results were thus consistent.

Attention was next given to producing TiC powders which have much less metallic inclusions. This

led to the use of the 'reversed polarity' for producing these powders. Spark machining is normally

done with the workpiece positive relative to the tool, because electron bombardment of the metal in

this mode leads to faster cutting rate. If the polarity is reversed the metal will be bombarded with

positive ions (carbon ions if kerosene is used as a dielectric). This should lead to better

carburization as was borne out by the experiments. The TiC powders, obtained by the reversed polarity

were not attacked visibly by the etchant. More significantly, the measured lattice parameter

indicated it to be almost pure TiC. Finally, TEM again showed that these particles were probably pure

Tic. A few of them were found to be single crystals (see Fig. 4).

The other processing variable is the frequency of the spark discharged. This parameter was varied and the TiC content of the specimens were ascertained using x-rays. The amount of a phase Ti

found in these powders decreased with decreasing frequency.

361
Figure 1. Titanium carbide powder prepared by spark machining of Ti in kerosene.

(a)

Powder obtained at 6000 Hertz with normal

polarity

(b)

Powder obtained at 12000 Hertz with

'reverse' polarity.

Figure 2(a) . A titanium carbide powder produced

by machining Ti with normal polarity at 6000

Hertz and then etched in Kroll's reagent which

attacked Ti but not TiC. Metallic titanium has

been removed revealing TiC dendrites.

Figure 2(b) . A titanium carbide powder obtained

under the same processing conditions as (a)

,

and

etched similarly. The dendrites in this case

are coarser, indicating a higher carbon content.
Figure 3. A TEM micrograph of a TiC particle obtained with normal

polarity showing grains in the particle and some free Ti.

Zirconium Carbide Powders

As is expected from the similarity of Ti-C and Zr-C phase diagrams the electron machining

behavior of Zr was similar to that of Ti. Many ZrC particles showed surface structure similar to TiC.

Some particles exhibited a typical surface morphology. TEM revealed that those surface features (many

small irregular patches) are each the base of elongated grains growing inward in the powder particle.

TiC-ZrC Alloy Powders

The technique of using these powders has been extended to powders of alloys of titanium and

zirconium also. Single phase TiC-ZrC has been produced. Heat treatment of these single phase alloy

leads to a phase separation into TiC and ZrC rich phases, both with NaCl structure but with different

lattice parameters.

Conclusions

Spark machining was found to be a simple and convenient tool to study rapid solidification

processes in metal carbides. With suitable processing conditions (reversed polarity and low frequency

discharge), single phase TiC can be produced. This study has also demonstrated that spark machining

provides a convenient tool for investigating the basic mechanism of rapid solidification in metal

carbides

363
Figure 4. A TEM micrograph of a TiC particle obtained with reverse polarity showing a single crytal TiC particle. The dislocation densities in these particles rarely exceeded

10^/ cm^

Acknowledgments

The authors are grateful to Miss Kathy Moore and Mr. Steve Smith for their contribution to this

work.

References

[1] Rawdon, H. S., AIME Trans., Vol. 70, 1924, p. 37 (Discussions).

[2] Ayers, J. D. and Moore, K. , (to be published).

1

364
PROCESSING, MICROSTRUCTURE AND PROPERTIES OF MELT

SPUN AND CONSOLIDATED SUPERALLOYS

R.H. VanStone*, J.R. Hughes**, R.E. Maxwell***,

H.H. Liebermann , and R.G. Rowe** 

*General Electric Aircraft Engine Business Group, Cincinnati, Ohio 45215

**General Electric Corporate Research and Development, Schenectady, NY 12301

***General Electric Corporate Research and Development, Schenectady, NY 12301

currently with General Electric Lamp Glass and Components, Cleveland, OH 44110

****Allied Corporate Technology, Parsippany, New Jersey 07054, formerly with General

Electric Corporate Research and Development, Schenectady, NY 12301

ABSTRACT

The single roller chill block melt spinning process, commonly used to produce amorphous

alloy ribbons, was modified to produce rapidly solidified 20-30 micrometers thick ribbon of

the nickel-base superalloys Rene'95 and the developmental alloy 1200/19. These ribbons had

cellular or columnar microstructure which were devoid of branched dendrites. Scanning transmission electron microscopy analysis showed no significant microsegregation. Several pounds

of narrow and 11 mm wide ribbon were produced.

Compacts of randomly packed narrow ribbon and wound spools of wide ribbon were hot isostatically pressed to full density. The randomly packed ribbon compacts were extruded to improve their toughness. After both HIP compaction and HIP plus extrusion, the prior ribbon

boundaries contained a relatively high density of aluminum and zirconium oxides and MC-type

carbides

The mechanical properties of consolidated Alloy 1200/19 were evaluated in both the HIP

plus heat treated condition and the HIP plus extruded plus heat treated condition. The tensile,

creep, rupture and fatigue crack growth properties of the HIP plus extruded ribbon were similar

to those obtained for rapidly solidified, helium atomized HIP-compacted powder of the alloy.

INTRODUCTION

The continuing demand for improved performance in aircraft gas turbine engines having longer

operating lives and improved reliability has encouraged the development of new turbine disk alloys

which are capable of operating at higher stresses and higer temperatures than conventional alloys.

A major advancement in enhancing the temperature capability of such alloys was the development of

nickel -base superalloys which are manufactured by powder metallurgy (PM) methods

This technology

allows more efficient material utilization and reduces mechanical property variations resulting from

compositional segregation and difficult forging operations associated with conventionally cast and

wrought superalloys. The reduction of segregation in j-y' superalloys has permitted the use of

higher y' volume fractions and higher levels of solid solution strengthening elements.

The next logical step in the quest for performance of superalloy disks is through the use of more advanced rapid solidification technology. One such rapid solidification process is chill block melt

spinning (CBMS), which has been used to fabricate amorphous alloy(^"^) and superal loy(^»^) ribbons.

The CBMS process, shown schematically in Figure 1, consists of a thin stream of liquid alloy which is

ejected under pressure from a nozzle onto a rapidly rotating substrate wheel which removes heat by con- duction. The simultaneous mass and heat transfer produces thin, nickel-base superalloy ribbon solidified at rates in excess of those obtainable using conventional gas atomization. This process provides

high yields of rapidly solidified material with uniform thickness and more uniform microstructures than

material obtained in rapidly solidified powders^^^).

This paper describes the development of melt spinning process parameters and compaction procedures

used for the production of advanced nickel-base superalloys. Selected mechanical properties of the com- pacted CBMS materials were also determined. These studies were performed using Rene'95 and experimental

Alloy 1200/19. Rene'95 is an advanced y-y' alloy which has been used in aircraft engine disks and other

high temperature, high stress applications. Alloy 1200/19 is a developmental alloy designed to simultaneously increase temperature capability and cyclic defect tol erancevl ^ ) . The compositions of these

alloys are listed in Table 1. Alloy 1200/19 has a y' content of 38 percent compared to 45 percent for

365
Rene'95. The y' content of Alloy 1200/19 was reduced primarily by reductions in aluminum and columbium.

To offset the loss in tensile, creep, and rupture strength associated with lower y' content, the amount of solid solution strengthening was increased by raising the tungsten content and adding rhenium.

Table I. Nominal Compositions of Rene'95 and Alloy 1200/19

Alloy Composition (Weight Percent)

Co Cr W Re M 11 Cb Zn i B

Rene' 95 8.0 13.0 3.5 3.5 -- 3.5 2.5 3.5 0.05 0.05 0.01 Balance

A1 loy 1200/19 14.0 17.0 — 8.0 2.0 1.75 3.5 0.65 — 0.02 0.01 Balance

Melt Spinning of Superalloy Ribbons

Davies, et. al.(®) and Wood, et. al.^^^ used chill block melt spinning to produce relatively small

volumes of nickel-base superalloys. The CBMS process used to fabricate the large volumes of ribbon re- quired for the mechanical property evaluation of the percent investigation was described in previous

publications ^

' and will be summarized here.

The quench rates achieved during rapid solidification are controlled primarily by the

I ( 1

, 1

) thickness of

the sample and the nature of interfacial contact between the wheel and the specimen^ '

Calculation,

by Ruhl(15)have shown that the average quench rate varies inversely with the square of specimen thicknes

for ideal cooling or inversely with specimen thickness for Newtonian interface-controlled cooling.

Therefore, careful control of the ribbon thickness is required to produce uniform rapidly solidified

microstructures . Ribbon thickness is primarily controlled by the liquid flow rate through the nozzle,

the ribbon width, and the wheel surface velocity. Liebermann, et. al.^*^^ demonstrated that the fluidribbon thickness relationships developed for amorphous alloys also apply to superalloys. Subsequent

experiements determined ribbon thickness as a function of wheel surface velocity for fixed mass flow

rates. The variation ribbon thickness as a function of substrate velocity is shown in Figure 2 for

Rene'95 melt spun on OFHC cooper and 304 stainless steel wheels. Ribbon thickness is strongly dependent on substrate velocity (V ) at low values of V ; however, attempts to obtain ribbon thickness significantly smaller than 20 microns by increasing wheel velocity will probably not be effective, as shown in Figure 2. Similar CBMS ribbon thickness-wheel velocity data have been obtained for amorphous

alloys (^6) and Alloy 1200/19(^2)^

The microstructures of Rene'95 CBMS ribbons were characterized using optical metallography and

scanning transmission electron microscopy (STEM). Direct evidence for the decreasing solidification

rate with increasing specimen thickness is shown in Figure 3, which shows Nomarski interference micrographs of ribbons with thicknesses of 16 and 52 microns. In these micrographs, the ribbon surface

which was in contact with the rotating wheel is shown at the bottom of each micrograph. Figure 3a shows the presence of branched, secondary dendrites near the free surface of the thicker ribbon. This

is direct evidence of reduced solidification rate with increasing ribbon thickness. The microstructure

of the 16 micron thick ribbon as shown in Figure 3 b is devoid of these branched, secondary dendrties.

It was concluded from several such observations that ribbon thicknesses less than 30 microns were sufficient to suppress the branched dentritic microstructure for the CBMS process used in this investigation. As shown in Figure 2, this thickness can be readily obtained using the melts spinning process.

All materials produced for evaluation of mechanical properties were sufficiently thin to suppress the

formation of secondary dendrite arms.

EJECTION

GAS

PRESSURE

ROTATING

SUBSTRATE

WHEEL

MELT CRUCIBLE

HEATING

SYSTEM

FREE MELT JET

MOLTEN ALLOY

PUDDLE

ICHILL BLOCK^^M

(CBMS)

Figure 1. Schematic Diagram Showing the Chill

Block Melt Spinning (CBMS) Process

Figure 2. Variation of CBMS Rene'95 Average

Ribbon Thickness with Substrate

Surface Velocity

366
j STEM was used to study phases and segregation in CBMS ribbons having several different thicknesses,

i A bright field TEM micrograph from the midplane of a 30 micron thick ribbon is shown in Figure 4a.

I This microstructure consists of small, equiaxed grains 0.2 to 0.6 microns in diameter separated by

I high angle boundaries containing second phase particles. These high angle boundaries are not the low

I angle boundaries usually associated with cellular instability of the liquid-solid interface. This

observation suggests that cellular solidification instabilities were suppressed in this region.

I Selected area diffraction of grain boundary regions was used to identify the grain boundary particles

! as MC-type carbides. Figure 4b shows a dark micrograph of the region shown in Figure 4a imaged with

I a MC carbide diffraction spot confirming their identification. STEM analyses of the carbide suggests

I that these carbides are primarily (Cb, Mo, Ti) c(^2)^

I It should be noted that other areas of this ribbon showed a tendency towards a more cellular

I microstructure^^ . In these regions, the grains tended to be larger; this presumably resulted because

the columnar grains were farther from their nucleation sites and therefore wider. Cell boundaries within the grains also contained MC carbides.

Microsegregation near cell, grain and dendrite boundaries was evaluated using STEM. All measurements were performed using a Hitachi H-600-1 electron microscope operating at 100 KV. Analysis of cell

wall and dendrite arm segregation was performed using beam spot sizes of approximately 100 a and 200

respectively. Due to scattering in the matrix and stage drift, the actual size of the regions analyzed

Figure 4. Thin Foil TEM micrographs of CBMS Rene'95 Showing (a) Bright Field, (b) Dark Field

of a Region of Equiaxed Grains.

367
-300 - 200-100 0 100 200 300 DISTANCE FROM BOUNDARY, nm

Figure 5. Results of STEM Compositional Profile (b) Across Grain Boundary (a) in CBMS Rene'95,

were approximately 250 A and 400 A, respectively. Variations in STEM specimen thickness, the absence

of exact multi-component standards, and differences in absorption characteristics make identification

of exact local compositions in complex superalloys like Rene'95 extremely difficult. As a result, the

spectra will be presented as intensities normalized to NiK^^ for a qualitative indication of composition.

No attempts were made to further quantify the results.

25

24

23

14 o° 13

The results of a STEM profile across the boundary between grains of a 15 micro thick sample are shown in Figure 5. This foil was taken near the free surface of the ribbon, the most slowly cooled

portion of the material. The superimposed line on the electron micrograph in Figure 5a shows the

location where the STEM trace was performed. The results of the STEM analysis are shown in Figure 5b. There was little observable segregation on the scale of the STEM probe.

Results of other probes occasionally showed increases in columbium, molybdenum, and titanium, the

strong carbide-forming elements, along the grain or dendritic boundaries. This result is presumably

due to MC carbides in the boundaries, such as those shown in Figure 4. No conclusive evidence of

segregation was observed in specimens with thicknesses up to 37 microns. A 100 micron thick Rene'95

ribbon was melt spun to evaluate the amount of segregation present in interdendritic regions. This

ribbon has a dendritic microstructure with a secondary dendrite arm spacing of approximately 1 micron.

The results of the STEM analysis of this sample are shown in Figure 6. The electron micrograph of the

region analyzed is shown in Figure 6a. The dashed line is the interdendritic boundary inferred from

the carbide distribution and the solid line shows the location of the STEM scan. The STEM results are shown in Figure 6b, where the dashed line is the position of the inferred interdendritic boundary. The'

titanium and columbium contents show a much larger increase than that previously shown in Figure 5. The gradient, shown in Figure 6, extended over a distance of approximately 0.75 micron and the segregation level .is expected to be approximately two times that of the bulk composition. Previous diffusion

calculations^^^) for Rene'95 have shown that during 1800°F HIP compaction, chemical segregation can be

homogenized over a distance of approximately 1.75 microns. The increased compaction temperatures used i to produce the mechanical test specimens in the present study and the reduced ribbon thickness which

prevented formation of branched dendritic microstructures would reduce the amount of segregation present

even further. These STEM results confirm that ribbon with thicknesses less than 30 microns will have

extremely limited micro-segregation in the matrix.

HIP Compaction of CBMS Ribbon

Two different forms of CBMS ribbon were compacted using a hot isostatic press (HIP) cycle. The j first form consisted of a narrow ribbon which was approximately 3 mm wide and 12.5 to 25,0 cm long.

The narrow ribbon segments were randomly packed and pressed into HIP containers. The other form of i CBMS ribbon studied in this program was approximately 11 mm wide having lengths ranging from 0.02 to | 25 m. These segments were spot welded together and subsequently would under light tension around a mandrel and placed in a HIP container.

368
Dendritic Microstructure.

The HIP containers of both forms of CBMS material were degassed and sealed. The HIP cycle

selected for the Rene'95 ribbon was a 3-hour, 2200°F exposure at 15 ksi. This temperature is approximately 65°F above the y' solvus temperature of Rene'95. This type of HIP cycle increases the healing

of voids by creep and diffusion bonding .Super y' solvus temperature have been shown to enhance the

cyclic defect tolerance for Rene'95v- ^ and AF115 n 8 )^ /\ total of 15 Rene'95 containers were HIP compacted during this program to densities equivalent to those obtained for PM Rene' 95 (* 2 )^ Another

important measure of the quality of HIP compaction is thermally induced porosity (TIP). TIP is the

porosity generated within the compacted material from entrapped gases during elevated temperature

exposure. The TIP values of the CBMS ribbon compacts as measured by a change in density after the

thermal exposure fell within the normal limits for PM Rene' 95 (^2)_

The microstructures of both the narrow, randomly packed and the wide, wound ribbons were characterized using optical metallography and transmission electron microscopy (TEM). Figure 7 shows optical

micrographs of both product forms taken from metallographic sections perpendicular to the axis of the

HIP container. Both materials have a grain size of approximately 25 microns and show no evidence of

lack of bonding; however, the prior ribbon boundaries are very heavily decorated in these microstructures. Comparison of the micrographs suggest that the ribbon thickness in the randomly packed specimen

was significantly greater than the wound ribbon, but this was not the case. In the compact containing

randomly packed ribbon, the ribbon thickness direction will tend to be aligned perpendicular to the

metallographic surface in Figure 7a. Thus, a metallographic plane would section the ribbon width

resulting in a large apparent spacing. In the wound ribbon micrograph in Figure 7b, the metallographic

plane is perpendicular to the ribbon width resulting in transverse viewing the ribbon so that the minimum ribbon thickness is observed.

The low magnification micrographs in Figure 7 appear to indicate that grain growth is impeded by

the prior ribbon boundaries. This was generally the observed situation, but higher magnification

optical metallography and thin foil TEM has shown occasional grain growth across prior ribbon boundaries

(PRB).

The nature of PRB regions in Rene'95 CBMS ribbon were studied using TEM extraction replicas. A

PRB region of HIP-compacted CBMS Rene'95 is shown in Figure 8 . The PRB region consists of second phase

particles and coarse . The second phase particles were identified, using electron diffraction techniques, as aluminum oxides, zirconium oxides, and MC carbides. Similar observations have been made for

Alloy 1200/19 CBMS ribbons. The source of these oxides is not fully known at this time. Eight oxygen

content measurements of the Rene'95 compacts were in the range from 67-227 ppm. Two of these measurements were outside the specification limits (150 ppm) of Rene'95 PPM product. Ritter and Henry(21)

have documented oxide and carbide precipitation on prior particle boundaries (PPB) in HIP-compacted PM Rene'95, but the density of boundary decoration in PPB was much lower than that shown in Figure 7 and

8 in CBMS ribbon.

369
Figure 7. Optical Micrographs of (a) Randomly Packed and (b) Wound CBMS HIP-Compacted Rene'95.

The tensile properties of HlP-^compacted CBMS Alloy 1200/19 ribbon were measured at 1000 F. _This

material was HIP-compacted in a three hour, 2050°F HIP cycle at 15 ksi, solution treated at 1925°F

for one hour, cooled with a helium purge, and aged at 1250°F for 8 hours. This processing was identical to that used on the PM version of Alloy 1200/19(^^). The Y' solvus temperature of this alloy

is approximately 2000°F, so the HIP temperature was above the Y' solvus temperature while the solution

temperature was below it. The results of individual tensile tests on Alloy 1200/19 are shown in Table

II for randomly packed CBMS ribbon and PM Alloy 1200/19. The major differences between the materials

are the extremely low tensile ductilities in the CBMS conditions relative to the PM product. The

specimens with low ductilities also have low tensile strengths due to specimen failure before necking

occurred. Fractography of the failed CBMS tensile specimens showed that the fracture occurred along

PRB. The lower tensile ductility in the longitundingly oriented specimen of the CBMS Alloy 1200/19

' most likely resulted from the applied stress being perpendicular to a majority of prior ribbon boundaries. A lower number of PRB were perpendicular to the axis of applied stress in transverse specimens.,

In an attempt to improve the ductility of CBMS Alloy 1200/19, the HIP'ed compacts were extruded. I

A

•'5

S

Figure 8. TEM Extraction Replica of HIP Compacted

CBMS Rene'95 Ribbon Showing Second Phase

Particles Along Prior Ribbon Boundaries.

1

i

370
Figure 9. Optical Micrograph (a),TEM

Extraction Replica (b),and TEM Replica (c) of extruded CBMS Alloy

1200/19.

Table II. Tensile Properties of HIP-Compacted PM and CBMS Alloy 1200/19

Material

Condition

Temperature

(F)

0.2% Offset

Yiel d Strength

(Ksi)

111 timate

Tensile

Strength

(Ksi)

Elongation

(%)

Reduction

in

Area

(%)

Specimen

Orientation

Relative to Axis of HIPPM 75 158.0 217.0 32.0 32.7 Longitudinal

1000 142.4 199.8 23.3 28.8 Longitudinal

1200 138.2 203.0 18.3 20.0 Longi tudi nal

CBMS 1000 93.1 0.0 2.0 Longi tudi nal 1000 158.9 0.06 2.0 Longi tudi nal 1000 160.8 168.4 1 .6 5.0 Transverse

1000 162.2 189.7 5.6 12.5 Transverse

*Sample Failed Before Reaching

0.2% Plastic Strain.

371
Extrusion of CBMS Alloy 1200/19

Extrusion was used to simultaneously impart more shear deformation and elongate the spacing W

between adjacent second phase particles along prior ribbon boundaries. The randomly packed ri bbon SI I* form was used for the HIP'ed extrusion billet instead of wound ribbon in order to obtain sufficiently

billet length for test specimen evalaution.

Approximately 0.8 kg of Alloy 1200/19 CBMS ribbon was loaded into a 2 in. dia. HIP container and I HIP'ed at 2050°F. The container was subsequently machined into 1.75 inch diameter extrusion billet,

[ loaded into an extrusion can and extruded at 1960°F with an extrusion ratio of 8:1. The Alloy 1200/19

portion of the resulting extrusion was approximately 0.6 inch in dia. and 11.5 inches long. The

extrusion temperature was below the Y' solvus temperature of this alloy and resulted in the extrusion

having a fine grain size. The extrusion was given a double solution treatment; a three hour, 2050°F

cycle followed by a one hour, 1925°F treatment. The first cycle was intended to restore the microstructure to the grain size observed in the material HIP'ed at 2050°F. The second solution treatment

and the subsequent aging cycles were the same as that used for the HIP-compacted CBMS ribbon in order

to form the same Y' distribution.

The microstructure of the extruded CBMS Alloy 1200/19 material was studied using optical metallography and TEM replicas. Figure 9a shows that the grain size is similar to that previously shown

in Figure 7 for HIP compacted Rene'95. Even after extrusion, the PRB regions were still decorated,

as shown in Figure 9a. Figure 9b shows a TEM extraction replica of a PRB region in the extruded

material. The PRB particles in these boundaries have a spacing approximately twice that observed in HIP'ed CBMS Rene' 95. This probably results from the lower reactive element content of Alloy 1200/19

relative to that of Rene' 95 and hence a lower oxide particle density. The increased oxide spacing

may also result from the extrusion process. Figure 9c shows the Y' distribution of this material as revealed in a TEM replica. This Y' distribution is similar to that obtained in HIP-compacted helium

atomized Alloy 1200/19 powder which received the same heat treatmentv^ ^ . The large amount of primary

Y ' resulted from the final solution treatment being 75°F below the Y

' solvus temperature.

The tensile, creep, rupture, and cyclic defect tolerance properties of extruded CBMS Alloy 1200/r

were determined in axisymmetric specimens with uniform gage lengths. The gage diameter of the creep/

rupture specimen was 0.2 inch, while all other specimens had a 0.25 inch diameter gage section.

Tensile tests were performed at room temperature, 1000°F, and 1200°F. These data are plotted in

Figure 10 along with previous results on the HIP-compacted PM Alloy 1200/19 and the average tensile

properties of PM Rene' 95. The yield and ultimate tensile strengths of both forms of Alloy 1200/19

are very similar, but both are approximately 15 percent lower than that of PM Rene' 95. Alloy 1200/19 was solution treated approximately 75°F below its Y ' solvus temperature, while the curve for Rene' 95 was determined for material solution treated only 30°F below its Y

' solvus temperature. The strength,

of Alloy 1200/19 could probably be substantially increased by increasing the solution treatment temperatureU8,20) _ jh-js would reduce the volume fraction of primary y' snd allow a larger fraction of

the Y' to be used for precipitation hardening. The tensile ductilities of the extruded CBMS Alloy

1200/19, as measured by the elongation and reduction in area, are inferior to those of the PM version

of this alloy but are substantially greater than those shown in Table II for the HIP-compacted CBMS

Alloy 1200/19. Previously, the lower tensile ductilities in the extruded CBMS Alloy 1200/19 relative

to those of the PM version of this alloy were caused by the higher density of oxides and carbides T along prior particle boundaries in the extruded material. The difference in ductilities between

Rene' 95 and Alloy 1200/1 9 is most likely due to their differences in strength(lS) , The good ductilities observed in the extruded CBMS superalloy are quite encouraging because this is the first time

that mechanical properties of relatively large volumes of a melt-spun superalloy have been reported.

Figure 11 shows scanning electron micrographs of CBMS tensile specimens fractured at 1000°F.

The HIP-compacted CBMS ribbon specimen shown in Figure 11a failed almost entirely along prior ribbon

boundaries. Higher magnification observations showed that most of the fracture surface was covered

by small dimples which apparently nucleated at the oxides and carbides within the PRB regions.

Other studies on cast and wrought product forms have shown this general type of intergranular dimple j rupture(22-24) ^ The extruded CBMS specimen also failed by dimpled rupture; however, the fracture

j path is not totally along the PRB regions. This apparently is the cause for the large improvement | in tensile ductility. Further reduction in the frequency of oxide and carbide in PRB regions may ! result in ductilities equivalent to those of PM Alloy 1200/19 shown in Figure 10. |

A single combined creep/rupture test was performed on the extruded CBMS Alloy 1200/19 at 1200°F | with an initial stress or 125 ksi, which is 87% of its 0.2 percent yield strength at 1200F. This

specimen had a 0.2 percent creep life of 1220 hours and a rupture life of 5146.6 hours. These

j properties are shown using a Larson-Mil ler normalization in Figure 12. The solid lines in this

plot indicate the 0.2 percent creep and rupture properties of Rene' 95. The arrow on the PM Alloy

1200/19 creep data point indicates that the terminal creep strain was less than 0.2 percent'^ '1. These results show that both PM and extruded CBMS Alloy 1200/19 have creep and rupture capabilities

372
TEMPERATURE (F)

Figure 10. Variation of the (a) Strengths and (b)

of Alloy 1200/19 With Temperature.

Tensile Ductilities or CBMS and PM Versions (b)

Figure 11. 5 ^;^ Fractographs of Tensile Specimens of (a) HIP-Compacted & (b) HIP'ed Plus Extruded

CBMS Alloy 1200/19.

373
which exceeded those of Rene' 95. The creep and rupture capabilities of Alloy 1200/19(12) are similar to those of AF115 (2, 25), a PM disk alloy with a 65°F creep and rupture capability improvement

over Rene' 95. i

i I The fatigue crack growth rates were determined using the DC potential drop techniques developed*

by Gangloff(^°> and VanStone, et. al . for axisymmetric specimen geometries containing 0.1 mm deep electrical discharge machined notches. The cyclic defect tolerance evaluations included both

strain controlled tests on specimens with 0.1 mm radius semicircular defects and load-controlled

tests on specimens with 0.1 mm chord-shaped defects identical to those used in other investigations

on nickel-base superalloys 19, 20, 26, 28)^ growth rate data determined in 1000°F.

0.33 Hz (20 cpm) crack growth rate tests were analyzed in the form of a linear elastic fracture

mechanics fatigue growth curve where fatigue crack growth rate (da/dn) is plotted against the range

in stress intensity factor ( A K). Figure 13 shows the fatigue crack growth curve for HIP'ed PM and

HIP'ed plus extruded CBMS Alloy 1200/19. The solid line in these curves is the baseline curve con- structed by VanStone and Gangloff ^ ' 91for the PM Rene' 95. Both sets of data show that Alloy 1200/19

has lower fatigue crack growth rates than Rene' 95 at 1000°F.

The similarity of properties between PM and CBMS Alloy 1200/19 indicates that this alloy has a wide range of processiblity. The high creep and rupture properties suggest that this type of low 9, high solid solution strengthened alloy has potential as an advanced disk material. In addition,

the low T' volume fraction and high ductility may permit heat treatments which i^efine y' size and

possibly retard the fatigue crack growth rates relative to those shown in Figure 13(^^^.

Pi scussion

This investigation has shown that CBMS process can be used to produce relatively large volumes

and high yields of rapidly solidified superalloys which are free of microsegregation and branched

dendritic microstructures. These ribbbons can be compacted to full density using several techniques,

including the winding of ribbon around a mandrel to fabricate axisymmetric shapes. This approach

was used to fabricate the six inch diameter HIP-compacted -shape shown in Figure 14. This process

has many of the advantages summarized by Breinan, et. al.'^^^ for layer glazing and avoids the

thermally induced cracking problems associated with layer glazing.

The primary technical barrier issue for the use of CBMS products for structural applications

is the extensive PRB decoration which occurs in HIP-compacted material which results in low tensile

ductilities. In this study, it was demonstrated that extrusion substantially improved

these properties to a level where they approached those of PM product forms. This approach cannot

be easily used to solve the ductility problems in large wound HIP-compacted shapes due to the large

diameters of such part. Thus, it is imperative to develop the melt-spinning and compaction procedures which can eliminate or substantially reduce the PRB decoration to more fully utilize the

potential advantages of rapidly solidified CBMS materials.

Acknowl edgements

The authors acknowledge the financial support of the Defense Advanced Research Projects Agency

(DARPA) which funded portions of this investigation and the guidance of Dr. E.C. VanReuth of DARPA

and Mr. A.M. Adair and Mr. W.T. O'Hare of the Air Force. We would also like to thank Mr. E.C. Nagy,

Mr. R.J. Zabala, and Mr. L.C. Perocchi of General Electric Corporate Research and Development, Mr. R.W, Smashey of the General Electric Aircraft Engine Group, and Mr. F.J. Rizzo of the Crucible Compacted Materials Operation for their assistance in producing and compacting the CBMS ribbon. Mr. L.T. Duvelius and Mr. E.D. Deaton were instrumental in conducting the mechanical property evaluations.

Dr. E.N. Hall performed the STEM microanalysis. The authors are deeply indebted to Mr. A.M. Johnson

and Dr. D.C. Williams for their encouragement, support and many helpful suggestions during this

i nvesti gation

374
CRACK

GROWTH

RATE

(mm/cycle)

Figure 14. Photograph of Six Inch Diameter

Wound CBMS Alloy 1200/19 Shape

Figure 12. Larson-Miller Plot of the Creep

and Rupture Properties of CBMS and

PM Versions of Alloy 1200/19

Relative to Those of Rene '95 STRESS INTENSITY (ksl lni'2) 10 20 50 10-2

10-3 -

io-< -

10-5 -

10-6 -

10-7

10-8

—I—I—l—I—]

ALLOY 1200/19 CBMS RIBBON

2050F HIP

1960F EXTRUSION

2050F SOLUTION

1925F SOLUTION

1250F/8HR AGE

T

1000F AIR

20 CPM

o STRAIN CONTROL

« LOAD CONTROL

V NO GROWTH

10-4

10 -

10 -

5 O

3D>

O

o

® O

$

10-7

3D

5im

3

O

><n

- 10-8 *

10-9

10 20 50 STRESS INTENSITY (MPa mi'2)100

STRESS INTENSITY (ksi ini/2) 10 20 50 10-2 r

10-3

ug 10-4 E

10-5

$

o

oco

o

<

oco

10-7

10-8

—l—I 1

1

1

1

' . 1

. . 1

I COMPACT H19 HELIUM ATOMIZED . -

- -270 MESH / 2050F HIP - 1925F SOLUTION X

- 1000F SALT QUENCH Z

- 1250F/8HR AGE

z

- f

*

- 1

n

] 1

1000F AIR

20 CPM

- STRAIN CONTROL

“ » LOAD CONTROL

-

r> 1 1 i 1— 1 1 1 _1 1 1 1-L- 10 20 STRESS INTENSITY (MPa mi'2)

10-4

10-5

10-6

10-7

10-8

10-9

(b)

Figure 13. Comparison of the Fatigue Crack O^rOwth of Alloy 1200/19.

Rate Curves of the CBMS (a) and PM (b) Versions

375

CRACK

GROWTH

RATE

(In/cycle)
REFERENCES

1. P.S. Mathur and J.L. Bartos, USAMRDL-TR-76-30 , May 1977.

2. J.L. Bartos, AFML-TR-74-187, December 1976.

3. R.D. Eng and D.J. Evans, NASA-CR-1 35409, March 1978.

4. R.D. Eng and D.J. Evans, NASA-CR-165549, May 1982.

5. D.F. Polk and B.C. Giessen, Metallic Glasses , ASM, Metals Park, Ohio, 1978, p.l.

6. R.W. Cahn, Contemp Phy , 21 \ 1980, p.43.

7. C. Sury and Rayanna, Suppl . Sci. Rep. RITU , A-7 .

8. H.A. Davies, N. Shohoji, and D.H. Warrington, Rapid Solidification Processing Principles and

Technologies II , Claitor's, Baton Rouge, 1980, p. 153.

9. J.V. Wood, P.F. Mills, A.R. Waugh, and J.V. Bee, J. Mat . Sci . , 1

, 1980, p. 2709.

10. H.H. Liebermann , R.E. Maxwell, R.W. Smashey, and J.L. Walter, "Process and Structural Aspects of

Melt Spun Rene'95 Nickel-Base Superalloy Ribbons", Submitted to Metallurgical Transaction .

11. R.H. VanStone, AFWAL-TR-82-4032 , Vol . II, Cincinnati, Ohio, June 1982

12. R.H. VanStone, H.H. Liebermann, J.R. Hughes, J.L. Walter, H.A. Davies, and R.W. Smashey,

AFWAL-TR-82-4032, Vol. Ill, June 1982.

13. H. Jones, Rep. Prog. Phys ., 36 , 1973, p. 1425.

14. H. Jones, Rapid Solidification Processing Principles and Technologies, Claitor's, Baton Rouge,

1978, p.28.

15. R.C. Ruhl, Mat. Sci. Eng ., 1, 1967, p. 313.

16. H.H. Liebermann, Mat. Sci. Eng ., 43 , 1 980, p. 203.

17. R.H. VanStone, F.J. Rizzo, and J.F. Radavich, Rapid Solidification Processing Principles and

Technologies II , Claitor's, Baton Rouge, 1980, p. 317.

18. R.H. VanStone, "Long Life Disks From Rapidly Solidifed Materials: Volume I, Rene'95 and AF115

Process Development", AFWAL-TR-82-4032, Vol. I, June 1982.

19. R.H. VanStone and R.P. Gangloff, Rapid Sol idificaton Processing Principles and Technologies II , Claitor's, Baton Rouge, 1980, p. 317.

20. R.H. VanStone, M.F. Henry, and A.M. Ritter, Final Report, Contr. No. NOOOl 9-79-C-0695 , Dec., 1980

21. A.M. Ritter and M.F. Henry, General Electric Corporate Research & Development, Schenectady, N.Y. 12301, Unpublished Research.

22. I. Kirman, Met. Trans ., 2, 1971, p. 1761

23. Jack Nutting, Physical Chemistry in Metallurgy , 1976, p. 489-495.

24. Gerald M. Ludtka and David E. Laughlin, Met. Trans . , 1 3A , 1982, p. 411.

25. D.M. Carlson, AFWAL-TR-81 -4078, August 1981.

26. R.P. Gangloff, Fatigue of Engineering Materials and Structures , 1981, p.l5.

27. R.P. Gangloff ASTM STP 738 , ASTM, Philadelphia, PA., 1981, p. 120.

28. R.H. VanStone, D.D. Krueger, and L.T. Duvelius, "Use of DC Potential Drop Crack Monitoring

Technique in the Development of Defect Tolerant Disk Alloys, to be published in ASTM STP 791 , ASTM, Philadelphia, PA., 1983.

29. E.M. Breinan, D.B. Snow, C.O. Brown, and B.H. Kear, Rapid Solidification Processing Principles

and Technologies II , Claitor's, Baton Rouge, 1980, p. 440.

376
I BORIDE IN MELT-SPUN SDPERALLOYS

S. C. Huang*, E. L. Hall*, and T. K. Glasgow**

* i

I *General Electric Corporate Research and Development

P.O. Box 8 Schenectady, NY 12301

**NASA Lewis Research Center

I Materials Division

*1 Cleveland, OH 44135

J ABSTRACT

I Superalloys with various levels of M^B^ boride have been studied by the melt-spinning

! rapid solidification process. The alloy compositions studied were chosen using the forj mula (matrix)^QQ_^(boride)^, where the matrix atom composition was

! 13"°0.02>0.85‘**0.7^i0.3>0.15' «'io.l‘^'o.3““o.5''‘o.l> 0.6®0.4

! ,„d ih. borld. fr.otion , = 0.15, 1.5 and 7.5. It «as found that tha botou content in tbe

i range of the present study did not significantly affect the ribbon castability. Inc^eas

, ing boride level decreased the bend ductility of the as-cast ^

was drastically improved after the step anneal of 1080 C/4 hrs and 900 C/10 hrs. TE

studies of an as-cast ribbon showed uniform gamma prime precipitates of about 10 nm in

I diameter in crystalline cells which were less than 1 pm in size. The cell boundaries were

decorated with boride particles smaller than 50 nm in diameter. Also studied was the

annealed ribbon microstructure, particularly in regard to the grain and boride coarsening,

boride-matrix coherency, and the boride chemistry. The boride composition was determined

after electrolytic extraction. Finally, the incipient melting temperatures of the rapidly

solidified ribbons were determined by DTA.

Introduction

Small boron additions, around 0.01 w/o have been found to improve the creep properties

stalline nickel base superalloys [1]. This improvement is generally attributed toron f^tti g mically into small holes in the grain boundaries. Significantly larger amounts of boron, around

w/o, form discrete boride phases (such as M^B^) which may also improve superalloy mechanical

iperties

The present paper reports on our investigation of the MgB

2 boride rising the ^apididification, melt-spinning process. In particular, preliminary results on the boride chemistry,

ibility, structure, and volume fraction, along with some ribbon microstructure, as

J

lealed, will be presented. Incipient melting as influenced by boride fraction and rapid solidif

:ion will also be discussed.

Experimental

Three alloys have been studied. The compositions were based on formula

(matrix) 100-x

(boride)

ere the MgB^- boride atom fraction, x, was 0.15, 1.5 or 7.5. The atomic composition of the matrix

s chosen as

matrix = (Ni^ ggCr^^^gMoQ q2^0.85^^^0.7^^0.3^0.15

377
representing approximately 50% gamma prime, and that of the boride was

boride

5 ^^0 . 1

^ 0 . 6 ®

0.4

selected according to the chemistry reported in Ref. (3) and (4). Alloy compositions in

weight percent are listed in Table 1. Ribbons of the alloys were fabricated in vacuum by the

rapid-solidification, melt-spinning process. Ribbon surface topography was studied by optical

metallography and scanning electron microscopy (SEM). This yielded information about the

ribbon-wheel contact, solidification structure and the ribbon fracture characteristics. Further detailed microstructural examination was conducted using analytical electron microscopy

(AEM) . The analytical electron microscope has transmission imaging and electron diffraction

capabilities, and is equipped with an energy dispersive X-ray spectrometer for chemical analysis with "10 nm rated resolution. Samples were prepared by jet-electropolishing in a solution

of 20% perchloric acid in methanol. Boride particles, after being extracted electrolytically

in a 15% HCL-methanol solution, were studied by X-ray diffraction, electron diffraction, energy

dispersive X-ray spectroscopy (EDS), to determine their structure and chemistry. Finally, differential thermal analysis (DTA) was performed on the as-cast ribbons. The DTA conditions

were: sample weight ~80 mg, heating rate 10°C/min., and the argon flow rate 150 cc/min.

Table 1 - Nominal Compositions of Alloys Studied (w/o)

alloy

number fraction Ni Cr Mo A1 Ti B

70 0.15 77.34 10.50 3.05 5.16 3.94 0.012

71 1.5 76.49 10.60 3.73 5.10 3.96 0.12

72 7.5 72.69 11.07 6.75 4.82 4.07 0.60

Results and Discussion

(A) Ribbon Formabilitv and Ductility

The three compositions studied were cast into ribbons having uniform dimensions and smooth

surfaces. No consistent variations in the r ibbon-to-wheel sticking distance or in the wetting

pattern observable on the ribbon wheel-side surface [5] could be found as a function of the

boride levels. Ribbons of all three compositions were obtained in fragments 2 to 20 cm long,

presumably fractured in flight after separation from the casting wheel surface or upon impact

with the vacuum chamber walls. A scanning electron micrograph of the fracture surface in

alloy-71 ribbon shows the intercellular fracture characteristics. Fig. 1. Hot tearing or intermediate-temperature embrittlement might have contributed to the ribbon fracture during

casting.

Fig. 1 An SEM micrograph showing a cross-sectional surface in the alloy-71

ribbon, exposed by ribbon fracture

during casting. The ribbon top surface

is shown at top.

378
Preliminary tests have demonstrated that the bend ductility of the as-cast ribbons

decreases with increasing boron content, but the ductility can be significantly improved after

a step anneal of 4 hours at 1080**C and 10 hours at 900°C.

(B) Microstructure in Allov-71 Ribbon

The solidification morphology developed in alloy-71 ribbon is also revealed in Fig. 1: cellular near the wheel-side surface and dendritic near the top surface [6,7] . An extremely

fine nucleation structure is barely perceptible at the wheel-side surface of the ribbon [7] . It has been further found that the alloy-70 ribbon is predominantly cellular while the alloy-72

ribbon becomes more pronounced in dendritic structure. Increasing boride level generally

decreases the ribbon solidification size-scale. In alloy 71 ribbon, the fiber crystals growing

unidirectionally from the bottom to the top surface have an average diameter of ~1 pm in the

cellular region and ~2 pm in the dendritic region. Fig. 1.

Alloy-71 microstructure in an area near the middle of the ribbon thickness has been studied by analytical electron microscopy (AEM) . The AEM sample geometry was such that the

transverse section of the fiber crystals were examined. Preliminary results on the as-cast

ribbon show a uniform grain (or cell) structure with a grain size of about 0.5 pm. Fig. 2.

Also clearly shown in this figure are particles along grain boundaries, ranging in size from 20

to 100 nm. These particles, as revealed by energy dispersive X-ray spectroscopy (EDS), are rich primarily in Cr and secondarily in Mo, with slight enrichment in Ti. Significant amounts

of Ni are also found, but the majority of it is attributable to the interference from matrix

which overlaps the particles. Accordingly, the grain-boundary particles may correspond to the

(Cr,Ni)

2 (Mo,Ti)B2 boride [4]. Further characterization of the as-cast boride will be carried

out. Finally, Fig. 2 also displays in the grain interior a mottled contrast which corresponds

to a uniform distribution of gamma prime precipitates which are about 20 nm in diameter. In

dark field images formed using a gamma prime reflection, some coarse y', up to 50 nm in diameter, is observed along the grain boundaries.

Fig. 2 Microstructure near the center of the alloy-71 ribbon, as cast.

The overall structure in alloy-71 ribbon, annealed in two steps at 1080°C for 4 hours and

at 900 C for 10 hours, is shown in Fig. 3. It can be seen that the grains are on the order of

a few microns in diameter and that the spheroidal borides are from 0.2 to 0.6 pm in diameter,

both of which have coarsened in size significantly from the as-cast structure shown in Fig. 2. Most of the coarsened grains have an internal dislocation structure, and in many cases dislocation tangles outline the gamma prime particles. Some of these may be misfit dislocations; some

others may be the as-cast grain-boundary dislocations not completely removed in the anneal process. Boride particles can be found at the grain boundaries as well as in the grain interior.

The intragranular borides, generally smaller in size, mark some previous boundaries.

379
Fig. 3 The overall microstructure near the center of the alloy-71 ribbon, step-annealed 1080**C/4

hrs. and 900*C/10 hrs.

Figure 4 is a bright field/dark field pair showing the gamma prime particles in the annealed

alloy 71 ribbon. In Fig. 4(b), a gamma prime reflection has been used to form the image, and thus

the Y* particles appear bright. In addition to the large y' particles which average 0.5 ^m in diameter, small spherical y' can be observed formed either during cooling from or at the lower heat

treatment temperature. Furthermore, in Fig. 4(a) apparent misfit dislocations can be seen at the

interfaces between the boride particles and the matrix. In Fig. 4(b) some boride particles are

Fig. 4 A bright field/dark field pair of TEM micrographs showing detailed microstrncture in

the annealed alloy-71 ribbon.

380
brightened along with the y' particles since a boride reflection has occurred near to the y' reflection used to form the image. Electron diffraction studies indicated that the boride/matrix orientation relationship is

(001) ^. matrix II <001),oride'

[001] ^

. matrix II ^®®^^boride'

and

[100] ^

. matrix II t^lO^boride*

The spheroidal annealed boride particles have been identified by electron diffraction as having a tetragonal structure, a = 5.71 and c = 3.0?X. An EDS of the annealed boride shows the

enrichment of primarily Mo and secondarily Cr — in reverse order to that found in as-cast boride.

Nickel and titanium also appear in the spectrum. The above results are consistent with the

molybdenum-rich (Mo,Ti)

2 (Cr,Ni)B

2 boride as reported in Ref. (4).

(C) Borides in Alloy-72 Ribbon

Efforts to characterize the alloy—72 ribbon structure were concentrated on the boride particles, extracted from ribbons as-cast or step-annealed 4 hours at 1080**C and 10 hours at 900°C. The

extracted as-cast boride, as shown by SEM, was fine in size, ranging from 0.006 to 0.6 pm in diameter. During annealing, the size became much larger, mostly between 0.2 and 0.5 pm. These results

are similar to those in alloy 71, suggesting that the boride size is quite independent of the boron

level. The boride particles in alloy 72 are thus conceivably larger in number and can be expected

to stabilize the grain structure more effectively. SEM study further shows clearly the spheroidization of the boride particles due to the heat treatment.

Energy (keV)

(b)

Energy (keV)

Fig. 5 The energy dispersive X-ray spectra of borides extracted from (a) as-cast, and (b) annealed

alloy-72 ribbon.

For alloy 72, the as-cast boride has a tetragonal structure with a = 5.67 and c = 3.05X, as determined by X-ray diffraction. An energy dispersive X-ray spectriun of the boride. Fig. 5(a),

shows strong Cr, medium-strong Mo, medium Ni, and weak Ti peak. These results suggest the as-cast

boride being of the chromium-rich, low-parameter (Cr,Ni)

2 (Mo,Ti)B

2 type [4]. On the other hand,

electron diffraction of the annealed boride yields lattice parameters of a = 5.8 and c = 3.lA, and

EDS shows strong Mo, medium Cr, weak Ni, and barely perceptible Ti peaks. Fig. 5(b). These results

suggest .the annealed boride being of the molybdenum-rich high-parameter Mo

2 (Cr,Ni)B

2 type [4]. A

similar boride composition change has been found in the as-cast and annealed^ribbons of alloy 71.

(D) Ribbon Thermal Analysis

Results of the differential thermal analysis (DTA) on ribbons of all three compositions studied

are compared in Fig. 6. In general, each on-heating thermogram exhibits two endothermal peaks: a high-temperature peak corresponding to the melting of the matrix, and a low temperature one corresponding to melting of the chromium-rich, as-solidified boride [4,8]. As shown in Fig. 6, with

the boron content increasing from 0.012 to 0.60 w/o, the liquidus temperature decreases from 1354 to 1285 C. More significantly, incipient melting of the boride phase occurs at 1223 C for the rapidlysolidified alloy-70 and 71 ribbons, with the boron content of the latter being larger by an order of

magnitude. For the alloy-72 ribbon, containing 0.6 w/o boron, an additional five-fold increase, the

incipient melting temperature decreases only slightly to 1212**C. These compare to an incipient

381
melting temperature decreases only slightly to 1212**C. These compare to an incipient melting temperature of 1201**C reported for the conventionally-cast B-1914 alloy which contains 0.1 w/o boron

[8]. As also shown in Fig. 6 , the intensity of the incipient-melting peak increases with boron

level, in proportion to the quantity of the boride phase formed during the rapid ribbon quenching.

Fig. 6 DTA thermographs of the alloy-70, 71 and 72 ribbons.

Conclusions

Superalloy ribbons with boride levels up to 7.5 atom fraction can be fabricated quite | readily by the rapid-solidification melt-spinning process. The ribbons, however, tend to fracture j during casting, perhaps due to hot-tearing or intermediate-temperature embrittlement. The as-cast

^ boride is identified as (Cr,Ni>

2 (Mo,Ti)B

2

, and after a step-anneal, becomes (Mo,Ti)„(Cr,Ni)B

2

, richer in molybdenum. The annealed boride has a semicoherent relationship with the matrix:

(001). .. 11(001) ^ . , [001], I I [001] ^ , and [310], ., 1 1 [100] ^

. The step- | boride matrix boride matrix , TM boride, ., matrix , , a anneal also coarsens the grains and the boride particles. The size of the as-cast and annealed

boride, however, seems to be unaffected by the boride level. While the bend ductility of the ascast ribbons decreases with increasing boride fraction, the ductility can be improved by the heat

treatment. Finally, the incipient melting temperature is 1223*C for alloy-70 and 71 ribbons and ! 1212®C for alloy-72 ribbon. '

Acknowledgments

We thank R.J. Zabala, L.C. Perocchi, R.P. Laforce, and P.C. Irwin for technical assistance. We also thank H.R. Gray and L.A. Johnson for stimulating discussions. This research is partially sup- i ported by the NASA Lewis Research Center.

References I

I

[1] R.F. Decker and J.W. Freeman, Trans. AIME, 218 (1960) 277.

j

[2] D.H. Maxwell, J.F. Baldwin, and J.F. Radavich, Metall. Met. Forming, Oct. 1975, p. 332. |

382
j [3] J.R. Mihalison, in Reviews on High-Temperatnre Materials. Ed. J.B. Newkirk, Frennd PnblishI ing, Tel-Aviv, Israel, Vol. II, No. 3, 1974, p. 244.

[4] W.J. Boesch and H.B. Canada, J. Met., April 1968, p. 46.

i [5] S.C. Hnang and H.C. Fiedler, Met. Trans. A, 12 (1981) 1107.

[6] H.A. Davies, N. Shohoji, and D.H. Warrington, Proc. 2nd Int. Conf. Rapid Solidification Processing. Ed. R. Mehrabian et al., Reston, Virginia, 1980, p. 153.

[7] S.C. Hnang and A.M. Ritter, to be published in Proc. Svmp. Chem. Phvs. Rapidly Solidified

Materials

TMS-AIME Fall Meeting, Oct. 1982, St. Louis, Missouri.

[8] C.J. Burton, in Proc. 3rd Int. Svmp. Sunerallovs; Metallurgy and Manufacture. Ed. Kear et al., Claitor's Publishing, Seven Springs, Pennsylvania, Sept. 1976, p. 147.

383
PREPARATION OF METALLIC GLASSES WITH SECOND PHASE PARTICLES

D.G.Ast and P . G . Z ielinski

Materials Science and Engineering

Cornell University

Ithaca N.Y. 14853

ABSTRACT

Ni[75]Si[8]B[ 17 ] metallic glasses with varying fractions of WC and

TiC second phase particles were prepared using two different techniques.

In the first method, the second phase particles were added to the melt

beiore the latter was extruded. In the second method, the particles were

injected immediately after extrusion into the still liquid melt puddle.

The volume fraction of second pnase particles incorporated into the cast

ribbon and their spatial distribution were studied as a function of the

preparation parameters by optical microscopy of sectioned specimens.

In ribbons prepared with the first method, the spatial distribution

of second phase particles depended critically on the geometry of the melt

puddle. As the volume fraction of second phase particles increased, the

particles tended to agglomerate in periodically spaced clumps along the

center plane 01 the ribbon.

In ribbons prepared with the second method, the distribution of the

second particles was homogeneous and independent of volume fraction.

The mechanical properties of the dispersed alloys and of undispersed

control specimens were studied by tensile and bending tests, and by

scanning electron microscopy observations of shear bands formed during

deformation. In ribbons with homogeneously distributed second phase

particles, yield stress and Youngs modulus increased relative to

undispersed control specimens. Ribbons containing agglomerated particle

clumps, on the other hand, had fracture stresses lower than undispersed

control specimens.

1 .Introduction

Dispersion of second pnase particles is commonly used to modify important

mechanical properties of materials such as yield strength (metals) or fracture

toughness (ABS polymers, zirconia dispersed ceramics). In metals, second phase

particles are usually formed by precipitations and this limits the chemical

composition of the strengthening phases. Powder metallurgical methods can overcome

this restriction (e.g.TD Nickel) but these alloys have not been commercially

successful

The properties 01 metallic glasses containing (externally added) second phase

particles have not been studied in detail. However, observations on partially

crystallized alloys [1] indicate that the yield strength should improve.

Allied Chemical has filed a patent on the preparation of metallic glasses with

second phase particles [2] and a preliminary report on the preparation and

384
properties of a Ni based metallic glass containing WC particles has been published

by Kimura et al [3J. The latter report that the yield strength improves relative

to undispersed matrix material.

The mechanical properties of dispersed materials depend critically on the

spatial distribution of particles. In the ribbons prepared by chill black casting,

the distribution of particles is controlled by the fluid flow in the melt puddle

which, under certain operation conditions, can be turbulent [4]. In such a case,

particles added to the melt prior to extrusion tend to agglomerate in the center of

vortexes which are shed periodically into the cast ribbon. This difficulty can be

overcome by adding the particles after extrusion but before solidification.

This paper discusses the microstructure of ribbons prepared by these two

different preparation techniques and their mechanical properties.

2 .Preparation techniques

In order to achieve satisfactorily results in a system in which second phase

particles are introduced into the liquid state, the following conditions must be

met; i) the density difference between the particles and the matrix should be

small to avoid segregation by gravity ii) the particles should be wetted by the

matrix ana should not dissolve in the melt before solidification iii) the particles

should not act as nuclei for the crystallization and iv) fluid flow should be

laminar in order to avoid segregation in turbulent flows. Obviously, many of these

conditions are easier to meet in a system which minimizes the time of contact

between the fluid and the particle.

In our apparatus, density differences are usually not a serious problem, even when the particles are added to the melt before extrusion since stirring by the RF

field of the heating coil keeps the particles suspended. A more serious problem is

the commencement oi matrix crystallization at the corners of second phase particles

which was occasionally observed in [3]. Thus, high quenching rates are required

for the matrix material. High quench rates require a good heat transfer to the

wheel. An analysis shows that the direct influence of particles (i.e. rule of

mixing contribution to the thermal conductivity and heat capacity of the matrix) on heat transfer is small [5] compared to the surface condition of the wheel which can influence heat transfer up to >60% [6]. However, investigations of the contact of

the liquid solid interface [73 show that the true area of contact depends

critically on the surface morphology of the solid, and that the liquid contacts the

solid continuously only when a critical wetting condition is satisfied. A

consequence of this condition is that the presence of particles which are not

wetted by the matrix can change this critical wetting condition and thus indirectly

greatly reduce heat transfer from the matrix to the quenching surface.

Ni[75]Si[8]B[ 17] was chosen as the matrix material. This alloy is a good

glass former which can be prepared in ribbons as thick as 225 urn [8]. The second

phase particles were WC and TiC, which are wetted well (contact angle about 0 (WC)

and 32*(TiO) by liquid Ni [93.

Ingots Of Ni[753Si[83B[ 173 were prepared by arc melting of high purity (4Nine)

nickel, nickel boride and silicon. Ribbons were prepared by chill block casting

using a quartz crucible with a slit orifice (0.2 mm) positioned close (0.5-0. 7mm)

to a rotating copper wheel (290 mm diameter, surface speed 21 m/sec). The second

phase particles were introduced either by adding the powders to the crucible prior

to heating (method 1j, or by blowing the particles into the melt puddle during the

extrusion process (method 2) using an arrangement similar to a small sand blasting

gun [103. The He operating the gun served simultaneously as a conventional quench

stabilizer of the liquid.

In either case, the resulting ribbons were 30 - 60 urn thick and 4 - 7 mm wide.

The amorphous structure of the matrix material was confirmed by x-ray diffraction

and transmission electron microscopy.

ts

3.1 Particle distribution

The distribution and volume fraction of second phase particles was determined

by optical microscopy of surfaces polished parallel either to the surface or to the

385
! edge of the ribbon. This polishing process was continued until the center of the

ribbon was reached in order to determine the three dimensional particle ' distribution. i

Fig. 1 Cross section (a) and

surface (b) of Ni [75 ]Si [ 8]B[ 1?]

ribbons containing various

fractions of WC particles

(Preparation method 1).

10 mm

Fig. 1 shows an optical micrograph of the cross section and the surface of ribbons

prepared by method 1 (i.e. addition of particles to the crushed ingot prior to

melting). At low volume fractions (Fig. la) the WC are nearly homogeneously

distributed, with the occasional formation of a particle agglomerate. As the

volume fraction of WC increases large and periodically distributed particle

agglomerates are observed along the ribbon (Fig. 1b). A ribbon with an average

volume fraction of 6% WC was selected for mechanical testing in order to evaluate i the influence oi these agglomerates on the mechanical properties. i

’ a

" I

.

' 2Qum

• ft,

Fig. 2

Homogenous distribution of

second phase particles

[WC - (a), TiC - (b)] in

Ni[75lSi[8]B[ 17] ribbons

prepared by method 2.

Ribbons prepared by method 2 (i.e. post extrusion addition) contained a i_; homogeneous distribution of second phase particles at all volume fractions j investigated. A typical cross section is shown in Fig. 2. The volume fractions j were estimated from several micrographs and averaged over the length of the ribbon, i For the samples shown in Fig. 2 the volume fractions are 4.5% WC (Fig. 2a) and j 6.4% TiC (Fig. 2b).

3.2 Mechanical properties

Ribbons prepared by method 1 were shaped into tensile specimens with reduced

cross sections and polished edges. All tensile specimens fractured in .the linear ' regime of the stress strain curve. The fracture stresses calculated from the load ‘j at failure and the measured cross sections are shown in Table 1. Inspection of

this table shows that the yield strength of dispersed ribbons prepared by method 1 '] is lower than that of undispersed ribbons. Scanning electron microscopy (SEM)

j fractography was carried out in a JEM 35 B and the results are indicated in Fig.

3. The failure pattern of the undispersed material (Fig. 3a) is typical for j' metallic glasses and shows the vein pattern characteristic of shear-tearing.

Inspection of Fig. 3b suggests that the agglomerates act as the initial site for }

decohesion. Since in a metallic glass, in the absence of work hardening, failure

in shear occurs at the weakest cross section a very small fraction of agglomerates S can cause premature failure in a tensile specimen.

386
Fig. 3. Fracture surfaces of undispersed (a) and dispersed (b) Ni [75 ]Si [ 8]B[ 1 7]

with 6 vol./S of WC prepared by method 1.

Ribbons prepared by method 2 were difficult to polish (because of the presence

of hard transition metal carbides) and were therefore investigated in bending. In

this test a ribbon is bent into a hairpin configuration, inserted between two

parallel platens, and then slowly compressed. The force exerted on the platens by

the ribbon is recorded as a function of the platen separation D, in an Instron

testing machine. Typical results of such tests are shown in Fig. 4 for an undispersed ribbon (curve A) and for two ribbons containing second phase particles.

Fig. 4 of the

(D ) as persed

Loa d

r ibbo recor

r ibbo

corresponding to unit width

n (P/w) vs separation distance

ded for undispersed (A) and disns (B,C) in bending test.

Inspection of this figure shows marked differences in the mechanical behavior of

the three ribbons. The dispersed ribbons show a serrated plateau (see curves B, C)

which is absent in the non dispersed ribbon. Since metallic glasses behave very

close to ideal elastic - plastic solids, bending tests can be analyzed

quantitatively to obtain the flow stress and the elastic modulus of the ribbon

under test [11,12]. The results obtained are generally in good agreement with

those obtained in tensile tests which require a more demanding specimen preparation

(reduced cross section, carefully polished edges) in order to achieve reliable

results. Table 1 lists the the yield strength and Youngs modulus of undispersed

and dispersed metallic glasses prepared as a function of the preparation and

testing conditions. Note the good agreement between tensile and bending tests in

the undispersed case. Table 1 indicates th

increased yield stress

the influence of the seco

i.e. the formation a materials were bent into

electron microscope at de

at ribbons prepared by method 2 have a significantly

and a high Oy/E ratio of about 1/40. In order to examine

nd phase particles on the plastic deformation process,

propagation of shear bands, both undispersed and dispersed

a hairpin configuration and examined inside a scanning

creasing platen spacings (see Fig. 5).

387
Table 1: Mechanical properties of Ni[75]Si[8]B[ 17] metallic glasses

with second phase particles.

Volume fraction Preparation method Oy(GPA) E(GPa) Test

0 standard 1 .59 tension

0

II 1.57 81 .5 bending

6% WC Method 1 1 .35 - ten Sion

4. 556 WC Method 2 2.15 107.5 bending

6.456 WC Method 2 2.65 107.6 bending

Fig. 5. SEM micrographs of shear bands formed during bending of undispersed (a)

and dispersed (b) [4.5 vol.% WC] Ni [75 ]Si [ 8]B[ 1 7] ribbons (Method 2).

The character of the shear bands observed on polished cross sections differed

markedly inaicating that the presence of second phase particles significantly

changes the microscopic deformation mechanism.

4 .Discussion

The formation of undesirable particle agglomerates in ribbons prepared by

method 1 can be explained in a semiquantitative fashion by the boundary layer

instability [4]. A large shear gradient exists in the melt puddle since fluid

elements in the vicinity of the melt nozzle have zero tangential velocity whereas

those in the vicinity of the wheel surface move with high velocity (about 20m/sec).

A qualitative evaluation shows that the Reynold number is such that turbulent flow

and boundary layer separation can occur. In the latter case the particle

accumulate in regions of low rotational flow. Periodic vortex shedding releases

the clumps of particles "down stream" where they are frozen in into the ribbon

cross section as snown in Fig. 1b. These agglomerates, which were observed

already at very low volume fractions are weak centers in the center of the ribbons

and degrade the mechanical properties of the glasses investigated.

Ribbons prepared by method 2 contain a homogeneous distribution of second

phase particles. Although the detailed mechanism of the elastic and plastic

deformation in the presence of these particles is very complex, one can try, in a zero approximation to model these materials as composites. In such materials it is

frequently found that Youngs modulus obeys the rule of mixture, i.e.

E = Ep V,p + EnVfn (1)

where Ep and En represent Youngs modulus of the particle and matrix material and

Vfp and Vfn their respective volume fractions. Taking E = 440 GPa (TiC) and 660

GPa (WC) [13], relation (1) is plotted in Fig. 6. A comparison with the

experimental data snows good agreement with the predictions of the rule of mixture.

Thus, it appears that the elastic properties of these reinforced glassy materials

are in agreement with the behavior of composites.

388
E

/GPa/

Fig. 6. Young modulus of dispersed

Ni[75]Si[8]B[ 17 ] glasses as a function

of volume fraction of second phase particles.

Solid line is the prediction of eq . (1).

The plastic deformation requires a more detailed discussion. When an undispersed Ni[75]Si[8]B[

7 ] metallic glass is tested in a hairpin configuration,

the load vs platen separation curve contains typically three stages. At large

platen separations (2.1 to 1.4 mm) the material deforms homogeneously without

formation of any shear bands. Below 1.4 mm, the deformation proceeds via the

formation of shear bands which are formed in a spatially homogeneous manner such

that the radius 01 curvature at the tip of the hairpin is roughly given by half of

the platen separation. At distances below about 1.1 mm, the deformation proceeds

only in a few shear bands close to the apex of the ribbon (kinking) until fracture

occurs or the material is doubled back on itself. These three stages of

deformation in bending are discussed in more detail in ref [14].

The plateau in the load vs distance curves (Fig. 4) which develops with

increasing amount of second phase particles corresponds to stage II, i.e. to the

formation and propagation of shear bands. The serrations reflect the interactions

of the shear bands with the second phase particles. These interactions include

delocalization of shear bands, cutting processes of particles, and nucleation of

shear bands at particles. The deformation mechanism is complex and requires

further investigations. From the technological point of view, the formation of the

plateau indicates an extended plastic deformation range, i.e. a desirable feature

in metallic glasses.

Acknowledgments

This research was sponsored by ONR under contract NR 039-151. Additional

support was provided through the use of central facilities of the Materials Science

Center at Cornell, which in turn is supported by NSF.

Ref er ences

1. H.Kimura a T. Masumoto, Acta Metall. 28(1980)1677

2. M. Narasimham - US Patent 4 330 027 (1982)

3 . H. Kimura, B. Cunningham a D. G. Ast - Proc. IV Int. Conf.

Rapidly Quenched Metals, Eds. T. Masumoto. K. Suzuki

(The Japan Inst. Metals, Sendai 1982) V. II, p.1385

4. P. G. Zielinski a D. G. Ast, submitted to Scripts Met.

5. P. G. Zielinski a D. G. Ast, submitted to Metall. Trans.

6. S. Huang a C. Fiedler - Metall. Trans A 12(1981)1107

7. R. S. Timsit - Appl . Phys. Lett. 40 (1982) 379

8. M. Hagiwara, A. Inoue a T. Masumoto - Metall. Trans. A 12(1981)1027

9. G. V. Samsanov - Handbook of High Temperature Materials, no 2

(Plenum Press, New York 1964)

10 P. G. Zielinski a D. G. Ast - patent pending

11 D. J. Krenitsky a D. G. Ast, J. Mat. Sci. 14(1979)275

12 L. Bresson, J. P. Chevalier a M. Fayard, Scripts Met. 16(1982)499

13 L, E. Toth - Transition Metals and Nitrides (Academic Press, New York 1971)

14 P. G. Zielinski a D. G. Ast - to be published

389
The Effect of Solidification Velocity

on the Microstructure of Ag-Cu Alloys

W. J. Boettinger, R. J. Schaefer, F. Biancaniello

Metallurgy Division

National Bureau of Standards

Washington, DC 20234

and

I

D. Shechtman* ^

Center for Materials Research \ The Johns Hopkins University J Baltimore, MD 21218 j

'I1 ABSTRACT

Electron beam surface melting and resolidification passes have been performed on a series of Ag-Cu alloys between 1 wt.% Cu and the eutectic (28.1 wt.% Cu) at speeds between

1.5 and 400 cm/s. At low growth rates conventional dendritic or eutectic structures are obtained. At high growth rates microsegregation-free single phase structures are obtained

for all compositions. The velocity required to produce this structure increases with com- position for dilute alloys and the values agree with the predictions of absolute stability

of the liquid-solid interface with equilibrium partitioning. For alloys between 15 and

28 wt.% Cu, the velocity required to produce the microsegregation-free extended solid

solution decreases with composition and is related to nonequilibrium trapping of solute at the liquid— solid interface. At intermediate growth rates for alloys with 9 wt.% Cu or greater, a structure consisting of alternate bands of cellular and cell-free material is

obtained. The bands form parallel to the local interface.

Introducti on

Rapidly solidified crystalline alloys are well known to exhibit a refined scale of microsegre- ! gation and interdendritic second phase particles. In some cases this microsegregation and the i i nterdendri tic phases can be completely eliminated by rapid solidification. These microsegregationfree structures can occur in alloys with compositions greater than the equilibrium solubility of the

' primary phase. Such structures provide an ideal starting microstructure for subsequent thermomej chanical processing. However, the solidification velocities required to produce these structures i are not wel 1 known.

The classic splat-quenching experiments of Duwez, Klement, and Willens [1] showed by x-ray dif- [ fraction that complete solubility of Ag and Cu in the solid state at all compositions could be s obtained despite the presence of a solid miscibility gap and eutectic reaction in the phase diagram.

Boswell and Chadwick [2] showed later by TEM examination that splat quenched Ag-Cu solidified from

the melt as a single solid phase free of microsegregation. Surface melting and resolidification « experiments using lasers or electron beams permit better control over solidification velocity than

other rapid solidification techniques. Such experiments involve no bulk undercooling due to the

' presence of the alloy substrate and nucleation processes do not therefore play an important role. ? Also, because the liquid is contained in its own solid, a knowledge of heat transfer coefficients is ji not required. Elliott, Gagliano, and Krauss [3] using a pulsed laser, and Copley et al .' [4-6] using

a scanned CW laser, showed that supersaturated solid solutions of Ag-Cu could be formed by surface

melting and resolidification experiments. The solutions however, contained a rather unique struc- ; ture consisting of bands formed parallel to the local solidification interface. In these papers no structures are reported which are free of the bands and the velocities at which the transitions of

microstructure occur are not described.

*Guest Worker, Metallurgy Division, National Bureau of Standards.

390
Il

I In the present paper, a series of electron-beam surface melting experiments has been performed

I on Ag-rich, Ag-Cu alloys. These experiments determine the transition velocity for changes from dendritic or eutectic growth to microsegregation-free solidification. This information permits the

evaluation of existing rapid solidification models and points to areas where new models are

• required. The results and discussion in this paper appear in more detail elsewhere [7].

Experiments

Alloys were prepared by induction melting of 99.99 percent pure components in a graphite cruciI ble and chill casting under argon in a carbon-coated copper slab mold. Ingots were cold-rolled to

3 mm thick sheets. For the surface melting experiments, the substrate must be prepared with either

I no microsegregation or microsegregation fine enough to ensure complete mixing during surface melting.

Alloys containing 1 and 5 wt.% Cu were homogenized for 20 h at 750 °C. Alloys containing 9, 15, and

23 wt.% Cu were surface melted using overlapping electron-beam passes to a depth of ~ 0.5 mm at a scan speed of 3 cm/s. This produces a dendrite arm spacing less than 1 pm. Alloys containing

28.1 wt.% Cu (the eutectic composition) required no pretreatment. Wet chemical analysis of the six

alloys yielded compositions of 1.07, 5.49, 9.23, 15.25, 23.50, and 28.61 wt.% Cu respectively. Rapid

solidification experiments were performed by single, one-dimensional sweeps of a focused 21 keV, ~ 50 mA, -V 1 mm diameter electron beam at speeds between 1.5 and 400 cm/s across the surface of a prepared alloy substrate which was attached to a water-cooled copper block.

Significant effects arise from the change in weld puddle shape as the scan rate is increased.

, The local solidification velocity is equal to the product of the scan velocity with the cosine of the

angle between the scan direction and the local growth direction. Because of this factor, growth

velocities even at the top center of the weld puddle for high scan speeds may differ from the scan

speed. By examining longitudinal sections through the weld centerline by optical microscopy and

assuming that grain boundaries propagate parallel to the growth direction, the angle, 6, between the

' electron beam scan direction and the growth direction at the top of the weld was measured. For scan

; speeds below 30 cm/s, the growth velocity near the weld surface usually differs from the scan speed

by 15 percent or less. At velocities of 200 cm/s, the measurements indicated that the growth rate

is about one-half of the scan speed.

Results

Figure 1 shows a summary of the microstructures observed at the center of the weld zone near the

free surface as a function of the electron beam scan velocity and alloy composition for all experiments performed. The figure is divided into four regions exhibiting four general types of microstructure: (a) dendritic (cellular), (b) eutectic, (c) banded, and (d) microsegregation-free. The

detailed results will be described by alloy composition.

1 Summary of microstructures observed in Ag-Cu

alloys as a function of composition and

electron-beam scan velocity.

391
Ag-1 wt.% Cu and -5 wt.% Cu .

I

Dilute alloys exhibit a transition from segregated to microsegregation*free structures at veloc- I ities which increase with alloy composition. These velocities agree within a factor of two with the I theoretical predictions of absolute stability of the liguid-solid interface as described in another ; paper in these proceedings [8]. ^

Ag-9 wt.% Cu, -15 wt.% Cu, and -23 wt.% Cu .

These alloy compositions are beyond the equilibrium solubility limit (8.8 wt.% Cu) of Cu in Ag

at the eutectic temperature. The 9 wt.% Cu alloy does not exceed a calculated [9] Ag-solidus metastable retrograde composition however. As seen in figure 1, the microstructure of each of these

alloys changes with increasing velocity from dendritic to banded and from banded to microsegregationfree.

The refinement of the primary arm spacing with scan velocity for Ag-9 wt.% Cu alloy is shown in

figure 2. The rate of change of spacing with the actual solidification velocity may however be some- what different due to changes in weld puddle shape as discussed previously. TEM examination of a Ag-15 wt.% Cu melt trail scanned at 2.5 cm/s showed a cellular microsegregation pattern with a relatively small volume fraction of eutectic in some areas and no eutectic in other areas (fig. 3). The

maximum velocity at which dendritic growth is observed decreases with alloy composition from 30 cm/s

for Ag-9 wt.% Cu to 4 cm/s for Ag-23 wt.% Cu (fig. 1).

3 cm/s 6 cm/s 1 5 cm/s

Fig. 2 Refinement of primary dendrite arm spacing with

scan velocity for Ag-9 wt.% Cu alloys. SEM.

392
A structure consisting of light and dark etching bands or striations is observed in the intermediate velocity range shown in figure 1. Figure 4 shows a transverse and a longitudinal section

through the weld centerline for a 15 wt.% Cu alloy scanned at 30 cm/s. The lower part of the melted

region, where the solidification velocity is lowest, is dendritic. The upper part of the melted

region is banded. As seen in figure 4b the bands are parallel to the expected shape of the trailing

edge of a melted zone. A transition from dendritic to banded structure, observed in the same sample

suggests strongly that the origin of the banding phenomenon is growth related rather than an artifact

of nonsteady beam power [3] or melt convection [5].

Fig. 4 (a) Transverse and (b) longitudinal sections of a Ag-15 wt.% Cu

alloy scanned at 30 cm/s, showing transition from (unresolved)

dendritic structure to the banded structure. Optical microscopy.

The scan direction is to the right in (b).

TEM examination of the banded structure (fig. 5) in a 23 wt.% Cu alloy scanned at 17 cm/s re- vealed alternate layers of cellular and cell-free alloy. Evidently the dark etching material seen

in optical micrographs of the banded structure corresponds to the cellular layer. X-ray fluorescence

measurement in the STEM with a probe size just large enough to average over the entire width of an individual layer (-'- 0.4 pm) revealed no statistically significant difference in average composition

between the cellular and the cell-free layers. Clearly microsegregation does exist however within

the cellular layers. Care was taken in the preparation of the TEM sample to document the growth

direction. The cellular layer appears to originate microstructurally from an interface instability

occurring in the cell -free layer. The cellular structure then propagates for a short distance until

a layer of solute marks the cell fronts. The continuity between the cell-free layer and the subsequently formed cellular region is seen clearly in the isolated subgrain (dark grain). For higher

scan velocities the widths of the cellular layers decrease compared to the widths of the cell -free

layers, until at high velocity the cellular layers disappear and the banded structure is absent. In eutectic alloys, the banded structure is identical with the exception of a finer cell size.

Fig. 5 High magnification view of the banded structure.

The growth direction is from lower right to upper

left. TEM.

393
The 9 wt.% Cu, 15 wt.% Cu, and 23 wt.% Cu alloys are free of the banded structure at velocities

of 300, 200, and 157 cm/s and above (fig. 1). As discussed in a previous section, actual growth rate may differ from the scan velocity because of changes in weld puddle shapes at scan velocities over

100 cm/s. Clearly however, the banded structure can be suppressed by high scan velocities to produce

a microsegregation-free structure. Furthermore, this velocity decreases with increasing concentration. Figure 6 shows micrographs of a 23 wt.% Cu alloy scanned at 157 cm/s. The transition from a banded structure at the bottom of the melted region, where the growth rate is low, to a band-free

structure at the top of the melted region, where the growth rate is high, is seen in figure 6b. TEM micrographs of this sample (fig. 6c) confirm the absence of bands but show significant solid state

(spinodal) decomposition of the microsegregation-free structures produced from the melt.

Ag-28 wt.% Cu .

Fig. 6 Microsegregation-free structure

obtained in Ag-23 wt.% Cu scanned

at 157 cm/s. Note the presence of

the banded structure at the bottom

of the weld zone in (b) and the

presence of spinodal decomposition

in (c). (a,b) optical microscopy

(severely etched); (c) TEM.

j

The eutectic composition, when electron beam melted at 2.5 cm/s, consists of a fine eutectic

O structure with a spacing of 200 A. The electron diffraction pattern from the eutectic yields latO O tice parameters of 4.05 A and 3.65 A for the two phases. These values are close to those obtained

[10] for Ag and Cu at their respective limits of equilibrium solid solubility at the eutectic temperature. Above 2.5 cm/s the eutectic composition does not solidify by a coupled growth mechanism. The

microstructure consists of the same banded structure observed for noneutectic alloys. Above 70 cm/s

the eutectic composition forms a structure without segregation similar to those shown in figure 6.
Di scussion

Coupled Growth .

The maximum observed growth rate for the Ag-Cu eutectic in the present work is 2.5 cm/s with

O a spacing of ~ 200 A. Cline and Lee [11] have directionally solidified the Ag-Cu eutectic up to a

O growth rate of 0.27 cm/s with an average lamellar spacing of 950 A. In their experiments growth

at higher velocities was prevented by the limitations of heat extraction in their solidification

furnace. They found the lamellar spacing A and the growth velocity, v, to be related by

A^v = 1.4 X 10 cm^/s. (1)

In the present experiments, the heat extraction rate is not limited by a furnace. Extending their

O relation to a growth velocity of 2.5 cm/s predicts a lamellar spacing of 237 A. This compares quite

favorably with the eutectic spacing seen in the present work.

The present results strongly suggest that there should exist a theoretical limit on the rate at

which coupled eutectic growth can occur. One theoretical limit can be found from the inclusion of a temperature dependent diffusion coefficient in eutectic growth theory [12]. The results of Eq. (1),

phase diagram information and diffusion data [13] can be used to show that for slow growth, v = 3.5 x -4 -1 -2 2 10 (cm s K )AT , where AT is the difference between the interface temperature and the eutectic

temperature. Incorporating the temperature dependence of D to extend this relation to large underO cooling, yields a maximum velocity of 4.9 cm/s at a spacing of 78 A. These estimates compare to

O the experimentally determined values of 2.5 cm/s at 200 A. Possible reasons for this discrepancy

and other theoretical limits on the eutectic growth rate are discussed elsewhere [7].

It should be emphasized that the temperature dependence of the diffusion coefficient becomes

important in the present analysis only because of the relatively large undercoolings required to accomplish the solute redistribution and surface area creation necessary for eutectic growth. For

dilute alloys the temperature dependence is much less important.

It is also interesting to note that the finest spacing seen by Boswell and Chadwick [2] in splat

O quench alloys was 300 A. This observation was used by them to assess the growth rate imposed by

the splat quench on the foil. Because this eutectic structure occurred only in small regions adjacent to the metastable extended solid solution, it seems more likely that a maximum growth rate for

the eutectic was reached locally in the small regions and surpassed in most of the foil to form the

metastable structure.

Dendritic Growth in Concentrated Alloys .

To be consistent with the idea that a temperature dependent diffusion coefficient leads to a maximum rate for coupled growth of a eutectic, a maximum rate for dendritic growth of concentrated

alloys must also exist. Although the theory of dendritic growth is continually being refined, especially for high rate [14], a simple expression formulated by Burden and Hunt for controlled growth

[15] with the inclusion of a temperature dependent diffusion coefficient demonstrates the qualitative features of this maximum. The calculated maximum velocity [7] increases for increasing deviation from the eutectic composition. The maximum thus parallels the experimental data for the boundary between dendritic and banded microstructure. The calculated values are a factor of 2 to 4 higher

than the experimental data, consistent with the disparity between the calculated and observed maxima

for eutectic alloys. For very dilute alloys this maximum is irrelevant because it occurs at velocities above those at which the planar interface is stable.

Microsegregation-Free Structure .

For Ag-1 wt.% Cu and -5 wt.% Cu alloys, microsegregation-free structures are produced by planar

growth at velocities predicted by absolute stability theory with equilibrium partitioning [8]. For

concentrated alloys beyond the solidus retrograde composition, equilibrium partitioning is not a possible mechanism to produce a microsegregation-free structure. The transition from banded structure

to microsegregation-free structure for concentrated alloys is most likely related to the onset of

complete solute trapping. The upper velocity boundary of the banded region (fig. 1) when extrapolated to pure Ag suggests a velocity for solute trapping of 5 m/s for dilute alloys. This velocity

is consistent with data obtained for solute trapping in dilutely doped Si [16]. The fact that the

395
eutectic composition can be produced as a single phase microsegregation-free material at lower

growth rates than, for example, 9 and 15 wt.% Cu alloys suggests a significant composition dependence

for the onset of solute trapping.

Banded Structure .

The origins of the banded structure are difficult to determine with certainty. Beck, Copley,

and Bass [6] have recently suggested that the bands are due to an instability proposed by Baker and

Cahn [17] related to growth with solid composition and interface temperature above the solidus.

However the observation of band-free structures in alloys for alloy composition beyond the solidus

retrograde, where single phase growth can never occur below the solidus, seems to negate this idea.

The preceding discussion has suggested that the lower velocity boundary of the banded region

in figure 1 is caused by a limitation of the maximum growth rate of dendritic and eutectic growth

in concentrated alloys. The discussion also has suggested that the upper velocity boundary of the

banded region in figure 1 is caused by significant deviation from interfacial equilibrium and the

concommitant increase of the partition coefficient to unity. Clearly the microscopic events which

occur during the formation of the banded structure must be related to these two phenomena. These

microscopic events are discussed in more detail elsewhere [7].

Acknowledgment

The authors appreciate support from DARPA for this research. We also thank S. R. Coriel 1 for

many helpful discussions.

References

[ 1] P. Duwez, R. H. Willens, and W. Klement, Jr., J. Appl . Phys. 1136 (1960).

[ 2] P. G. Boswell and G. A. Chadwick, J. Mat. Sci. 12, 1879 (1977).

[ 3] W. A. Elliott, F. P. Gagliano, and G. Krauss, Met. Trans. 4, 2031 (1973).

[ 4] S. M. Copley, M. Bass, E. W. Van Stryland, 0. G. Beck, and 0. Esquivel, Proc. Ill

Int. Conf. Rapidly Quenched Metals , 1, p.l47. The Metals Society, London (1978).

[ 5] D. G. Beck, S. M. Copley, and M. Bass, Met. Trans. A 12A , 16 (1981).

[ 6] D. G. Beck, S. M. Copley, and M. Bass, Met. Trans. AIM, 1879 (1982).

[ 7] W. J. Boettinger, R. J. Schaefer, F. Biancaniel lo, and D. Shechtman, Met. Trans. A., to be

publ i shed.

[ 8] W. J. Boettinger, S. R. Coriell, and R. F. Sekerka, these proceedings.

[ 9] J. M. Murray, Met. Trans. A, (1983) to be published.

[10] R. K. Linde, J. Appl. Phys. 934 (1966).

[11] H. E. Cline and H. Lee, Acta Met. 18, 315 (1970).

[12] W. J. Boettinger, in Rapidly Solidified Amorphous and Crystalline Alloys , B. H. Kear,

B. C. Giessen, and M. Cohen, Eds. , Elsevier, p. 15 (1982).

[13] T. Yamamura and T. Ejima, J. Jap. Inst. Met. 901 (1973).

[14] R. Trivedi, to be published.

[15] M. H. Burden and J. D. Hunt, J. Cryst. Growth 22, 109 (1974).

[16] M. J. Aziz, J. Appl. Phys. 1158 (1982).

[17] J. C. Baker and J. W. Cahn, in Sol idi fi cation , ASM, Metals Park, p. 23 (1971).

396
NONDESTRUCTIVE CHARACTERIZATION OF RAPIDLY SOLIDFIED A1 -Mn ALLOYS

R. J. Schaefer*, M. Rosen**, J. J. Smith**, D. Shechtman** and R. Mehrabian*

*National Bureau of Standards

Washington, DC 20234

**The Johns Hopkins University

Baltimore, MD 21218

ABSTRACT

Supersaturated A1 -Mn alloys were prepared by melt spinning and by electron beam surface melting of bulk samples. The samples were evaluated by several methods, in the as- prepared condition and after annealing for several different times at 450° C. The measurements gave consistent indication of the high degree of supersaturation attained in the

rapidly solidified materials, and showed the changes of several properties as precipitation

progressed. At low supersaturation the precipitates were localized at the grain boundaries

and the supersaturation of the lattice disappeared only slowly, whereas at higher manganese

concentrations the precipitates were abundantly dispersed throughout the matrix and the lattice supersaturation decreased more rapidly.

Introduction

Nondestructive methods are needed to characterize the modified layers which can be generated by

the action of directed energy sources on the surface of metals. The information generally required

includes the depth of the layer and its microstructural character. In many cases the surface layer

will not differ in composition from the bulk material, i.e. no surface alloying is present, and the

surface layer is detected through the effect of the modified microstructure on the probing signal.

Manganese has a maximum equilibrium solid solubility in aluminum of approximately 1.55 wt.% at the eutectic temperature [1], which is only slightly below the melting temperature of aluminum. It has long been known [2], however, that this solubility can be increased by rapid solidification to

compositions far beyond the eutectic (2%) between A1 and AlsMn. In response to heat treatment,

precipitation may occur as either the equilibrium orthorhombic Al 5 Mn or other phases [3] depending

upon the temperature and impurity content. Nondestructive evaluation of the state of precipitation

may either detect some direct influence of the precipitates or it may primarily detect the decrease

of Mn in solid solution as the process progresses.

Results and Discussion

Sample Preparation

Melt-spun ribbons with Mn concentrations of 0.1, 1, 2, 3, 7, 9, 12, and 15 wt.% Mn were prepared from 99.999% A1 and 99.95% Mn. The ribbons were typically 2 mm wide and 20 pm thick. For annealing, the ribbons were sealed in borosilicate glass ampoules with 1/2 atmosphere of helium.

Annealing was at 450°C for 5 minutes, 1 hour and 3 hours.

A sample of A1 -7 wt.% Mn was prepared for electron beam surface melting by chill casting, rolling at 600°C to a thickness of 9 mm, and annealing for 1 hour at 450°C. Surface melting was carried out by electron beam scans along the length of the sample at a velocity of 50 cm/s, with

successive scans offset 0.4 mm laterally and with a time delay of 15 seconds between scans to allow

cooling. In this way a rapidly solidified surface layer 8 cm long, 2 cm wide and 0.2 mm deep was

formed.

Sample Evaluation

The melt-spun ribbons were evaluated by measurements of lattice parameter, electrical resistivity, and sound velocity, and by transmission electron microscopy. The measurements indicated that

397
Fig. 1. Variation of the lattice parameter

of the A1 phase as a function of Mn concentration in the as-spun and

annealed conditions.

Electrical Resistivity of

Al-Mn Alloys

Fig. 2. Variation of the electrical resistivity i in Al-Mn alloys as a function of Mn i concentration in the as-spun and an-

| nealed conditions.

at concentrations up to 12 wt.% most of the manganese was in solid solution in the as-spun ribbons. I At 450°C, precipitation occurs heterogeneously at the grain boundaries at low Mn concentrations \ but at many more locations throughout the sample at high concentrations. As a result, the supersaturation of manganese in the aluminum lattice diminishes rapidly in the high concentration samples i but more slowly in the more dilute alloys where longer-range diffusion is required.

j Lattice Parameter Measurements

i

! Lattice parameters were measured on an X-ray diffractometer using CuK^^ radiation, plotting • calculated values for individual peaks against f = (cos^e/sin e) and extrapolating to 'i' =0. In the as -spun ribbons, diffraction peaks were seen only from the aluminum phase at compositions up to I 12 wt.% Mn, and the lattice parameter (Fig. 1) decreased linearly with increasing Mn content, with a ; slope in agreement with prior measurements [2,3]. At 15 wt.% Mn, additional peaks, not corresponding;

to the equilibrium Al 5 Mn phase, appeared in the pattern and the lattice parameter of the A1 phase

indicated a composition of approximately 4 wt.% Mn in solution. Measurements of the top and bottom

surfaces of the 12% alloy showed essentially identical results, indication that the supersaturated

phase was present throughout the thickness of the ribbons.

Annealing of the A1 -7 wt.% ribbon for 1 hour at 350°C showed little change in the diffraction

pattern. After 1 hour at 450°C however, A1 gMn was present and the peaks from the aluminum phase were significantly broadened, with the average value of the aluminum lattice parameter having increased approximately one half of the way to its equilibrium value at that temperature. At 9 and 12

398
wt.% Mn, the 1 hour 450°C anneal gave an aluminum phase lattice parameter close to the equilibrium value but at 3 wt.% Mn or less the same anneal gave no perceptible change in the lattice parameter.

Annealing for 3 hours brought the lattice parameter of the 7wt.% Mn sample close to its equilibrium

value and also caused significant changes in the low composition samples.

The diffraction chart of a 9 wt.% Mn sample annealed for 5 minutes at 450°C showed extreme

line broadening with a line shape indicating that much of the aluminum phase was still close to its

original composition but that localized regions had already become significantly depleted in manganese.

The A1 -7 wt.% Mn sample prepared for surface melting showed initially a lattice parameter for

the aluminum phase indicating only a small concentration of Mn in solution, and strong AlsMn peaks.

After electron beam surface melting, only a single faint peak indicative of a second phase was present and the measured lattice parameter of the aluminum phase lay well below the line for as-spun

ribbons in Fig. 1. The value of the lattice parameter in this case, however, may have been seriously

affected by the large residual stresses which are inevitably present in such surface melted layers.

The surface layer showed a pronounced orientation texture as revealed by the complete absence of the

(111) and (222) peaks in the diffraction pattern.

Electrical Resi sti vity

The electrical resistivity of the ribbons (Fig. 2) was determined by means of a standard four

probe potentiometric method with an accuracy of about 1 part in lO^. Considering the overall uncertainty in the geometrical shape factor of each sample, the estimated error in the absolute resistivity

values is about ±5%. The as-spun ribbons exhibit a strong increase in electrical resistivity with

increasing manganese content. An apparent leveling of the resistivity is observed at concentrations

above 9 wt.% Mn. Annealing at 450°C for even short periods (5 minutes) has a large effect on the

electrical resistivity of the higher concentration alloys. By means of a nucleation and diffusion

controlled process, Al 5Mn particles are precipitated out of the aluminum matrix, thus denuding the

lattice of effective scatterers of electrons. This "purification" process is apparently responsible

for the observed decrease in electrical resistivity relative to the supersaturated, as-spun, alloy.

The difference in the absolute values of the electrical resistivity between the as-spun and annealed

condition, Ap, increases with increasing manganese concentration up to 9 wt.% Mn, and subsequently

levels off. The electrical resistivity of the annealed 12 wt.% Mn alloy is almost identical to that

of the 2 wt.% Mn alloy, indicating that the contribution of the fully grown, and incoherent, AlsMn

precipitates to the measured total resistivity is almost negligible. Such a behavior is expected

from relatively large isolated precipitates, about 5xl0~^m in size, in a continuous matrix phase, in

which case the resistivity of the sample is determined primarily by the resistivity of the matrix.

Electron Microscopy

The microstructures of as-spun and annealed ribbons of the 3 wt.% and 9 wt.% Mn alloys are shown

in Fig. 3. The 3 wt.% Mn alloy shows only dislocations in the as-spun condition, and after 1 hour

at 450°C it shows a few AlgMn precipitates localized at grain boundaries. In the A1 -9 wt.% Mn alloy, the as-spun alloy consists of a cellular structure with very fine precipitates at the cell

boundaries. It is conjectured that the fine precipitates were formed in the solid state during

rapid cooling after solidification. It is not clear whether the' fine precipitates are direct precursors of the large ones that form upon annealing at 450°C. After 1 hour at this temperature,

the 9 wt.% Mn sample shows abundant Al 5Mn precipitates throughout.

X-ray, TEM and electrical resistivity observations all indicate that at high supersaturations

precipitation progresses most rapidly because of the large number of nuclei formed and the correspondingly short diffusion distances. The leveling of the electrical resistivity at concentrations

above 9 wt.% Mn may be the result of the presence of the cellular structure. It is suggested

that the concentration of Mn in the A1 lattice adjacent to the cell boundaries is very low, and in

this case these regions would act as a continuous network of low-resistivity material which would

decrease the electrical resistivity of the sample to a larger extent than a simple volume fraction

would indicate. The x-ray diffraction peaks of the Al-12 wt.% Mn ribbons have distinct tails on the

low angle side, confirming that some aluminum is present with considerably less than 12 wt.% Mn in

solution.

Sound Wave Vel ocity

The velocity Vg of ultrasonic extensional waves is (E/p)^/^, where E is the Young's modulus

and p is the density. The velocity of such waves propagating along the ribbon was determined by

measuring the transit time of a single pulse generated by a laser and detected by a piezoelectric

transducer. The detector was about 200 mm from the spot on the ribbon irradiated by the Q-switched

Nd:YAG laser (Fig. 4). The transit time of the extensional sound wave propagating along the ribbon

399
AI-3 Wt. % Mn

No segragation is observed in the as-spun rib- bon. Following heat treatment, large AleMn

precipitates form at the grain boundaries.

As-Spun Ribbon Heat Treated 450°C-1h

AI-9 Wt. % Mn

Fine particles of metastable compound form

at the grain boundaries of as-spun ribbon. Fol- lowing heat treatment, precipitation of AleMn

is observed.

As-Spun Ribbon Heat Treated 450'’C-1h

Fig. 3. TEM of A1 - 3 wt.% Mn and A1 - 9 wt.% Mn in the as-spun and annealed conditions

BEAM STEERING crystal

Fig. 4. Laser generation and piezoelectric detection

of ultrasonic waves in thin ribbons.

could be determined to within one part in 10^. The separation between the impinging laser pulse

and the quartz receiver was determined to within ± 0.2 mm. The total estimated error in the calculated value of the sound velocity was ± 0.1%. The measurements were carried out at room temperature in the as-spun condition and after the ribbons were subjected to the appropriate heat treatments.

The propagation velocity of Rayleigh surface waves was used to characterize the supersaturated

solid solution generated by electron beam surface melting of the slab of A1 -7 wt.% Mn. The Rayleigh

velocity in the electron beam melted and rapidly solidified layer was found to be 3039 m/s. On the

untreated side of the slab, the microstructure of which consisted of Al 5Mn precipitates in an aluminum

phase matrix, the Rayleigh wave velocity was 3091 m/s, substantially higher than the velocity in the

supersaturated solid solution. This behavior was expected since the specific sound velocity of the

AlsMn particles is substantially higher than that of the aluminum matrix. A similar behavior was observed in A1 -Cu alloys after aging [5].

Figure 5 shows the variation of the sound velocity in A1 -Mn alloys as a function of Mn content

for both as-spun and annealed alloys. The salient feature in Fig. 5 is the fact that the sound

velocity, and thus the Young's modulus, of the heat treated alloy is always higher than that of the

400
Sound Velocity vs. Wt % Mn

in Al-Mn Alloys

Fig. 5. Variation of the extensional sound wave velocity in A1 -Mn alloys as a function of Mn concentration in the

as-spun and annealed conditions.

as-spun state. The difference between the as-spun and the annealed conditions increases with increasing Mn concentration, up to 9 wt.% Mn. This relatively large difference is significant when

nondestructive techniques are considered for monitoring on-line and in real-time the progress of the

metallurgical process leading to a rapidly solidified single-phase supersaturated alloy.

The extensional sound wave velocity initially decreases with increasing Mn content up to about

2 wt.% Mn. Subsequently the velocity increases, followed by a plateau starting at 9 wt.% Mn. The

density, based on the measured lattice parameters of the as-spun ribbons, is expected to be a linear

function of composition. Reported values of the static elastic modulus as a function of composition

[6] show a slight suggestion of nonlinearity, but not enough to account for the pronounced minimum

in sound velocity seen here. It is thus apparent that either the density of the melt-spun ribbons

differs from the value derived from X-ray measurements, or the elastic modulus of the ribbons differs

from that measured statically on bulk samples.

CONCLUSIONS

Nondestructive ultrasonic and electrical resistivity measurements were found to be feasible for

in-process dynamic monitoring and determination of the supersaturation range of rapidly solidified

alloys. Up to 12 wt.% Mn can be incorporated into an A1 lattice by rapid solidification, and the

subsequent precipitation can be monitored by any of several methods.

ACKNOWLEDGEMENTS

The authors thank F. S. Biancaniello for sample preparation and C. Bechtold for consultation

on X-ray diffraction. Thanks are also due to DARPA for financial support.

401
REFERENCES

[1] F. A. Shunk, Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York (1969)

[2] G. Falkenhagen and W. Hofman, Z. Metallkunde 43, 69 (1952)

[3] R. Hoier, S. E. Naess, and E. Nes, Z. Metallkunde 64, 640 (1973)

[4] I. Obinata, E. Hata, and K. Yamagi , J. Japan Inst. Met. 17, 496 (1953)

[5] M. Rosen, E. Horowitz, S. Fick, R. Reno and R. Mehrabian, Mater. Sci. Eng. 53, 163 (1982)

[6] N. Dudzinski, J. R. Murray, B. W. Mott, and B. Chalmers, J. Inst. Metals 74, 291 (1948)

402
EXPERIMENTAL STUDIES OF THE DYNAMICS OF RAPID SOLIDIFICATION

USING ACOUSTIC EMISSION

R. B. Clough, H. Wadley, and R. Mehrabian

Center for Materials Science

National Bureau of Standards

Washington, DC 20234

ABSTRACT

One- and two-dimensional heat flow models for rapid melting and resolidification of

metallic substrates subjected to high energy laser and electron-beam sources are now available, including both stationary and moving heat sources. Methods are accordingly

needed for monitoring these processes in real time during solidification. In this paper

we discuss the use of acoustic emission methods to monitor rapid solidification of precisely controlled electron beam surface melts of commercially pure aluminum and 2219

aluminum alloy. These results are discussed in relation to heat flow calculations and

resol idification microstructures.

The driving forces for the emission are the density changes due to phase transformations and the stresses induced by thermal contraction during and after solidification.

For a restricted range of values of the product 3a, heat flux times the radius of the

circular region over which it is absorbed, steady-state thermal profiles and melt depths

are established. Increasing this product results in increased melt depth and increased

amounts of acoustic emission on resolidification when the beam is switched off. The

amount of emission is in addition much greater in the 2219 aluminum alloy than in the

commercially pure 1100 aluminum. Micrographs show that the 2219 aluminum alloy exhibits

extensive hot tearing while the 1100 aluminum has no cracks, so that we attribute part of

the enhanced signal to hot tearing in the 2219 aluminum alloy. Additional experiments in

which the substrate microstructure was varied were found to have small effect upon solidification cracking but a significant effect upon acoustic emission. This suggests both

plastic deformation and hot heating generate detectable acoustic emission in 2219 aluminum

alloys.

Introduction

Directed high energy sources (laser and electron beams) are currently being evaluated for

rapid surface melting and resolidification of a range of engineering alloys [1]. In these processes,

the bulk substrate, in intimate contact with the molten surface layer, acts as the quenching medium

resulting in high solid/liquid interface velocities and short solidification times during cooling.

These rapidly solidified alloys exhibit enhanced wear and corrosion resistance, arising, it is

believed, from improved chemical homogeneity and the fine scale of microstructure. Such materials

offer promise for applications such as bearing surfaces, cutting tool faces, and corrosion resistant

parts.

At present, critical process variables such as solid/liquid interface velocity or the cooling

rate can only be deduced using one- and two-dimensional heat flow models for melting and resolidification. Experimental methods to measure, i_n situ , process variables are required so that confirmation of microstructure-process variable relationships may be made. Furthermore, surface modification of some alloys is severely limited by the occurrence of cracking during solidification.

In this work, we report the use of acoustic emission techniques to monitor rapid solidification

of electron beam surface melted 1100 commercially pure aluminum and a 2219 aluminum alloy in several

different heat treated conditions. Heat flow theory has been used to set up controlled experiments

in which shallow melts were formed at the surface of the aluminum plates. After attainment of

steady state conditions, the energy source was terminated, resulting in rapid solidification for

which the acoustic emission response was measured. Effects of varying both the substrate composition and electron beam parameters on the acoustic emission response have been investigated and

linked with post-mortem metal lographic observations of microstructure to deduce the origin of the

elastic waves.

403
Heat Flow Calculations

The problem considered was the rapid melting and resolidification of an alloy substrate subjec- 1 ted to a high-intensity stationary heat source over a circular region on its bounding surface. The

^ numerical solutions were developed for an Al-4.5 wt.% Cu alloy which is close to the composition of

the 2219 aluminum alloy (whose principle solutes were 6.6 wt.% Cu, 0.34 wt.% Mn) used in the experiments reported in this paper. It was assumed that the substrate had previously been subjected to conventional solidification-- it segregated and precipitated a weight fraction of eutectic calculated i from the Scheil equation [2].

The generalized two-dimensional heat flow model and solution technique developed in an oblate i coordinate system [3] can readily take both space and time variations of the heat flux into consid- I eration. For the purpose of this paper, it is reasonable to assume that a step function heat flux , is applied at time zero with a uniform circular distribution. The shape and location of the "mushy"

zone, bounded in space by the liquidus and solidus isotherms, are then calculated for various

combinations of process variables. Special emphasis is placed on conditions that result in steady

state heat flow. That is, conditions under which the total amount of heat absorbed at the surface

circular region is exactly offset by the conduction of heat into the substrate interior. This | results in a steady- state thermal field with a maximum temperature at the center of the circular I surface region. This enables a quiescent condition (i.e., with static stresses) to be established i prior to solidification. Two dimensionless variables are used to describe heat flow:

k (T(0,0) - T^)/£a and a/V4o(^t (1)

where k is thermal conductivity, T(0,0) is the temperature in the center of the circular region; T^

is initial substrate temperature, q and a are absorbed heat flux and radius of the circular region,

respectively. a/V4a^t denotes dimensionless inverse time, is thermal diffusivity of the solid,

and t is time. Values of the variables are given in table I.

Table I

Properties of Al-4.5% Cu Alloy Used in Calculations [4,5]

Notation Name Val ue

k

s Thermal conductivity of solid 180.6 Wm-^k-i

Thermal conductivity of liquid 100.8 Wm-^k-i

a

s Thermal diffusivity of solid 7.6 X 10-s m^s-i

Solidus temperature 821 K

Tl Liquidus temperature 923 K

Tv Vaporization temperature 2744 K

In order to account for the release of heat of fusion, the first term of expression (1) was enlarged to:

(k^(T(0,0) - T^) + AC + kg (T^ - T^))/ga (2)

where the subscripts £ and s denote liquid and solid, respectively and Tj^ and T^ are the liquidus

and solidus temperatures, respectively. AC is an equivalent term accounting for heat conduction in j the "mushy" region and is calculated to be 6.2 x 10“^ W/m for the Al-4.5 wt.% Cu alloy [4].

It has been verified that for a given value of expression (2) all isotherms in the alloy

substrate are located at the same dimensionless distance for identical dimensionless times. Figure

1 shows a plot of results obtained from computer calculations relating the two dimensionless variables noted above. A number of important facts are deduced from this figure. First, there is a minimum value of the product £a required if the center of the circular region is to reach a given

temperature T(0,0), e.g. the liquidus T|^ or the vaporization T^ temperatures, of the alloy. That

404
is, when the power is maintained for long times (for small values of a/ V4a^t) the vertical axis

approaches unity—steady state conditions then prevail if T(0,0) < T^. The sum of the second and

third terms in the numerator of the ordinate At + “ T^) = 1.56 x 10^ W/m, where is ambient

temperature. Therefore, for the Al-4.5 wt.% Cu alloy, a minimum value of the product £a 1.56 x 10^ W/m is required if the center of the circular region is to reach the liquidus temperature of

the alloy, T(0,0) = T|^. Similarly, ~ 1.84 x 10^ W/m, a minimum value ga 3.4 x 10^ W/m

is deduced for T(0,0) = T^. The temperature at the center of the circular region for any dimensionless time can also be readily deduced from Figure 1. For example, for a heat flux of 3 = 2.7 x 10® W/m^ absorbed over a circular region of radius a = 10“®m ( 3

a = 2. 7 x 10® W/m) temperatures of

T(0,0) = 1384 K and T(0,0) = 2054 K are reached in t ~ 9.2 ms (a/\/4agt ~ 0.6) and t = 52.7 ms (a/\/4a^t 0.25), respectively. Note that for t > 52.7 ms (a/-\/4a^t < 0.25), the plot in Figure 1 approaches unity on the vertical axis, and steady state conditions are obtained, i.e. a maximum

temperature T(0,0) = 2054 K is reached. Therefore, for products 1.56 x 10® ^ 3a ^ 3.4 x 10® W/m,

a = 10“® m, and times t 52.7 ms temperatures are likely to be in the range T|^ S T(0,0) ^ T^ and

steady state conditions are expected to prevail for the Al-4.5 wt.% Cu alloy substrate.

Figure 1. Temperature at the center of the circular region during melting of a semi-infinite Al-4.5 wt.% Cu alloy substrate as a function of uniform

absorbed heat flux, radius of the circular region and time [4].

Figure 2 shows the location of the "mushy" zone (the region bounded by liquidus and solidus

isotherms) during the approach to steady-state conditions, followed by rapid solidification. The

isotherms here represent the situation in which heat was applied until almost a steady-state condition was achieved, the solidus isotherm reached a plateau, at which point beam power was turned off

resulting in rapid solidification. The vertical distance between the two curves denotes the size

of the "mushy" zone. Upon attainment of a steady-state condition the "mushy" zone extended ~ 1/4

of the maximum melt depth z/a = 0.4 or -v 0.1. Loss of superheat and shallower temperature gradients

during solidification result in significantly larger "mushy" zone sizes. For example, when the

substrate surface reaches the liquidus temperature, the solidus isotherm is at approximately

z/a -v- 0.3--the size of the "mushy" zone is three times larger than when solidification started.

405
0.4 0.8 1.2 1.6 2.0 2.4 2.8 3.2 3.6 4.0

at

FOURIER NUMBER,

a2

Figure 2. The dimensionless size and location of the "mushy" zone along the z-axis versus time for product £a = 2.0 x 10 ^ Wm~^. T923 and Ts 2 i denote the

locations of the liquidus and solidus isotherms, respectively [4].

Average solidus and liquidus interface velocities and solidification times can be readily

calculated from the data shown in Figure 2. Assume that a = 10~^ m, 3 = 2 x 10® W/m^. At for the*

liquidus and solidus to move from their maximum depths, z^^ = 3 x 10“'* m and z^ = 4 x lO”** m, to the*

surface are 3.7 ms and ^ 4.75 ms, respectively. This translates into average liquidus and solidus , velocities of 8.1 x 10 “^ m/s and 8.5 x 10“2 m/s, respectively. Note that while the two average/

interface velocities are almost equivalent, initially the solidus moves much slower than the liqui- | dus resulting in the large "mushy" zone size noted above.

The steady-state shapes and locations of the "mushy" zone, liquidus and solidus isotherms as . a function of both dimensions r and z are readily obtained from the computer calculations. In | Figure 3, we show the positions of the steady-state "mushy" zones for two values of the product £a.

Fractional Distance, r/a 0 .2 .4 .6 .8 1.0 1.2

6 .6

a 2.0 X 105Wm-i

a 2.7 X lO^Wm-i

Comparison of Experimentally and

Theoretically Obtained Steady State Profiles

... Theoretical (Al-4%Cu)

Experimental

(2219 A1Alloy)

Figure 3. Comparison of predicted and experimentally measured steady-state melt

profiles for aluminum-copper alloy

at two val ues of £a.

1

406
Both values are within the range noted above for which S T(0,0) ^ T^. While the theoretically

calculated size of the "mushy" zone does not change significantly with increasing value of 3a, from

2 X 10^ to 2.7 X 10^ Wm“^, Figure 3 shows that the melt depth almost doubles. Experimental melt

profiles of the 2219 aluminum alloy are plotted on the same figure for comparison. Detailed experimental procedures are described in a later section. Agreement between the two is reasonably good

considering the various assumptions of the calculations including a slightly different alloy composition, and constant but different thermophysical properties for the solid, the "mushy" zone, and

the liquid. The calculations are based on complete absorption of the incident energy by the substrate. This is a reasonable assumption for electron beam surface melting.

Acoustic Emission Theory

Acoustic emission is the elastic radiation from changes of stress in a body. Such changes may

be either positive (e.g., elastically constrained thermal expansion during absorption of a laser

pulse) or negative (e.g., the growth of a crack through an already stressed region). The propagation of elastic waves through an observation point within/on the body results in transient elastic

displacements of the point. These are the basic physically observable quantities in an acoustic

emission experiment.

Calculations [ 6

- 8 ] using dynamic elasticity are able to predict acoustic emission waveforms.

To demonstrate the calculations, the transient displacement vertically above four point sources, a vertical dipole, a horizontal mode I loaded crack, a thermal expansion (dilatation), and an inclined

shear loop are shown in Figure 4 [9]. If such waveforms due to actual sources are measured, the

dynamics and other source properties can be deduced [ 11 ].

Experimental

Material

In the present investigation the acoustic emission from two aluminum alloys has been simultaneously studied. One alloy was 1100 aluminum, a 99.0 percent commercial purity aluminum. The second

alloy studied was 2219 aluminum containing 6.6 wt.% Cu, 0.08 wt.% Si, 0.3 wt.% Mn, and 0.104 wt.%

Zr. The 2219 aluminum alloy was in the -T87 condition. Rectangular plate samples were cut from

each alloy with dimensions of 4.0 x 1.5 x 0.4 cm.

Electron Beam Facility

Samples were mounted in the facility shown in Figure 5 and placed in an electron beam heating

unit. To insure attainment of steady-state conditions, the samples were heated with a 100 ms duration stationary electron beam pulse with a circular cross section having a 1 mm radius a. The

electron beam was accelerated through 25 kV and the beam current deduced from the voltage developed

across a resistor placed in series with the line to ground. The heat flux was calculated for each

pulse, and the current varied so that heat flux could be investigated as an experimental variable.

The product ga was systemmatical ly varied between 1.6 x 10^ and 3.4 x 10^ W/m to produce steadystate melt profiles at the long times t 100 ms noted above.

Acoustic Emission Measurements

A schematic diagram of the acoustic emission instrumentation is shown in Figure 6 . A piezoelectric transducer was used to convert the elastic displacements of the acoustic emission into a voltage signal that could subsequently be electronically processed. To optimize sensitivity

(signal-to-noise ratio) signals were bandpass filtered between 0.02 and 1 MHz and a transducer with

high sensitivity in this passband (see inset in Fig. 6 ) was used to measure emission. The gradual

build-up of stress could activate rapid stress drops of ± 1 ms risetime accompanying deformation

and fracture events. The coarse sampling interval required to follow the more gradual contraction

stress results in effectively only random discrete sampling of the fast risetime (high frequency)

signals. In effect, the bandwidth is limited by a 5 ps sampling rate to about 100 kHz. Studies

showed, by varying the low pass filter, that approximately half the typical signal was below

100 kHz. Much of this can be attributed to the filtering effect of the transducer; its sensitivity

(Fig. 6 ) drops about 40 dB between 100 kHz and 1 MHz. Signals were amplified 50 dB for the 2219

aluminum alloys and 70 dB for 1100 aluminum.

407
E

i|

E

8 => 2

Q

c

<D

E

30

= E I

c^l 20

<D SS ^

E *c 0) ® = O

8 o g 10

0.0- 5 tn

|U3

30

15

¥

(c) Dilatation

10

5

f\0

2 16 8 1

3 ^ £

CO

CO 3 I

=

E ©

C

® .= </) E E ©

a> 3 O u »_ t (0 © S

Q.

(0 O

O. O

(0

- .U3

- (b) Horizontal

-

Microcrack

" ‘ * 4 6 8 10

Time//-ts

TimeZ/iiS

Figure 4. Transient epicenter vertical displacements above four point sources: (a) Vertical

dipole; (b) Horizontal microcrack; (c) Dilatation; and (d) Inclined shear loop [9].

Figure 5. Schematic diagram of the experimental facility to measure

acoustic emission during the rapid

solidification of electron beam

surface melted aluminum alloys.

408
Pulsed

Electron Beam

Dual Channel

Transient Recorder

Figure 6 . Schematic diagram of acoustic

emission instrumentation together

with the transducer plus preamplifier response determined by the

breaking capillary method. (Transducer calibration by F. Breckenridge).

The acoustic emission signal and the beam current were digitally recorded on separate channels

of a 2-channel 8-bit transient recorder. For the examination of emission throughout the entire

pulse, a coarse sampling interval of 200 ps (giving a record length of 400 ms) was used for a few

pulses on each alloy. In this study our interest was focused on the acoustic emission during

solidification and for the majority of measurements, digitization was performed at 5 ps intervals

(10 ms record) from just prior to the moment the electron beam was cut off. The digitized data was temporarily stored on floppy diskettes at the remote electron beam installation and transferred,

off-line, to a 1.5M Byte minicomputer for data analysis. In these experiments we are not attempting

to deduce dynamic source information. Wave propagation within the sample is too complex to allow

this at present. Rather, we are attempting to follow the kinetics of solidification through the

temporal variation of acoustic emission activity. This "activity" was arbitrarily (but in a rigorously standardized manner) measured by computing the mean square voltage of the emission as a function of time and then integrating this over time until the signal became lost in noise. This

provided a useful single number (related to acoustic emission energy) with which to describe the

relative "activity" of the different electron beam pulses/samples.

Results

Microstructure

The effect of increasing the product 3a, absorbed heat flux q(Wm~^) times radius of the circular region a, upon melt size and microstructure is shown in Figure 7 for both the 1100 aluminum and

the 2219 aluminum alloy. Figure 3 shows a comparison between the calculated and experimentally

measured depth profiles for the 2219 aluminum alloy. Agreement is good considering various assumptions of the heat flow model including constant thermophysical properties for the alloy. The size

of the melted region for £a values in the steady-state regime increased with 3a. For the same £a,

the alloy, with its lower liquidus temperature, had a slightly deeper melt depth. The 2219 alloy

also had a less clear interface between the rapidly solidified region and the unmelted substrate.

This is marked the "mushy" zone in Figure 8 . It denotes a partially melted region between the melt

pool and the unmelted substrate. Figure 9 shows the effect of 3a upon melt profile for 3

a values

in the vaporization regime. Deep penetration can be observed in both alloys.

Detailed metal lographic observations indicate considerable plastic deformation to accompany

melting and resolidification, and, in the case of the 2219 alloy, considerable hot tearing also

occurred (see Fig. 8 for example). The effect of 3a upon melt depth along the z-axis for both

alloys is shown in Figure 10. Melting occurred at slightly lower £a values for the 2219 alloy and,

for a given 3a, melt depths were always greater for 2219.

409
1100 Aluminum

Figure 7. Optical micrographs of 1100 and

2219 aluminum alloys showing the

effect of £3 upon melt profile

and ensuing microstructures in the

"steady state" region.

I—

0.5mm —Melt

-—Mushy Zone

-—Substrate

I 1 100 p<m

2219 Alloy

qa = 1.9 X 10* Wm-’

qa = 2.7 x 10= Wm*’

qa = 3.3 x 10= Wm-’

Figure 8. Optical micrographs of 2219

aluminum alloy surface melted at £a = 2.7 X 10^ Wm~^ showing

detailed microstructure of the

region of the "mushy" zone at the

bottom of the melt pool.

410
Fractional

Melt

Depth,

zla

2219 Alloy 1100 Aluminum

qa = 4.1 X lO^Wm-’

qa = 5.7 x lO'Wm-’

qa = 6.3 x lO^Wm-’

Figure 9. Optical micrographs of 1100 and

2219 aluminum alloys showing the

effect of 3a upon melt profile

and ensuing microstructure of 3a such that T[0,0] > T^.

Figure 10. The effect of £a upon melt depth

for 1100 and 2219 aluminum alloys.

Melt depth is measured from the

liquidus isotherm to the substrate

surface for the 2219 alloy.

Acoustic Emission

The temporal behavior of the acoustic emission signal, together with absorbed heat flux, are shown in Figure 11 for ga = 2.7 x 10^ Wm“^ for both metals over the entire beam pulse. It can be

411
3.0 X 1Q5

Figure 11. Acoustic emission generation

during pulsed electron beam

heating (£a = 2.7 x 10^ W/m) and

subsequent cooling of 1100 and

2219 aluminum alloys. Note that * 1100 aluminum measurements were

made at a lOx higher gain, but

even so generated orders of magni-;

tude less emission.

seen that very much more acoustic emission is observed from the 2219 alloy than from 1100 aluminum

both while melting was occurring and during rapid solidification after beam cut-off.

The acoustic emission energy during the period of solidification (during and after beam cut- off) is shown, as a function of qa, in Figure 12. Figure 12(a) and (b) focuses on the AE behavior

of 2219 material in a -T87 condition in the thermoelastic region, where no melting occurs, (12(a),

qa = 1.36 x 10^ W/m) and the melt region (12(b), £a = 2.70 x 10^ W/m), showing the beam current

(dotted line) and AE response (solid line). Note that melting in this case often produces an extra

pulse of AE during solidification. In Figure 13 the cumulative energy of the AE signals during

solidification is shown for both the 2219 and 1100 materials. Each of the data points are the

average of 15 separate electron pulses on three different, but nominally similar samples. It can

be seen that the acoustic emission is typically a hundred times more energetic in the 2219 aluminum

al loy.

Hi

_IUJib< CO z z O liJ co^

CO ~ SS 5 <

TIME IN MILLISECONDS

TIME IN MILLISECONDS Figure 12. Effect of 3a level on the type and quantity of AE during rapid solidification of 2219

aluminum, (a) £a = 1.36 x 10^ W/m. (b) 3a = 2.70 x 10^ W/m.

412
|-*-T<T|_ I"*— T|_<T<Tv l*^T>Tv —

qa, W/m

<

Figure 13. The effect of ^ upon acoustic

emission after electron beam cutoff for steady-state conditions

in (a) 2219-T87 and (b) 1100 aluminum alloys. Note the different

scales for the two alloys.

Acoustic emission during solidification was measured from samples with different substrate

microstructures. Table II shows that the amount of AE during rapid solidification is dependent on the microstructure of the substrate. In this experiment, 2219 aluminum was tested in the solutionized condition (75 min. at 535 °C, ice water quench), and -T6 conditions (aged to peak hardness at

177 °C for 18 hours). The grain size was 50 pm. The table indicates that the T6 condition emitted

the greater AE.

Table II

Acoustic Emission After Beam Cut-off as Affected by Substrate Material Condition

Material Rockwell B Energy (Relative Units)

Condition Hardness qa g 1.5x10^ W/m £a > 1.5x10^ W/m Total

2219-

Sol utioni zed

39

1 0.008 0.013 0.021

2219-T6 77 0.056 0.117 0.173

Di scussion

During electron beam heating and subsequent cooling, both 1100 and 2219 aluminum alloys generate detectable levels of acoustic emission but that from the 2219 alloy is 1 to 2 orders of magnitude greater in intensity (Fig. 11). Here, we would like to deduce the physical origin of this

emission. To do this we shall use our understanding of the differences in the behavior of the two

alloys during solidification, in conjunction with elementary elastodynamic calculations, to deduce

the stress change mechanisms that generate detectable signals.

413
Every change in stress radiates elastic waves. Due to noise and to sensitivity limitations

not every stress change generates detectable elastic waves. Many of the stress changes occurring

during solidification of the 2219 aluminum alloy are clearly detectable, while very few in 1100

aluminum are of sufficient strength to satisfy detectability criteria. To develop this argument

further, it is essential to elucidate the detectability criterion for this system. That is, what

is the relation between the measured acoustic emission energy and the source stress change?

We have devised a rather simple method for this calibration in which we use the predictable

thermal stress induced by the absorption of a short duration laser pulse to generate elastic waves

at the natural source position. For the experiments we used a 25 ns duration 12 mJ energy Nd:YAG

pulse from a Q-switched laser as the artificial source.

This source generated a signal whose energy, when measured using the same conditions as for

the 2219 aluminum alloy experiments, was 0.30 units. From this, since the smallest detectable

energy was about 0.0005 units, we estimate the smallest detectable dipole source to be about 1 pNm

[7], Differences in the radiation patterns of various sources are likely to be reduced because of

the multiple paths of reflected rays that contributed to measurements.

Origin of Acoustic Emission During Cooling

From above, it is clear that only those stress changes for which D^.j 1 pNm will give detectable acoustic emission in these experiments. During the electron beam treatments on 1100 aluminum

we have observed very weak acoustic emission signals: presumably, the stress changes accompanying

melting/resolidification here have dipole components < 1 pNm. The 2219 aluminum alloy has generated

intense acoustic emission signals: the addition of 6.6 wt.% Cu has resulted in many stress changes!

during processing for which D.. >> 1 pNm. We can use these differences to aid our deduction of the

acoustic emission sources.

Since the solidification interface velocity, cooling rate, and thermal constants do not vary, appreciably between the two alloys we would expect the development of thermoelastic contraction

stress to be similar in both alloys. Since little or no acoustic emission was observed from the

1100 alloy it is reasonable to assume this potential acoustic emission is undetectable in these'

experiments. Very high contraction stresses undoubtedly develop, but in solidification from a' steady state condition the cooling rate is presumably too slow to generate detectable elastic waves

above 20 kHz frequency. The more rapid cooling of short duration, high flux density laser pulses

might be measurable however.

The maximum potential thermal stress that could develop upon cooling can be estimated from the

relation:

Aa = (A. + 2/3p)pAT

where AT is the temperature drop, about 600 °C in aluminum. This results in a stress of 900 MPa,

many times the yield strength of both alloys. Thus, during cooling, we expect plastic deformation

to relax the thermal stress. Each dislocation that slips will radiate elastic waves and could be a detectable emission source.

Suppose a unit of slip consists of the simultaneous expansion of n dislocation loops each

expanding over an area dA on the same slip plane. Then the dipole matrix [8] is given by

D . . = n C . . b. dA. ijk£ k £ s np b dA

For detectability this must exceed 1 pNm leading to the criterion ndA ^ 0.13 mm^.

Previous studies reviewed in reference [10] have indicated the fraction of slip events during

uniaxial plastic deformation able to generate detectable signals is highly sensitive to composition

and microstructure. For grain sizes below about 100 pm it is unlikely the 1100 aluminum alloy

could generate detectable signals because n would have to be unrealistically large to satisfy the

detectability criterion. In the alloy, during solidification a spectrum of microstructures exist

ranging from hot solid solution to the peak hardness precipitate distribution. Previous studies

O have indicated that deformation of aluminum alloys containing small (g 100 A) diameter shearable

precipitates generates intense acoustic emission due to cooperative slip phenomena. Variation of

the substrate microstructure, which we expect to have only a small effect upon processes other than

414
plastic deformation, resulted in large changes in acoustic emission. The most active sample con- dition (T6 ) was the one containing a mixture of 0

" and 6 ' precipitates. Thus, for 3

a > 1.56 x 10^ Wm“^ a part, at least, of the extra acoustic emission from the 2219 aluminum could be due to plastic deformation. While for 3a < 1.56 x 10^ Wm~^ (for which no melting occurred) probably all

the emission is from plastic deformations.

For 3a > 1.56 X 10^ Wm“^ melting resulted, and upon subsequent resolidification hot tearing

occurred raising the possibility of an extra emission source. The dipole strength for formation of

a circular microcrack under mode I stress a is approximately given by

D. . = (A + 2p) 2(l-v^) — 7t E

where v = Poisson's ratio, r = crack radius, and E = Young's modulus. Substituting physical property

values for aluminum and assuming a = 100 MPa yields the detectability criterion r ^ 1 pm. Many

cracks clearly satisfied this criterion. For these cracks to generate detectable signals they

should also propagate at speeds greater than 1 ms~^. Data on the micromechanisms of crack extension

during hot tearing are sparse. Metal lographic evidence indicates cracks have propagated beyond the

prior melt boundary into unmelted material suggesting fast fracture may indeed have occurred.

However, such evidence is far from conclusive and, because of the pronounced substrate effect, for

the moment, we may only speculate that a part of the emission from samples for which qa > 1.56 x 10 ^ Wm"^ was from cracking.

Acknowledgments

We are grateful for the advice and assistance of R. Schaefer, C. Brady, J. Chang,

F. Breckenridge, and C. Turner and the financial support of the Defense Advanced Research Projects

Agency under DARPA Order No. 4275.

REFERENCES

[ 1] R. Mehrabian, Int. Metals Rev., 1982. ^(4), 185.

[ 2] M. C. Flemings, Solidification Processing (McGrawHill, New York, NY, p. 142, 1974).

[ 3] S. C. Hsu, S. Kou, and R. Mehrabian, Metall. Trans., 1980, IIB, 29.

[ 4] J. A. Sekhar, S. Kou, and R. Mehrabian, accepted for publication Metall. Trans.

[ 5] S. C. Hsu, S. Chakravorty, and R. Mehrabian, Metall. Trans., 1978, 9B, 221.

[ 6 ] J. A. Simmons and R. B. Clough, "Theory of Acoustic Emission," Dislocation Modeling of Physical

Systems, J. Hirth and M. Ashby, Eds., Acta/Scripta Metallurgica, Pergamon Press (1981).

[ 7] H. Wadley, C. Stockton, J. Simmons, M. Rosen, S. Ridder, and R. Mehrabian, "Quantitative

Acoustic Emission Studies for Materials Processing," DARPA/AFML Annual Review of Quantitative

NDE, held in 1982.

[ 8 ] H.

[ 9] C.

[10] H.

[11] H.

N. G. Wadley, C. B.

B. Scruby and H. N.

N. G. Wadley and R.

N. G. Wadley, C. B.

Scruby, and J. E. Sinclair, Phil. Mag., 1981, 44, pp. 249-274.

G. Wadley, Accepted for publication in J. Phys. D, 1983

Mehrabian, to be published.

Scruby, and G. Shrimpton, Acta Met. 1981, 399.

415
CRYSTALLIZATION KINETICS OF AMORPHOUS ALLOYS BY ULTRASONIC MEASUREMENTS

1

t

I

M. Rosen*, E. Horowitz*, H. N. G. Wadley**, and R. Mehrabian**

*The Johns Hopkins University, Baltimore, MD 21218 and

the National Bureau of Standards, Washington, DC 20234

**The National Bureau of Standards, Washington, DC 20234

ABSTRACT

The crystallization kinetics in amorphous PdCuSi and CuZr ribbons, prepared by the

melt spinning technique, were determined from ultrasonic wave propagation measurements,

microhardness and optical metallography. Ultrasonic waves were generated using a laser

pulse. The waves were piezoelectrically detected after propagation over a predetermined

distance. From the variation of the elastic moduli with crystallization time at isothermal

holding temperatures the activation energies and the growth laws of the processes were determined. The elastic moduli of the alloys increase by about 40% during the transition

from the amorphous to the crystalline state. Furthermore, the sound velocity and the

elastic moduli were found to be sensitive to structural relaxation processes occurring in

the disordered amorphous state prior to crystallization. The structural relaxation is

accompanied by a significant increase in hardness, as revealed by microhardness measurements.

It is conjectured that the frequency dependence of the ultrasonic attentuation, as interpreted within the framework of the theory of viscoelasticity, may explain the hardening

process in the amorphous state. Metallographic examination showed that the crystallization

process in the ribbons proceeds by rapid nucleation at the free surfaces followed by

frontal linear growth of the crystallized phase towards the center of the ribbon. In the

interior, however, random nucleation and growth occur. The observed growth kinetics are in

agreement with theoretical considerations. The effect of composition on the Young's moduli

of the amorphous and crystalline states has been investigated in the Zr-Cu system. In contrast to the crystalline structure the amorphous state exhibits a smooth S-curve in the

variation of the moduli with copper content. The inflection point at 50% Cu is interpreted

as the point of maximal configurational entropy of the amorphous alloy with respect to the

Cu-Zr atom pairs.

Introduction

Formation of metallic glasses through rapid solidification techniques and their subsequent

crystallization upon heating has been the subject of numerous investigations. The extensive interest

in the materials is due to their unusual mechanical, physical and chemical properties. While some practicing engineers are vigorously pursuing new alloys and processing techniques to develop

commercially attractive amorphous products for use in electronic and power industries, others are recrystallizing specific alloys through selected heat treatments into microcrystalline structures

which exhibit high coercivltles or lend themselves to consolidation into bulk products. Characterization of the range of microstructures thus produced using nondestructive techniques is an attractive

proposition both in terms of its contribution to the understanding of the kinetic phenomena at play

during the various transformations, as well as its ultimate commercial utilization for on-line feedback

control of process variables.

Metallic glasses are not thermodynamically stable, and they tend to structurally relax and

finally crystallize upon appropriate heat treatment. Waseda et al. (1) showed that annealing metallic

glasses increased the short-range order. Egami (2) used energy-dispersive methods to confirm the logtime kinetics of the relaxation process. His work indicated that the activation energy for relaxation

continuously increased with time, which may correspond to the removal of quenched-in defects. The process of crystallization into a stable state involves drastic variations in the properties of the material

416
Although changes in density and electronic structure associated with the crystallization process are rather minute (3), appreciable variations were observed in several physical properties, including the

elastic moduli (4-6). The crystallization behavior of metallic glasses was examined by several

workers (7-10). It was found that small additions of noble metals (Cu, Ag) to Pd-Si greatly stabilizes

the amorphous phase, and particularly in PdQ. 775^'^0.06S^0. 165 glass ( 11 ) which became the subject of numerous investigations. The effect of isothermal annealing on the crystallization kinetics of

P‘^0.775^'^0. 06^10. 165 determined by differential scanning calorimetry ( 11 ).

Associated with relaxation and crystallization of metallic glasses are variations in the elastic

and mechanical properties. Young's modulus E and shear stiffness y generally increase by 20-40%, but

the bulk modulus K by only about 7% upon crystallization (5). Kursumovic and Scott (12) found that E for Cu5oZt40 increased by about 10% due to structural relaxation and another 15% upon crystallization.

In contrast, the density of glassy metals is generally only 0.3-1. 5% less than the corresponding

crystalline structures. As-quenched metal glasses contract approximately 0.5 volume percent upon

relaxation. Freed and VanderSande (13) found that after aging Cu45 Zr54 at 375°C, the tensile fracture

stress Of and Knoop microhardness numbers (KHN) rose to a maximum from the as-spun state and then

fell to a minimum before finally increasing again. They ascribed the initial "humps" in Of

and KHN to be due to structural relaxation, and the final increases to crystallization. Davis (14)

used the simple relation Hy = 30y, where Hy is the Vickers micorhardness and Oy is the yield

stress, to estimate strengths of metallic glasses. He found that observed Hy/Oy values for

Ni^gFe 29P]^ 4 B^Si

1^2 (Metglas"^^ 2826B) and PdQ 775 CUQ q^SIq agreed reasonably well with

theory.

Straightforward elasticity and wave propagation theories enable one to calculate E (15). For one-dimensional extensional wave propagation in a homogeneous, isotropic, linearly elastic solid.

Young's modulus is given by E = Vg • p where Vg is the extensional wave velocity and p is density.

Objectives

The objectives of the present investigation were as follows:

• Determination of the isotheirmal transformation kinetics from the amorphous to the crystalline

state by means of monitoring changes in sound wave velocity (elastic modulus)

• Calculation of the kinetic parameters: activation energies and crystallite growth regime.

Comparison with experimental evidence.

• Information about the relaxational changes occurring in the amorphous state, prior to crystallization.

• Corroboration with optical and electron microscopy, x-rays and microhardness concerning the

relationship between microstructure and properties.

Experimental

Metallic glass ribbons were prepared using the melt-spinning technique (16). A quartz shooter

with 0.5 mm diameter nozzle, contained a PdCuSi or CuZr crystalline button. An induction coil melted

the button, and the shooter was then pressurized with helium to 70 KPa to inject the molten alloy

onto a balanced oxygen-free, high-conductivity copper disk spinning at 7,000 rpm. A 2 mm wide,

30 pm thick ribbon was produced. Ultrasonic, x-ray, and STEM examinations indicated that the ribbon

was completely amorphous. The as-spun ribbon was cut into 200 mm sections and placed into Vycor

tubes. After evacuating the tubes to approximately 5 x 10~^ torr, they were sealed. The encapsulated

ribbons were then heat-treated in a forced-air convection furnace. For each Isothermal heat treatment,

the temperature was held constant to within ±20°C. Three heat treatment temperatures were chosen.

The first was near the crystallization temperature Tc , the second near the glass transition temperature

Tg, and the third intermediate temperature Tf. The experimental technique. Fig. 1, for measuring vg,

and thus E, has already been described elsewhere (17).

The velocity of the ultrasonic extensional wave Vg (Vg = (E/p) , where E is the Young modulus

and p is the density) propagating along the ribbon was determined by measuring the transit time of a single pulse generated by, a laser and detected by a piezoelectric quartz-crystal transducer located

at a distance of about 200 mm from the spot on the ribbon irradiated by the laser. The transient

load was applied by rapid deposition of energy from a single-pulse of a Q-switched neodymium glass

laser with a wavelength of 1.06 pm. The laser-pulse duration was 15 ns, and the pulse energy, for

this specific series of measurements, was about 20 mJ. The deposition of the laser energy onto the

417
absorbing specimen produces a state of unbalanced compressive thermal stress as the locally heated

material tends to expand but is restrained initially by its own mechanical equilibrium. The elastic

stress wave propagates along the specimen and may be detected by a piezoelectric transducer. The

laser pulse was line-shaped to produce a nearly plane wave. A photodetector was used to trigger a transient pulse recorder as the laser pulse was generated. Thus, the transit time of the extensional

sound wave propagating along the ribbon could be determined to within better than 1 part in 10^. The

separation (about 200 mm) between the impinging laser pulse and the quartz receiver was determined to within ±0.2 mm. The total estimated error in the calculated value of the Young modulus was ±0.2%.

The sound velocity measurements were carried out at room temperature.

Laser Generation and Piezoelectric-Sensor

Detection of Ultrasonic Waves

in Thin Ribbons

SEAM STEERING CRYSTAL

TRIGGER

Fig. 1 Laser generation and piezoelectric

detection (PZT transducer) of ultrasonic waves in thin ribbons.

Fig. 2 Variation of the extensional wave

velocity, Vg, in PdCuSi ribbons

as a function of crystallization

time at 380, 390 and 400°C.

In this work, ultrasonic extensional waves were detected using a National Bureau of Standards

conical PZT transducer having a flat response between DC and 2 MHz. Signal amplification was provided

by an EMI 6306 charge-sensitive amplifier powered by an EMI Power Supply Type 6396*. A silicon

photodiode triggered a Nicolet Explorer III digital oscilloscope so that the elastic wave could be

recorded and transit times read directly from the screen of the oscilloscope.

Finally, small sections of each heat treated specimen were mounted in epoxy for Knoop microhardness

number (KHN) measurements. To minimize anvil effects and to ensure the KHN values were taken on the

bulk material, each specimen was cut perpendicular to its length vector, and then was mounted with

its cut face exposed. After polishing the mounted ribbons to 0.05 pm alumina, KHN values were determined using a lOg load and averaging four symmetrical indentations.

Results and Discussion

Fig. 2 shows the variation of the extensional wave velocity, Vg, as a function of crystallization

time at three isothermal holding temperatures, 380, 390, and 400°C. The initial sound-wave velocity,

in the amorphous state, was found to be 2970 m. s“l

, whereas in the fully crystallized state it reached

an asymptotic value of 3490 m.s~l. Taking 10.52 and 10.69 g.cm"^ as the density (5) of the amorphous

and crystalline PdQ^ 775 CuQ^Q5 SiQ^]^ 55 , respectively, the Young moduli could be calculated. They were

found to be 9.28 MPa for the amorphous state and 13.0 MPa for the crystalline state, i.e. AE of

about 40%. The present values of Vg and E for the amorphous state are somewhat higher than those

quoted by Davis et al. (18) for Pd-Cu-Si.

Fig. 2 exhibits a relatively sharp increase in the sound velocity during the first 5 minutes of

the isothermal holding. Subsequently, typical sigmoidal curves are developed, depicting the

crystallization kinetics.

*A11 references to commercial equipment in this paper are for identification purposes only and in no way constitute any endorsement or evaluation of the relative merits of such equipment.

418
The initial increase in the extensional sound velocity, of about 3.7%, i.e. 7.4% in the Young

modulus, appears to be related to a structural relaxation that precedes the crystallization process

in the amorphous state (19). Structural relaxation was found to be particularly enhanced in glasses

obtained at high quenching rates, where greater structural disorder is to be expected. The lower

elastic stiffness in the glassy state has been attributed to the interatomic displacements inherent

in a disordered structure (20). The elastic stiffness should, therefore, increase with increased

local ordering as a consequence of the structural relaxation. This is compatible with the increase

in the sound velocity, as shown in Fig. 2. Within the framework of the present investigation, the

variables of this thermally-activated relaxation process were not studied in detail.

Fig. 3 Crystallization sequence of amorphous PdCuSi as a function of time

at a constant temperature, by optical metallography.

The crystallization process in Pd-Cu-Si involves a thermally-activated nucleation and growth

process dominated by the frequency of nucleation, and a diffusional growth. Turnbull (21) adopted

the classical nucleation theory (22) for glass formation and crystallization in the amorphous state.

Formation of critically-sized stable nuclei capable of further growth is characterized by a series of

concentration fluctuations. A steady state nucleation rate can be expressed as I = ke“(0+A)/RT^ where

k is a constant, A and Q are activation energies of formation and growth, respectively. In terms of

the reciprocal rate theory Int = (K + Q/RT + A/RT) or 91nt _ Q ^ i ^ 9A . By assuming Q >> A,

9(1/T)“ R R RT 9(1/T)

then 91nt _ Q. Thus, if t is the crystallization time for a specific fraction of volume

9(1/T)~ R

transformed from the amorphous into the crystalline state (Int) vs. (1/T) gives a value proportional

to the activation energy of the process. The activation energy for 50% crystallization of amorphous

^‘^0.775^'^0.06^^0.165» obtained from the ultrasonic measurements. Fig. 2 was found to be 232 kJ/mole.

419
This value is significantly lower than 400 kJ/mole determined by calorimetry (10), but is in rather

close agreement with values of activation energies for recrystallization in other glasses (23.)

Optical metallography indicates, however, that the crystallization process in Pd-Cu-Si ribbons

is not homogeneous in the bulk. Fig. 3 is a view from the thin edge of a ribbon heat treated for

various times at 390°C. Two coexisting growth processes are apparent. One is rapid nucleation on

' the free surfaces, followed by a frontal growth from the external surface of the ribbon towards the

interior. The other growth process involves nucleation of crystallized regions in the bulk, followed

by spherical growth. The growth rate was determined from the metallographic observations. It was found that the frontal growth is linear with respect to crystallization time, in agreement with ^ theoretical considerations due to Cahn (24). It should be mentioned in this context that the ultrasonic

measurements detect the kinetics of the bulk growth, rather than the specific linear growth of the fl advancing front, as was determined by metallography. Furthermore, this study demonstrates the

importance of the specimen geometry in the analysis of the crystallization kinetics in thin ribbons.

The Johnson-Mehl analysis (25) should be used with care by considering the possibility of simultaneous

operation of different nucleation and growth processes that exhibit different reaction kinetics. i Therefore, experiments specifically designed to isolate nucleation and growth processes are necessary i for a better understanding of the contribution of each to the total crystallization process.

Figure 4 shows the effect of copper content on the Young's modulus of Zr-Cu alloys in the

crystalline and amorphous state. The lowest value of the Young's modulus is for pure crystalline

zirconium. A maximum and a minimum are found for alloys with compositions near the intermetallic

compounds Zr2Cu and ZrCu, respectively. The variation of the Young's modulus with composition of the' amorphous alloys, between 30 and 70% copper, is a smooth S-curve with an inflection point at 50 ^/o

copper. It is conjectured that the amorphous state can be interpreted in terms of the Bragg and

Williams approximation for short range order where the mixing pairs Zr-Zr, Cu-Cu and Cu-Zr are of

equal probability at the composition of 50 ^/o Cu. The inflection point on the Young's modulus versus composition curve will then represent the point for maximal configurational entropy of the amorphous

alloy. The theoretical model should be corroborated by appropriate X-ray analysis in order to establish its validity.

Heat Treatment Time, Min.

Fig. 4 Variation of Young's modulus

with copper content of amorphous and crystalline Zr-Cu

alloys.

Fig. 5 Variation of microhardness

during crystallization of

amorphous Zr 5 oCu 5o alloy at

a constant temperature.

The initial increase in the extensional sound velocity, of about 7.4% in the Young modulus of _ PdCuSi and a similar increase in the Zr-Cu amorphous alloys, appears to be related to a structural If relaxation that precedes the crystallization process in the amorphous state. Structural relaxation ^ was found to be particularly enhanced in glasses obtained at high quenching rates, where greater

structural disorder is to be expected (26). The lower elastic stiffness in the glassy state has been S attribtured to the interatomic displacements inherent in a disordered structure (20). The elastic

stiffness should, therefore, increase with increased local ordering as a consequence of the structural!

relaxation. This is compatible with the increase in the sound velocity, as shown in Fig. 2.

420
Microhardness measurements. Fig. 5, on the ribbon edge, perpendicular to the length vector of

the ribbon, revealed that while in the amorphous, precrystallization state, the material undergoes

some significant changes in its mechanical properties. Fig. 5 depicts a plot of KHN (Knoop Hardness

Numbers) log-time for Zr 5 QCu 5Q heat treated at 455°C. This isothermal treatment significantly

alters the hardness of the amorphous alloy before crystallization ensues. In agreement with Freed

and VanderSande (13) for Zr 54Cu 45 ribbons, microhardness first increased from the as-spun value to a relative maximum of 88%. With continued aging, Knoop microhardness fell to a shallow minimum before

rising to a final asymptotic value of approximately 720 kg. mm~^. This last point contrasts results

by Freed and VanderSande in that their KHN values continued rising even after 600 minutes at 375°C

(15°C below Tg). Using DSC and TEM, they determined that the volume fraction of crystalline material

was between 0.80 and 0.90 after 600 minutes. The asymptotic Knoop hardness values found in this

work, however, represent the fully crystalline state.

1 ! The embrittlement mechanism in the amorphous alloy can be related to the removal of the quenchedin free volumes (vacancy clusters). Initially, the amorphous alloy has a disordered packing of the

j| constituent atoms. The reduced coordination of the atom complexes, due to the present of excess free

volumes, allows shear transformations to be sustained. The free volumes are analogous to dislocation

I core segments. Consequently, the material is very ductile in the disordered amorphous state.

! Heat treatment in the amorphous region at temperatures below the glass transition temperature,

i Tg, or for relatively short periods of time at temperatures above Tg causes structural and compositional

I relaxation. The free volumes are redistributed by short range shuffling of the atoms, thus a more

i! rigid close packed atomic distribution of high coordination is achieved. The dense random packing of

ji hard spheres (Bernal structure) leads to a significant embrittlement of the amorphous phase.

The topological and compositional relaxation processes occurring in the amorphous state can be

; investigated by analyzing the behavior of the shear and compressional components of the ultrasonic

i attenuation, and their frequency dependence. From the theory of vicsoelasticity > fot a shear

I relaxation:

U = 2 2

1+0)2 T

2 + iha 0)T,

1+0)2 ^

2

where, p and are the relaxed and unrelaxed shear

s s moduli, respectively, o) is the vibration frequency, and Xg is the shear relaxation time.

Ts = hg/V® and = ^”^s where, and Pg are frequency dependent and frequency independent

1+0)2 Tg 2 shear viscosities, respectively.

In addition, the compressional viscosity can be expressed in terms of the relaxation time for a compressional process, Tq, the relaxation time for a shear process Xg, and the bulk modulus,

B, of the material: _ ®x„

The compressional and shear viscosities of the medium, and Pg, respectively, are related

to the experimentally determined attenuation, a, in the following manner: ' ' ' < / -

where v is the sound velocity and p is the density of the amorphous material.

Thus, from the frequency dependence of the ultrasonic attenuation, obtained by means of the duallaser interferometer, the behavior of the bulk and shear viscosities can be characterized. Consequently,

it is possible to establish the functional relationship between the appropriate relaxation times and

the diffuslonal reactions leading to topological and compositional changes in the amorphous state.

This specific phase of the investigation is currently in progress. It can be realized through the

capability of the laser generation-laser interferometric detection of propagating sound waves, and

the subsequent analysis of their frequency dependence (27).

Summary and Conclusions

Structural relaxation and crystallization processes, induced by isothermal heat treatment above

the glass transition temperature in CuZr and PdCuSi amorphous alloys was studied by means of

ultrasonics, microhardness and metallography.

421
• Sound wave velocity versus log-time of isothermal heat treatment exhibited isokinetic sigmoidal

behavior. Young's modulus increased by about 40% when the alloys transformed from amorphous to crystalline state.

• Structural relaxation, before the crystallization process ensues, is manifested by an increase

in Young's modulus and by a drastic increase in microhardness.

• The eutectic-like decomposition of the amorphous phase in melt-spun ribbons occurs by rapid

nucleation and linear growth at the free surfaces, and by homogeneous nucleation and linear

growth in the bulk of the ribbon.

• The activation energies for crystallization were for PdCuSi and Zr5oCu5Q were found to be 232

and 360 kJ/mole"^, respectively.

Acknowledgment s

The authors are thankful to F. Biancaniello for preparation of the specimens. Fruitful discuss

with S. Fick are greatly appreciated. Support from Defense Advanced Research Projects Agency (DARPA)

under DARPA Order No. 4275 is gratefully acknowledged. Technical monitor of this contract was Lt. Col. L. A. Jacobson.

References

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20 . D. Weaire, M. F. Ashby, J. Logan and M. S. Weins, Acta Met. 19 , 799 (1979).
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27. H. Palmer, S. Pick and M. Rosen (in press).

423
OXIDATION RESISTANCE OF A FINE-GRAINED

RAPIDLY SOLIDIFIED 18-8 STAINLESS STEEL

i

G. J. Yurek

The H. H. Uhlig Corrosion Laboratory

Massachusetts Institute of Technology

Cambridge, MA 02139

ABSTRACT

A fine-grained, rapidly solidified 18-8 stainless steel (AISI 303) exhibited superior

resistance to oxidation under both cyclic and isothermal conditions compared with a conventional wrought alloy of the same composition. The enhanced resistance to oxidation is at- tributed to three factors: (

1 ) the formation and growth of protective scales is promoted

by the fine alloy grain size C^lOym) and by the presence of a dispersion of fine MnS precipitates in the steel matrix, ( 2 ) scale adherence is increased by the formation of intrusions

of oxide from the external scale into the alloy, which form along closely-spaced alloy grain

boundaries and probably around the dispersed phase, and which act to key the scale mechanically to the alloy, and (3) growth stresses in the scales are reduced by the formation of

scales of uniform composition, which are promoted by the unifoi^m composition of the alloy.

Modification of the microstructures of 18-8 stainless steels through rapid solidification

processing appears to allow an increase in operating temperature of at least 25°C under cyclic oxidation conditions and about 75°C under isothermal conditions.

INTRODUCTION

Resistance of an alloy to oxidation at elevated temperatures requires the development of a protective scale on the alloy surface. A protective scale is one that provides an effective barrier between the alloy and its environment. To be an effective barrier requires that solid-state diffusion

of cations and anions through the scale be very slow. Because scales that meet the latter requirement

usually comprise Cr203 , AI 2 O 3 or Si 02 , many structural alloys and coating compositions are designed to form surface scales of these oxides at elevated temperatures. A second requirement that must be ful- \ filled for an oxide scale to be protective is that it adhere tightly to the alloy substrate both during ' isothermal exposure to a high-temperature environment and under conditions wherein the temoerature is cycled between low and high temperatures.

Although most studies of the oxidation behavior of alloys focus on the effects of the environment

and alloy composition on oxidation resistance, the structure of an alloy can also play a very important

role in the development of protective scales. As described in a recent paper [1], the combination of

a stable, fine alloy grain size, a fine, uniform dispersion of nonmetallic or intermetal 1 ic inclusions

and a uniform alloy composition, all of which may be achieved through rapid solidification orocessing, | should provide an excellent microstructure for oxidation resistance. The objective of this paper is to describe the oxidation resistance of a fine-grained solidified 18-8 stainless steel that contained i a dispersion of fine nonmetallic precipitates.

EXPERIMENTAL

A rapidly solidified 303 stainless steel powder, which is designated 303RS in this paper, was prepared by Pratt & Whitney Aircraft by forced convective cooling in He of centrifugal ly atomized droplets:

Particles that were -140 mesh were consolidated by hot extrusion at 900°C with a 10:1 reduction in area Samples of the consolidated bar were prepared for oxidation experiments by annealing for 6 h at 1000°C

to stabilize the microstructure. The same heat treatment was given to specimens of conventional,

wrought 304 and 310 stainless steels (designated herein as 304W and 310W). The composition of the

three alloys are given in Table 1.

424
Table 1. Compositions of Steels Employed in this Investigation

Composition, Wt Pet

Steel C Mn P S Si Ni Cr Mo Co Cu N

304W 0.04 1.62 0.033 0.016 0.13 8.78 18.28 0.32 0.19 0.26 0.045

303RS 0.059 1 .60 0.028 0.34 0.62 8.68 17.31 0.37 0.15 0.78 0.032

31 OW 0.06 1.64 0.028 0.006 0.52 19.90 25.0 0.18 0.16 0.16 0.074

After annealing, the specimens were polished through 6ym diamond paste. Cyclic oxidation experiments were conducted by inserting the specimens into a furnace at 900°C in pure oxygen for 20h_, quenching the specimens to room temperature in air, measuring the change in weight (Am) and reinserting

the specimens into the furnace. Isothermal oxidation tests at 900°C in pure oxygen were conducted

using a thermogravimetric apparatus that employed an Ainsworth semi -microbalance as the weighing unit.

ALLOY MICROSTRUCTURES

The 303RS after the 1000°C anneal had a fine grain size (5.6ym average grain size) and contained

a uniform dispersion of fine MnS precipitates (0.2 to 0.5ym diameter). The 304W and 310W alloys had

grain sizes of 170 and 19ym, respectively. After 300h at 900°C, the 303RS had an average grain size

of 8ym, which indicates considerable resistance to grain coarsening. The size and distribution of the

MnS precipitates were virtually unchanged.

OXIDATION RESULTS

Results of the cyclic oxidation experiments for the three alloys are shown in Figs. 1-3 as plots

of the weight change per unit area (Am/A) versus time. Three separate experiments were conducted for

each alloy.

The results in Fig. 1 for 304W appear quite variable; however, each specimen exhibited considerable

spallation during these tests. Hence, the specimens that showed no apparent weight change after about

40h of exposure actually lost weight by spallation and regained weight upon subsequent oxidation. All

specimens exhibited a decrease in cross-sectional thickness, which is perhaps a better indication of

the degradation of the alloy by oxidation. The variability of the results presented in Fig. 1 is probably due to inhomogeneities in the 304W.

The 303RS exhibited considerable resistance to cyclic oxidation. Very little spallation was observed during cycling, thereby demonstrating that the plots in Fig. 2 yield a true indication of the

oxidation resistance of this alloy; i.e., the oxidation rate decreased as time increased.

Data for the 31 OW alloy are shown in Fig. 3. Although the 31 OW contains more Cr and Ni than 303RS,

the latter alloy exhibited similar, if not better, resistance to cyclic oxidation.

Data for the isothermal oxidation tests, which were obtained only for 303RS and 304W, are shown in

Fig. 4. The test temperatures were 800, 900, 1000 and 1100°C. There was little difference in oxidation resistance at 800 and 900°C. Both alloys exhibited relatively good resistance to oxidation at these temperatures, although the 304W did exhibit some oxide breakdown at 900°C. At 1100°C, both alloys oxidized rapidly. Although the 304W also oxidized rapidly at 1000°C, the 303RS exhibited good

resistance to oxidation at this temperature (the latter result was reproducible; the curve shown in

Fig. 4 for the 303RS at 1000°C represents the fastest oxidation kinetics observed for this alloy under

these conditions).

The maximum recommended operating temperature for an 18-8 stainless steel under cyclic conditions

is usually specified to be 875°C[2]. The 303RS alloy used in this study exhibited excellent resistance

to cyclic oxidation at 900°C, and thus, it appears that the maximum operating temperature of an 18-8

stainless steel under cyclic conditions can be increased by at least 25°C through modification of the

alloy microstructure. The maximum recommended operating temperature for an 18-8 stainless steel under

isothermal conditions is usually specified to be 925°C. The 303RS steel exhibited very good resistance to oxidation under isothermal conditions up to 1000°C, which is an increase of 75°C in the maximum operating temperature.

425
(Arn/A),

mg

/cm

.0

Figure 1. Kinetics of the cyclic oxidation of a 304 stainless

steel. Data for three separate experiments are depicted.

Figure 2. Kinetics of the cyclic oxidation of a fine-grained,

rapidly solidified 303 stainless steel. Data for

three separate experiments are depicted.

42b
(Am

/A

),

rng/cm^

Figure 3. Kinetics of the cyclic oxidation of a 310 stainless steel.

Data for three separate experiments are shown.

Figure 4. Isothermal oxidation kinetics for 304W and

303RS stainless steels.

427
The mechanisms of the oxidation of the 303RS were discussed in detail in a recent paper [1]. In ' brief, the fine grain size of this alloy appears to promote the formation of protective, CraOa-rich i scales by preferential formation of CraOa over grain boundaries and by rapid grain boundary diffusion

of Cr from the interior to the surface of the alloy. Although the alloy contains only 0.6wt% Si, Si02

forms in a relatively short time along the alloy/scale interface, which acts as a secondary barrier

between the alloy and the gas phase. The SiOa also forms as intrusions from the external scale into

the alloy along the closely spaced alloy grain boundaries. These oxide intrusions may act to mechanically key the oxide scale to the alloy surface, thereby explaining the very good adherence of the

oxide scale to the alloy. The major role played by the MnS precipitates is probably pinning of the

alloy grain boundaries which stabilizes the fine grain size at high temperatures. The MnS may also

contribute to scale adhesion by helping to promote a convoluted alloy/scale interface over alloy grains.

Finally, the uniform composition of the alloy appears to promote formation of external scales of unicomposition, which would tend to decrease stresses and stress gradients in the scale, thereby enhancing

the ability of the scale to resist spallation.

ACKNOWLEDGEMENTS

This work was supported by a NASA grant (contract #NSG-7645) through the Materials Processing

Center at M.I.T.

REFERENCES

G.O.Yurek, D.Eisen and A.Garratt-Reed, Met Trans A , 1982, vol . 13A, p. 473-.

L. A. Morris, Handbook of Stainless Steels , D. Peckner and I.M. Bernstein, eds., McGraw-Hill, N.Y.,

1977, p. 17-1.

[

1 ]

[ 2 ]
CHARACTERIZATION OF METASTABLE PHASE POWDERS PREPARED BY ELECTROLESS DEPOSITION

L. V. Mieisel, P. J. Cote, and D. M. Quinn

U.S. Army Armament Research and Development Command

Large Caliber Weapon Systems Laboratory

Benet Weapons Laboratory

Watervliet, NY 12189

ABSTRACT

Methods are described for preparing fine powders of amorphous and other metastable

alloys by rapid deposition from electroless baths. Five different alloys have been prepared by

these methods and characterized by means of x-ray diffraction, SEM, and DTA techniques. The

properties and transformation kinetics of these alloys are described. In particular, some features of non-isothermal transformation kinetics and incubation effects are discussed.

Introduction

One of the principal areas of interest in the field of rapid solidification processing is the

preparation of metastable phase alloy powders and their consolidation into bulk components. A. Wurtz

[1] in 1844 was apparently the first to report on the preparation of metastable phase powders; his

method involved the chemical reduction of nickel ions in aqueous solutions of sodium hypophosphite.

Later work on sodium hypophosphite solutions by Brenner and Riddell [2] at NBS led to the development

of the electroless plating process; related work on electrolytic plating by Brenner et al [3] led to the preparation of the first metallic glass, NiP. Electroless deposition of nickel, cobalt, copper,

and palladium is now routine in the plating Industry. Commonly used reducing agents include

hypophosphite, formaldehyde, hydrazine, borohydride, and amine borane.

In this paper we describe some characteristics of powders which we prepared by inducing

decomposition of a variety of electroless baths. Five different metastable alloy powders were obtained in this manner: amorphous and crystalline phases of NIP, and crystalline phases of NiWP, CoP,

CoWP, and NiB.

Our most extensive results concern the transformation kinetics of NiP alloy powders using

differential scanning calorimetry. Isothermal transformation data are in excellent agreement with

Johnson-Mehl-Avrami (JMA) theory [4]; i.e., the degree of reaction x(t) at time t is given by

x(t) = 1 - exp[-(k(T)t)’^] (1)

where the rate constant k(T) at temperature T is given by

k(T) = A exp(-0/T), (2)

with A constant, E = RT is an activation energy for the transformation process, R is the gas constant,

and the JMA exponent n is determined by the details of the transformation process.

Constant heating rate results are well described by

x(T) = 1 - exp[-W'^] (3)

where

k(T)T2 „ „ 2 W(T) = Wo + -r [1 - 2 ^ + 6(2) ] (4)

T0 0 0

with T = Tq + Tt, T is the (constant) heating rate, Wq = [-£n( 1 - x(Tq))]^/’^ which is zero if no reaction has occurred at t = 0. Equations (3) and (4) are deduced in Reference [5] for systems

governed by JMA theory in the isothermal case and where the degree of reaction is determined by the

number of atom movements; equivalent results were obtained by Henderson [6] on the assumption

that the reaction rate in the general case is given by the JMA isothermal theory. The Kissinger

equations, which form the basis of the Kissinger method [7], can also be deduced from Eqs. (3) and

429
(4). Thus, References [5] and [6] provide a theoretical basis for the analysis of constant heating

rate transformation kinetics in many systems and, in particular, for the powders analyzed here.

Interesting incubation effects, apparently unrelated to the formation of critical nuclei, are observed in the isothermal and constant heating rate transformation of metastable crystalline NiP

powders, while no such effects are apparent in the transformation of the a-NiP powders.

Experimental

Sample Preparation

Powder samples of NiP [2], NiB [8], CoP [9], NiWP [10], and CoWP [11] were prepared from

solutions based on standard electroless baths operated under conditions which produce spontaneous

decomposition (i.e., deposition throughout the entire volume of the bath). The concentration of the

reducing agent was doubled relative to standard baths, high bath temperatures were maintained (T > 96° C) and additions of bath stabilizers were omitted in the electroless baths.

Powder samples of the same alloys were also prepared by seeding the baths with low concentrations

(~ .001 M) of PdCl 2 following the method of Randin et al [12]. The seeding process apparently results

in a colloidal suspension of palladium metal from the reduction of palladium ions by the

hypo phosphite, which in turn leads to rapid deposition throughout the entire bath, as with spontaneous

decomposition. We find that this deposition can be interrupted by filtering out the palladium

suspension. No significant differences were found between samples prepared by either method.

Structural Characterization

Sample compositions were determined by standard wet chemistry methods. Debye-Scherrer x-ray

analysis was used to determine alloy phases. Powder morphology was examined by scanning electron

microscopy.

Transformation Kinetics

The kinetics of the phase transformations were studied using isothermal and constant heating rate differential scanning calorimetry (DSC) employing a Dupont 990 Thermal Analysis System. Aluminum

sample pans and an argon atmosphere were used in all runs. An empty sample pan served as a reference.

The isothermal transformation kinetics data were obtained from the time dependence of the DSC

output for a series of temperatures. The DSC output is assumed to be proportional to the reaction

rate dx/dt (IngC “ cdx/dt, c = constant) so the degree of reaction x(t) is given by x(t) = s(t)/S

where s(t) = J IqsC “ c[x(t) - x(0)] and S = s(“). The technique for extracting the time exponent

n and k(T) in an isothermal experiment described by Eq. (1) is to plot £n[-Jln(l - s(t)/S)] vs. Ant so that n is given by the slope and k(T) = tg"^ where Jlntg is the int intercept.

The constant heating rate data were analyzed by the Kissinger method based on Eqs. (3) and (4) as described in References [5] and [6]. Thus, the activation energy for the transformation process is

given by the slope of the plot of £n(T/T]^) vs. T^"^ where T is the heating rate and Tf^ is the

temperature at the maximum reaction rate. Corrections are required to account for incubation effects

and to account for thermal lag effects which are inherent with this method. The thermal lags were established empirically using the Curie temperature of Ni and the melting point of Sn.

Results and Discussion

Structural Characterization

Sample compositions for the powders prepared from electroless baths are given in Table 1.

Table 1. Compositions in at% for powders prepared from electroless baths. a-NiP and x-NiP refer to amorphous and metastable crystalline phases, respectively. The other as-prepared powders are in metastable crystalline phases.

a-NiP x-NlP NiWP CoWP NiB CoP

Ni:80 Ni:85 Ni:84 Co; 92 Ni:80 Co: 92

P:20 P:15 W:5 W:3 B:20 P:8

P:ll P:5

430
We find that the compositions of the NiP powder samples are extremely uniform and can be

controlled by maintaining constant bath pH. A measure of the uniformity in composition is given by

the positions of the DSC transformation peaks which are the same for all NiP samples from a given

bath. In addition, the width of the DSC peaks are close to the theoretical widths, which further

reflects the homogeneity of the powder compositions. (Examples of DSC peaks are shown in Figure 5.)

On the other hand compositions of the other powders were not precisely controlled by bath pH so that

inhomogeneous powders, which exhibit complex DSC traces, were obtained frequently.

X-ray diffraction data indicate that, except for a-NiP, the as-deposited powders are metastable

solid solutions of the various constituents. The as-deposited nickel and cobalt alloys are found to have fee and hep crystal structures, respectively.

Scanning electron microscopy (SEM) reveals that the powders have similar morphology; being

comprised of clusters of attached spheres with diameters ranging from ~ 1 p to 10 p, and having a wide

distribution of cluster sizes, ranging from individual spheres to clusters with overall dimensions up

to ~ 100 p. Figures 1 and 2 show samples of amorphous and metastable crystalline NiP powders which

are typical of all the powders and were taken from the samples that were used for the detailed studies

of transformation kinetics.

Figure 1. SEM micrograph of amorphous NiP powder Figure 2. SEM micrograph of metastable

sample prepared from an electroless crystalline NiP powder sample

bath. prepared from an electroless bath.

Figure 3 is an optical micrograph showing sections of typical x-NiP spheres. There are laminations evident within the spheres. Such features are characteristic of alloy deposition from

aqueous baths although no satisfactory explanation yet exists to their origin [13]. This may be due

in part to the fact that the detailed structure of the laminations varies widely from one alloy to another. As noted earlier, the compositions of the NiP powders appear remarkably uniform which seems

to rule out composition fluctuations as the origin of the laminations. Changes in crystallographic

orientation, on the other hand, offer a consistent explanation for the observed features in these

powders.

Figure 4 is a SEM micrograph of a section x-NiP spheres which was etched with modified Fry's

etchant (HNO 3 , HCl, and CUCI 2 ). The boundaries between the laminations shown in Figure 3 appear to have been completely dissolved. This suggests that a thin impurity layer, possibly containing

hydroxides or hydrated oxides, are present at the interface of the laminations, since such impurities

are commonly found in electrodeposits and are soluble in acids.

431
Isothermal and constant heating rate studies of transformations in amorphous and metastable

crystalline NiP powders were performed using DSC methods. The isothermal data are well described by

the JMA equation, deduced from Eqs. (1) and (2). The constant heating rate data are well described by

Eqs, (3) and (4) and were analyzed by the Kissinger method. Activation energies, preexponential

factors, and JMA exponents n describing the transformation of a-NiP, as-prepared x-NiP, and

preannealed (as described later) x-NiP powders, deduced from isothermal and constant heating rate DSC

studies, are summarized in Table 2.

Table 2. Kinetic parameters for amorphous (a-NiP) , as-prepared metastable crystalline NiP (x-NiP)

,

and preannealed powders, obtained from isothermal and constant heating rate DSC measurements. h represents the mean value of the JMA exponent.

Isothermal Constant T

Sample

E (kcal/mole)

An A

(A in sec“^) n

E (kcal/mole)

An A

(A in sec“^

)

a-NiP 38.6 30.4 2.3 40.7 32.5

x-NiP 52.7 39.2 3.09 55.5 41.0

Preannealed 54.2 40.4

x-NiP

A feature of transformations in the as-prepared x-NlP powders which is not observed in the a-NiP

powders is an Incubation time, tj^nc» representing a time delay before any indication of a transformation is observed in an isothermal anneal. We find that is apparently controlled by an activation process with activation energy E = 66.6 kcal/mole for x-NiP.

In the isothermal experiments, the DSC output (dx/dt) is zero for t < t^^c increases from

zero for t > 0, proceeds through a maximum, and returns to zero. Following standard procedures, the

DSC traces for a series of temperatures were integrated and yield straight lines over most of the

range in plots of £n [-Zn( 1-s/S) ] vs. £n(t-tinc)* values range from 2.8 to 3.2 with a mean n of

3.09 in the x-NiP powders and range from 2.1 to 2.5 with a mean h of 2.3 in the a-NiP powders.

In the constant heating rate experiments, we found that a modification of the standard Kissinger

method is required to account for the effects of incubation times. A demonstration of these effects

in the crystalline NiP powders is given in Figure 5 which is a plot of the DSC output vs. temperature

for a heating rate of 50K/min. The theoretical curve was obtained from Eqs. (3) and (4) with values

for the parameters E, JlnA, and n obtained from the isothermal measurements. The difference between

the theoretical curve and the data on the as-prepared powders is substantial. The data show a maximum

transformation rate at T^ = 636K and a full width at half maximum (FWHM) of approximately 8K while the jV' corresponding theoretical values are 631K and 11. 7K.

432
Figure 5. Experimental and theoretical DSC trace at 50K/min for metastable crystalline NiP powders.

Parameters used in the theoretical curve were obtained from isothermal data. Area of

theoretical curve is normalized to unity. Peak heights of experimental curves were scaled

to match the theoretical curve.

The small observed value for FWHM reflects a high degree of uniformity in sample composition as we suggested earlier. The fact that the FWHM is actually less than predicted and Tjvj higher than

predicted suggests that incubation effects are responsible since they act to delay the onset of the

transformation. The isothermal data are routinely corrected for the incubation process while no such

corrections are made in the standard Kissinger method. We found that if we preannelaed x-NiP samples

at T = 568K for 15 minutes (i.e., the Incubation time at 568K) before a constant heating rate study,

incubation effects were eliminated, i.e., the data on the preannealed x-NiP sample are seen from Table

2 and Figure 5 to be in excellent agreement with the theoretical predictions. The elimination of

incubation effects in the constant heating rate runs by means of isothermal preanneals suggests that

the standard model for the formation of critical nuclei [14] cannot explain incubation in the x-NlP

powders

Another illustration of incubation effects is given in Figure 6 which shows Kissinger plots of

data from as-prepared and preannealed samples. The parameters E and £nA were obtained from least

squares fits to the data. It can be seen from Table 2 that preannealing through t^ji^ leads to improved agreement with the isothermal parameters.

The kinetic parameters of the NIP powders reveal considerable differences in the nature of the

transformations of the amorphous metastable crystalline phases. The JMA exponent of the crystalline

powders is near 3 which is associated with activated rate spherical growth on pre-existing nuclei

[15]. The JMA exponent is closer to 2.5 in the amorphous powders, which is associated with diffusion

controlled spherical growth with activated non-zero nucleation rate [15].

Besides yielding consistent values of the kinetic parameters describing the transformation of

amorphous and metastable crystalline NiP, the excellent agreement between the results of the analyses

by the Isothermal method and the Kissinger method provides evidence of the appropriateness of the

generalized form of the JMA theory as given in References [5] or [6] to non-isothermal heterogeneous

reactions. Greer [16] obtained similar results in a detailed study of transformation kinetics in

a-FeB ; the basic assumptions in his numerical analysis are equivalent to those used in the generalization of JMA theory given in Reference [5]. A more complete presentation of the experimental results

in x-NiP as well as an alternative derivation of Henderson's equations describing non-isothermal solid

state reactions with more general starting assumptions are given in Reference [5].

433
10® K / T„

Figure 6. Kissinger plots for as-prepared and preannealed x-NiP powders.

T unit is Ksec“^ and is in K.

Acknowledgements

We wish to thank many of our colleagues for their generous assistance: Mr. G. Capsimalis

obtained the x-ray data, Mrs. T. V. Brassard, Mr. L. McNamara and Mr. C. Rickard obtained the

micrographs; Drs. E. Chen and F. K. Sautter provided many helpful discussions on electrodeposition

and Ms. E. Fogarty was instrumental in the timely completion of the manuscript.

References

[1] A. Wurtz , Compt. Rend. J^, 702 (1844); 149 (1845).

[2] A. Brenner and G. E. Riddell, J. Res., Natl. Bur. Std. 1 (1946).

[3] A. Brenner, D. F. Couch, and E. K. Williams, J. Res., Natl.. Bur. Std. 44, 109 (1950).

[4] W. A. Johnson and R. F. Mehl , Trans. AIME, 135 , 416 (1939) and M. Avrami, J. Chem. Phys. 1_^ 1103 (1939).

[5] L. V. Melsel and P. J. Cote, Acta. Met., to be published.

[6] D. W. Henderson, J. Non-Crystalline Solids 3^, 310 (1979).

[7] H. F. Kissinger, J. Res. Natl. Bur. Std. 57_, 217 (1957); Anal. Chem. 2^, 1702 (1957).

[8] T. Watanabe and Y. Tanabe, Mat. Sci. , and Eng., _2^ 97 (1976).

[9] A. Brenner and G. E. Riddell, J. Res., Natl. Bur. Std. 39_, 385 ( 1947).

[10] F. Pearlstein and R. Weightman, J. Electrochem. Soc. 427 (1968).

[11] F. Pealstein and R. Weightman, Plating, 1 (1967).

[12] J. P. Randin and H. E. Hintermann, Plating, 54, 523 (1967).

[13] A. Brenner, Electrodeposition of Alloys , Academic Press, New York, (1963), Chapter 7. [14] J. H. Holloman and D. Turnbull, Prog. Metal. Phys. 333 (1953).

[15] See for example, V. Raghavan and M. Cohen in Treatise on Solid State Chemistry , Ed. N. B. Hannay, Plenum Press, New York (1975).

[16] A. L. Greer, Acta Met., 30., 171 (1982).

434
PREPARATION OF GLASSY METALS AS POWDERS VIA HYDROGEN EMBRITTLEMENT*

B.C. Giessen, S.H. Whang and F. Dabkowski

Materials Science Division, Institute of Chemical Analysis

and Department of Chemistry

Northeastern University, Boston, MA 02115

ABSTRACT

Preparation of metallic glasses by melt spinning readily yields metastable materials

in the form of a homogeneous product; however, the alloy ribbons are frequently ductile

and are thus not directly amenable to the pulverization process required to produce

powders suitable for further P/M processing. Extending the well-known metals processing

technique of pulverization by reversible hydrogen embrittlement to metallic glasses

(and probably other RSR produced metastable alloys), such alloys can be prepared as powders by loading them with hydrogen at pressures of 5-20 MPa in an autoclave or high

pressure ball milling chamber. Hydrogen charging is generally done at elevated temperature to provide activation while comminution is carried out at room temperature. The

sample is then unloaded by evacuation of hydrogen before exposure to air. Particulates

are screened and studied by XRD, DSC and microscopy. This approach has also been used

and described in a recent paper by Maeland and Libowitz. The paper describes preliminary

results on the pulverization of glassy Fe 78 Bi 2 Siio.

Introduction

Glassy alloys can be prepared in ribbon form as well as in the form of atomized powders using

rapid liquid quenching techniques. RSR processing by solid-substrate quenching methods such as melt

spinning is generally preferable to quenching methods which utilize heat conduction by gases, such

as quench-atomization, due to the higher cooling rates and greater product homogeneity which distinguish the former techniques over the latter; ribbon material tends to have a more uniform microstructure than powders which have microstructures with wide qualitative and quantitative differences,

depending on the droplet size and other aspects of their quenching history [l].

However, in order to reap the benefits of melt spinning while at the same time being able to apply well-established P/M techniques to the consolidation of melt spun ribbon alloys, the quenched

material must first be converted to powder without destroying its desirable properties such as ductility and toughness. To accomplish this and to pulverize ductile metallic glasses or other

ductile, RSR processed alloys into powders with unchanged mechanical properties appears to be a problem that, at present, is still solved only incompletely.

The high bend ductility of many RSR alloys makes it very difficult to pulverize them into fine

particulates via a mechanical treatment unaided by embrittlement. The following methods are therefore candidates for a solution to this problem:

a) Melt spinning into small particles, e.g., by quenching on a profiled substrate configurated to produce flake shaped particles. It is not clear whether a large scale preparation facility based on this

principle is viable on a long-term operation basis.

b) Cutting or other size reduction by shearing action. An approach to prepare alloy flakes by this

method is under study.

c) Pre-comminution embrittlement by physical means. Low temperature embrittlement is an obvious

candidate method; however, it is not known which ductile glasses, if any, can be embrittled enough

to permit sufficient size reduction.

d) Reversible chemical embrittlement by hydrogenation. This method, patterned after the well known

process for producing powders of ductile metals such as tantalum by reversible hydrogenation, is a major candidate for powder production and its application is discussed here.

435
Ill

Attempts at comminution in hydrogen atmosphere are based on the prior observation that some Fe '

' based glassy alloys become brittle when charged cathodically with hydrogen [2, 3] and on many studies i of hydrogen absorption in glassy metals [4, 5] where reversible embrittlement was reported. However,

j| •in these studies [5] it was also found that even in hydride forming glassy alloys such as Zr-Ni there

is no pressure plateau in the pressure-concentration isotherms similar to that in hydride-forming

crystalline matal -hydrogen systems. In addition, Fe based alloys do not form equilibrium hydrides i at low pressures. As a result, the concentrations of hydrogen in Fe base glassy metals that are ! needed for embrittlement require application of high hydrogen equilibrium pressures. Recently, em- brittlement of Fe base glassy metals by low hydrogen concentrations added before melt spinning has

been reported [6]; also, Maeland and Libowitz [7] have demonstrated that several Fe based glassy

alloys could be pulverized under high hydrogen pressures (10-15 MPa). In the following, we will des^il

cribe a similar experiment conducted in this laboratory in an effort to produce powders from glassy /!

ribbon material. K|

Experiments

Figure 1 shows the laboratory scale hydriding system used for comminution experiments. It con-B|

sists of a commercial autoclave with a magnetically activated hammer (originally intended in this

system for use as a stirrer), a vacuum system and a hydrogen pressurizing system. The autoclave can

be operated at a maximum pressure level of 5000 psi (34.5 MPa) at 343°C and at correspondingly lower

pressures at higher temperatures. Glassy metal samples are charged with hydrogen in the autoclave by 1

a series of cyclic heatings under high pressure. Samples activated and hydrogenated at high tempera- | ture are subsequently comminuted at room temperature. After comminution, hydrogen is removed from

the samples by evacuating the autoclave. Following removal from the autoclave the powders are screened using graduated sieves. Parameters and typical results are presented in the next section.

Results and Discussions

A small amount (about 0.5g) of glassy Fe 7 eB 22Siio alloy ribbons was cut into pieces 0.5 to 1.5cm

in length and placed in the autoclave. (Only small quantities were used as the hammer system presently

used becomes clogged by larger sample quantities.) At the start of each thermal cycle hydrogen with j pressures of about 1000 psi (6.9 MPa, as measured at room temperature) was admitted and the temperature was then raised to 400°C (i.e., to 127°C below the glass transition temperature which is 527°C

for this alloy [8]). Two or three thermal cycles, each lasting about 30 minutes, were then made. j| Starting during the last cycle and continuing after cooling, the sample was comminuted for a specified

time, typically of the order of 20 hrs.

To test whether the resulting embrittlement was indeed due to hydrogenation, similar cyclic ?

|

heatings were carried out under vacuum or, alternatively, under argon gas, using the same loading

J I,] conditions; no alloy embrittlement was found in the ribbons after these treatments. This rules out f

|!i)

[ the possibility of embrittlement being due to partial crystallization or structural relaxation during lli-i the cyclic heating and demonstrates the need for the presence of hydrogen during the embrittlement i heating cycles. It can be assumed that cyclic heating in hydrogen generates microcracks and hence

produces fresh surfaces and facilitates hydrogen absorption. | The powders produced were checked by X-ray diffraction and found to be amorphous. After screen- 1 ing with standard sieves the powders exhibited the size distributions shown in Table 1. It was noted

that this size distribution is not Gaussian but that alloy fraction quantity is related to sieve size

by a simple linear relationship. This unusual size distribution was also found by Maeland and Libowitz

[7], and appears to be independent of the hydrogen pressure which was 7-10 MPa in the present work, i i.e., lower than that used by these workers (12-15 MPa) [7]. This unfavorable size distribution is assumed to be due to the relatively soft hammering mechanism used which is clearly insufficient; a ball milling system operating under high hydrogen pressure can be expected to permit the production I of RSR materials in larger quantities and hopefully, with better size distributions of the RSR

powders (such a system is now under construction).

|

It appears that Fe alloy glass powders of small particle sizes adequate for P/M work can be

prepared by this method; the major question remaining to be answered is how the toughness and ductility^

of metallic glasses thus pulverized and subsequently consolidated compares with the corresponding

properties of the original, as-spun materials. This study must await the availability of adequate

amounts of powdered material.

Acknowledgements \

j We are grateful for support of this work by the Army Research Office, Durham, Contract DAAG-29- T 80-0088 under subcontract with M.I.T. '

436
Fig. 1 Schematic diagram of system used for hydrogenation and comminution of Fe-based amorphous

ribbons.

Run +250(um) -250 -125 -63 -45 -38

No. +125 +63 +45 +38

1 54.7 30.7 9.9 2.8 0.7 1.2

2 56.9 27.7 9.6 2.9 2.9 ---

Table 1

Size distribution (in wt%) of Fe 7 gBi 2 Siio glassy metal powders after hydrogen charging at 1000 psi at 400°C and comminution for 20 hrs.

437
References

[1] R.H. Patterson, II, in Rapidly Solidified Amorphous and Crystalline Alloys, MRS Symposia Proceed.

Vol . 8, B.H. Kear, B.C. Giessen and M. Cohen, Eds., North-Hol land N.Y. (1982) p. 123.

[2] M. Nagumo and T. Takahashi, Mat. Sci. Eng., 23 (1976) 257.

[3] S. Ashok, N.S. Stoloff, M. Glicksman and T. Slavin, Scripta Met., 15 (1981) 331.

[4] F. Spit, K. Blok, E. Hendriks, G. Winkels, W. Turkenburg, J.W. Drijver and S. Radelaar, Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai, Japan, 2 (1981) 1635.

[5] K. Aoki, A. Horata and T. Masumoto, Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai,

Japan, 2 (1981) 1649.

[6] H. Shahani, H. Soderhjelm and M. Nygren, Rapidly Solidified Amorphous and Crystalline Alloys,

MRS Symposia Proceed., Vol. 8, B.H. Kear, B.C. Giessen and M. Cohen, Eds., North-Hol land, N.Y. (1982) p.289.

[7] A.J. Maeland and G.G. Libowitz, Mater. Lett., 1 (1982) p.3. fl

[8] I.W. Donald and H.A. Davies, Phil. Mag. A42 (1980) 423.

*Communication No. 150 from the Institute of Chemical Analysis, Northeastern University, Boston, MA 02115

i
AN ARC FURNACE MELT SPINNER FOR THE RSR PROCESSING

OF REFRACTORY AND REACTIVE ALLOYS*

S. H. Whang and B. C. Giessen

Materials Science Division, Institute of Chemical Analysis

and Department of Chemistry

Northeastern University, Boston MA 02115

ABSTRACT

Melt spinning is a convenient method for the preparation of RSR processed alloys;

however, its most common forms use oxide crucibles such as quartz which introduces

melting temperature limitations due to the softening temperature of the crucible material

and crucible-melt reactions.

Cold hearth techniques are very desirable for the RSR processing of molten refractory

or highly reactive elements. Following a design described in the literature, we have

found arc melting to be convenient for small sample quantities. A two-chamber system

has been built, with a small, non-consumable electrode arc melter in the upper chamber and

a rotating copper quenching wheel located in the lower chamber directly underneath the

cold hearth which is provided with an appropriate hole of small bore. After melting the

alloy, a pressure differential is produced between the top and bottom chambers by a ballast tank and a thin stream of molten metal is quenched into a ribbon upon impingement

on the wheel. Among alloys prepared as ribbons by this technique have been alloy glasses

with high titanium content and Nb 3 (Nb,Si) alloys prepared as the metastable A15 type phase.

Possible modifications of the basic design are discussed, including the use of alternative

heat sources such as energy beam or induction melting.

Introduction

Melt spinning has become the technique of choice for the continuous production of rapidly

solidified metals and alloys in ribbon form, and in the last decade melt spinning technology has

substantially advanced in accordance with increasing requirements concerning the quality of ribbon

or sheet materials. The basic melt spinning technique goes back to the method of Strange and Pirn [1]

in which a molten jet impinges onto a cold, rotating disk; most alloys of interest in RSR technology

can be spun by this technique. An important component of the melt spinning system is the crucible;

various crucible materials are in use depending on alloy properties. The preferred crucible material

is silica; graphite, alumina, zirconia etc. have also been used. However, reactive alloys such as alloys rich in Ti and Zr or alloys with very high melting temperatures such as refractory alloys

containing Nb, Ta, Mo, W, Re, Ru, Os, Rh, Ir, etc. cannot be readily melted and superheated using

these crucible materials. Problems associated with these crucible limitations can be solved using

a water cooled copper hearth as a crucible. Such a cold copper hearth and a rotating quenching

wheel were first combined into an operational arc furnace melt spinner by Tutzauer et al. [2]. Use

of cold-hearth arc melting of small alloy pellets combined with quenching by the hammer-and-anvil

method as a way of avoiding crucible contamination has been described by Wang [3] and was subsequently

used and discussed [4]; such a unit, provided with a pressure driven piston for alloy quenching, has

been in continuous operation in our laboratory for a number of years. The following is a description

of an arc furnace melt spinner constructed and in experimental use in this laboratory and a review

of some experiments carried out with this unit.

Design Features and Operation

The Arc Melt Spinner

Figure 1 shows a schematic diagram of the laboratory scale arc melt spinner constructed at Northeastern University. The spinner system is made up of a main chamber, a vacuum system, inert gas

supply system, a D.C. power system and an A.C. motor. The main chamber consists of two compartments:

the top compartment, i.e., the melting section, contains an inert gas plasma arc electrode, a removable

copper crucible, a gas inlet and a vacuum outlet. The bottom compartment includes a spinning disk made

of copper of molybdenum, a ribbon guide pipe, a gas inlet and a vacuum outlet.

439
Fig. 1 Schematic diagram of the arc furnace melt spinner constructed at Northeastern University.

A motor which drives the spinning wheel is located behind the main chamber.

440
Melting

A premelted alloy button ('^^0.5g) is placed at the bottom of the copper crucible such that it covers the orifice centered in the crucible as shown in Fig. 1 and is melted using the arc electrode

while both compartments are kept under argon atmosphere at equal pressures. The molten alloy is superheated to facilitate formation of a fine molten jet. This is particularly important for refractory

alloys where the surface tension at the melting temperature is still too high to allow formation of a fine melt stream.

Spinning

The superheated melt is formed into a fine molten jet flowing through the orifice of the copper

crucible by a pressure differential applied between the two compartments, either by rapidly pressurizing

the melting compartment or depressurizing the spinning compartment. A stationary pressure gradient can

be maintained during spinning as the molten alloy "plugs up" the orifice which is the only opening

between the two compartments. In the design shown in Fig. 1 the spinning compartment is rapidly depressurized by opening a gate valve connecting an evacuated ballast tank to the spinning compartment.

Spinning proceeds at that point and the spun ribbon is caught in a retrieving pipe dimensioned such

that impingement of the ribbon on the back wall does not affect the quenching process.

Crucial Components and Dimensions

The operation of this system is determined by the same parameters critical in a melt spinning

system: flow rate (determined for a given material by the orifice size and the pressure differential

between the two chambers), wheel speed, crucible-to-wheel distance, angle of impingement and degree

of superheating. Additional parameters specific to an arc melting system are the shape of the

copper hearth and the ratio of the length to width of the orifice; this aspect ratio must be kept below

a critical value to prevent freezing of the melt to the cooled orifice wall.

Experiments

To date many Ti rich alloys have been spun into ribbon material using this unit. The quenched

compositions include alloys from the binary systems Ti-Fe, Ti-Cu and Ti-Ni and the ternary systems

Ti-Ni-Si [5], Ti-Zr-B(-Si ) , Ti-6A1-4V, etc. For these alloys the following spinning parameters were

used: orifice diameter, 0.4mm; pressure differential, about 10 psi; disk surface speed, '^^20 m/sec,

angle of impingement, 90°. The resulting ribbons had a thickness of 20um. The ribbons spun from

Ti-Cu, Ti-Ni and Ti-Ni-Si were amorphous as checked by X-ray diffractometry. In another experiment

NbszSiis having a high melting temperature ('^-2000°C) was successfully spun by this technique using

an increased melting current of '^>75A and a higher pressure differential of 18 psi. According to X-ray analysis, the spun ribbon consisted predominantly of a textured A-15 structure (as found earlier

[6] for arc furnace quenched [4] foils of Nb-Si alloys). However, the ribbon was narrow and thick,

probably as a result of the high surface tension of the melt.

Discussion

Current Status

The reported results show that this method is promising for its intended purpose; however, in

its present experimental form, the unit has some operational drawbacks. Thus, the reproducibility of

the spun ribbons is not yet satisfactory since melting, superheating and timing of pressurization

for spinning rely on the operator's skill and reactions. Also, prolonged heating of a small alloy

button in the arc may result in vaporization of a significant amount of volatile component from the

alloy. It should be noted that the arc spinning technique is not favorable for alloys having a very

high thermal and electrical conductivity such as alloys rich in A1 or Cu; the fact that their melt - crucible interface resistance to heat loss as well as the thermal gradient in the alloy are low

makes it difficult to superheat them sufficiently for ready flow.

Further Developments

Looking beyond the present, we believe that this technique can be upscaled and modified to produce RSP refractory materials in the amounts needed for alloy development and processing and,

ultimately, production. This will require, in the first place, a means of supplying adequate amounts

of alloy materials in a continuous fashion. This can be accomplished by continuously providing feed

material to the melt pool from a reservoir in particulate form, e.g., as small pellets or a coarse

powder. If discontinuous operation is acceptable, use of a consumable electrode prepared from the

alloy under study may be considered; however, this approach is less economical considering the

higher manufacturing cost of such electrodes. In view of the control problems inherent in cold-hearth

arc furnace operation (which are aggravated by the need of keeping the orifice size constant, i.e..

441
of avoiding any freezing) one may also consider (or re-consider) alternative forms of heating, e.g.

by induction or laser as discussed in the following.

An induction heated cold hearth system making use of an rf field concentrator coil and a slitted

silver hearth with a central orifice has been described by Willens and Buehler [7] (reviewed in

Ref. [8]), Although this unit was intended for use with a "gun" type splat cooling unit (i.e. using

atomization and particle acceleration by shock wave) combined with a stationary copper substrate, it

could be adapted to melt spinning by use with a gas pressure system and a copper wheel to provide a thin stream of metal flowing onto a rapidly rotating substrate.

Turning to laser heating, use of this energy source to melt pellets of materials prior to quench

ing by the hammer-and-anvil method has been described [9]; with a sufficiently powerful laser, it

would be possible to maintain a melt pool fed by a continuous supply of pellets (as described above)

and drained by a continuous stream through an orifice onto a quenching wheel.

Acknowledgement

Support of this work by the Army Research Office, Durham, under Contract DAAG29-80-C-0147 is

gratefully acknowledged. The authors also thank Y.Z. Lu and Y.Q. Gao for their assistance.

References

[1] Strange, E.H., and Pirn, C.A., 1908, U.S. Patent No. 905758.

[2] Tutzauer, H. , Esquinazzi, P., De La Cruz, N.E., and De La Cruz, F., Rev. Sci. Instrum. 51 (1980)

546.

[3] Wang, R. , Rev. Sci. Instrum., 41 (1970) 1233.

[4] Fischer, M. , Polk, D.E., and Giessen, B.C., Rapid Solidification Processing, Principles and

Technologies, Mehrabian, R. , Kear, B.H., and Cohen, M. , eds., Claitor's Publishing Division,

Baton Rouge (1978) 140.

[5] Polk, D.E., Calka, A., and Giessen, B.C., Acta Met,, 26 (1978) 1097.

[6] Waterstrat, R.M., Haenssler, F., and Muller, J., J. Appl . Phys,, 50 (1979) 4763.

[7] Willens, R.H., and Buehler, E., Met. Trans. TMS-AIME 236 (1966) 171.

[8] Giessen, B.C. and Willens, R.H., in Phase Diagrams, Alper, Allen M. , ed., Vol . Ill (1970) 103.

[9] Krepski, R. , Swyler, K. , Carleton, H.R. and Flerman, H. , J. Mat. Sci. 10 (1975) 1452.

*Communication No. 149 from the Institute of Chemical Analysis, Northeastern University, Boston, MA 02115

442
A NEW HIGH STRENGTH NICKEL-BASE MICROCRYSTALLINE

ALLOY VIA RAPID SOLIDIFICATION TECHNOLOGY

V. PANCHANATHAN

Marko Materials, Inc.

No. Billerica, MA 01862

S. SASTRI and S. HAHN

The Gillette Company

Boston, MA 02106

ABSTRACT

A new high strength nickel -base alloy with microcrystalline microstructure has been

developed by the rapid solidification technique. The composition of the alloy designated

as Markomet 8 is as follows (weight percent):

Bal Ni, 21.8 Cr, 2.8 Fe, 8 Mo, 9.3 W, 1.5 Cb, 0.7 C, and 0.7 B

The method of melt-spinning-pulverization was used to prepare powders which were con- solidated into bar stock by hot extrusion. The alloy exhibits excellent oxidation resistance at elevated temperatures upto 1800°F. Typically, at 1800°F the oxidation rate is

1.37 mg/cm2. The alloy also has excellent corrosion resistance in nitric acid, formic

acid, and phosphoric acid. The mechanical properties (tensile and stress rupture strength

values) of the consolidated microcrystalline alloy at different temperatures are reported.

Typical ultimate tensile strength values are:

175 KSI (1205 MPa) at room temperature and 97 KSI (667 MPa) at 1400°F.

The microstructural features of the alloy investigated by TEM and SEM are reported.

INTRODUCTION

Rapid solidification processing (RSP) has been recognized in recent years as having outstanding

prospects for creation of new engineering alloys with improved service performance characteristics.

The major R & D efforts on RSP alloys in the past several years have been geared towards enhancement

of mechanical properties and high temperature capabilities suitable for many aerospace/defense oriented hardware application [1,2,3]. One field where RSP can make significant contribution is in the development of new alloys having high corrosion resistance combined with superior mechanical strength.

These improvements in corrosion resistance and mechanical properties are anticipated to result from

one or more of the following characteristics of these rapidly quenched materials:

(1) unique chemical compositions and novel microstructures, and

(2) excellent chemical homogeneity and uniformity of structures.

Recent research efforts at Marko Materials led to the development of a new family of rapidly solidified nickel base alloys possessing high strength and excellent corrosion resistance suitable for potential technological applications. The present paper describes the structure and properties of one alloy

belonging to the aforesaid category.

ALLOY-COMPOSITION, PREPARATION TECHNIQUES & STRUCTURES

The new alloy was designed based on nickel for general corrosion resistance. It contains as primary

constituents, chromium for oxidation resistance and molybdenum for pitting corrosion. The alloy is

designated as Markomet 8 and has the following composition in wt.%.

BalNi, 21.8 Cr, 2.8 Fe, 8 Mo, 9.3 W, 1.5 Cb, 0.7 C, and 0.7 B

Markomet is the registered trade mark of Marko Materials Inc.

443
The alloy was produced by the method of melt spinning as rapidly solidified ribbons, typically

.002 inch thick and .06 inch wide. Incorporation of certain critical amounts of metalloid elements

[e.g. B and C] as alloying elements enables the alloy to be rapidly quenched into single crystalline

solid solution phase with fCC structure as revealed by x-ray diffraction analysis. The as quenched

structure consists of an extremely fine grain structure with grains, subgrains or crystallites typically ranging 100-200 °A in size as evidenced by TEM examination (Fig. 1).

The as quenched ribbons were found to be hard ['^100 DPN] and brittle. The brittle ribbons were readily pulverized into powders under 100 mesh by a standard hammer mill. The powders were subsequently

consolidated into bars by hot extrusion at 2100°F using an extrusion ratio of 12:1. A hot-extruded bar specimen (3/8" dia.) was cut, polished, and electrolytically etched in 10%

oxalic acid solution for SEM study of the microstructure. A Hitachi S-700 field emission scanning

electron microscope was used with operating conditions as follows: 20 KeV acceleration voltage, 20°

specimen tilt, 5mm W.D. SEM micrographs of the specimen (Fig. 2) show a fine dispersion of second

phase precipitates on the general matrix. The size of these multi-faceted precipitates vary from 0.5fjm to 3.0pm. Larger precipitates, however, consist of two or more precipitates, i.e. precipitates of

lighter and darker contrasts. An optical micrograph (Fig. 3)taken at 2,000X does not show this differ-] ance in contrast, which arises from the atomic number difference.

EDX (Energy Dispersive X-ray) analysis of these precipitates show that the lighter contrast precipitates contain very high concentration of Mo and W, while the darker precipitates contain much high- er Cr concentration than that of the matrix. The matrix composition has been found by EDX analysis to be as follows:

65.38 Ni, 17.7 Cr, 5.4 Fe, 0.66 Cb, 3.72 Mo, and 5.74 W

MECHANICAL PROPERTIES:

Table 1 lists the tensile properties of the as extruded alloy at room and elevated temperatures.

TABLE - 1 ROOM AND ELEVATED TEMPERATURE TENSILE PROPERTIES

Temperature Ultimate tensile

Strength

0.2% Offset yield

Strength

% Elongation 1

1

; °F °C MPa 1 KSI MPa KSI 75 24 1205 ! 175 T059 r~l54 8

600 316 1072 1 155 933 135 9 800 427 1052 152 891 129 9 1200 649 951 , 138 886 128 10

1400 760 667 i 97 638 93 18 ; 1600 871 369 ; 54 296 43 44 ;

A study of the table indicates that this alloy has good mechanical properties at room and elevated

temperatures. It is of interest to compare the properties of this alloy with a conventional Alloy 625 [ (Bal Ni, 20-30 Cr, 8-10 Mo, 5(max) Fe, 3.15 Cb and Ta). Alloy 625 has got UTS of 138.8 KSI at room

| temperature and 88.4 KSI at 1400°F. [4] .The rapidly solidified alloy shows nearly 25% higher tensile ! strength at room temperature and 10% at 1400°F. i Table 2 lists the stress to cause rupture in 100 hrs. at high temperatures. It is nearly 60%

of the yield strength at 1200°F and drops to 25% at 1400°F. The stress rupture values are slightly l| lower compared to Alloy 625 (71 KSI at 1200°F and 27 KSI at 1400°F). ij

TABLE - 2 |i STRESS TO CAUSE RUPTURE IN 100 HRS. I

Test Temperature Rupture Strength

°F °C MPa KSI 1200 649 413 60

1400 760 158 23

1600 871 48 7

CORROSION BEHAVIOR:

Immersion corrosion testing was carried out in different solutions and Table 3 lists the corrosion

rate of Markomet 8 alloy in various solutions. A study of the Table indicates its good corrosion resistance in various acid media.

444
TABLE - 3 CORROSION BEHAVIOR

Solution Concentration

(weight percent)

Test

Temperature

°F °C

Test Time,

hours

Corrosion Rate,

mils per year

Nitric Acid 50 176 80 48 2.6

65 212 100 48 10

Formic Acid 20 158 70 48 0.1

40 158 70 48 0.1

Phosphoric Acid 70 212 100 48 6 Sulfuric Acid 10 158 70 48 2 Hydrochloric Acid 5 158 70 48 157

The corrosion rate of this alloy is comparable to Alloy 625 in all the acids excepting hydrochloric acid.

OXIDATION BEHAVIOR : Table 4 lists the oxidation behavior of Markomet 8 alloy. The sample was heated to the desired

temperature and cooled to room temperature every 50 hrs. and then reheated. From the weight changes,

the oxidation rate was determined after 250 hrs. The oxidation resistance is excellent as evidenced

by less oxidation rate on prolonged exposure. However Alloy 625 on exposure at 1600°F for 250 hrs.

had an oxidation rate of 0.85 mg/cm2, indicating superior oxidation resistance of Markomet 8 alloy compared to Alloy 625.

TABLE - 4 OXIDATION BEHAVIOR

' Test Temperature

°F °C

Exposure Time,

hours

Oxidation Rate,

mg/cm2

1600 : 871 250 08 ^ 1800 i 982 250 1.37 ; 2000 ! 1093 250 2.16

SUMMARY:

A new bulk microcrystalline nickel-base alloy, designated Markomet 8, has been developed by Marko

Materials, Inc. using rapid solidification technology and powder metallurgical consolidation. This

alloy is characterized by high strength room and elevated temperature properties in combination with

good corrosion and oxidation resistance. The microstructural features of the alloy investigated by TEM

and SEM are reported.

REFERENCES:

[1]

. Rapid Solidification Processing II, R. Mehrabian, B.H. Kear, and M. Cohen, Eds, Claitors Publishing Division, Baton Rouge, Louisiana, 1980.

[2]

. Rapidly Quenched Metals III, 3rd International Conference, Vols. 1 and 2, Ed, B. Cantor, The

Metals Society, London, 1978.

[3]

. Rapid Solidification Processing, R. Mehrabian, B.H. Kear, and M. Cohen, Eds., Claitors Publishing

Division, Baton Rouge, Louisiana, 1978.

[4]

. Bulletin Cabot Alloy 625, Published by Cabot High Technology Materials Division.

445
Fig. lb

Fig. 1: a) Transmission Electron Micrograph of Markomet 8 and

b) The Corresponding SAD Pattern

446
40 um

» ( 8 um

I i

Fig. 2: Scanning Electron Micrographs of Markomet 8

Fig. 3: Optical Micrograph of Markomet 8

447
DYNAMIC WETTING/SOLIDIFICATION PHENOMENA

R.P.I. Adler and S.C. Hsu

GTE Laboratories, Inc. Waltham, MA 02254

ABSTRACT

A key to better understanding and control of Rapid Solidification Processing is the

characterization of the dynamic wetting interaction and the solidification behavior that

takes place between the liquid metal supply and the moving heat sink. The approach taken

has been to design simple screening procedures which reduce the experimental effort but

still allows the identification of significant parameters for the dynamic interaction between the solidifying liquid, a gas boundary layer, and a solid-heat removing substrate

without evaluating every process combination for this coupled mass/heat transfer phenomena.

A simple dynamic contact test procedure is now being used to generate phenomenological information and empirical results on the wetting and solidification of a droplet having a glancing impact with a metallic surface in a vacuum. The present methodology of impacting

fixed volume droplets from a series of metals and alloys near their respective melting

points can discriminate between the limiting dynamic wetting classes as well as systematic

trends as a function of impact angle within each grouping.

Introduction

The objective of this study is to identify the significant operational variables and thermo-physical parameters that control the efficiency of the melt spinning process. This involves understand- b|

ing the dynamic interaction between a solidifying liquid, a gas boundary layer, and a solid heat mj

removing substrate. l.i

Attention has been paid to the dynamics of the solid-melt-gas interaction [1-5]. The effect of 1'

surface free energy (liquid-vapor, solid-vapor, and liquid-solid) on the wetting/spreading of liquid *[

on a solid substrate has been well established through the conventional stationary (sessile drop) wet- f ting experiments [6,7]. The influence of substrate roughness, temperature and vibration, and inter- " facial tension and liquid-solid mutual solubility on wettability has been reported for these experiments [8-10]. There have also been several investigations in which a metal/alloy was rapidly quenched [ onto a moving substrate. The evaluations have separately considered the effects of: substrate rough- ‘ ness [2,4,11,12]; increased substrate temperature [13] (using a heated substrate to improve spread- i ing); and liquid nitrogen cooled substrates for enhancing cooling rates [14] or related to the increased solidification rate [15]. Others studied the influence of liquid characteristics such as the /

effect of : superheat on the foil width [16,17]; the impact forces on wetting [4,17]; and the in- *, volvement of quenching stress [18] on the foil structure. Finally, discussions about gas entrapment I

between the liquid and substrate leading to reduced effective areas of thermal contact and thus lower ':i heat extraction rates [2-4, 19-23] for the liquid, as well as the dependency on alloy composition [2] .j have been reported. ,

1 Since the spatial transients for this unique combination of phenomena are complex, this multiphase [ interaction has been difficult to comprehend and then correlate with the structure and morphology of

the resulting product. However the systematic classification of the variants of this kind of transient phenomenon not only could provide fundamental information but can also be usefully applied to optimizing productivity for the melt spinning process. In essence then we are trying to develop em-

^ pirical tests and then generate predictive guidelines that can be used to indicate appropriate com- Sr

binations of material and process conditions for any desired melt spun product. Tj

Experimental -

The experimental approach for this program involves developing a simplified dynamic wetting sys- I* tern that allows the screening of a range of materials in a consistent manner. Essentially a falling,

molten metal droplet is made to impact on an inclined substrate; by running experiments in a vacuum , environment it is possible to isolate the primary phenomenon from effects that may be induced by a

j gas boundary layer as illustrated in Figure 1. ]

I Initial activity has been directed at evaluating whether the Falling Droplet/Inclined Substrate

448
I

I

II

i;

VACUUM

Figure 1. Experimental Setup

Plane test produces results than can identify differences in dynamic wettability between various material combinations of melt and heat extracting substrate. In this test a fixed volume liquid metal

droplet (0.23 inch diameter) having a constant impact velocity of 66 in/sec strikes an inclined plane

of a "standard" OFHC copper with a surface prepared with 320 grit emery paper. These substrates have

: reproducibl e surface roughness values of 16 micro inches CLA and normally had a 280C surface temperature .

Resul ts

I Observations of the impact phenomena and any solidification that occurs were made using high

speed cinematography. The substrate inclination angle (0) was the primary experimental variable for

a representative group of five melt compositions having a wide range of melting temperatures and surface tension values. The initial impact and spreading phenomena at any given inclination angle were

surprisingly independent of melt composition. Thus each impacting molten metal sphere spreads into

a tear drop shaped disc with a maximum width that was inversely proportional to sin 9.

The subsequent phenomena could be classified as splat solidification or droplet recollection.

There were systematic dimensional trends that represented various types of dynamic wetting behavior.

As will be demonstrated, these effects would not be evident using classical, sessile drop wetting procedures.

The characteristic of forming solidified splats during dyanmic droplet testing has been used to identify the wetting class of materials (Cu, Zr, and A1 ) . The solidified splat width, taken as the

largest lateral dimension, is usually upwards of twice the droplet diameter but tends to fall off as

0 is increased (Figure 2). The elongated length of the splat tends to increase faster than the re-

' spective width decreases for increasing values of 9 so that the aspect ratio (Table 1) also increases

I with 0. Comparatively the magnitude of these effects are more pronounced for Cu, followed by Zr, and

1 least for A1

I

i Even though impacting droplets of the non-wetting class of materials also form a spreading liquid

I film on impact, dominant restraint forces cause this film to retract into a droplet that rebounds

1 away. Experiments using greater inclination angles (0=7OO) were useful in reducing the thickness and

! doubling the aspect ratios of the primary splats of most materials. However for Ni - 20% Cr

i rebounding droplets formed even at 0=70° because a sufficiently thin liquid film was not produced so

j that solidification could prevent droplet formation. On the other hand there is a transition in the

I wetting behavior of the AMS 4778 alloy (Ni-2.8% B-4.5% Si) droplets as indicated on Figure 2. At in-

' clination angles below 40O there is a characteristic non-wetting behavior while at angles of 45° and

I above the wetting behavior is intermediate between Zr and A1

I

I The systematic differences between materials exemplified by the trends in Figure 2 can not be

I exclusively related to the relative magnitudes of the repsective surface tensions. More significantly

i the transitional behavior of the nickel -base brazing alloy (AMS 4778) may be important because this

alloy can be readily melt spun on a copper substrate wheel. Further testing, analysis, and correlations are required to substantiate these results and correlative inferences to "melt spinnabil ity"

i but certainly the designed test procedure appears to have the resolution to discriminate between variI 449

i
Figure 2. Relationship Between Final Splat Width and Inclination Angle 0

MATERIAL ASPECT RATIO AT 0=30° 0=45° O

O

Cu 1.3 2.1 4.9

Zr 1.2 1.8 4.5

A1 1.0 2.1 4.6

AMS A778 DROPLET 1.7 4.0

Table 1. Systematic Dimensional Relations for Solidified Splats

ous classes of dynamic wetting behavior.

Concl usions

1) Splat dimensions represent balance between:

• spreading/solidification phenomena

• retractive surface energy forces.

2) Molten droplet/inclined substrate test discriminates between dynamic wetting classes.

3) This procedure to identify positive/negative classes of melt/substrate combinatipns for RSP

remains viable.

Acknowl edgment

The authors would like to thank the Office of Naval Research for the support of this research

under contract NOOOOl 4-82-0434. The technical assistance of Walter Allen and Neil Schauder is also

appreciated

450
References

1. T.R. Anthony and H.E. Cline, J. Appi . Phys . 50 (1 ) , 239 (1 979).

2. C.E. Mobley, R.E. Maringer and L. Oil linger, Proc. International Conference Rapid Solidification

Processing: Principles and Technologies, R. Mehrabian, B.H. Rear and M. Cohen, eds . , Claitor's

Publishing Division, Baton Rouge, LA, p. 222 (1977).

3. J.V. Wood and R.W.K. Honeycombe, J. of Materials Science 9, 1183 (1974).

4. S.C. Huang and H.C. Fiedler, Conf. Proc. "Continuous Casting of Small Cross Sections," Y.V. Murty,

F.R. Mollard, eds.. The Metallurgical Society of AIME, p. 127 (1980).

5. R.B. Pond, R.E. Maringer and C.E. Mobley, New Trends in Materials Processing, C.S. Hartley, T.E. Tietz and B.H. Rear, eds., American Society for Metals, Metals Park, OH, p. 124 (1976).

6. Lawrence E. Murr, Interfacial Phenomena in Metals and Alloys, Addison-Wesl ey Publishing Company,

Reading, MA (1975).

7. Arthur W. Adamson, Physical Chemistry of Surfaces, Interscience Publishers (1967).

8. S.J. Hitchcock, N.T. Carroll and M.G. Nicholas, J. Materials Science 16 , 714 (1981).

9. J.F. Oliver and S.G. Mason, J. Materials Science 15 , 431 (1980).

10. R.J. Rlein Wassink, J . of the Institute of Metals 95 , 38 (1967).

11. H. Jones, Material Sci. Eng. 5 , 1 (1969).

12. P.J. Wray, Met. Trans 12B , 167 (1981).

13. R. Togano, H. Rumakura and R. Tachikawa, Appl . Phys . Lett. 40 (1 ) (1982).

14. D. Pavuna and D. Hainsworth, J. Non-Crystalline Solid 37 , 417 (1980).

15. P.H. Shingu and R. Ozaki , Met. Trans 6A , 33 (1975).

16. F.E. Luborsky and H.H. Liebermann, Materials Sci. Eng. 49 , 257 (1981).

17. C.A. Williams and H. Jones, Material s Sci . Eng. 1

, 293 (1975).

18. A.R. Yavari , P. Desre and P. Hicter, Scripta Metal lurgica 1 5 , 503 (1981 ).

19. Ranjan Ray, Carl F. Cline, D.E. Polk and L.A. Davis, U.S. Patent 4,154,283, (May 1979).

20. M.C. Narasimhan, U.S. Patent 4,142,571 (March 1979).

21. Davor Pavuna, J. Materials Science 16 , 2419 (1981).

22. Makoto Matsuura, M. Rikuchi, M. Yagi and R. Sizuki, Japanese Journal of Applied Physics 19, 9,

1781 (1980).

23. H.H. Liebermann, Proc. Third International Conference on Rapid Quenched Metals, B. Cantor, ed..

The Metals Society, Vol . I, p. 34 (1978).

451
RAPID-SOLIDIFICATION PROCESSING OF TITANIUM ALLOYS

T. C. Peng, S. M. L. Sastry, and J. E. O'Neal

McDonnell Douglas Research Laboratories

St. Louis, Missouri 63166

ABSTRACT

Rapid-solidification processing (RSP) of titanium alloys was investigated using surface melting, splat quenching, and spin atomization techniques. Rapidly solidified 0.05-

0.25 mm thick flakes, and 100-500 pm diameter spherical powders were produced, and the

dependence of cooling rate on flake thickness and powder diameter was determined. Microstructures of several commercial Ti-alloys and novel Ti-alloys produced by RSP were determined. Novel alloy compositions and beneficial microstructures that can be obtained by

RSP of Ti alloys were determined.

Introduction

Rapid solidification processing (RSP) of titanium alloys results in fine, homogeneous metastable

microstructures in the as-solidified state [1-3]. These microstructures may correspond to desirable

property combinations, or they may be decomposed in a controlled manner by heat treatment to produce

microstructures desirable for improved properties [4,5]. Although RSP has had a major impact on the

metallurgy of nickel-, iron-, and aluminum-based alloys [6-9] , there has not been wide application of

RSP to Ti-alloy metallurgy. Recent studies on the rapid solidification of Ti, modified by the additions of rare-earths or boron, as well as titanium/high-aluminum alloys containing rare-earths

,

indicated that significant microstructural refinement can be achieved by RSP technology [10,11].

Preliminary studies of these systems indicated the persistence of the microstructural refinement

during subsequent heat treatment and suggested a correlation between this refined microstructure and

significant increases in strength and hardness of the alloys. The results obtained indicate promise

of RSP titanium alloys with higher modulus, strength, and creep resistance. The effectiveness of

different RSP techniques for producing the desired high cooling rates and unique microstructures in

Ti-alloys was determined in this investigation.

Experimental Procedure

The methods used for RSP were laser surface melting [3,12], electron-beam-melting/splat-quenching

and laser-melt/spin-atomization.

Laser surface melting was used principally for determining the dendrite-arm-spacing/cooling-rate

correlation for Ti-alloys [13,14] and for screening large numbers of potential compositions during

the development of new RSP Ti alloys. Prealloyed Ti alloys produced by arc-melting as well as Ti- | substrates coated with different alloying elements were used as the starting materials for determiningj

microstructural refinement and novel-alloy formation.

Electron-beam-melting/splat-quenching illustrated in Figure 1(a) was used to produce 50-500 g i lots of Ti-alloy flakes for microstructural examination by transmission electron microscopy and j j mechanical property evaluation. The electron beam is focused onto a prealloyed Ti-alloy rod. The j

; molten drops hit a rotating copper disk, and the centrifugal force stretches the droplets into thin

j

i flakes. The melt-drop size and drip rate are controlled by varying the power into the alloy rod. The|-

flake thickness (and hence the cooling rate) are controlled by varying the rotational speed of the

copper disk.

In the laser-melt/ spin-atomization technique, illustrated in Figure 1(b), a continuous-wave CO2 laser beam with an incident power of 6.2 kW was used to melt the top surface of a cylindrical, prealloyed titanium rod rotating at 10 000-30 000 revolutions per minute. The molten droplets expelled

by the centrifugal force from the circumference of the rod are cooled by a helium gas flow.

452
Inert gas

flow

Figure 1. Schematic diagrams of (a) electron-beam-melting/splat-quenching and (b) laser-melt/spinatomization methods for rapid solidification processing of titanium alloys.

The RSP Ti—alloys were analyzed by optical microscopy and scanning and transmission electron

microscopy to characterize the microstructural modifications effected by rapid solidification.

Results and Discussion

Morphologies and Cooling Rates of Rapidly Solidified Ti Alloys ; The morphologies of rapidly solidified Ti alloys produced by electron-beam-melting/splat-quenching and laser-melt/spin-atomization are

shown in Figures 2(a) and 2(b). The dendrite-arm/ cell spacings in the rapidly solidified alloys were

1-10 pm corresponding to cooling rates of 10^-10^ K/s. Although laser surface melting is not readily

Figure 2. Particulate morphologies of rapidly solidified titanium alloys produced by (a) electronbeam-melting/ splat-quenching and (b) laser-melt/spin-atomization.

453
adaptable for bulk production of material, it is an effective technique for producing high cooling

rates, and it is a rapid and inexpensive means for screening large numbers of potential alloy compositions. The electron-beam-melted and splat-quenched flakes are particularly suitable for a detailed

study of as-rapidly-solidified microstructures and microstructural evolution upon subsequent

annealing

The success of the laser-melt/spin-atomization technique for the production of fine (< 100 pm diam)

spherical powders with cooling rates > 10^ K/s indicates promise of solving some of the problems which

hitherto impeded titanium rapid-solidification technology. The reactivity of molten titanium with

crucible materials and the high affinity of Ti for interstitial elements such as carbon, nitrogen, and

oxygen have been the major obstacles in producing rapidly solidified Ti-alloy powder. These problems

are overcome in the laser-melt/spin-atomization technique. Furthermore, sporadic and local overheating, sputtering, electrode contamination, and preferential vaporization of low-melting components

of Ti-alloys encountered in spin-atomization using electric arc-discharge and electron-beam melting

are significantly reduced in the laser-melt/spin-atomization technique.

Microstructures of Rapidly Solidified Ti Alloys ; The novel Ti-alloys that can be obtained by rapid

solidification processing include dispersion-strengthened Ti alloys, intermetallic compound forming

alloys, eutectoid formers, and fine-grained conventional Ti alloys.

Dispersion Strengthed Ti Alloys ; Homogeneously distributed, fine, incoherent dispersoids were obtained in Ti-rare earth alloys. Figures 3(a)-3(c) show microstructures of conventionally processed,

rapidly solidified, and annealed Ti-2.0Er. The conventionally cast alloy has coarse, equilibrium,

constituent particles (Figure 3(a)) which produce low ductility and toughness. Rapid solidification

results in single-phase microstructures consisting of 0.2-0. 5 pm martensite plates (Figure 3b).

Although the equilibrium solubility of Er in both a-Ti and p-Ti is low [15], the single-phase microstructure in rapidly solidified Ti-2.0Er alloy indicates large increases in solubility by rapid

solidification. Annealing the splat-quenched flakes at 700-800°C produces homogeneously distributed

fine (< 0.1 pm) incoherent dispersoids (Figure 3(c)). Similar results were obtained in the other

Ti/rare-earth alloys. The incoherent dispersoids produce significant improvements in room- and

elevated-temperature tensile strength and creep resistance [16].

(a) 10 /im

Figure 3. Microstructures of Ti-2.0Er alloy: (a) conventionally processed, (b) rapidly solidified

and, (c) rapidly solidified and annealed at 700°C for 100 h.
Compound Formers ; Boron and carbon additions to conventionally processed Ti result in coarse titanium

boride and titaniuia carbide dispersions (Figures 4(a), and 5(a)). However, rapid solidification of

Ti-l.OB and Ti-l.OC alloys results in significant grain refinement and large volume fractions of fine

dispersoids (Figures 4(b) and 5(b)). The presence of such dispersions have been shown to increase the

strength by 100% and modulus by 30% of titanium [11].

Figure 5. Microstructures of (a) conventionally processed and (b) rapidly solidified Ti-l.OC alloy.

Eutectoid Formers ; The segregation at grain boundaries of eutectoid forming elements such as Ni, Cu,

and Fe in conventionally processed Ti alloys has hampered their development for high-strength application. The microstructures of rapidly solidified Ti-3Ni and Ti-7Ni alloys in Figures 6(a) and 6(b)

demonstrate the fine-scale homogeneous microstructures that can be produced by rapid solidification

processing. Ti-3Ni exhibits a fine martensitic structure, while Ti-7Ni has a retained-beta cellular

microstructure with fine Ti

2 Ni at the cell boundaries. Annealing of Ti-3Ni at 700°C for 16 h resulted

in growth of the alpha phase from the martensite and the formation of Ti

2 Ni at the boundaries of

impinging alpha grains. The fine-grained microstructure consisting of homogeneously distributed Ti

2 Ni

precipitates is conducive to increased strength. The results demonstrate that rapid solidification

processing of Ti-alloys containing eutectoid formers allows greater flexibility in controlling the

entectoid decomposition products.

Combined Precipitation- and Dispersion-Strengthened Alloys ; In rapidly solidified Ti-alloy systems

such as Ti-8Al-2.0Er a combination of shearable, coherent, ordered Ti^Al (u

2

) precipitates and incoherent erbium-containing dispersoids can be produced. These microstructures are particularly effective for increasing the room-temperature tensile strength and high-temperature creep resistance.

455
Ij Fine-Grained, Conventional Ti Alloys ; Rapid solidification processing of conventional near-alpha and f! alpha-beta titanium alloys such as Ti-8Al-lMo-lV, T1-6A1-4V, and Ti-6Al-2Sn-4Zr-2Mo produces martensitic structures (Figure 7(a)) containing dislocation substructures. Annealing these alloys at 800-

1000°C results in equiaxed, fine-grained microstructures (Figures 7(b) and 7(c)). Alloys having such*

equiaxed, fine-grained microstructures have superior room-temperature mechanical properties and hightemperature superplastic formability. J

Figure 7. Microstructures of rapidly solidified (a) Ti-6A1-4V, (b) T1-6A1-4V annealed at 900°C for 2 hours and, (c) T1-6A1-4V annealed at 1000°C for 1 h.

456
Acknowledgements

This research was conducted in part under the McDonnell Douglas Independent Research and Development program and in part under the Office of Naval Research Contract No. N00014-76-C-0626.

References

[1] S. M. L. Sastry, J. E. O'Neal, and T. C. Peng, Dispersion Strengthening and Rapid Solidification

Processing of Titanium Alloys, 110th AIME Meeting, Chicago, IL, 22-26 Feb 1981.

[2] J. E. O'Neal, T. C. Peng, and S. M. L. Sastry, Microstructural Studies of Laser-Melted TitaniumRare Earth and Titanium-Boron Alloys, in Proc. Electron Microscopy Soc. Amer ., ed. by G. W. Bailey (Claitor's Publ. Div., Baton Rouge, LA, 1981), p. 66.

[3] T. C. Peng, S. M. L. Sastry, and J. E. O'Neal, Exploratory Studies of Laser Processing of

Titanium Alloys, in Annual AIME Meeting, Chicago, IL, 22-26 Feb 1981.

[4] R. E. Maringer, C. E. Mobley, and E. W. Collings, Preparation and Properties of Compacts of Melt

Extracted Staple Fibers of Ti-6A1-4V Alloy, AIChE J. 180 , 111, (1978).

[5] S. M. L. Sastry, T. C. Peng, J. E. O'Neal, and L. P. Beckerman, Dispersion Strengthening of

Titanium Alloys, TMS-AIME Fall Meeting, 26-28 Oct 1982, St. Louis, MO.

[6] M. Cohen, B. H. Rear, and R. Mehrabian, Rapid Solidification Processing - An Outlook, in Proc

of 2nd Int . Conf . on Rapid Solidification Processing , ed. by R. Mehrabian, B. H. Rear, and

M. Cohen (Claitor's Publ. Div., Baton Rouge, LA, 1980), p. 1.

[7] J. J. Rayment and B. Cantor, The Rapid Quenching of Steels, in Proc. of 2nd Int. Conf. on Rapid

Solidification Processing , ed. by R. Mehrabian, B. H. Rear, and M. Cohen (Claitor's Publ. Div.,

Baton Rouge, LA, 1980), p. 165.

[8] A. R. Cox, J. B. Moore, and E. C. Van Reuth, On the Rapid Solidification of Superalloys, in

Superalloys; Metallurgy and Manufacture , ed. by B. H. Rear, D. R. Muzyka, J. R. Tien, and S. T. Wlodek (Claitor's Publ. Div., Baton Rouge, LA, 1976), p. 45.

[9] V. Anand, A. J. Raufman, and N. J. Grant, Rapid Solidification of Modified 7075 Aluminum Alloy

by Ultrasonic Gas Atomization, in Proc. of 2nd Int. Conf. on Rapid Solidification Processing , ed. by R. Mehrabian, B. H. Rear, and M. Cohen (Claitor's Publ. Div., Baton Rouge, LA, 1980),

p. 273.

[10] P. J. Meschter, J. E. O'Neal, and S. M. L. Sastry, Characteristics and Thermal Stability of

Dispersed Phases in Rapidly Solidified Titanium-Rare Earth Alloys, TMS/AIME Meeting, St. Louis,

MO, 24-28 Oct 1982.

[11] T. C. Peng, S. M. L. Sastry, and J. E. O'Neal, Microstructures and Mechanical Properties of

Rapidly Solidified Ti-B and Ti-6A1-4V-B Alloys, TMS-AIME Annual Meeting, Atlanta, GA, 6-10 March

1983.

[12] E. M. Breinan, B. H. Rear, and C. M. Banas, Processing Materials with Lasers, Physics Today, 29 , 44 (1976).

[13] H. E. Cline and T. R. Anthony, Heat Treating and Melting Material with a Scanning Laser or Electron Beam, J. Appl. Phys. 48 , 895 (1977).

[14] T. C. Peng, S. M. L. Sastry, and J. E. O'Neal, A Correlation Between Dendrite-Arm-Spacing and

Cooling Rate for Laser-Melted Ti-13V-3Al-3Sn-3Cr , Int. Conf. on Application of Lasers in

Materials Processing II, Los Angeles, CA, 24-26 January 1983.

[15] R. p. Elliott, Constitutions of Binary Alloys, (McGraw Hill, New York, 1965) 1st Supp

[16] S. M. L. Sastry, T. C. Peng, J. E. O'Neal, and L. P. Beckerman, Consolidation, Thermomechanical

Processing, and Mechanical Properties of Rapidly Solidified Dispersion-Strengthed Titanium

Alloys, Third Conf. on Rapid Solidification Processing, Gaithersberg, MD, December 6-8, 1982.

457
THE MICRO-PARTICLE PROCESSOR TO PRODUCE FINE POWDERS AND COATINGS*

j- Julius Perel , John F. Mahoney, Scott Taylor, Zef Shanfield**, and CarloS Levi

Phrasor Scientific, Inc. Duarte, California 91010

ABSTRACT

The Micro-Particle Processor has been fabricated at the Phrasor Scientific Laboratory after extensive research and development of the Electrohydrodynamic (EHD) technique of producing very fine droplets of various alloy metals. The Processor produces

small quantities of liquid metal droplets for use in studies of rapid solidification

processes and in the production of coatings. EHD has been used to produce tiny spheres

of alloys that are amorphous, single crystalline, bi-crystalline, etc. Adhesive coatings

have been produced such as aluminum on titanium and silicon on graphite. The Processor

could be used to produce small quantities of fine powders for material science studies

and for low volume applications. The method provides for rapid coating rates and

stoichiometic coatings because the coatings are made by impinging droplets rather than

atoms . o EHD produces fine droplets (<30A) by the application of an intense electric field

to a liquid surface. The liquid is produced by the heating of alloy metals in an alumina

reservoir and delivered to an alumina capillary nozzle held at high potential . Heating

of the droplet source, control of high voltage and all the other operations are performed

by a micro-processor via a video terminal. The software contains the sequence of operations that automatically controls the production and collection of sub micron powders.

The operator inputs parameters specific to the alloy tested by means of a keyboard.

The Micro-Particle Processor was fabricated and tested at the Phrasor Laboratory

then partially dismantled and shipped to the National Bureau of Standards. The Processor was in operation a few days after being uncrated by both Phrasor and NBS personnel.

Several minor modifications were made after delivery to allow for rapid turn around.

The Processor is being used at the NBS to further studies of fundamentals of rapid solidification, extended solubility and new alloy designs, such as, Ail-Mn and AJl-Fe combinations. Results of these tests will be presented by the NBS.

* This work was sponsored by DARPA and monitored by ONR under contract #N00014-77-C-0373.

** Present address: Northrop Research and Technology, Palos Verdes, California

t Present address: Mexico City

1. INTRODUCTION | In May 1982, Phrasor Scientific, Inc. delivered a Micro-Particle Processor (pPP) to the National

Bureau of Standards, thus completing the Table Top Powder Generator Program supported by ARPA and X

monitored by ONR. The program was initiated to (1) develop the method of high temperature Electrohydrodanamics (EHD) , (2) apply this technology to rapid solidification and other materials-oriented i areas, and (3) produce a unique laboratory instrument capable of generating micron and submicron

Jj particles of various alloy systems (Ref. 1-5). Before the initiation of this contract little, if any, J

technical information existed in the field of high temperature EHD to draw upon the experience of « previous research of state-of-the-art development. This program has succeeded in introducing into I Materials Science a unique material processing and research tool based on advances made in the •! understanding and construction of molten metal EHD devices during the past few years. The integration > of these two very distinct and separate technologies , EHD and Materials Science, for rapid solidification;

studies has stimulated interests in such diverse applications as nucleation studies, magnetic properties of fine powders, powders composed of refractory metal and ceramic materials, rare earth and

precious metal applications, copper powders for dispersion in inks used with hybrid circuits, photovoltaics, novel coatings and others. The capability of the EHD process to deposit thin films and

458
coatings has been evaluated as a process that could produce low-cost polycrystalline silicon sheet

for solar photovoltaic applications (Ref. 6).

The main objective of this program was the development of an automated Micro-Particle Processor,

capable of producing rapidly solidified powders in laboratory quantities for use by material scientists.

This work was motivated by the possibility of processing new material structures and compositions by

rapid solidification. Amorphous powders and new alloy compositions were originally anticipated and

have been produced by EHD.

Much more work in metallurgy is required to increase our understanding of the inter-relationship

between various process variables involved in rapid solidification. These variables, such as alloy

composition, cooling rate, microstructure, etc. can only be obtained through careful experimentation

and imaginative theoretical considerations. The Micro-Particle Processor is able to provide controlled and reproducible scientific experimentation affordable in a relatively small research laboratory.

It also provides a means for producing submicron droplets of both low and high temperature alloys

which are rapidly quenched in flight or by deposition on various substrates (conductive cooling).

The EHD technique, initially developed to produce charged liquid droplets of glycerol for space

propulsion applications, requires the application of an intense electric field to a liquid surface.

The intense field is obtained by applying high voltage to the tip of a capillary nozzle to which

the liquid is transported. The electric field draws the liquid out to a spike releasing charged

droplets which are further accelerated by the field. These charges can be as small as atomic ions

and can be as large as twice the capillary inner diameter, depending upon the electric field and the

liquid flow rate. For the rapid solidification application, alloy metals heated in an alumina crucible

are delivered by gas pressure to an alumina capillary nozzle, entirely operated under high vacuum

conditions (10 ^torr). The droplets cooled by radiation and solidified to powders, are collected by

various substrates for examination by electron microscopy. Powders ranging from 30nm to lOOym have

been observed with estimated cooling rates by radiation alone ranging from 10 K/s to 10 K/s. When

coatings are made the conditions are arranged so that the droplets remain molten when they impinge

upon the collecting substrate.

2. CAPABILITIES AND ACCOMPLISHMENTS

Capabilities of the Micro-Particle Processor

The capabilities of EHD to produce rapidly quenched metallic powders for the materials scientist

are listed below:

° Solidification studies relating average cooling rate, prior to solidification, and

particle size to the formation of amorphous and crystalline structures.

° Rapid solidification studies performed to determine the effect of particle size and

cooling rate on increases in solute solubility of alloyed elements.

° Rapid solidification and heat transfer studies on extremely fine crystalline and glass

metal powders.

° Production of electron transparent metal powders. The direct observation using TEM of

a complete casting is a unique capability heretofore not available.

° Nucleation studies that relate particle size to the number of crystals (number of

nuclei) in each particle.

° Electron diffraction study of fine powders to establish the sequence of formation of

stable and metastable phases in a given alloy system as a function of cooling rate.

° The electron transparency of the fine powders produced permits studies of the growth

morphology (plane front, cellular, or dendritic) during solidification. Effects of

particle size, cooling rate, and solute content upon liquid-solid interface morphology

can be studied.

The above listed capabilities have been demonstrated in several publications (Ref. 2, 3, 4, 7).

459
Results of Powder Studies

In the course of developing the yPP, several materials were used to extend the EHD technique to higher temperatures and to include a wider variety of material types. The major differences between

the alloys experienced in EHD involved such liquid parameters as melting point, surface tension and

wetting angle between liquid and nozzle. Table I lists the pure elements and alloy combinations tested at Phrasor. The work at NBS extends this list even further. These studies resulted in observations

of amorphous, single crystalline, bi-crystalline and tri-crystalline powders. Studies of extended solubility segregation, nucleation and metastable phases were made, greatly aided by the fact that the

powders were sufficiently small to be electron transparent.

A few examples of the powders produced by the EHD technique and analyzed by TEM are shown here.

They illustrate interface breakdown due to segregation and nucleation behavior producing bi-crystals

j

in smaller powders and micro-crystals in larger ones.

The Ail-6% Si powders below 200nm solidify in a plane front mode without segregation even for Si

content as high as 6% when the solubility is 1.65%. This extended solubility indicates the achievement of large undercoolings prior to nucleation. In the larger particles, there was usually a tran- sition from plane-front to cellular (segregated) growth as the droplet size increased and the achievable

undercooling was consequently decreased. Fig. 1 shows an example of the interface break-down due to segregation in an A£-6% Si powder.

A large fraction of the powders in the range of 200nm or smaller are bi-crystals as shown in

TABLE I

Figure 1. TEM image and diffraction pattern of an A£-6%Si powder, showing the transition between plane

front and cellular growth modes due to segregation.

LIST OF ELEMENTS AND ALLOYS SPRAYED

Elements Binary Amorphous Alloys

Pb Pb-60% Sn Fe-15% Si-10% B Sn A5.-4.5% Cu Fe-40% Ni-14% P-6% B Al A£-28% Cu AJl-30% Ge

Cu AH-3% Si

Ge AH-6% Si

Si AI-12Z Si

AJl-25% Fe

Al-1% Mn

A£-30% Ge

460
Fig. 2. Some preliminary diffraction and micro- chemical information suggests that one of the crystals

may be A£-rich and other, Ge-rich but this is not conclusive.

Most of the larger powders, above lOOnm, have a structure similar to that in Fig. 3 where a large

number of crystals in the range of lOnm can be distinguished by diffraction contrast. Some of the powders that could he properly tilted in the TEM revealed a microstructure reminiscent of that produced

by the plane-front to cellular transition in Fig.l.

3. RESULTS AND APPLICATIONS

A micro-structure study was made by Levi and Mehrabian (Ref. 6) using the electron transparent submicron aluminiam alloy powders produced at Phrasor Scientific by the EHD process. Various microstructures of several hundreds of powder samples were observed including homogeneous plane-front solidified single crystals, cellular crystals, block segregates, multiple grains and twins. With the aid

of thermodynamic concepts, thermal history and solidification modes were deduced. It was found that

segregation decreased with smaller particle sizes. This presumedly was caused by the greater undercooling experienced by the smaller particles. Nucleation studies also showed that larger undercooling

promotes multiple nucleation and the occurence of twins. These and other observations were made showing how electron transparent powders generated by EHD are very useful in fundamental solidification

studies.

Among the most interesting applications of EHD produced powders are their use as Analytical Electron Microscopy standards and their suitability for re-melt and re-solidification experiments in the

electron microscope (Ref. 4 and 8). These uses are made possible because of the small size of the EHD particles and because they are generally spherical in shape.

The capability of the EHD generation method to produce and control charged liquid droplets leads

to a wide range of applications in materials Science. Some of the important features of EHD useful in

these applications involve: copious production of submicron droplets, the high speed to which they can

be accelerated, the vacuum or gas environment selected, the choice of selecting that either droplets

or powders can be impinged upon the substrate. Different uses considered for applications of powders

include micron sized copper powders for conductive inks, silicon spheres for photovoltaics and

461
special chemical compounds produced by the projection of molten droplets through carbide, nitride or hydride producing atmospheres. Fine metal (or non-metal) flakes can be produced by splat impingement

on a non-adhering surface.

Surface properties can also be altered by coatings to increase hardness, prevent corrosive contact

retard oxidation, increase wear resistance, and in general provide a coating which lowers the overall

cost of the raw material or production such as the production of solar cells (Ref. 6). Spongy coatings

have been produced. Special coatings on a reactive substrate may be obtained that is not achievable

in a bath or gaseous atmosphere.

4. THE MICRO-PARTICLE PROCESSOR

The production of the Micro-Particle Processor was the final goal of the DARPA/ONR Program (Tabletop Powder Generator) (Ref . 9) . The Processor is able to produce about 20 grams of alloy material in

about two to three hours. This limitation is due to the size of the reservoir volume (as listed in

Table II, Specifications) in this particular model. It is possible to increase powder yields by using

a larger reservoir or by operating multi-nozzle droplet sources. The production rate increases nearly

linearly with the number of nozzles. Table II also lists some of the other pPP specs.

A photograph of the pPP delivered to NBS, is shown in Figure 4 and a diagram showing its dimensions are given in Figure 5. Within the Main Console are the vacuum pump system, the high voltage

system, the source heater controls, and the electrical control relay system. On the left side on top

of the Main Console is the Computer Console containing the computer, floppy disc drive, video screen

and the vacuum gauge (controller) . In front of the Computer Console is the printer and the control

keyboard. On the right side in the vacuum chamber which contains the EHD source visible through the

top view port and the powder (sample) collector which can be observed through the bottom view port.

Behind the vacuum chamber is a cylindrical column routing the high voltage and source heater cables to the top of the chamber. These cables are shielded by the High Voltage Cover on top of the column and

the chamber. Figure 6 shows a close-up photograph of the Computer Console and the chamber showing the EHD source operating at high temperature. The video screen is shown displaying parameters of the ex- perimental run. Figure 7 shows the video screen in the Standby Mode, as it appears when the pPP is

first turned on. It provides the date and time, inserted by the operator, and describes the states of

the various vacuum valves, the pumps, the pressure, whether the cooling water is on and whether the

view ports are open or closed. The different modes of the yPP, shown on the screen determine the se- quence of the operation that is all computer controlled. The modes of operation are listed below:

1. Standby Mode. 3. Start Vacuum Pumps 5. Heat up & Operation

2. Set Parameters 4. Pumpdown Chamber 6. Shutdown

TABLE II SPECIFICATIONS

SOURCE

Voltage 1 to 20kV

Current 0 to 3mA

Nozzle Heater Power 500W Max. Reservoir Heater Power... 500W Max. Operating Temperature. ... To 1350°C (2462°F)

Reservoir Pressure To 2000 Torr Max. Reservoir Volume 0.3in^ (4.8cm^)

Nozzle Orifice 002" DIA (.005cm)

MATERIALS:

Reservoir High Purity Alumina

Nozzle High Purity Alumina

Source Gasket GRAFOIL (graphite)

Extractor 030" Graphite

Nozzle Heater.... .020" Ta

Reservoir Heater. .020" Ta

Thermocouple Type K (Chromel-Alumel)

CONSOLES

Main Console 40" high by 47" wide

by 40" deep (101 x 119 x

101cm)

Computer Console. ... 27" high by 21" wide by

21" deep (68x53x53cm)

Vacuum Chamber 12" i.d. x 27.5" high

(30 X 70cm) Figure 5. Diagram of the Micro-Particle

Processor.

462
Figure 7. Close up of ttie Standby

Mode on the Video Screen.

Figure 6. Photograph of the Computer Console and the Vacuum Chamber.

When the Set Parameter Mode is selected, the screen changes to display all the important parameters to !dentiS ?he parSSlar run such as date,time,operation,alloy,temperatures,etc. After these are complete, the display returns to the Standby Mode with the cursor on the next mode xn the sequence. When

Start Vacuum pLps is selected, the operator can turn on the pumps and open the proper valves. The

coSterprogram also prevents the operator from making mistakes such as turnxng on the vacuum pump and

opeLng tL Lte valve while a view port is open. The next mode allows for the chamber pumpdown. The

nLtup and Operation Mode is the one in which the alloy is sprayed. After completion of the run, the

Shutdown steps are sequenced.

The Micro-Particle Processor produced on the AKPA/ONR Program was delivered to the National Bureau

of Standards in May 1982. In a co-operative effort between Phrasor Scientific and the Bureau, the p wL Jufinto operSion in a few days'and the A.-Fe alloys were processed The NBS Solidification Program

using the pPP includes fundamental studies of extended solubility, new alloy design, comparison

theory with experiment and development of predictive models of micro-structure formation Results from

the iLtial experiments on the yPP as part of the extensive NBS program are reported by Shechtman a d Ridder in this Third Rapid Solidification Processing Conference (Ref. 10).

ACKNOWLEDGEMENTS

The authors are grateful for the support provided by Dr. E.vanReuth and Col. L. Jacobson of D

and Dr. B. MacDonald of ONR. Dr. Pol Duwez, well known in Rapid Solidification, was extremely help

REFERENCES

1 J F Mahoney, J.Perel,Z.Shanfield and B.E.Kalensher, "Development of a Tabletop Model for the Generation of AmLphous7Microcrystalline Metal Powders", Interim Tech.Report #80-l,ONR,Arlington,VA(1980)^

2. J.Perel,J.F.Mahoney,B.E.Kalensher,K.E. Vickers & R.Mehrabian, Rapid Solidification Processing. Prin

cinles and Technologies, edited by R.Mehrabian,B.H.Kear,M.Cohen(Claitor ' s P^b.Div. ,BatonRouge,^,1978)

3. j.Perel,J.F. Mahoney, B.E.Kalensher & R.Mehrabian,Advances in Metal Processing , edited by J.J. Burke,

R.Mehrabian & V. Weiss, (Plenum Pub. Corp. 1981) ,p. 79. u

4. J.Perel,J.F.Mahoney,S.Taylor,Z.Shanfield & C .Levi, Amorphous and Crystalline Allo^,B.H.Kear,

B.C. Giessen and M. Cohen, editors , (Elsevier Science Pub. Co. , North Holland 1982),p.l31.

5. J.F. Mahoney, "Tiny Particles Aid Research on Microstructures , Indus. Res. &Dev. ,24, 150 (January 1^

6. J.F. Mahoney & J .Perel, "Electrostatic Spray of Silicon for Photovolataic Applications ,IEE/IAS 1981

Annual Meeting, Philadelphia, PA, (5-9 October 1981).

^ c v • v aoT-c"

7. G.C.Levl & R.Mehrabian, "Microstructures of Rapidly Solidified Aluminum Alloy Submicron Powders , Metallurgical Trans., 13A, 13, (1982).

8 H v-rncp>r TTniversitv of Illinois, private communication.

I'. jiF.Mahoney.J.Perel.Z.Shanfield.B.E.Kalensher S, s. Taylor, Development ot a 0ma982)

Generation of imorphous/Microcrystalllne Metal Powders”. Final Technical Report #82-1.0NR(1982)

10 D Shechtman.S.D. Ridder, & R. Mehrabian,"Microstructural Studies on Supercooled ^ub-Micrometer

Powder Particles. "Third Conference on Rapid Solidification Processing: Principles _& Technologic,

NBS, Gaithersburg, MD(6-8 December 1982).

463
RAPID SOLIDIFICATION SURFACE TREATMENT

OF

DISCS AND RINGS BY ELECTRON BEAM PROCESSING

M. TULI

Pilgrim Materials Corp. South Windsor, Ct. 06074

P.R. STRUTT

Uni. of Connecticut, Storrs, Ct.

ABSTRACT

The technological requirement of surfaces with enhanced wear and rolling contact fatigue resistance frequently involves treatment of cylindrical

surfaces. As a consequence, the electron beam surface melting technique

has been applied for treating discs and rings with various dimensions. The

technique has been successfully applied for the mi cros tructua 1 modification

of hard materials.

INTRODUCTION

Methods for producing hard, wear and corrosion resistance facings on materials

are important in terms of:

1. Improved performance

2. Cost reduction

3. Rejuvenation of high technological components.

For this reason various types of plasma deposited coatings are being used to a ever increasingly scale in variety of industrial applications.

Directed energy beam processing now provides a unique approach to forming sur- face layer having an unusual degree of microscopic refinement and homogeneity (1,2).

Most frequently, these layers are perfectly bonded to the substrate material. This

particular aspect can be of extremly importance in specific application where the

lack of perfect bonding in the case of plasma deposited layer is frequently regarded

as a significant draw- back.

Electron beam processing is a particularly versatile technique since the beam

may be readily programmed to treat large area of complex geometry (3).

EXPERIMENTAL TECHNIQUES AND RESULTS

In laser/electron beam glazing momentary melting and ensuing rapid solidification

proceed as the high energy beam traverses the material surface. This results either

in the formation of microcrystalline or amorphous layer. The fact that some alloys

undergo transformation into the amorphous state has led to the term "laser glazing"

and "electron beam-glazing". The actual cooling rates achieved and the thickness of

the rapidly solidified layer critically depends upon the power density and the velocity of the beam relative to the substrate material. In contrast to other rapid

solidification techniques the cooling rate may be varied over a particularly wide

range, from 10^ to IQoOc.s"!, for example, with melt depths ranging from 2mm to a micron ( 4 ).

Rapidly solidified surface layers of high demensional tolerance and homogeneity

are attainable using ultra-rapid beam deflection system specifically developed by

Pilgrim Materials Corporation. Precision control of electron motion is achieved by use of a low inductance/high current deflection yoke driven by an ultra-rapid response

464
! current control amplifier. This circuit is designed so that the beam can follow any

j type of oscillation voltage waveform,

j Fig. la shows the actual current triangular waveform through the coil and Fig. lb

shows the actual input voltage waveform which is produced by current feed-back cir-

{cuitry. Negligible disturbance of triangular current waveform is maintain upto a 'frequency of 15KH. The complete circuitry involved in this application is contained

! within a single unit, as shown in Fig. 2. A cylindrical bar (M-42 tool steel)

1 diameter, was treated using el ectromagneti cal ly oscillated beam generated by this

unit (Fig. 3).

Topographical and Mi crostructural Characteristics

Fig. 4a and b are low and medium magnification micrograph of a 1 ongti tudi nal

view of a 1.25 inch (3.18 cm) wide zone produced by an electron beam laterally os- cillated at 1000 hertz. Clearly evident in these figures is the uniform thickness of

I the rapidly solidified layer; the melt depth is approximately 250 urn. In addition,

j the actual surface is particularly smooth and does not display any evidence of ripple

structure. The entire zone etches in a homogeneous manner, which is indicative of a

I uniform microstructure and an absence of significant segregation of alloying elements.

In Fig. 4c the zone thickness seems to taper-off very smoothly at the edge.

Obviously, this highly desirable feature results from an almost instantanous reversal

of the electron beam at the maximum amplitude.

Finally, a high magnification scanning electron micrographs in Fig. 5 shows a homogenous cellular solidification structure in the rapidly solidified M2 high speed

j steel. This structure is typical of the hard martensi ti c/ansteni ti c microstructure

observed both in laser and electron beam glazed M2 high speed steel. (6).

CONCLUSIONS

This study shows that a much higher quality, homogenous electron beam glazed

surface ( 1 micron) can be formed on a cyl i ndri cal /di sc surface using "line heat

source"

An important industrial application of the technique just described is for the

surface treatment of high performance bearing used in gas-turbine jet.

REFERENCES

1. P.R. Strutt, H. Nowotny, M. Tuli and B. Kear, Mat. Sci . & Engr. 217 (1978)

2. P.R. Strutt, Metal. Forum, 4 117 (1981)

3. M. Tuli, P.R. Strutt, Materials Research Society Annual Meeting (1981)

4. R.J . Schaefer , T.R. Tuker and J.D. Ayers, Laser/E. B. Processing of Materials, Academic Press, N.Y. 749 ^980)

5 . P.R. Strutt, Mat. Sci. & Engr. 44 , 239 (1979)

465
F i g . 1

a. Shows the actual current triangular waveform through the coil.

Fig. 2. Single unit with complete

ci rcui try.

b. Shows the actual voltage waveform

which is produced by current

feed-back circuitry.

Fig. 3. A cylindrical bar (M-42 tool steel) was treated

using el ectromagneti cal ly oscillated beam

generated by unit shown in fig. 2.

466
(C)

F i g . 4

a&b are low & medium magnification

micrograph of a 1 ongti tudi nal view

of a 1.25 inch (3.18 cm) wide zone produced by an electron beam laterally

oscillated at 1000 hertz. Melt depth

is approximately 250um.

In 4c the zone thickness seems to

taper-off very smoothly at the edge. F i g . 5

Shows a homogenous cellular solidification structure in the rapidly

solidified M2 high speed steel.

467
PLASMA DEPOSITION EFFECTS ON STRUCTURES AND PROPERTIES

OF RSPD SUPERALLOYS

R. W. Smith

General Electric Gas Turbine Division

Schenectady, NY

D. V. Rigney, J. S. Smith

General Electric Aircraft Engine Group

Evendale, OH

J. R. Rairden

General Electric Corporate Research and Development

Schenectady, NY

ABSTRACT

The low pressure plasma deposition process has evolved from research and development

to manufacturing during the past six years for the application of environmentally resistant coatings to superalloy components. Recently, research and development studies have

extended this technology to the fabrication of rapidly solidified plasma deposited (RSPD)

structural bodies. This report documents the results of a coordinated study, funded by

DARPA, (Project 4289) in three organizational components of GE to evaluate several com- mercially available plasma spray systems for deposition of superalloy components.

Introduction

The low pressure plasma deposition process has evolved from research and development to manufacturing during the past six years for the application of environmentally resistant coatings to superalloy components. Recently, research and development studies have extended this technology

to the fabrication of rapidly solidified plasma deposited (RSPD) structural bodies. [1] This is a progress report on the results of a coordinated study to evaluate and further develop several com- mercially available plasma guns for deposition of superalloy components.

Three important factors in plasma spray deposit quality are: the plasma jet, powder distribution/injection, and environment.

These factors control the heat content of the gases, the heat transfer to the powder particles, the location of the particles in the jet, the particle velocity, and the environmental reactions. Therefore, the degree of particle melting, the deposition pattern, the deposit density,

the deposit efficiency, and the extent of deposit inclusions are determined by the processing

factors. The arc gas flow rate and composition, the nozzle design (manufacturer or manufacturer

model), and the power level contribute most to the energy content of the gas. The nozzle design

and the jet pressure ratio (Po/Pchamber) will determine how the energy is distributed in the

plasma jet, either in heat content (enthalpy, H) or velocity (V) and determine the exit Mach

number (Ma #). Powder injection parameters, such as powder size distribution, injection gas flow rate, powder flow rate (P.F.R.), injection location and injection angle, will all affect

powder particle velocities and trajectories. Hence, how well powder particles will be melting

become controlled by the plasma jet and the powder injection.

The intent of this paper is to present a progress report on a coordinated study, under contract to DARPA, to evaluate several commercially available plasma guns, operated in accordance

with manufacturer's instructions and modifications discussed with the manufacturer. These evaluations were conducted on two Ni-base superalloy deposits, Rene 80 and Hastelloy Alloy X, in an effort to determine plasma deposition parameters.

468
Experimental

Different vendors' guns were studied as a function of processing parameters i) power level

ii) arc gas composition iii) arc gas flow rate iv) nozzle design v) powder injection location,

internal and external vi) pressure ratio (chamber pressure) vii) and relative gun/substrate

motion. A y' strengthened superalloy, Rene 80 and a non-y', solid solution/carbide strengthened

superalloy, Hastelloy Alloy X were selected as the candidate alloys for this study. The initial

processing conditions used were either those recommended by the gun vendor or those which had been

developed in-house prior to the start of the project. Table 1 summarizes the parameter ranges

evaluated for each of the plasma gun designs.

TABLE 1 : PLASMA DEPOSITION PARAMETER RANGES

Arc Gas Flow Chamber Powder Pow der Iniection TA Spray

Anode Power

(kw) Ar

(slm) He Hz Pressure

(torr) Feed Rates

(g/min)

Flow

(slm) Location Angle (KW) Distance

(mm)

Vendor

Ar/He

A

Mach 3 Mach 2 (std)

(new) 40-80

50-72

108-180; 90-120;

28-140;

23-64; —

14 - 120

20 - 300

100

100

15 4-15Int. & Ext.

Int. & Ext. 20°

20°

0-3

0 400

400

Vendor

Ar/Hz

A 74 169 — 23 60 190-380 17 Int. 20°, 25° 30° 1 381-457

Vendor

Ar/Hz

B 40-56 50 — 6-9 30 - 235 76 0.7-

2.9

Ext.

8 - 22mm

-2° 0 260-410

Stationary (sweet spot) or traversing deposits were made onto 150 x 150 mm plates in thicknesses from 3mm to 15mm for deposition times from 1 1/2 to 30 minutes, depending on the particular

evaluation or specimen requirement.

The alloys, Rene 80 and Hastelloy-X, were selected for study because they are used extensively

in gas turbines. RSPD data can be compared with a large data base for both alloys as conventionally processed.

Results and Discussions

Evaluations of deposits made during processing parameter studies were spray pattern (sweet

spot and traversing), deposit efficiency, as-deposited density, heat treated density, as-deposited

metallography, heat treated metallography, and selected tensile testing.

Spray Pattern

The deposit spray pattern is a measure of the particle trajectories in the plasma jet measured

at the substrate. The spray patterns, in general, are characterized by bell shaped mounds which

build up during deposition. The patterns have been found to be affected by the gun/substrate

motion, spray distance, chamber pressure powder size distribution, nozzle design, and powder feed

injection parameters (i.e. location, number of ports, injection angle, etc.). The spray pattern

spreads, shrinks, shifts axis, or changes thickness distribution with these process variables.

The thickest, central region of the deposit is the most dense, however, the periphery of the

deposit rapidly decreases in density. This drop off in density has been attributed to a "fringe"

of the plasma jet where powders are heated less efficiently and may be completely unmelted yet

become entrained in the deposit. The "fringe" region has also been found to be sensitive to the

aforementioned process variables.

In general the spray pattern has been found to increase in diameter and become more distributed with i) decreasing chamber pressure ii) increasing power iii) increasing spray distance

(S.D.), while the "fringe" becomes larger with i) decreasing power ii) increasing chamber pressure

iii) increasing arc gas flow rates iv) and gun/substrate motion.

Deposit Density

The as-deposited density is a measure of the melting efficiency of the plasma jet and as such

is a good indication of the process variables effects on deposit quality. The density has been

469
found to be quite sensitive to deposit location (center vs. fringe) and substrate/gun motion where

the "fringe" of the plasma jet has to be included over the entire deposit, unlike "sweet spots"

where the fringe is not included in the deposit center.

The deposited structures have all been heat treated from 1200°C to 1250°C for one or more

hours to complete the consolidation process, in accordance with the results reported in Ref 2. Density improvements from 1 to 3 percentage^^ points were observed by heat treatment. For this

study, all specimens received a 2 hour 1250°C densif ication heat treatment in vacuum or argon

before the final alloy heat treatment.

Table 2 shows the range of as-deposited and heat treated deposit densities over a wide range

of processing parameters measured under the conditions of this evaluation for the Rene PO and

Hastelloy-X sweet spot and traversing deposits. It was found that densities varied from 94.0 to 100.4 as-deposited to 98.9 to 100. S'*" for heat treated samples.

TABLE 2: DENSITY OF DEPOSITS

% Density”

Rene 80 Hastelloy-X”

Gun As Deposited Heat Treated 0 1250°i: As Deposited Heat Treated 0 1250“C

Vendor B, Ar/H 2 Large Plates

96.5 - 96.1 - 100.5 98.4

99.1 - 100.9 99.4 - 100.6 97.6 - 97.0 - 100.5 99.5

99.8 - 101.5

99.3 - 100.7

Vendor A, Ar/He

Large Plates

96.3 - 94 99.9 98.7

98.9 - 100.3 N/A 94.4 - 95.4 - 99.7

97.7

96.3 - 99.8

N/A

Vendor A, Ar/H 2 Large Plates

95.9 - 97.8 - 99.3 98.5

99.4 - 100.6 99.4 - 100.8 96.7 - 96.1 - 99.6

99.4

99.1 - 100.3 98.4 - 100.6

•For Rene 80, 100% density =8.17 gm/cm^ •’For Hastelloy X, 100% density = 8.23 gm/cm*

Microstructure - As-Deposited j

I The as-deposited microstructures of Rene 80 and Hastelloy-X were most affected by RSPD proces- ; sing variables. The deposits are formed by a rapid succession of powder particle impacts built up on top of one another. The impacting powder particle population has been found to be in various ; stages of melting and/or solidification. Hence the deposits include features of unmelted parti- I cles, partially (interdentritical ly) melted powder particles, smaller pieces of larger dendritic ! powder particles, or completely melted particles which splat out into thin lamellae. Figures 1 | thru 4 are examples of these microstructures. It has been found that the plasma jet enthalpy

level most controls the as-deposited microstructures described above, but other variables can also

effect microstructure. The plasma jet heat flux and high substrate temperature was found to induce directional grain growth mostly in the Hastelloy-X, coarsen y', or precipitate M 23 C 5 ! carbides in Rene 80. These changes occurred both through the thickness (function of time vs. | temperature) or radially. The radial changes in microstructure were attributed to both heat flux , and the degree of entrainment of unmelted powder particles from the plasma jet "fringe". Figure la illustrates the typical unmelted particulate nature of the "fringe" area Rene 80 deposit. I

j Figures la through Ic illustrate the range of deposit microstructures of Rene 80. Figure la / microstructure represents low enthalpy deposits made (either low power or high gas flow). Fig- ! ure lb shows a typical microstructure of Rene 80 or other high y' nickel-base superalloy showing a

| moderate volume fraction of unmelted (spherical with dendrites or cells from atomization) or partially melted particles in a melted matrix. These type deposits were obtained from the recom- I mended plasma jet parameters from Vendor A Mach 3+ conditions. Vendor B jets or further changes f by lowering the Vendor A jet Mach number while keeping the power high have higher enthalpy "hotter" jets which produce more complete melting seen in the deposit microstructure Ic. The sweet

j spot deposit has very few unmelted particles, has a high density, and also 1-3 ym thick "splats" f showing very good melting.

| Selected jet conditions were chosen for spraying the Hastelloy-X, and Figures 2a and 2b show

the representative microstructure of the deposits. It can be seen that in all cases recrystal li- l. zation has occurred, removing any of the "splat" microstructure. It is believed that the low ' alloy content and lack of y' or any other precipitate in these Hastelloy-X deposits has removed . the barrier for grain growth that in most cases maintained the splat nature of the Rene 80 depo-

|j sits. The "sweet-spot" deposit plasma jet heat flux or substrate temperature in all cases was I' enough to induce grain growth in all the Hastelloy-X deposits. Relative differences in density |:

'‘'Note: Greater than 100 percent density values probably result from compositional difference

from reported values. j 470
and uniformity of grain size across the deposits can be used to judge the optimum plasma jet for

depositing Hastelloy-X. Again, the higher enthalpy jets produced the densest, most uniform

Hastelloy-X deposits.

Traversing of the plasma jet has been found to reduce the grain growth that occurs in the

Rene 80, and more substantially, the Hastelloy-X deposits. In addition, traversing the substrate

forming deposits which include the "fringe" has been found to cause interpass layering of most

deposits.

Figure 3 shows the typical structure of a traversing Vendor A Mach 3+ Rene 80 RSPD deposit in this evaluation. A layer line is seen in these deposits and its origin lies in the traversing of

the plasma jet over an area larger than a sweet spot. Note the high unmelted particle content of

the layer lines in Figure 3. It is believed that during the traversing, the fringe of plasma jet

deposits both less than ideal ly' melted particles and condensed vapor from vaporized powder on to

the deposit surface.

Figure 4 is an example of the interpass layering that also occurred in this evaluation with

the Vendor B plasma deposits. The structures are well melted; however, bands every 30-45ym delineate pass interfaces.

In summary, under the condition of this evaluation, plasma nozzle design, arc gas composition

and flow rate, and power level varied the jet enthalpy the most. Hence, most of the as-deposited

major structural differences resulted from these changes. Chamber pressure, which changes the

nature of the jet expansion and powder feed (location and feed rate) can affect the as-deposited

microstructure, density, and degree of melting. The balance of parameter changes, spray distance,

and transferred arc did not significantly change the structure of these deposits, except where

heat flux to the deposits may have induced grain growth. The larger traversing deposit plates

produce a very uniform deposit along the centerlines and variations in deposit location are sensitive to deposit condition or alloy composition. Large deposits are also sensitive to interpass

layer lines believed to be entrapment of fringe "debris" (unmelted or vaporized and condensed

powder particles)

Microstructure - Heat Treated

The most important processing aspect of the heat treatment on microstructure is the influence

that the as-deposited microstructure has on the final-heat treated structure. The effect of plasma processing on the as-deposited microstructure has been discussed previously. Features in the

Rene 80 deposits were more varied than in the Hastelloy-X and include porosity, completely

unmelted (possibly resolidified) deformed particles, fine grain melted matrix, interpass lines,

lamellar structure from splatting of particles, and recrystallized grains. The Hastelloy-X deposits showed mostly recrystallized structures and the other features observed in Rene 80 were not typically discernible.

Figures 5a and 5b illustrate the range of microstructures observed for the heat treated

(1250°C/2 h FC + 1204°C/2 h RQ + 845°C RQ) Rene 80 deposits. One example shows a fine grained

structure (2 to 20 ym) with limited grain growth being possibly controlled by the volume fraction

of unmelted particles; in fact, the prior particles are decorated similar to P/M alloys. Another

example shows better grain growth with equiaxed grains, more uniform size (25 to 75 ym) and no prior unmelted particles.

In summary, it has been found that for the highly alloyed y' strengthened, Rene 80 grain sizes were limited to maximum 50-75ym, with high heat treat temperatures, while for the simpler non-y'

strengthened alloy Hastelloy-X, grain growth was not as limited resulting in uniform structures.

It is believed the stable MC carbide distribution in the Rene 80 is controlling grain size and

that the as-deposited structure of the deposit will determine this carbide distribution either

along particle, splat interfaces or interpass interfaces.

Tensile Tests

Tensile tests for both Rene 80 and Hastelloy-X sweet spot and large plate deposit produced

over a wide range of processing conditions are in progress. Results to date for tests conducted

at 910°C on small pin loaded flat tensile specimens, machined from these sweet spot deposits, are shown in Table 3. Prior to machining, the deposits were given a high temperature heat treatment

(1250°C - 2 hrs) to further density the deposits and increase grain size.

471
TABLE 3: 910“C TENSILE PROPERTIES SUMMARY SHEET SPECIMENS MACHINES FROM SWEET-SPOT DEPOSITS

Rene 80 Hastel loy-x

0.2% YS UTS Elongation 0.2% YS UTS Elongation

Process (MPa/ksi

)

(MPa/ksi

)

(%) (MPa/ksi

)

(MPa/ksi ) (%)

Vendor B Gun 268-436 492-568

Ar/Hz 38.8-63.3 71 .4-82.4 0.5 - 3.5

Vendor A Ar/He Gun Sweetspot

331-527 405-582 119-172 154-190 48.0-76.4 58.8-84.4 1.6 - 3.4 17.3-24.9 22.3-27.6 30.3 - 39.6

Vendor A Ar/Hz 324-428 337-428 159-179 168-188 47.1-56.6 49.0-62.2 0-1.0 23.1-26.0 24.4-28.5 11.0 - 50.2

Avg Conv. * processed

430.9

62 604.7

87 13 162.3 23.5 185.6 26.7 42.0

The 910°C test temperature was selected because it has been reportedfll that a significant

ductility loss was observed in this range for y' alloys in their study. This study then evaluated

the effect processing variables may have had on this ductility minimum.

As can be seen from Table 3 the low ductility of Rene 80 was not strongly affected by variations in processing. The yield strengths and ultimate tensile strengths observed in this evaluation for Rene 80 ranged from 267 MPa/38.8 ksi - 528 MPa/76.4 ksi and 338 MPa/49.0 ksi

581 MPa/84.4 ksi, respectively. The upper end of these test results compare favorably with those

attained for conventionally cast Rene 80. The variation in yield strength was found to correlate

with the as-deposited density, with the lower strength specimens corresponding to deposition

conditions which resulted in lower density deposits.

The Hastelloy-X tensile properties exhibited much less variation than was observed for the

Rene 80. This tensile property insensitivity to process variables for Hastelloy-X was also seen

in the deposited microstructures for this alloy. These property levels, including ductility, are comparable to those observed for wrought Hastelloy-X sheet.

The conclusions from these tensile tests is that large variations in deposition parameters can influence the tensile strength results for Rene 80. The same variations show relatively little

effect on Hastelloy-X, as was observed in both the as-deposited and heat treated microsturctures.

Summary

This paper presents a progress report on the continuing evaluation of the plasma spray process

for producing rapidly solidified alloys. The authors have made a sensitivity study of two superalloys, Rene 80 and Hastelloy-X. It has been found that the plasma jet enthalpy as function of

equipment or by selection of operating parameters most affects the deposited structure of these

materials. However, deposition parameter effects for Hastelloy-X are largely normalized by post

deposition heat treatment as reflected in negligible differences in microstructure and tensile

properties. For Rene 80, a more complex y' strengthened alloy, a more restrictive range of parameters are required to acheive similar results. Tests are continuing to further evaluate the

effect of some of the other features introduced by the processing variables such as substrate size

or gun/substrate motion. Other deposit parameters such as powder feed variables, and chamber

pressure, most affect deposit shape and efficiency with no discernable difference in "sweet spot"

structure. Finally, substrate temperature and plasma jet heat flux were found to control grain

growth and/or grain orientation while deposit structure (i.e. splat vs. unmelted) dictated carbide

or other stable phase alignment or location.

The advantage of using high enthalpy processing conditions to improve deposit density has been

demonstrated. It was shown that the control of powder injection is particularly important in

using the Vendor B Ar/H2 gun.

In order to increase the density of bodies formed using gun and/or substrate motion, improved

guns which minimize or eliminate the fringe region in the spray patterns will be required.

*Conventional ly cast Rene 80 or wrought sheet Hastelloy-X

472
Acknowledgements

The research reported herein was sponsored by DARPA Project 4289, Lt. Col. L. Jacobson and the

authors and Air Force Project Monitor Dr. H.C. Gegel. We would also like to thank the General

Electric Co. for their support and the many people there who have contributed to this effort.

References

[1] M.R. Jackson, J.R. Rairden, J.S. Smith, R.W. Smith, J. of Metals, Nov. 1981.

[2] J.R. Rairden, M.R. Jackson and M.F. Henry, To be published in the Proceedings of the 10th

International Conference on Thermal Spraying, Essen, FRG, May 1983.

WP1008/C326/2072A

473
b) high heat flux.

a) lower heat flux. Figure 2 Typical as-deposited microstructures

of RSPD Hasten oy Alloy X; (etched).

Figure 3 Mach 3 large traversing deposit microstructure of RSPD Rene 80. Figure 4 High enthalpy large traversing deposit microstructure of RSPD Rene 80.

474
b) well melted deposit.

Figure 5 Heat treat microstructure of RSPD

Rene 80 deposits.

475
THE MICROSTRUCTURE OF RAPIDLY SOLIDIFIED NiAl-Cr QUASIBINARY EUTECTIC

D. Shechtman*, T. Z. Kattamis^, F. S. Biancanieno^"*^ and W. J, Boettinger^^

*Center for Materials Research

The Johns Hopkins University

Baltimore, MD 21218

^Dept. of Metallurgy

University of Connecticut

Storrs , CT 06268

^^Metallurgy Division

National Bureau of Standards

Washington, DC 20234

ABSTRACT

The microstructure of melt-spun ribbon of the NiAl-Cr quasi binary eutectic composition

has been characterized by optical and transmission electron microscopies. The eutectic

composition is Ni-38. 5wt%Cr-19. 4wt%Al and is of interest because of the similarity of crystal

structures (CsCl for g-NiAl and BCC for a-Cr) and lattice parameters of the two phases in

the eutectic. The rapidly quenched microstructure consists of 0.5 pm diameter columnar

grains of the e-NiAl phase supersaturated with chromium to the eutectic composition. Between

these grains a fine rod- type eutectic structure of the B-NiAl and a-Cr phases is observed

with eutectic spacings as fine as 12 nm. The composition of the phases in the eutectic

portion of the microstructure were found to be close to the equilibrium solubilities for

these phases. A rationale for the appearance of a supersaturated B-NiAl phase at the

eutectic composition, rather than a supersaturated a-Cr phase, will be presented based on the T curves for this alloy system.

Introduction

The directional solidification of NiAl-Cr quasibinary eutectic alloy (Ni-33at%Al-34at%Cr) was previously investigated by Walter, et al. [1-4], who saw in it potential for high temperature

applications. The eutectic consists of fine, nonfaceted fibers of a-Cr BCC phase embedded in a 6- NiAl matrix, which is an ordered BCC phase (CsCl-type) containing some chromium in solid solution.

From previous work by Kornilov, et al. [5], the NiAl-Cr quasibinary eutectic temperature is 1445 °C. The a-Cr phase at 1200 °C contains 20.6 mole% NiAl and, hence, has a composition of 17.1at%Ni17.1at%Al-65.8at%Cr. The 6-NiAl at 1200 °C phase contains 84.5 mole% NiAl, hence has a composition

of 45.8at%Ni-45.8at%Al-8.4at%Cr. The volume fraction of fibers is about 0.34. In the directionally

solidified eutectic all directions and planes of the two phases are parallel. The solidification

microstructure consists of eutectic colonies of a-fibers of circular cross-section embedded in <100>-

oriented columnar grains of B-NiAl and parallel to the growth direction, except near the cell boundaries where they are inclined to this direction. Despite this inclination the crystallographic relationship between fibers and matrix remains the same [1]. Both cubic phases are crystallographically

related with only a small lattice mismatch, as evidenced by the large distance between dislocations

observed at the interphase interface.

Walter, et al. [2] studied the effect of growth velocity on inter-fiber spacing. They quantitatively established the decrease in cell size and inter-fiber spacing with increasing growth velocity.

They also investigated the effect of alloy additions on microstructural changes, in particular the

rod-to-plate transition [3].

It appeared interesting to extend the study of solification microstructures of the NiAl-Cr

quasibinary eutectic to those achieved by rapid solidification, anticipating that the small difference in lattice parameter between the two phases would lead to substantial extension in solid solubility. Melt-spinning was used in this investigation. The growth rates achieved in this process

are difficult to quantify, because of the possibility of bulk undercooling and an imprecise knowledge

of the interface heat transfer coefficient. However, melt-spinning provides a simple way to assess

476
the alloy's response to rapid solidification processing.

Results of the work on microstructure of melt-spun ribbon of NiAl-Cr quasibinary eutectic alloy

are reported herein and are also included in another publication [6].

Experimental Procedure

Rapidly solidified ribbons of NiAl-Cr quasibinary eutectic were melt-spun on a room temperature

wheel in a helium atmosphere. Ribbons of about 25 ym thick and 2.5 mm wide were obtained at a linear

velocity of 24 m/s. The ribbons were studied by optical microscopy, X-ray diffraction and analytical

electron microscopy. Specimens for electron microscopy were electropol ished and studied in a scanning transmission electron microscope equipped for nondispersive X-ray fluorescence analysis. The

composition of the phases^in the foil was evaluated by using the nominal composition of the specimen

as standard. The Cl iff-Lorimer coefficients were determined by using the intensities obtained from a thin area representing the general microstructure with a probe size of about 1 ym. Analysis of individual phases was then conducted in the same area of similar thickness. The C-L coefficients were redetermined for each set of analyses in different regions of the foil. This procedure eliminates

the need for additional corrections of the measured X-ray intensities.

Results

The microstructure of an arc-melted chill-cast cylindrical specimen, 6.3 mm in diameter is

illustrated in Figure 1. The rod-like a-Cr-rich phase is revealed as the dark etching phase.

Figure 1: As-cast microstructure of arc-melted Figure 2: Columnar grains in melt-spun ribbon,

NiAl-Cr quasibinary eutectic (optical micros- edge-on section (optical microscopy).

copy).

The cast microstructure of the ribbons consists mainly of a columnar zone with columnar grains

parallel to the heat flow direction. The intergranular spaces were not optically resolved. The av- erage columnar grain width increased slightly with distance from the chill. Figure 2, from about 0.5

to 1.3 ym. The average columnar grain width also increased slightly with ribbon thickness.

The microstructure of the ribbon as observed by electron microscopy is shown in Figure 3, which

represents cross-sections of columnar grains that grew perpendicular to the ribbon surface. The

columnar grains average 0.5 ym in diameter and are surrounded by a fine eutectic. The boundary between a columnar grain and a region of particularly fine eutectic is shown at a higher magnification

in Figure 4. The columnar grain is decomposed into a fine microstructure typical of conditional

spinodal decomposition [7] with an average wavelength of 10 nm (lower right. Figure 4). The eutectic

rods grow roughly perpendicular to the columnar grain and their spacing ranges from 12 nm. Figure 4, to 60 nm. Figure 3. The microstructure is not uniform and in some areas the volume fraction of the

columnar grains by far exceeds that of the eutectic. Figure 5. In other regions of the same foil the

volume fraction of the eutectic is much higher. Figure 6. These variations may correspond to different distances from the heat sink, however this was not confirmed in the present study.

A large proportion of columnar grains are single ordered domains. However, some of them contain

two or more domains. Anti-phase domain boundaries are shown in Figures 3 and 7. These APBS run across the grain and in most cases end in a chromium-rich rod of the eutectic. The significance of

these APBS to the understanding of the solidification process will be discussed later in this paper.

A diffraction pattern taken from a columnar grain. Figure 8, identified the structure as CsCl

(B2). A diffraction pattern taken from a columnar grain and its surrounding eutectic. Figure

9, shows no extra reflections besides the CsCl . This indicates that both phases in the eutectic and

477
Figure 3: Transverse section of columnar grains

and surrounding eutectic in melt-spun ribbon

(TEM). Note APBS in center grain.

Figure 5: Region of foil consisting primarily

of columnar grains (TEM).

Figure 4: Columnar grain-eutectic interface.

Inter-rod eutectic spacing: 12 nm (TEM).

Figure 6: Region of foil consisting primarily

of eutectic (TEM).

the columnar grain have the same orientation and almost identical lattice parameters. An X-ray diffraction scan taken from the ribbon shows that the reflections from the CsCl lattice and those from

the BCC lattice are indistinguishable. Lattice parameter measurements resulted in a figure of

a = 0.28882 nm. Microdiffraction patterns were taken from a rod and the adjacent matrix of the

eutectic, as shown in Figure 10. The rod is identified as BCC, Figure 10a, the matrix as CsCl

structure. Figure 10b. A composition analysis of the columnar grain cores has shown that they have the same composition

478
as the original alloy, Figures 11a and b. They are, therefore, 3 -NiAl supersaturated with chromium

to the eutectic composition. The BCC phase in the eutectic was also analyzed, Figure 11c and was found to be composed mainly of chromium of near equilibrium composition at the eutectic temperature

(13.9 at%Ni and 7.5 at%Al). The CsCl phase in the eutectic was not analyzed due to experimental difficulties, but based on an estimated volume fraction of the phases in the eutectic the phase is estimated to contain about 13 at% chromium, i.e. near the expected equilibrium composition for the g-NiAl

phase at the eutectic temperature.

Discussion

Two features of this rapidly solidified structure are unique and will be discussed separately;

The columnar grains of spinodally decomposing alloy with the B-NiAl structure and with composition

near the eutectic, and the fine rod-type eutectic structure. It is important for the analysis of

rapid solidification processes to establish whether the 3 -NiAl phase was a direct product of solidification or whether it has ordered in the solid state from a supersaturated a-Cr phase.

Columnar Grains

Because of the extension of solubility required for the formation of the columnar grains, an examination of the possible Tg curves for the NiAl-Cr system is in order. Figure 12 shows the stable

phase diagram as solid lines [5] and a possible schematic set of metastable extensions of the solvus

curves (dashed) and the Tg curves (long-short dashed). The figure assumes that the two solvus curves

meet at a metastable tricritical point above which the transition from the BCC phase to the CsCl

phase is second order. This transformation can be second order according to the symmetry restrictions of Landau [8]. This metastable extension upward to the tricritical point, where the crystal

structures differ by an ordering reaction, is formally similar to the metastable extension of the

solvus curves between identical crystal structures (FCC) in the Ag-Cu system upward to an ordinary

critical point. The proposed metastable tricritical point is thermodynamically identical to the

stable tricritical point in the Fe-Al system where the ordering is between Fe (BCC) and FeAl (B2)

[7]. The schematic free energy curves for the solid phases shown in Figure 12 are identical to those

proposed by Allen and Cahn [7]. With the addition of the liquid free energy curve to generate the

stable eutectic phase diagram the structure of the Tq curves becomes apparent. In this figure we

have assumed that the tricritical temperature is relatively low due to the similarity in the lattice

parameters of 3 -NiAl and a-Cr. The lattice parameters would be identical above the tricritical

point. In Figure 12 the dotted curve is the metastable extension of the second order phase transition below the tricritical point and is the locus of points where the free energy curves of the

disordered and ordered phases join tangentially. This curve is generally thought to lie much closer

to the solvus for the disordered phase rather than that for the ordered phase [7]. The Tq curve for

the disordered phase is seen to split into two branches when it crosses this dotted line. The lower

branch is the continuation for the disordered phase, while the upper branch is for 3-NiAl phase.

This upper branch continues to the NiAl side of diagram where it is the normal Tg curve for 3-NiAl.

The relatively small depression of the Tq curve for 3 -NiAl below the eutectic temperature is a

479
Figure 8: Selected area diffraction (SAD) Figure 9: Diffraction pattern taken from a pattern of a columnar grain. columnar grain and its surrounding eutectic.

Figure 10: (a) Microdiffraction from rod in eutectic.

(b) Microdiffraction of matrix in eutectic.

consequence of the relatively low tricritical temperature. This fact again is formally similar to the T„ curve for Ag-Cu.

This diagram leads to two important conclusions which are consistent with the experimental

results. First, the Tq curve for 6-NiAl phase at the eutectic composition occurs at a higher tem- perature than the Tq curve for the a-Cr phase. Second, if the 3-NiAl phase was formed with the

eutectic composition directly from the melt, it would undergo subsequent solid state spinodal decomposition into ordered and disordered phases. This type of spinodal has been termed a conditional

spinodal [7]. Typical decomposition structures of FeAl by conditional spinodal show coarse antiphase domains of the FeAl along with a fine scale compositional separation into disordered Fe and

ordered phase.

The structure of the columnar grains in the present work is identical to the above mentioned

structure. The grains contain coarse antiphase domains and fine spinodal structure. This strongly

suggests that the columnar grains solidified from the melt as supersaturated 3 -NiAl rather than

supersaturated a-Cr. No antiphase domain boundaries would be visible, if the supersaturated Cr phase

had solidified from the melt. The solidification of the ordered phase directly from the melt is also

suggested by the fact that at the eutectic composition Tq for the ordered phase is higher than for

the disordered phase. Hence, at the eutectic composition partitionless solidification of the 3 -NiAl

phase can occur with less undercooling than that required for the partitionless solidification of the

Cr phase.

Eutectic Structure

The solidification of the ribbon is envisioned in two steps. First, the formation of the

columnar grains of supersaturated 3-NiAl by partitionless solidification followed by the formation of

a fine rod eutectic structure in the regions between the columnar grains by near equilibrium solidification. This eutectic solidification would presumably result from the reduction in growth rate

caused from the evolution of latent heat during growth of columnar grains. The nucleation of the BCC

480
290 200

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II

II

II 1

1

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ir 1

1 —

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I » (a)

a- "ICHNI PHI 4 B LIST

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Figure 11: (a) General composition of the ribbon (used as standard).

Composition of a columnar grain (b) and of a Cr rod in the eutectic (c).

phase from the CsCl phase seems to be difficult and is demonstrated by the fact that in most cases

the eutectic originating from one columnar grain impinges directly upon an adjacent columnar grain.

Only- rarely do eutectics originating from adjacent grains meet. Once the eutectic has nucleated

it rapidly consumes the liquid that remains between the two columns.

The eutectic solidification apparently occurs with near equilibrium compositions for the two solid phases. Although a wide variety of eutectic spacings are seen ranging from 60 nm down to 12 nm the minimum spacings are of particular interest. As described by Boettinger [9] the temperature

dependence of the liquid diffusion coefficient leads to a predicted maximum solidification rate for

eutectics typically around 10 cm/s. In Ag-Cu eutectic alloys the maximum observed growth rate is

2.5 cm/s with a eutectic spacing of about 20 nm [10]. Hence, the present observation of 12 nm spacing is very interesting. Cline and Walter [2] have measured the rod spacing, X, as a function of

solidification velocity, v, up to 0.021 cm/s. They found these parameters to be related by

x2v=7xl0“^2cm3/s , deduced from their graph. The value of the X^v constant for NiAl-Cr is smaller

than for Ag-Cu (1.4xl0"l^ cm^/s) and for most other eutectics [11], presumably because of the low

interface energy between 3-NiAl and a-Cr phases, that can be attributed to the strong orientation

relation and similarity of lattice parameters of the two phases in the eutectic. An experimental

determination of the maximum growth rate of the NiAl-Cr eutectic structure will be performed in the

future, using electron beam melting experiments.

Acknowledgments

Thanks are due to Dr. Frank Lemkey for pointing out to WJB the similarity of phases in this

alloy system and to Dr. John W. Cahn for interpretation of ordered structures. The authors

acknowledge the support of DARPA.

References

[1] J. L. Walter, H. E. Cline and E. F. Koch, Trans. TMS-AIME, 245 (1969) 2073.

[2] J. L. Walter and H. E. Cline, Met. Trans., 1 (1970) 1221.

[3] H. E. Cline and J. L. Walter, Met. Trans., 1 (1970) 2907.

481
M0LE7oNiAI

Figure 12: (a) NiAl-Cr phase diagram (solid lines), metastable

extension of solvus curves (dashed) to tricritical point and Tq curves for a-Cr and 3-NiAl . (b) Schematic free energy vs composition curves for NiAl-Cr system.

[4] J. L. Walter and H. E. Cline, Met. Trans., 4 (1973) 33.

[5] I. I. Kornilov and R. C. Mintz, Dokl . Akad. Nauk USSR, 94 (1954) 1085.

[6] D. Shechtman, T. Z. Kattamis, F. S. Biancaniello and W. J. Boettinger, to be published,

J. Mater. Sci

[7] S. M. Allen and J. W. Cahn, Acta Met., 24 (1976) 425.

[8] L. D. Landau and E. M. Lifshitz, Statistical Physics , Pergamon Press, London, 1958, p. 434.

[9] W. J. Boettinger, Rapidly Solidified Amorphous and Crystalline Alloys , B. H. Kear,

B. C. Giessen and M. Cohen, ed., Elsevier Publishing Company, 1982, p. 15.

[10] W. J. Boettinger, R. Schaefer, F. Biancaniello and D. Shechtman, these Proceedings.

[11] W. Kurz and D. J. Fisher, International Metals Reviews, Nos. 5 and 6 (1979) 177.

482
WEAR PROPERTIES OF RAPIDLY SOLIDIFIED SURFACES

OF A PSEUDO-BINARY Fe-TiC ALLOY

P. R. Strutt*, B. G. Lewis*, S. F. Wayne*, S. L. Rice*, B. FI. Kear**

*University of Connecticut, Storrs, CT

**Exxon Research and Engineering Co., Linden, NJ

ABSTRACT

The wear behavior of the steel -base cemented carbide material Ferro-TiC (SK) has been

studied following electron beam glazing and tempering treatments. A distinction is made

between wear of SK against a moderately hard martensitic stainless steel (17-4PH), and a hard molybdenum-base high speed steel (M42). The effect of glazing is to produce a threefold increase in the wear life in tests performed using a fully hardened M42 high speed

steel counterface material.

FIG. 1. Microstructure of Fe-TiC Alloy (a) As Sintered, (b) After

e.b. -Glazing, Showing Refined Dendritic Structure

Introduction

When the cemented carbide Ferro-TiC, otherwise known as SK, is subjected to a rapid solidification

electron-beam surface melting (glazing) treatment, the original blocky TiC carbides are transformed

into a highly refined dendritic structure [1,2], H'gure 1. This feature coupled with a high hardness

suggests a possible benefit in terms of improved wear properties. The present study was undertaken to test this proposition. As will be shown, under appropriate conditions, significant improvements in wear

resistance due to e.b. -glazing can be achieved.

Specimen Preparation

Wear-pin specimens of SK material were glazed using a Carl Zeiss electron beam unit. A beam power

density of 10^ watts/cm^ was achieved by setting the power level at 300 watts and by adjusting the

focusing to give a measured spot diameter of 0.3 mm. Under these conditions, and with a beam scanning

rate of 100 cm/sec, a narrow glazed band about 100 ym in depth was obtained, which experienced an average cooling rate of 5 x 10^°C/sec. Continuous surface coverage of glazed material was attained by

483
applying two nearly equal frequency signals to the X and Y deflection coils, so as to trace out a series of overlapping Lissajous figures [3], Figure 2. This particular scanning mode was originally

developed for the purpose of solid state transformation hardening of steels, without actual melting [4]

All the material used, whether it was subsequently tested in the unglazed or glazed conditions

was initially austenitized at 1050°C for 5 min., and then quenched into oil at 50°C. Following this

treatment, a few samples were doubled tempered at 450°C for 1 hr.

Table 1. Major Alloying Elements (Wt.%)

in M42, 17-4PH, and SK (Ferro-TiC)

Fe Cr W V Mo Co Ti Mn Si C Ni Cu P Nb/Ta

M42 Bal 3.70 1.36 1.12 9.62 7.88 0.24 0.40 1.09

17-4PH Bal 16.5 - - - - - 0.06 1.0 0.05 4.0 4.0 0.04 0.35

SK Bal 3.5 - - 3.0 3.0 21.0 - - 5.0 0.5 - - -

All other elements <0.01 wt.%

Wear Testing

Wear tests were carried out in a standardized test machine [5], in which the end of a small cylindrical SK wear pin (3mm dia.) both impacted and slid across a steel counterface. Tests were performed

using two different counterface materials - a moderately hard (400 DPN) martensitic stainless steel

(17-4PH), and a hard (930 DPN) molybdenum-base high speed steel (M42)., Table 1. In all tests, the

transverse sliding velocity was 8 m. sec“l and the nominal maximum impact stress was 69 MPa at a repetition rate of 800 cycles/min. To determine the extent of wear, pins and counterface samples were weighed before and after testing. Wear pins of SK material in both the unglazed (500 DPN) and glazed

(950-1000 DPN) conditions were tested against the two different counterface materials at ambient temperature.

Wear Against a Moderately Hard Steel

Figure 3 shows the general appearance of the microstructure of a glazed SK wear pin after testing.

The surface is covered with a thin layer of counterface material (Zone A), which has been transferred . onto the pin during testing. Below this layer is a much thicker dendritic carbide region (Zone B), i which represents the full extent of the homogeneous glazed material. Zone C is a' transition zone of

partial glazing, involving melting and re-solidification of the steel matrix only i.e. without complete

dissolution of the original blocky carbides. Zone D is representative of the unmelted parent material.'

Close inspection of Zone B reveals that it is separable into an upper heat-affected (tempered) deeply

etched region 3-4 ym below the surface, and a lower more evenly etched dendritic carbide structure.

These microstructural features were rather typical of all glazed and tested wear pins.

Figure 4 compares the microstructures of unglazed and glazed SK wear pins after testing against 1

17-4PH. The surface of the unglazed sample is rough. Figure 4(a), and there is evidence for patchy |

regions of transferred counterface material. Below this there is a region of intense plastic deforma-

^ tion in the sliding direction, which contains small fragmented carbides. The actual amount of plastic _ strain is estimated from measurements on several micrographs to be >50 pet. In complete contrast, in J

the glazed sample. Figure 4(b), there is a much greater amount of visible transferred material, but no a

evidence for significant plastic deformation in the upper part of Zone B. In both cases, there is I

evidence for cracking in the surface layers, arrows in Figure 4. These are limited to the transferred I

material in the unglazed sample, whereas in the glazed sample they extend into Zone B, terminating nearm

its base. Post-glaze tempering treatments to increase martensite ductility did not reduce the propen- M

sity for crack formation in the glazed material. I

No significant weight changes were recorded for glazed and unglazed pins. On the other hand, "

metal lographic examination provided indirect evidence for more wear of unglazed pins than glazed pins. ' This is deduced from Figure 4(a), where carbide particles have clearly worked their way into the I

adhered material (Zone A). In contrast, the corresponding counterface materials (impact discs) under- 1

went substantial wear, as determined by actual weight losses. Thus, the overall situation can be

described as essentially 'one-body wear'. Figure 5 shows plots of the actual experimental weight change

data. For completeness, the effect of tempering the glazed pins is included. Tempering reduces the wear rate of the counterface material, which may be the result of more material adhering to the ; tempered pin.

484
ill ELECTRON

ill BEAM

FIG. 2. Schematic of Electron Beam

Glazing, Utilizing Electromagnetic

Beam Deflection

FIG. 3. Electron Beam Glazed Ferro-TiC(SK) Wear

Pin Sectioned Parallel to the Sliding Direction,

Showing Adhered Counterface Material (Zone A),

and Sub-Surface Zones of Fully Glazed (Zone B),

Partially Glazed (Zone C), and

Unmelted Substrate (Zone D)

0. 02 -

E

o»

0.01

\9>

/ /

/ / if Specimen Counterface

‘fo Ferrotic SK, glazed 17 4PH —

• Ferrotic SK, glazed, 17 4PH —

double tempered

^no detectable weight loss for glazed pins, 25 50 75 100

Repetitive Impact Cycles x 10'

FIG. 4. Comparison of the Surface Micro- FIG. 5. Weight Loss Data for Glazed

structures of Unglazed (a) and Glazed (b) Ferro-TiC (SK) Worn Again 17-4PH

Ferro-TiC (SK) Wear Pins, After Wear

Against a Moderately Hard 17-4PH Steel

Counterface

485
Wear Against a Hard Steel

Figure 6 compares the microstructures of unglazed and glazed SK wear pins, after testing against

M42 high speed steel. The unglazed sample, Figure 6(a), exhibits some of the features characteristic

of wear against the softer 17-4PH material. Figure 4(a), such as surface roughness, and near-surface

plastic deformation and carbide fragmentation, but there is no indication of pick-up of counterface

material. On the other hand, the glazed sample exhibits quite distinctive features, not previously

seen in the 17-4PH case. As shown in Figure 6(b), the wear surface is smooth, contains no detectable

counterface material, and there is no evidence for cracking of the glazed material. Careful examination reveals a small amount of carbide fragmentation, limited to the top 2-3 ym of the glazed region.

Figure 7(a) shows that wear of unglazed SK pins against M42 produces significant weight losses

for both pin and counterface materials, i.e. two-body wear, in contrast to the one-body wear characteristic of 17-4PH. Both tempered and untempered samples wear at a similar rate up to 50,000 cycles,

beyond this there is a continued weight loss only for the untempered material. A possible explanation

for this behavior is that a greater amount of M42 counterface material adheres to the tempered wear

pin, which gives rise to wear of M42 against M42 in the later stages of testing. A similar two-body wear is found also in the case of the glazed material. Figure 7(b), except that to obtain a weight

loss equivalent to that of unglazed material. Figure 7(a), it is necessary to increase the test time

by a factor of 2.5 (i.e. to 250,000 cycles). In other words, e.b. -glazing in this case increases wear

life by 2.5x.

Pi scussion

i The present work shows that electron-beam glazing is capable of improving the wear resistance of a

; hard cemented carbide material. Although the underlying mechanisms are imperfectly understood, it is ! believed that the most important factors contributing to improved performance are (i) the high degree

of homogeneity of the rapidly solidified material, and (ii) the high yield strength of the ultrarefined microstructure.

Before considering wear of glazed material, it is instructive to consider the degradation process

of untreated material. The microstructure of unglazed SK material. Figure 4(a), after wear against a 17-4PH steel counterface, is indicative of a situation where fragmented carbide particles tend to aligns , themselves along curved flow lines within a plastic deformation slip field produced by shear along the wear direction. X-ray diffraction analysis of wear debris confirms that small TiC particles are indeed I carried from the SK material into the adhered layer.

'

Observations during actual wear tests indicate that an adhered layer of 17-4PH forms at an early

stage on the wear-pin surface. It is the continual breakaway and renewed build-up of this layer that . produces the wear debris. The 17-4PH readily shears from the counterface due to the plastic softening i resulting from localized heating during sliding and impact. When, however, a sufficiently thick layer ( has been transferred to the wear-pin the adhered zone exhibits a hardness of about 800 DPN. The

increase in hardness arises from a combination of precipitation hardening, mechanical mixing of alloy t phases and oxides, and the presence of untempered martensite formed as a result of wear test induced ,1 heat treatments. In addition to being hard, mi crostructural observation of severe cracking shows Zone II A structures to be rather brittle. Interestingly, the cracks in the adhered 17-4PH layer do not propagate into the cemented carbide material. Presumably this is because the matrix phase readily

accommodates the intensified stress field at the crack tip by plastic deformation. Although there are

no detectable cracks within the near-surface region of the SK material, a small amount of SK neverthe- i less is removed. Presumably, this arises from the small amount of near surface material, which is ; firmly bonded to the fragments of adhered 17-4PH, detaching itself during the wear process.

From the experimental observations, and the preceeding considerations, it is evident that glazing

is advantageous in improving wear resistance provided that an appreciably thick, brittle and tenaciousl

adherent layer does not form. Such a conclusion is well supported by the v/ear experiments using a hardened M42 steel counterface. Little or no detectable adhered layer is formed, presumably because

the good hot hardness prevents plastic shearing. In this case, the material degradation process appear^' to be surface abrasion involving the continued removal of fine particles. The creation of such

particles must involve surface microcracking, but such cracks are highly localized since they do not penetrate the glazed surface layer. This type of wear process occurs on both the SK wear pin and the " M42 counterface. In the glazed condition the fine particle fragmentation is reduced, thus sfiowing the

beneficial effect of electron beam glazing in reducing wear.

486
FIG. 6. Comparison of Surface Microstructures of Unglazed (a) and Glazed (b)

Ferro-TiC (SK) Worn Against a Hard M42 High Speed Steel Counterface

Repetitive Impact Cycles x 10*

(a) FIG. 7. Weight Loss Data for Unglazed (a)

and Glazed (b) Ferro-TiC (SK) Worn Against M42

(b)

References

1. Lewis, B. G., Gilbert, D. A. and Strutt, P. R. (1980). Laser and Electron Beam Processing of

Materials (eds. C. W. White and P. S. Peercy) , Academic Press, 747.

2. Strutt, P. R. (1977). Metals Forum, Pergamon Press, 117.

3. Strutt, P. R. (1981). Mat. Sci . & Eng

, 49 , 87.

4. Mayer, R., Dietrich, W. and Sundermeyer, D. (1977). Weld. J. , 56 , 35.

5. Rice, S. L. (1979). Wear , M, 291.

487
FABRICATION OF ORIENTED SINGLE CRYSTAL WAFER STOCK

FROM Ni-Al-Mo-X ALLOY POWDERS

R. E. Anderson, H. A. Chin and S. E. Hughes

PRATT AND WHITNEY AIRCRAFT/GPD

WEST PALM BEACH, FLA 33402

B. B. Rath

NAVAL RESEARCH LABORATORY, WASHINGTON, DC

ABSTRACT

Utilizing a rotary atomization device to produce prealloyed, rapidly-cooled, metal powder,

P&WA has developed alloys which exhibit uniaxial creep strength significantly greater

than that of the best available directionally solidified nickel-base superalloy. However,

the grain boundary strength of this material is poor necessitating a single crystalline

approach to allow application of this class of materials in many advanced gas turbine airfoils. An extensive effort has been undertaken to optimize alloy composition and processing parameters for producing bulk quantities of seed oriented single crystalline sheet

stock of Ni-Al-Mo-X compositions (X representing W, Ta, Cr, Hf, Y, C) for use in subsequent

turbine blade application. The properties of the alloy as it is currently configured will

be reviewed. Further, a description of the processing sequence as it is currently arranged,

and the effects of the various processing steps on the microstructure and texture of the

material will be presented. Furthermore, the influence of the varying microstructural

parameters on the relative perfection and growth selectivity of the crystals grown in the

solid state are discussed.

INTRODUCTION

The efficiency of a gas turbine engine is directly tied to the maximum gas stream temperature. The

need for further increases in turbine inlet temperature has not abated. Before the findings of re- cent programs at P&WA, the point of diminishing returns was being reached in terms of conventional

airfoil materials and cooling technology. The conclusion was reached that a radical departure from

established turbine airfoil materials and design approaches would have to be taken if further gains

in airfoil operating temperature capability and hence, engine performance were to be attained. To meet this challenge, P&WA in conjunction with the Defense Advanced Research Projects Agency (DARPA)

and the Air Force undertook a program aimed at applying the principle of rapid solidification to nickel base superalloys in general and to turbine airfoils in particular. This program, reported on previously ^ »2,3 , aimed at developing an alloy with creep and rupture capability 100°F higher than

the best available production airfoil material of that time. This effort was so successful that it was mated with the radial wafer cooling concept in an attempt to produce an airfoil with the potential for withstanding a 25 QOf increase in turbine inlet temperature in a device like the P&WA FIDO

turbine engine.

The fabrication effort aimed at producing prototype airfoils was also discussed previously. ^ This

effort led to a brief, but successful, engine test two years ago. Since that time, P&WA, the Air

Force, and DARPA have been aggressively pursuing a follow on program aimed at (1) producing these

types of airfoils at a reasonable cost, (2) further improving the durability of these components

through design and process modifications, and (3) substantially improving the material environmental

resistance and thermal stability through compositional modifications. A significant portion of this

effort is associated with the adoption of a wrought, directionally'recrystallized (D.R.) single crystalline component approach to enhance alloy and blade durability and properties. The technique for

the fabrication and the properties of these single crystalline assemblies represent the topic of

this discussion.

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BACKGROUND

GENERAL

The potential benefit of an airfoil fabricated from radially oriented wafers made from a very high

strength superalloy has been described previously. 4,5,6 ^35 ^gen speculated that such a component could lead to a 40% reduction in specific fuel consumption or to a 15 to 20% increase in thrustto-weight ratio in an engine like the P&WA FI 00 turbofan. However, these airfoils (a total of 20 were made and two were engine tested) contained several non-optimal features. First, the cooling

scheme lead to thermal inhomogeneities. Second, the alloy exhibited only marginal environmental re- sistance as well as thermal instabilities which made its strength time-dependent. Third, the traverse ductility of that material /component was marginal. Fourth, the cost of the assembly was extremely high. In short, the concept needed a significant amount of optimization to reach the level

of excellence needed to be considered as a prime candidate for the airfoil system in production turbine engines.

One of the most serious problems to be solved was the inherently poor transverse strength of the

grain boundaries in directionally recrystal 1 ized alloy 185 (Ni-14.5Mo-6.8Al .04C-6.1W) . This strength

in this direction was so low that it eliminated the material from consideration for use in some of

the more advanced applications. One solution to this problem was simple; eliminate all grain boundaries.

WHY RAPID SOLIDIFICATION?

To maximize the high temperature capability of nickel base superalloys, the matrix phase is commonly

strengthened with large additions of compatible refractory elements such as tantalum and tungsten.

The theoretical limits of solubility for the various additions is set by the phase equilibria, beyond which deleterious phases such as o and jU are rejected. In practice, the solubility limits of

the alloying elements are significantly lower than the theoretical limits because of the sluggish

kinetics of conventional solidification processing.

At the time this effort was initiated, the most advanced production airfoil material was directionally solidified alloy MAR-M200+ Hf, (nominal composition Ni-9CR-10Co-2Ti-5Al-12.5W-2Hf-lCb-0.14C0.015B). Several flight airfoils of this material were found to exhibit Laves - phase precipitation

but only in the dendrite cores of the castings. Microprobe analyses of these cored regions indicated

tungsten concentration in excess of 40% even though the nominal composition of the airfoils was within specification.

The implications of these findings were twofold. An inordinate amount of tungsten was being concentrated in the dendrite cores by the relatively long range segregation associated with the slow directional soldification process used in casting complex turbine airfoils. The suppression of this

precipitation could only be accomplished by suppressing solidification induced segregation. Now,

since the refractory alloying elements were being rejected from the 7 phase at the current level of

alloying, further alloying additions aimed at additional strengthening would be totally frustrated by

the rejection of these additions culminating in further deleterious third phase precipitation.

The simplistic idea of adding refractory elements to superalloys until saturation is reached to strengthen them is an appealing one because it is a relatively straight-forward approach. It is

easy to speculate on what the creep strength of MAR-M200+Hf would be if the composition of the dendrite

cores was the uniform composition of the alloy. Obviously, some of the effectiveness of this

approach would be lost on a self-loaded high speed rotation part like a turbine airfoil because the

density (and hence, the airfoil stress) would increase as the alloy's nominal refractory element content increased. If most of the tungsten is in the dendrite cores, then little is in the interdendr'

tic regions leaving these areas relatively weak in creep and stress rupture loading, where elements

like W, Ta, and Mo etc., show their biggest beneficial effect on superalloys.

There are two obvious methods for reducing constitutional segregation: reducing solidification induced segregation by increasing the solidification rate, and homogenizing the structure through the

use of long-time, high-temperature solutioning cycles. The effectiveness of the solutioning cycle is an obvious function of the maximum temperature as well as the time. The maximum temperature is, however, also a function of the solidification rate because it is limited by the incipient melting

point, a temperature which is developed in the structure by the solidification induced segregation

of low melting point alloying phases and elements to the grain boundaries and interdendritic regions.

With MAR-M200 and normal directional solidification techniques, this temperature (circa 2175°F) was

so lov: that full homogenization without local melting would require solutioning for inordinately long

times (weeks or months).

489
It is possible to produce simple geometry cast test bars with a fairly uniform composition by increasing the directional solidification rate (typically mold withdrawal rate). This approach would

allow both shorter homogenization times to complete the task (because of the higher homogenization

temperatures usable without melting) and/or the use of higher effective strengthener contents. How- ever, in complex airfoil castings where solidification rates must be reduced (withdrawal rate re- duced) to improve mold fill and grain structure, this approach too is frustrated. Consequently,

while it may be possible to improve uniformity and, hence, strength in test bars ( and perhaps in

simple airfoil castings) these gains are not translatable to high performance geometries with more

highly alloyed compositions using production directional solidification techniques.

Recognizing that uniformity in alloying element distribution can enhance strength In the metals

needed to make the turbine's most performance limiting component, it became necessary to develop a technique to use this approach to allow the practical fabrication of hardware. P&WA chose the most

rapid solidification technique which was available at the time for producing bulk quantities of homogeneous material, the powder metallurgy approach. Theoretically, the interdendritic distances in a powder particle are so short that even if dendritic solidification were the primary solidification

mode atomistic uniformity in composition could be reached by high temperature homogenization in

a reasonable engineering time. It was speculated that if MAR-M200 could be made uniform by the

powder metallurgy technique then sufficient tungsten could be added to drive the temperature capability up by well over an effective 100°F without either exceeding the solubility of tungsten or raising the density beyond reasonable limits.

Using commercially available powder from Special Metals Corporation, this theory was tested. Alloy

MAR-M200 powder was procured, consolidated by extrusion, forged to a pancake shape, DR'ed, and tested

in creep and rupture. Although the alloy was found to D.R. to a <11 0> orientation, a relatively

weak orientation for superalloys, this material was found to be clearly superior to the directionally solidified alloy in 1% creep strength and roughly equivalent to the cast material in stress

rupture resistance. Furthermore, the powder metallurgy approach appeared a viable technique for producing strong airfoil materials. What was needed to fuel the development process was a source of

large quantities of segregation free, or at least segregation suppressed, powdered material on which to do alloy content and optimization studies.

THE POWDER RIG

To supply large quantities of rapidly solidified powder, P&WA built the rotary atomization device

discussed previously. ^ Rotary atomization was selected because this method seemed to offer the best

probability of producing a controlled particle size distribution. Helium gas quenching was selected

for simplicity of handling, high heat extraction capability, and inertness. Conventional metal

melting and powder handling equipment was used whenever possible to minimize complexity.

The powder making device now has the capability of pouring 300 or more pounds of powder per atomization run. A recirculator has been added to reduce the helium consumption, and the turbine now has

the capability of turning at 35,000 rpm (thus producing a finer, more uniform droplet distribution).

This device has been used to produce in excess of 20 tons total of rapidly solidified metal powder

in assorted nickel, aluminum, copper, and iron based alloy systems.

ALLOY MODIFICATION

The alloy developed in the previous program, directionally recrystallized alloy 185 (Ni-6.8A1-

14. 4Mo-6. OW-0. 02C) , although inordinately strong in high temperature creep and rupture loading, ex- hibited several deficiencies. The columnar grained structure of the DR alloy possessed inadequate

rupture strength across high energy boundaries. The base composition is supersaturated in Mo and W and rejects the excess as CL-Mo, W. The oxidation resistance of alloy 185, particularly at high

temperatures, leaves much to be desired. Finally, alloy 185 in its most creep resistant orientation

( <111> axial) exhibits poor strain controlled and thermal fatigue resistance. This is a fundamental characteristic of the relatively high elastic modulus of nickel based alloys in this orientation ( =44 X 10^ psi at R.T. for the <111> as opposed to = 18 X 10^ psi for the <100> orientation).

A major alloy and process development effort has been conducted during the last two years in an attempt to solve these problems. The phase instability problem was essentially eliminated by re- ducing the Mo level from 14.5 weight percent to 10.0 or less. The high temperature oxidation

problem was resolved through the addition of minor alloying elements (most importantly Yttrium). The

problem of fatigue resistance and grain boundary strength has been resolved through an oriented

single crystalline approach. This effort has led to the development of PWA 1121 (single-crystalline,

directionally recrystallized Ni-6.8Al-3.0Cr-10.0Mo-6.0W-l .5Ta-0.15Hf-0.02C-0.1Y).

490
Oriented single crystals can improve fatigue life because of the inherent anisotropsy of the nickel

based superalloy system. It was discovered early in the effort that the <11 2> crystal lographic

direction in the Ni-Mo-Al system exhibits nearly as much creep and rupture strength, and substantially more rupture ductility, than the <111> orientation. Furthermore, this orientation has a characteristic elastic modulus which represents a compromise between the two prime orientations

(E <;ii2> — 34 X 10^ psi at RT.) Consequently, the strain controlled low cycle fatigue resistance

of single crystals of this material, oriented in the <112> direction, also should exhibit a compromise. Test results have recently substantiated this premise. However, <112> D.R. crystal

growth is not readily producible from material consolidated by high temperature extrusion, the technique used previously for consol idating Ni-Mo-Al-X powder. Apparently, D.R. orientation is a complex function of the input material deformation and microstructure texture as described elsewhere

in this conference. ^ Consequently, the technique most appropriate for producing a texture with a <11 2> direction in a useful dimension of the material appears to be a combination of hot and cold

rolling.

THE COMPONENT AND THE DEFORMATION SEQUENCE

Early experiments indicate^ that consolidation of Ni-Mo-Al powder by extrusion at high temperatures

(circa 2300°F) produces a <111> fiber texture. If this material is subsequently rolled to sheet

stock by cross rolling techniques, a |l 1 Cj <1 1 2> sheet type texture Is produced. This texture is

compatible with <112> D.R. growth parallel to the primary working direction. Obviously, the

product of this process is sheet stock. This is fortuitous because the component being used as a demonstration vehicle for this alloy is a radial wafer concept turbine blade which uses sheet stock

(wafers) exclusively for its input material as shown in Figure 1. The design of this component has

been upgraded to improve the cooling effectiveness and to make it compatible with the projected

properties of the material.

STATEMENT OF THE PROBLEM

Having solved most of the problems associated with the use of this material from a theoretical standpoint, all that remained was to produce single crystalline stock from prealloyed rapidly solidified

powder of the appropriate composition. Described below are the steps taken in consolidating the

material, shaping it, D.R.ing the subcomponents, and fabricating the final assemblies. Many of the

techniques are similar to those described previously,* while others have had to be changed significantly. What is discussed are the critical steps in the sequence and some of the rationale for their

choice.

FABRICATION SEQUENCE

The remelt stock for atomization was produced at PWA/GPD by conventional vacuum melting techniques.

However, since the atomization device has grown in capacity to a current level of 300 lbs, the 100 lb

capacity laboratory melting facility has been outstripped. Consequently, commercial vendors such as Special Metals are now producing the remelt stock for this effort. The material under consideration

here came from a 4,000 lb Special Metals master heat of the composition listed previously as PWA

FIGURE 1 FINAL BONDED ASSEMBLY OF SINGLE CRYSTAL RADIAL WAFER BLADE

1121

491
FKjUAE 2 CONSOClOATtON TEMPERATURE EFFECTS ON THE AS-HIP MtCROSTRUCTURE OF PWA 1120 BILLETS

2200*F

HIP IS ktl/2 hr 2250*F 23S0*F

L I

This material was atomized in the GPD powder device in heats ranging in sizes from 250 lb. to over

300 lbs. The heats were cross blended in an attempt to normalize any unforeseen anomalies such as undetected leaks in the atomization system, spot cleanliness problems, etc. Each step of the atomization sequence, sizing, storage and canning was all accomplished under either vacuum or positive

helium pressure in an attempt to all -inertly handle the powder. Immediately prior to canning, the

powder was passed through the GPD hot dynamic outgassing system into Inconel 600. Typical pressures

during canning and sealing were in the 10"5 torr range at 1000°F. The cans for this effort ranged ini

capacity from 145 to over 300 lbs. Inconel 600 is used as can stock because it also doubles as a hot, , rolling can. '

\ Consolidation of the powder was accomplished at IMT in Boston, Mass. The powder was consolidated to full density at all temperatures studied above 2100OF. Consolidation at the temperatures studied ledto a 50% or greater level of recrystallization as shown in Figure 2. Subsequent formability was found to increase with percent recrystal 1 ization so that HIP temperatures were kept intentionally

high. However, hypersolvus (T ^jp) > T-y' solvus — 2325°F) cycles were found to lead to reduced i formability resulting from excessive grain growth. Consequently, 2250° was used for the bulk of the j material processed in this effort. i

Hot rolling was accomplished at Cabot, Cyclops, AMAX, and GPD. Rolling schedules were designed to yield a final decanned product of approximately 0.400 in. thick using a reduction of 5% initially I increasing to 20% per pass with reheats at 2250°F on every pass. A typical hot rolled and decanned

j billet is shown in Figure 3. Losses due to edge cracking, tearing, etc. were minimal with the total | yield through hot rolling approaching 85%. The hot rolled plates were subjected to a normalization i heat treatment comprised of a two hour isothermal anneal at 2275°F followed by controlled cooling to | fully recrystallize the structure and normalize the 7 (second phase) morphology. ^

FIGURE 3 HOT ROLLED AND DECANNED PWA 1120 BILLET ILLUSTRATIVE MACRO- AND MICROSTRUCTURES

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j: Following hot rolling, normalization, and decanning (either chemically or mechanically) the billets

j| were cold rolled to final thickness. The schedules were adjusted to yield a final thickness of

nominally 0.145 in. using a 65% total reduction, the final 25% of which was accomplished inthe cross ' rolling direction. Reductions were taken at 1% per pass with 2200°F/10 min/AC anneals applied after

^ each 5% initial to 15% final reduction. Typical cold rolled products from these operations are shown I; in Figure 4. The yields through cold rolling, after trimming, approximated 80% of theoretical. <

492
FIGURE 4 TYPICAL COLO ROLLED AND OECANNED MACRO- AND MICROSTRUCTURES OF PWA 1120 SHEET

Following rolling and trimming to size, the wafers were subjected to surface preparation for diffuj sion bonding. This operation was generally accomplished by mechanical lapping. Sufficient material

I was removed to clean the surface and size the wafer to the proper thickness. In instances where

I large quantities of material had to be removed (for example when an 0.080" final thickness was dej sired) the wafers were double disk ground to rough size, and lapped to finish dimensions.

I Since a multiple piece assembly was required as the final product, two processing sequences were

I evaluated in component fabrication. The first technique involved the assembly of pregrown single

I crystalline stock. Alternatively, it was shown feasible to diffusion bond fine grained stock and

I then, to directionally recrystallize the monolithic assembly. Both techniques were demonstrated in

this effort, and since both have unique advantages (and disadvantages), both will be outlined here.

FIGURE 5 WAFER AND TOOLING DESIGNED FOR ORIENTATION CONTROLLING SEED BONDING PRIOR TO DIRECTIONAL RECRYSTALLIZATION

w»m and S«ed TZM Den* <t Tooling

To produce a pregrown assembly, the fine grained and lapped wafers were trimmed to the shape shown in Figure 5 to facilitate the attachment of an orientation controlling seed crystal. The joining operation was accomplished using differential expansion tooling in vacuum during heating to 2050°F. The

lapped surfaces were carefully cleaned, electropolished, and loaded into the TZM tooling and placed

The assembly was then placed into a vacuum furnace, and the chamber was evacuated and heated to the appropriate bonding temperature. Typically, three (3) samples were bonded in each

die stack per load. Bonding pressures varied , and successful bonds were produced with pressures as

igh as 10 torr. Bonding yields were very high with 90% of the seeded stock exhibiting grain

transferral across the lap bond joint.

After bonding, the preseeded fine grained samples were trimmed to remove any potential grain nucleating sharp edges and submitted for D.R. The D.R. heat treatments were conducted in the GPD vacuum

induction gradient furnace. The samples were loaded into the furnace either one, two, or three at a time and grown to a point where sufficient material remained unheat affected to generate seed stock.

Vacuum levels were maintained sufficiently low to minimize graphite hot zone and sample oxidation but

high enoughto minimize sample chromium and aluminum depletion. Traverse rates ranged from 1/8 to 3/4 in/hr with gradients measured at the 7 'solvus temperature approximating 25QOF/in. The yield of

single crystals from this approach was found to be highly orientation dependent but was above 50% on

an overall basis.

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FIGURE 6 COOUNG PATTERNS FOR TWO TYPES OF RADIAL WAFER DESIGNS

The D.R.'ed sinyle crystals were finish machined and transferred to assembly bonding. To facilitate

ready bond closure around the somewhat intricate internal airfoil passages of the type shown in

Figure 6, a compliant layer of pure nickel was applied to each mating face. This layer (typically

0.0005 thick) was applied by either sputtering (the preferred method) or plating. Bonding was accomplished in a vacuum hot press at pressures above 15 Ksi. Upsets were held below the 1% level to guarantee bond closure without inducing sufficient strain to recrystallize the assembly upon subsequent full solution heat treatment.

FIGURE 7 MICROSTRUCTURE OF BOND REGION OF PREGROWN S'NGLE CRYSTAL ASSEMBLIES

Ni Interlayer 2S«m

(0 0002 in/SiOe)

The final assembly (previously shown in Figure 1) contained typical bond interfaces illustrated

metal lographical ly in Figure 7. These assembl ies were to be final heat treated prior to delivery for

final machining by a variety of conventional and non-conventional machining techniques.

FIGURE 6 MACRO- AND MICROSTRUCTURE OF A BONDED ASSEMBLED WAFER STACK

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To produce a prebonded block, the processing sequence was reversed The

^ef^lhen

ilsion machined to the final thickness plus w1tf>out

machined into the wafers which were then electropolished and delivered directly y

the application of an inberlayer.

494
Bonding was accomplished in a vacuum hot press at 2050°F. Stop blocks were used to control deformation, and since the material is inherently highly superplastic, applied stresses were typically in

the i to 3 Ksi range. An assembled prebonded block is shown in Figure 8.

These blocks also need a seed to control D.R. grain growth direction. This seed was applied by

differential expansion diffusion bonding in a manner similar to that described previously. The full

width seeds were generated from the end croppings of the pregrown wafers described above. The

doubly bonded blocks were then trimmed to regular cross-sectional geometries and D.R.'ed at conditions

similar to that described for the preqrown wafers. A sample from this technique is illustrated in

Figure 9. For blade fabrication, a series of these blocklets (four in all) will subsequently be

bonded together to produce the full airfoil cross section (shown schematically in Figure 10) and

machined to final shape.

FIGURE 9 MACRO- AND MiCROSTflUCTURE OF DR HEAT TREATED PREBONOEO BLOCK

FIGURE 10 BLOCKLETS TO BE BONDED TO PRODUCE THE FULL AIRFOIL CROSS SECTJOH

DISCUSSION

The complexity of the radial wafer blade brings the engineer immediately to the problem of process

economics. This is doubly true if the required component must be single crystalline in nature. The

early radial wafer blades were made by slicing oversize extrusions, machining the radial cooling

passages into the slices, diffusion bonding the assembly, and D.R.'ing the block; This sequence is illustrated in Figure 11. This approach was simple from a technical standpoint but led to an apparently

poor material utilization factor. Furthermore, when the base alloy was found to be grain boundary

weak, some technique for minimizing or eliminating grain boundaries or grain boundary angles was re- quired. The process economics are also strongly effected by yield factors which, for this application, are largely controlled by heat treatment (D.R.) response.

FIGURE 11 MANUFACTURING SEQUENCE FOR RSR RADIAL WAFER BLADE

FD 254161

495
It was discovered early in the program that strong textures greatly enhance D.R. growth response.

Early data also indicated that deformation textures controlled D.R. orientation. Furthermore, since

the properties of the alloy under, consideration are anisotropic, it appeared feasible to control the

final properties of the fabricated assembly by controlling the processing/deformation sequence.

Finally, if D.R. orientations were controlled by deformation texture, a singular texture would, of

necessity, produce an effectively single crystalline structure.

A comparison of the properties of alloy X as a function of orientation indicated the following: the

<11 1> orientation is superior to all other orientations in creep loading; the <001 > orientation

offers the best strain-controlled low cycle fatigue resistance because of its inherently low elastic

modulus; and the 112 orientation seems to offer the best combination of the above mechanical and physical properties extremes for use in the radial wafer design. Figures 12, 13, and 14 compare the most relevant properties for turbine airfoil as a function of orientation (where appropriate for this material). Because of these orientation property variations, the <11 2> orientation

was selected for blading applications.

FIGURE 12 STRESS RUPTURE DATA AS A FUNCTION OF ORIENTATION FOR ALLOY PWA 1121

Larson-Miller Parameter, C-20.0

FD 254162

821311 0945B FIGURE 13 LOAD CONTROLLED THERMAL MECHANICAL FATIGUE DATA OF ALLOY PWA 1121

FO 2541&3

621211 0952B

496
With the <112> having been selected as the prime axial airfoil direction and with the need for

strong texture to enhance directional grain growth, cold cross-rolling was selected as the prime

material finish forming technique. Cold rolling appeared to offer the potential for developing the

strongest deformation texture, and cross rolling offered the potential for developing a truly singular

|110|<112> texture. Early cold cross rolling yields were less than satisfying (Figure 15). However, recent experience has been far more enjoyable with rolling yeilds approximately approaching 90%

of theoretical. The texture produced by these operations is depicted in Figure 16.

FIGURE 15 EFFECT OF TECHNIQUE ON COLO-CROSS ROLLING YIELDS

FIGURE 16 X-RAY DIFFRACTION POLE FIGURES MEASURED ON COLO CROSS ROLLED PWA 1120

Contour Values

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FIGURE 17 OmCNTATlONS DEVELOPED OLTRINQ UNDtRECTED DR Of COLO^iROSS ROLLED ALLOY PWA 1120

I I

This material was exposed to the D.R. cycle with no intentional anneals after rolling. As stated

above, early studies indicated that texture steered grain growth. Apparently, however, texture intensities of the level developed here are insufficient to fully control growth orientation. A sampling of the D.R. orientations developed when as-rolled material is directionally recrystallized

parallel to the primary working directions is illustrated in Figure 17. Note that the texture seems

to preclude the appearance of random orientations rather than to guarantee the presence of selected

orientations. Furthermore, since the <11 0> direction appears with such frequency, it appears that

this is the preferred growth direction in this alloy system, and this habit growth will overcome the

effects of texture of this level of strength. Again, it is speculated that a stronger texture may

offer the potential for greater orientation control, but to date, testing of this hypothesis has not been possible.

An alternate technique is required, therefore, to control orientation. Two techniques were tried.

The first technique was statistical in nature and involved initiating a large number of grains at the

top of the plate, selecting an appropriately oriented grain (if possible) by Laue' back diffraction

techniques, and isolating and propagating the selected grain through the wafer. This method was i

found too tedious for general applicability but, in fact, was used to generate seed material. The S

more preferred method, and the technique used for the growth of most of the crystals in this effort,

was to bond on a growth directing seed crystal. 1 ;

(^1 Initial attempts at bonding a truly single crystalline seed onto the fine grained stock gave spotty f'j results. It was not found feasible using the techniques described above to drive a grain with an :.'H inappropriate orientation through the pre-established texture. Grain growth would cross the bond ; plane, terminate, and new appropriately oriented grains would nucleate, grow, and consume the fine 'i grained stock. Even with appropriately oriented seeds, grain growth would often not continue from the,y) I seed through the matrix. The reasons for these growth breakdowns are unclear, but the "hard seed

technique" was abandoned because of its low success ratio. ^

The success ratio of the "soft-seed technique" has been more than adequate. The superplastic, fmeqrained region of the soft seed apparently complies well with the fine grained sample, allowing

local bond joint oxide disruption and/or grain boundary migration/recrystallization during

Whatever occurs, subsequent growth down the seed, across the joint, and into the matrix is en anc . Seed transferral using this method and differential expansion loading to mate the surface is

above 90%. Using this method, growth orientation control resolves down to mechanically aligning tne appropriate seed orientation with the desired primary growth axes of the plate. To date this alignment has been easily held to within about 8 degrees of the nominally required orientation. A evel

which has been somewhat arbitrarily selected as adequate for mechanical property predictability.

There appears to be preferred growth directions in the NiMoAl X alloys evaluated as cross rolled

plate. The <110> orientation has exhibited growth in this system at greater than /hr. The

growth rate of the <112> orientation has, at times, approached this >^ate but

Lhibit a growth limit nearer l/2"/hr. The<l 1 1> di rection seems to exhibit a ing near 3/8 /hr.

Normally, when one of the higher index grains breaks down, the new gram ^hich replaces it i ented such that a <11 0> type axis is parallel to the principal rolling direction with the UUUl

Diane in the plane of the sheet (cube on face, rotated 45° texture). This orientation unfortunately^

does not exhibit outstanding creep or fatigue strength and has been the secondary gram orientation

which has claimed most of the rejected samples.

I

498
A final statement will be made concerning the economics of the technique as outlined above. The

current technique calls for approximately 7 lbs of powder to be consolidated to yield each pound of

actual airfoil block material. This relatively low material utilization factor results from the need

to roll oversized product to protect dimensions and finished surfaces and the apparent need to cross

roll to enhance texture strength. Various schemes have been advanced to raise this ratio including

straight, bright finish rolling to 0.005" over finish thickness. Such an improvement in yield (to

about 30% overall) appears technically feasible, but not without a large amount of technical effort

to guarantee that the final product will be appropriately heat treatment responsive. A large percentage of the cost of this component is associated with the grain growth yields and with the

material utilization factor. If both of these factors can be improved through improved rolling technique designed to enhance texture strength and improve material utilization then a substantial re- duction in the cost of these components (several hundreds of dollars per component) could be

realized.

CONCLUSION

The feasibility of producing single crystal stock out of NiMoAl X alloys for use in radial wafer

turbine blade fabrication has been demonstrated using a variety of processing routes. Rolling to produce directly heat treat responsive material has also been shown to be a viable approach. "However,

reliance on texture to control grain growth orientation has been found unfeasible, and a supplemental orientation controlling technique (solid-state crystal seeding) has been shown necessary.

The properties of the single crystals produced by these techniques fully meet expectations and represent an excellent compromise between strength and durability. Effort should continue, however, to improve the economics of the approach.

REFERENCES

1. Bourdeau, R. G. and J. B. Moore, Thermomechanical Processing of RSR Superalloy Powders, in

Proceedings of the International Conference on Rapid Solidification Processing, Reston, Va, Nov. 1977, 334-346.

2. Cox, A. R., J. B. Moore and E. C. van Reuth, Rapidly Solidified Powders, Their Production,

Properties and Potential Application, AGARD Conference Proceedings No. 256, 1978.

3. Holiday, P. R. , A. R. Cox and R. J. Patterson, II, Rapid Solidification Effects on Alloy

Structures, Proceedings of the International Conference on Rapid Solidification Processing,

Reston, VA, Nov. 1977, 245-257.

4. Anderson, R. E., et. al . , Use of RSR Alloys for High Performance Turbine Airfoils,

Proceedings of the Second International Conference on Rapid Solidification Processing,

Reston, VA, 1980, 416-428.

5. Cox, A. R., and R. J. Ondercin, Manufacture of Advanced Turbine Airfoils Using RSR Powder

Alloys, Manufacturing Technology Advisory Group (MTAG), Phoenix, AZ, 1979.

6. George, D. B., B. T. Brown, and A. R. Cox, The Application of Rapid Solidification Rate Superalloys to Radial Wafer Turbine Blades, AIAA/SAE/ASME 15th Joint Propulsion Specialist

Conference, Las Vegas, Nevada, 1979.

7. Chin, H. A., S. E. Hughes, J. S. Andrus and J. L. Larson, On the Texture Development-Directional

Recrystallization Interaction of Rapidly Solidified Ni-Mo-Al-W and Ni-Cr-Al-W Alloys, paper

presented at the International Conference on Rapid Solidification Processing, Reston, VA,

December 6-8, 1982.

ACKNOWLEDGEMENTS

This work is jointly sponsored and monitored by the U.S. Defense Advanced Research Projects Agency

and the Air Force Wright Aeronautical Laboratories.

The views and conclusions contained in this paper are those of the authors and should not be interpreted as necessarily representing the official policies, either expressed or implied, of the U.S.

Defense Advanced Research Projects Agency, U.S. Air Force, or the U.S. Government.

499
DEFECT ANALYSES OF P/M SUPERALLOYS

W. H. Chang, H. M. Green and R. A. Sprague

General Electric Co., Evendale, OH 45215

ABSTRACT

This paper discusses the characteristics of manufacturing defects in a typical

high-strength P/M superalloy, Rene 95. The nature, origin and size range of four

types of defects found in the as-HIP material are described and their influence

on the LCF behavior analyzed. Isothermal forging of the HIP material is shown to improve the LCF behavior considerably. The improvement is attributed to dispersion of large defects, especially of the PPB (Prior Particle Boundary) type.

As a result of defect dispersion, the LCF life of the forged material is also

benefited by a greater tendency toward internal fatigue initiation and initiation

by crystallographic cracking or from small voids.

INTRODUCTION

The need for high-performance aircraft-engine disks has during the past decade spurred

the development of P/M superalloy technology. Among the challenges accompanying this development, none are more critical than first, the availability of high-quality powders or billets and second, the design and management of engine components for safe and reliable

service introduction. P/M superalloys can be consolidated by hot isostatic pressing (HIP)

,

HIP followed by forging, or other combinations of compaction and working. Depending on the

powder cleanliness and processing procedure, each type of material can contain defects

differing in nature, size and distribution. A thorough knowledge of the defect characteristics is a prerequisite to predictability of material behavior and hardware performance.

This paper discusses the defects which have been found or Induced in specimens of a typical high-strength P/M disk alloy, Rene 95, both in the as-HIP condition and after isothermal forging. Rene 95 in P/M form has a nominal composition (in wt %) of 0.065C, 13Cr,

8Co, 3.5A1, 2.5Ti, 3.5CB, 3.5Mo, 3.5W and balance Ni. It normally consists of coarse, medium and fine gamma prime in a highly-alloyed gamma matrix which also contains a fine

dispersion of Cb, Ti-rich carbides. Unless indicated otherwise, the materials discussed in

this paper were produced from -150 mesh powders by argon atomization. The heat treatments

consisted of 2065F/1 hr, oil quench, 1600F/1 hr, A.C. and 1200F/24 hrs, A.C. for the as-HIP

condition and 2055F/1 hr, A.C. and 1400F/16 hrs, A.C. for the HIP and isothermally forged

(H+F) condition.

Defects in As-HIP Material

Defect Characteristics

Table 1 summarizes the characteristics of four types of defects identified at the

failure Initiation sites of as-HIP Rene 95 LCF specimens. Representative appearances of

these defects are illustrated in Figure 1. Types 1 and 2 are ceramic inclusions which exist

as discrete and often chunky particles (Type 1, Figure lA) or as agglomerates of fine particles (Type 2, Figure IB) . Both types apparently originate from the melting crucible, pouring

tundish or the atomizing nozzle during powder production. Aluminum is usually the major

metallic element with Zr, Mg or Ca also present in various amounts depending on the particular ceramics involved. The sizes of Type 1 and Type 2 defects are limited by the screening

opening through which the powders have been sieved and are not altered significantly by

reaction during the HIP cycle. As indicated in Table 1, the ceramic defects average about

10 mils2 with very few reaching 50-100 mils^.

In direct contrast to Type 1 and Type 2 which are "condensed" defects. Type 3 or the PPB

(Prior Particle Boundary) defects are "diffused" in that they generally spread over a larger

volume of material. As shown in Figure 1C, Type 3 defects are characterized by a semi500
continuous network of fine oxide or carbide particles around the host powder surfaces. A

residual core of the initial defect-forming source may or may not be present at center of

the affected area. Figure 1C illustrates an example of PPB defect with a core of refractory remaining. Formation of Type 3 defects appears to require a source of oxygen or carbon

which reacts with powder surfaces during the HIP cycle. Potential sources of PPB defects

include refractories containing water glass (sodium silicate) as binder, organic contaminants inadvertently introduced during powder handling (e.g. Buna N and vacuum grease), mill

scale (rust) and leakage during powder canning. Depending on the nature and amount of the

contaminant and the rate of powder consolidation during the HIP process, PPB defects can range from a few mils^ to several orders of magnitude larger than the maximum (250 mils^)

shown in Table 1. The PPB defects are thus potentially the most detrimental to P/M superalloys .

Since the powders are produced by argon atomization, argon is unavoidably entrapped in some of the powders, giving rise to voids or Type 4 defects in the consolidated product.

Individual voids seldom exceed 2 mils^. Clusters of voids, however, often exist to yield

substantially larger combined sizes as indicated in Table 1.

Effect of Defects on LCF

Defects are of major concern to P/M superalloys primarily because of their adverse

effect on LCF. The total LCF life (Nj-) of high-strength disk alloys can be considered to consist of three parts: life spent in crack initiation (Nq)

life in Stage-1 propagation

(N]^) and life in Stage-2 propagation (Nji) which ends essentially in final failure. Within

certain temperature and strain ranges, the sum of No+Np can consume over 90% of total LCF

life. Yet micros tructurally , No+Np corresponds to but a very small crack size or area. This means that defects whose sizes exceed the corresponding No+Np crack size would reduce

LCF life drastically by depriving the material of resistance to crack initiation and

Stage-1 propagation. Figure 2, for example, shows the effect of EDM holes of prescribed

sizes intentionally introduced to the gage section of LCF specimens made from as-HIP, CAP

(Consolidated at Atmospheric Pressure) and rolled, and cast and wrought Rene 95, tested at lOOOF, A=l, and Ae^=0.66%. In the absence of significant defects, the intrinsic LCF life

is well over 100,000 cycles. Induced surface defects exceeding 5 mils^, however, are sufficient to reduce life to less than 10,000 cycles by eliminating Nq+Nj.

Defects affect LCF life to a degree which varies with both defect size and location of

failure initiation. This is illustrated in Figure 3 for as-HIP Rene 95 at lOOOF, A=1 and

Aet=0-66%. The data shown were obtained from specimens which contained surface or internal

defects sufficiently large to permit only Stage-2 crack propagation. Within this range, the

LCF life is linearly related to defect area on a log-log basis. Note that for a given size,

a defect is much less detrimental when situated internally than when it is located at the

surface. The less severe effect of internal defects is due in part to the reduced stress

range intensity amplification but in a large measure probably results from the absence of

adverse environmental interactions. Since surface failures are less likely the fewer the

number of defects. Figure 3 reinforces the simple concept that the LCF life of as-HIP P/M

alloys can be improved to the extent the defects are reduced both in number and in size.

This concept is particularly pertinent to PPB type defects.

Table 1: Major Defects in As-HIP Material

2 Major Size, Mils (*)

Type Description Typical Elements Avg. Max,

1 Discrete, chunky ceramic Al, Mg, Zr , Ca, 0 8 50

2 Ceramic agglomerates Al, Si, Mg, Ca, 0 10 no

3 Reactive agglomerates

forming PPB

Al, Zr,

C, Ca,

Cr , Ca , 0

(Fe)

15 250

4 Voids Ar 6 10

(*) Based on defects at Initiation sites of LCF test

specimens made from -150 mesh powder.

501
502

Figure

1:

Defects

in

As

HIP

Rene

95:

A,

B,C

and D

Reoresent

Types

1,

2,3

and

4,

Respectively

(See

Text)
Cycles

Figure 2: Effect cf Surface Cafect Size cri LCF Life of Various iypes of Rare 95

503
Nf,

Cycles

I

I

I

504
Defect Modification by Isothermal Forging

Results on Doped Material

The potential of isothermal forging in defect modification was initially assessed by

doping Rene 95 powders with Type 1 (chunky ceramic) and Type 3 (PPB) defect formers. The

dopants selected were AI2

O

3 (Type 1, -60+80 mesh, 1,000 particles per pound of Rene 95),

Buna N (0 ring material. Type 3, 500 particles per pound) and vacuum grease (Type 3, 1 cc per pound). Both doped and undoped materials, after HIP at 2050F, were isothermally

forged at 2050 for 0 to 80% reduction. Figure 4A shows that forging had little effect on AI2 O

3

, other than occasional fracturing. The larger PPB defects, on the other hand, were

greatly reduced in size, generally to stringer or cigar shapes, lying randomly in the

forging plane (Figure 4, B and C) . Consistent with these characteristics of defect modification, there were corresponding changes in mechanical properties. Table 2 shows that both

tensile ductility and LCF life of doped Rene 95 were improved by the forging operation.

Further, the improvements were much more pronounced in materials doped with Type 3 than

with Type 1 defects.

While as-HIP materials (doped or undoped) invariably initiated LCF failure at one of

the four types of defects described previously, some forged specimens failed by crystallographic cracking without impurity-related defects. An example is shown in Figure 4D. The

initiation was usually associated with one or more somewhat larger grains which are designated as Type G "defects". Note that since the forged structure generally had very fine

grain sizes (<10y) and that even occasional larger grains seldom exceeded 30y, Type G

failures preserved their initial processes of crack initiation and Stage-1 propagation, as evidenced by the small facets delineating the Initiation site in Figure 4D. Consequently,

Type G failures were generally associated with long LCF lives.

Effect of Forging on Undoped Material

Defect Type and Size Distribution ; Based on examination of initiation sites of LCF

specimens. Table 3 compares the defect type, defect size and location of failure between

as-HIP and H+F (HIP+isothermally forged, 60-80% reduction) Rene 95. Several trends are discernible pertaining to the influence of forging on defect modification:

1. The forged specimens were much less prone to surface initiation, being almost immune

at lOOOF (Table 3, Column D) . Under conditions more conducive to surface initiation

(750F/0.7% Ae^) , the frequency of such failure was 33% vs. 75% for the as-HIP specimens.

Table 2: Effect of Isothermal Forging on Ductility

and LCF Life of Seeded Rene 95

Material

RT/1200F

Tensile

El., %

lOOOF/0.

LCF Life,

X

66% Ae^- Cvcles

1T-3S

1 . As-HIP (Unseeded) 17/15 36,840 11,800

2 . H+F (Unseeded) 20/22 128,848 -

3. As-HIP (AI2 O

3 Seeded) 10/5 3,411 1,758

4. (3)+F 14/15 8,219 4,101

5. As-HIP (Buna N Seeded) 10/3 2,561 1,200

6 . (5) +F 19/19 134,001 22,241

*A11 -150 mesh powder. Forging reduction: 80%

505
’(i

D_ c ^ o c +-> • 1— to O 3

+->

C -rO

ra +-> -C

•1- c o

4-> 3 • r—

•r- “O C CO 0) 3

H-» CL QC CD

Q_

C

<D CT> •1— s- -o C 3 <D ‘r- o

1—

O CD 03

•1- 3 S- s. rO -O O

Lu

C Li_ CJ

t-H

u

U. r— O OJ ro -C

—I E cx

fO s- fO “D CU O) S- C S- -C CD

rO CD 4->

O

O r—

CO E to

+-> 3 »—03

O

o>

O ^ to

03

O >>

<D > CO s- Q o --- >> C CO

-Q “O

CD »* a; Ll c Ll + + M- CD^ •»- JZ

S- “O c o c o •f— to u_ ‘f- +-> c o

rC ^ O O)

COO CO •r- M- E +-> <D

S- to 03

Q

0) 3 C •r— -Cl— *1- -«-> *o

-MU •f- CU

O

C +->

C 4-^ CO t-H C> HH 03 1—1 dJ 03 M- Ll CD 4- C <D O QC O Q —1 1 E >>

to 3 CO ^ 4-> r— CL CO •!- CJ <: Cl S- 0) 3

“O CL

C

O c E

(U

Sk-i

-1|I

506
Table 3: Comparison of Defect Type and Size at LCF Initiation Site - As HIP vs H+F

A

Temp

B

^^t

C D Overall

Surface

Failure

E Frequency of^

Defect Type

%

F

Defect

G

Defect Size

“As HIP

, Mils

H +

2T

oF Material % 1 1 3 4 Type Ave. Max, Ave. Max.

1000 0.78 As HIP 76 21 74 4 1 0 1 6.7 ”3T 5.3 TO

H+F 5 66 5 0 0 29 2 9.5 50 6.1 11.3

3 18.7 242 None None

4 6.3 10.3 None None

0.66 As HIP 31 15 77 7 1 0 1 11.2 50 3.9 12.0

H+F 5 29 28 3 14 26 2 13.1 111 4.8 13.0

3 19.3 83.6 4.2 4.5

4 5.3 6.8 0.8 1.4

750 0.70 As HIP 75 33 58 9 0 0 1 7.5 9.1 6.4 13.7

H+F 33 13 24 0 30 33 2 8.6 13.1 4.2 9.3

3

- 6.4 None None

4 None None 1.1 1.5

* Type G represents crystallographic cracking without impurity-related defects

507
2. In contrast to as-HIP specimens which invariably failed from impurity-related

defects (Types 1 through 4), a substantial number (26 to 33%) of H+F specimens failed from

Type G defects (Column E) . Of these, about 5% at lOOOF and about 50% at 750F were surface

related, constituting a major source of surface failures at the lower temperature. As

explained before. Type G failures, even when initiated externally, were associated with

long LCF lives

3. Types 1 and 2 ceramic inclusions constituted the most frequently observed

defects in both materials. The average size of each type (5-10 mlls^) did not differ significantly between the two materials although the maximum sizes (especially Type 2) appear to be somewhat smaller in the H+F specimens (Column G) , suggesting that the Type 2 defects

may have been dispersed by the forging reduction.

4. Although Type 3 (PPB) defects were only infrequently found in both materials, the

frequency of occurrence was consistently higher in the as-HIP specimens (4, 7 and 9% vs.

0, 3 and 0% for H+F, under the three testing conditions. Column E) . Further, both the

average and especially the maximum sizes found in the H+F specimens (4.2 and 4.5 mils^,

respectively) were considerably smaller than those found in the as-HIP specimens (19 and

242 mils^, respectively).

5. Void-induced (Type 4) failures became much more frequent in H+F specimens, reaching 14% at lOOOF/0.66% Ae^ and 30% at 750F/0.7% as compared with 1% or less for the as-HIP

material. The lack of Type 4 failures in the latter case is not surprising since the

relatively small voids would be bypassed as initiation sites in favor of larger defects

which were present. Once the larger defects had been dispersed by forging, then the most

aboundant and uniformly distributed small defects would be favored for crack Initiation,

especially those located at or near the surface. These conditions were fulfilled by voids

in the H+F material, thus accounting for the propensity toward Type 4 failures in the H+F

material, from voids located only slightly below surface. Note from Table 3, Column G that

the voids in the H+F specimens were very small (<1.5 mil^), with correspondly minor detrimental effect on LCF life.

Effect of Forging on LCF Life ; The preceding comparisons indicate that through its

effects in eliminating or modifying large defects (especially PPB) , isothermal forging

facilitates (a) internal crack initiation, (b) Initiation by crystallographic cracking and

(c) void-induced initiation. As shown in Figure 5 and Table 4, these characteristics in

concert resulted in several fold improvement in LCF life (both average and -3S) over as-HIP

Rene 95. The Improvement was consistent with that shown previously in Table 2 for doped

materials.

Table 4 : LCF Life Comparison Between AsHIP and HIP+Isof orged Px.ene 95

Temp.

°F % Material

Nf, Cycles

X X-3S

750 0.70 As-HIP 21,003 3,920

H+F 104,148 34,103

1000 0.78 As-HIP 11,676 1,564

H+F 61,693 12,961

0.66 As-HIP 37,000 4,500

H+F 181,052 70,752

508
Cycles

Figure 5: Overall LCF Life Comparison at lOOOF

(As HIP vs H+F)

Concluding Remarks

Given clean powder and manufacturing care, the as-HIP process is capable of producing

high-quality components. The results presented in this paper indicate that further

improvement in LCF behavior can be achieved by isothermal forging through defect modification. The fine-grained structure effected in a forged product provides an additional

advantage in overall defect control by allowing more discriminating nondestructive inspection. The benefits of forging reported in this paper, while based primarily on laboratory

specimens, are achievable in engine components provided, again, a high degree of powder

cleanliness and manufacturing care is maintained.

Acknowledgments

The authors are indebted to the following individuals of the General Electric

Company; Cliff Shamblen for his work on defect characterization on USAF Contract

F33615-78-C-5225, David Mourer for his contribution in isothermal forging evaluation and

Ken Wright for reviewing the manuscript.

509
INTERPRETIVE REVIEW OF COMPARATIVE MICROSTRUCTURES OF

POWDERS AND CONSOLIDATES OF P/M SUPERALLOYS PRODUCED

BY MODERN ATOMIZATION PROCESSES*

Robert D. Kissinger and John K. Tien

Henry Krumb School of Mines

Columbia University

New York, NY 10027

ABSTRACT

This paper reviews in a comparative manner cooling rates, powder morphology and con- solidated microstructure and some properties of superalloys produced by the commercial and

the new processes. The review strategy will be to give preference to the comparison of

the same alloy produced by different atomization techniques. For example, the surface mor- phology of the RSR powder is smooth with occasional fracture, whereas the AA powder surfaces

are connected to many small satellite powders. Depending on powder size, the AA powders

have approximately twice the number of particles containing gas pores as the RSR powder.

From secondary dendrite arm spacing (DAS) measurements, it was estimated that argon atomized superalloy powders show a cooling rate with the same order of magnitude as the RSR

superalloy powders. Comparisons and conclusions involving other atomization processes

will also be presented and discussed, as will some mechanical properties.

* An invited review paper

510
Introduction

Commercial production of superalloy powders is primarily by inert (usually argon) gas atomization

(AA) and vacuum, or soluble gas, atomization (VA). A third commercial process, the rotating electrode

process (REP), has been used to produce superalloy powders, but has found greater success in titanium

powder production. Still in the developmental stage are the rapid solidification rate (RSR) and the

ultrasonic atomization (UA) processes. A detailed description of the above atomization processes can

be found in the literature [1] and will not be discussed in this compacted review dealing mainly with

powder metallurgy (P/M) superalloy microstructures. Emphasis will be given to comparisons of the same

alloy produced by different atomization processes and, more specifically, to a comparative study of VA,

AA, and RSR superalloy powders currently in progress at Columbia University. For this comparative

study, RSR and AA powder material of IN 792 (AA material only), IN 792 + Hf (ClOl), and IN 792-boron,

carbon modified (B1925) produced from the same ingot material were supplied by Pratt & Whitney Aircraft

through DARPA contract AF 33615-76-C-5136 and Special Metals Udimet Powder Division, respecti vely . UA

modified IN 100 was supplied by Homogeneous Metals, Inc.

Powder Size Distribution

Typical powder size distributions [2] of similar superalloy powders produced by VA, AA, and RSR

processes are given in Fig. 1. The RSR process apparently produces a narrower size distribution of

powder particles as compared to the broader distributions of the VA and AA processes. Although mean particle diameters of the RSR and VA powders are slightly smaller than the AA powders, these values

can be varied to some extent by changing certain processing parameters. For example, increasing the

speed of the rotating disk in the RSR process [3] or increasing the relative gas to liquid velocity

[4,5], or changing the nozzle configuration in the AA process will result in a decrease in the mean particle diameter. For the REP process, powder size is, on the average, coarser than the above pro- cesses with the powder particles usually falling within a narrow size range.

Surface and Interior Power Structures

An example of superalloy powders produced by AA, RSR, and VA is shown in Fig. 2. Although the

three atomization processes yield relatively spherical particles, there are slight differences in

surface morphology. The majority of RSR powders are smooth with occasional fractures caused by

particle-particle or particle-apparatus wall collisions. Both AA and VA powders exhibit small satellite particles connected to powder surfaces. All three types of powders occasionally show particles

enveloped by metal shells, occurring when a liquid droplet impinges on and surrounds an already so- lidified particle. The RSR powders generally show the greatest percentage of these splat cooled

structures. Not shown in Fig. 2 are REP powders, which are characterized by smooth, satellite-free

particles similar in surface character to the RSR powders.

Analysis of sectioned AA, RSR, and VA powders reveals the presence of porosity in the form of

entrapped gas. A simple qualitative calculation determining the percentage of particles containing

gas pores shows that at least twice as many AA powder particles contain gas pores as do RSR powders

of the same particle diameter (Fig. 3). For the one VA superalloy powder studied, the percentage of

powder particles containing gas pores was similar to AA powders.

Individual powder particles can be considered as small castings with cellular, dendritic, and/or

microcrystalline microstructures characteristic of each particle's thermal history. While cooling

microstructures of atomized powders differ for different powder particle sizes, powder particles of

the same diameter also exhibit varying cooling microstructures. For example. Fig. 4 shows both a dendritic and a microcrystalline microstructure for two 12 ym-diameter AA ClOl powder particles.

Relative Cooling Rates and Retained Carbon in Solution

For those particles which are dendritic, it is possible to measure secondary dendrite arm spacings (DAS) and correlate these measurements with particle diameter. It is found that DAS varies with

particle diameter according to the empirical relationship [6]:

DAS = Ad'^ (1)

where d is the particle diameter. Table I lists values of A and m calculated for similar superalloy

powders produced by VA, AA, and RSR. In graphical form (Fig. 5) the data show that for the same alloy

and average particle diameter, the AA process produces a coarser average DAS as compared to the RSR

process. Data for VA IN 100 are included in Fig. 5 for comparison with the AA and RSR IN 792-type

superalloys.

511
Fig. 1. Powder size distributions for RSR ClOl, argon atomized (AA) ClOl

and vacuum atomized (VA) IN 100, (Ref. 2.)

I

f:H

Fig. 2. Superalloy B1925 powders produced by (a) argon atomization and

(b) RSR and (c) Superalloy IN 100 produced by vacuum atomization.

(-140+150 mesh powders)

512
TABLE I

Secondary Dendrite Arm Spacing vs . Particle Diameter

DAS =

Process Alloy A m

VA IN 100 0.225 .525

AA IN 792 .164 .545

AA C 101 .130 .605

AA B 1925 .159 .540

RSR C 101 .116 .574

RSR B 1925 .105 .512

DAS has also been shown to vary with cooling rate according to the empirical relationship:

DAS = BT-i^ (2)

where T is the local cooling rate. Unfortunately, experimental values of B and n have been determined

for only a few alloy systems within a restricted range of cooling rates. However, it is possible to calculate approximate relative RSR/AA cooling rates by coupling Equation (2) with Equation (1). If

the ratio of DASrsr/DASaa is calculated for the same alloy from data listed in Table I,, and the con- stant B in Equation (2) is considered only alloy dependent, the relative cooling rate, TRSR/tAAs can

be expressed as:

IrSR = °^^RSR ”n (3)

\l\ '^^^AA .

Calculated relative cooling rates for RSR and AA B 1925 and ClOl as a function of particle diameter

are shown in Fig. 6 for n = 0.3, which is the experimental value determined for a Maraging 300 steel

from data spanning four orders of magnitude up to lO^oc/sec [7]. For both alloys the AA powders sur- prisingly show cooling rates within the same order of magnitude as the RSR powder cooling rates,

which are quoted to be on the order of 10^°C/sec for superalloy particles in the 100 ym range [8].

With these extremely high cooling rates available for superalloy powder production, it is possible to extend the solid solubility limits of various alloying elements. For example, a detailed

analysis of carbon partitioning in AA Rene 95 [9] has shown that as particle size decreases (and

cooling rate increases), the percent of total carbon in the powder present as MC type carbides de- creases. Figure 7 shows the hyperbolic nature of the carbide partitioning function.

Consolidated Microstructures

The most common consolidation techniques used in P/M superalloy processing are hot isostatic

pressing (HIP) and hot extrusion. Early attempts at P/M superalloy consolidation resulted in the

deleterious precipitation of carbide films at prior particle boundaries (PPB), as well as the formation of precipitates at PPB's due to oxygen and nitrogen gas adsorbtion on powder particle surfaces.

These problems have been solved to some extent by alloy modifications, inert handling of powders, and

changes in consolidation heating cycles. For example, carbon levels of P/M superalloys reduced to

0.02 weight percent [10,11] and the addition of strong carbide formers [12,13,14] , e.g., Hf, Nb, and

Ta, have reduced or eliminated PPB carbides and/or modified the type and distribution of carbides.

Table II lists the composition of modern P/M superalloy chemistries, which makes use of one or more

of the above composition modifications. With the combination of alloy redesign, improved powder

handling, and appropriate heat treating schedules, the integrity of as-HIPed P/M superalloys has

dramatically increased, as witnessed by a recent Pratt & Whitney Aircraft report [15], which states

that as-HIPed VA and AA MERL 76 can be used for commercial jet engine disks with considerable saving

in weight and machining costs.

The response of superalloy powders to HIP has been reported to vary with respect to the atomization process. Centrifugal ly atomized (RSR) iron-nickel base powders have shown a greater tendency,

as compared to AA powders produced from the same piece of bar stock, to retain the as-solidified

powder structure during HIP [16]. We have also found this to be the case for as-HIPed and HIP + heat

treated RSR and AA B1925 and ClOl. A comparison of consolidated microstructures of alloy ClOl is

513
PARTICLE DIAMETER, d (urn).

Fig. 3. Percent of argon atomized (AA) and centrifugally .. atomized (CA) powder particles containing entrapped ’ gas pores as a function of particle diameter.

1

Fig. 4. Two 12 ym-diameter argon atomized ClOl powder

particles exhibiting (a) dendritic and (b) microcrystalline microstructures.

I

I

I

514
TABLE II

Nominal P/M Superalloy Compositions (Wt. Percent)

Alloy Ni Cr Co Mo W Cb A1 Ti C B Zr Hf Other

Astroloy 55.4 15 17 5 4.0 3.5 .06 .03 _ _ IN 100 54.5 12.4 18.5 3.2 - - 5.0 4.3 .07 .012 1.0 - 1.0 V Rene 95 62.4 13 8 3.5 3.5 3.5 3.5 2.5 .06 .01 .05 - - IN 792 60.7 12 9 1.9 3.8 - 3.4 4.0 .11 .016 .08 .79 4.1

B 1925 60.3 11.6 8.7 1.7 5.1 _ 3.4 4.0 .01 .10 - * 5.1

C 101 59.7 12 9 1.9 4.3 - 3.4 4.0 .10 .017 .06 1.1 4.4 Ta

MERL 76 54.5 12.5 18.5 3.2 - 1.4 5.0 4.4 .02 .02 .06 .4 AF 115 55.3 10.7 15 2.8 5.9 1.7 3.8 3.9 .05 .02 .05 .75 -

Rene is a trademark of the General Electric Company

MERE 76 is an improved turbine disk alloy developed by Pratt & Whitney Aircraft

AF 115 is also an advanced turbine disk alloy developed by the General Electric Company for the

Air Force

shown in Fig. 8. The -104 ym (-150 mesh) RSR material seems to retain more of the pre-HIP powder

structure, with the average y' precipitate size in the large particle interiors remaining much smaller

than the y' precipitate size in the interparticle regions, even after heat treatment. There also

appears to be a higher percentage of thermally induced porosity (TIP) in the -230 ym (-60 mesh) AA

HIP + heat treated material.

The as-HIPed RSR and AA consolidated materials were studied using transmission electron microscopy, and it was found that the RSR material had a tendency to retain not only original grain size

within the larger prior particles, but also to retain the smaller scale DAS microstructure, as shown

in Fig. 9a. This tendency was not observed nearly as often in the AA material (9b) and is thought to

be related to the slightly higher cooling rate, and accompanying increase in homogeneity, of the RSR

superalloy powders as compared to the AA powders.

Conclusion

Following HIP or hot extrusion consolidation, P/M superalloys are generally hot forged or rolled.

These P/M materials show excellent hot working characteristics due to their small grain size, and the

lateral flow of material associated with forging or rolling improves low and intermediate temperature

mechanical properties by breaking up PPB's and the larger gas pores. Unfortunately , available me- chanical property data for P/M RSR nickel base superalloys are limited to creep rupture test results

[17,18], making comparisons with other P/M atomization processes, of which there are large inventories

of stress rupture and tensile data, very difficult. Comparisons must therefore be limited to commercially produced as-HIP, HIP + forge, cast, and cast + wrought materials. The literature shows that

as-HIP and HIP + forge P/M superalloy mechanical properties are as good, if not better, than the cast

and wrought versions [19,20,21].

Improved low cycle fatigue (LCF) resistance of P/M superalloys has been obtained by screening

out the larger non-metal lie particles introduced during atomization. Consolidated materials still

contain P/M processing related defects, however, which consist primarily of the smaller non-metallic

inclusions and thermally induced porosity (TIP). The effect of these defects on LCF properties of

P/M superalloys has been well documented [20,22,23]. Despite the unpredictable nature of these defects, P/M superalloys have shown good LCF properties [24,25,26] and, in some cases, equal or surpass

cast and wrought alloys.

P/M processing of superalloys has achieved the benefits of increased chemical homogeneity and

improved workability. Still to be fully realized is the ability to produce reliable near net shape

as-HIPed superalloy components with the accompanying savings in materials and cost. This will re- quire a better understanding in, and control over, the atomization and consolidation processes in

P/M superalloy production, e.g., minimizing or eliminating entrapped gases in powder particles and

maximizing material flow of all powder particles during HIP.

515
DAS

(i

Fig. 5. Secondary dendrite arm spacing (DAS) vs. particle diameter for similar nickel

base superalloys.

Fig. 6. Relative cooling rates for RSR and argon I atomized (AA) Superalloys B1925 and ClOl.

% OF c AS MC

0.04C 8

0.1 6C

INGOTS

(20 cm

'DIAMETER)

2x10 '

Fig. 7. Percent carbon as MC-type carbide as a function

of particle diameter for Ren^ 95. (Ref. 9).

51
Fig. 8. Consolidated microstructures of superalloy ClOl: (a) -60 mesh as-Hip argon atomized,

(b) -60 mesh HIP + heat treated argon atomized, (c) -150 mesh as-HIP RSR and (d) -150

mesh HIP + heat treated RSR.

Fig. 9. (a) Retained secondary dendrite arm microstructure in as-HIP RSR Superalloy ClOl.

(b) Typical microstructure of the interior of a prior particle of as-HIP argon atomized

Superalloy ClOl.

517
Acknowledgment

The authors wish to acknowledge sponsorship of this research by a grant from the Air Force Office

of Scientific Research, monitored by Dr. Alan Rosenstein.

References

1. N.J. Grant: Advances in Powder Technology , 1981 ASM Materials Science Seminar, G.Y. Chin, ed.

,

ASM, Metals Park, Ohio, 1982, Ch. 1. 2. J.K. Tien and T.E. Howson: ibid, p. 161.

3. S.J. Friedman, F.A. Gluckert and W.R. Marshall, Jr.: Ch. Eng. Prog. 48, No. 4, (1952) 181.

4. K.Y. Kim and W.R. Marshall: AIChE J. 17_, No. 3, (1971) 575.

5. D. Bradley: J. Phys. D.: Appl . Phys. (1973) 1724 and 2267.

6. P.A. Joly and R. Mehrabian: J. Mat. Sci . (1974) 1446.

7. R. Mehrabian: Rapid Solidification Processing--Principles and Technology , Proc. Int. Conf. on Rapid Solidif. Processing, R. Mehrabian, B.H. Kear and M. Cohen, eds.. Cl ai tors. Baton Rouge,

1978, p. 9.

8. A.R. Cox, J.B. Moore and E.C. van Reuth: Superalloys: Metallurgy and Manufacture , Proc. Third

Int. Symp. on Superalloys, B.H. Kear, D.R. Muzyka, J.K. Tien and S.T. Wlodek, eds., Claitors,

Baton Rouge, 1976, p. 45.

9. J.A. Domingue, W.J. Boesch and J.F. Radavich: Superalloys 1980 , Proc. Fourth Int. Symp. on Superalloys, J.K. Tien, S.T. Wlodek, H. Morrow III, M. Gell and G.E. Maurer, eds., ASM, Metals

Park, Ohio, 1980, p. 335.

10. L.N. Moskowitz, R.M. Pelloux and N. Grant: Superal loys— Processing , Proc. Second Int. Symp. on Superalloys, MCIC Report 72-10, 1972, p. Z. 11. S.H. Reichman and J.W. Smythe: Mod. Dev. in Pow. Metall. (1971) 73.

12. R.D. Eng and D.J. Evans: Superalloys 1980 , Proc. Fourth Int. Symp. on Superalloys, J.K. Tien,

S.T. Wlodek, H. Morrow III, M. Gell and G.E. Maurer, eds., ASM, Metals Park, Ohio, 1980, p. 491.

13. G.J. Lewis, D.M. Parkin and F.A. Thompson: Proc. Conf. on Forging and Props, of Aerospace Mats.,

Metals Soc. , London, 1978, p. 399.

14. M.J. Blackburn and R.A. Sprague: ibid, p. 350.

15. R.D. Eng and D.J. Evans: NASA Report #N82-26439, May 1982.

16. J.E. Smugeresky: Met. Trans. 13A , No. 9, (1982) 1535-1546.

17. A.R. Cox and E.C. van Reuth: Rapidly Quenched Metals III , Proc. Third Int. Conf. on Rapidly

Quenched Metals, Vol . 2, B. Cantor, ed.. The Metals Soc., 1978, p. 225.

18. R.J. Patterson II, A.R. Cox and E.C. van Reuth: J. Metals 32, No. 9, (1980) 34.

19'. J.K. Tien, W.J. Boesch, T.E. Howson and W.B. Castledine: M^. Dev. in Powder Metall., Vols. 12-

14, MPIF, Princeton, 1981.

20. J.L. Bartos : Proc Conf. on Powder Metall. in Defense Technology, Vol. 5, MPIF, Princeton, 1980,

p. 81.

21. W. Betz, B. Borchert and W. Track: Superalloys 1980 , Proc. Fourth Int. Symp. on Superalloys,

J.K. Tien, S.T. Wlodek, H. Morrow III , M. Gell and G.E. Maurer, eds., ASM, Metals Park, Ohio,

1980, p. 643.

22. D.A. Jablonski: Mats. Sci. Eng. (1981) 189.

23. R.V. Miner and R.L. Dreshfield: Met. Trans. A 12A , (1981) 261.

24. M.M. Allen, R.L. Athey and J.B. Moore: Met. Eng. Q. 1£, No. 1, (1970) 20.

25. D.J. Evans, D.N. Duhl and R.B. Slack: Mod. Dev. in Pow. Metall. 8, (1974) 473.

26. J.F. Barker and E.H. Van Der Molen: Superal loys --Processing , Proc. Second Int. Symp. on Superalloys, MCIC Report 72-10, 1972, p. AA.

518
CONSOLIDATION AND FORMING OF P/M POROUS BILLETS

H. L. Gegel , J. S. Gunasekera, S. M. Doraivelu, J. C. Malas,

AFWAL Materials Laboratory

AFWAL/MLLM

Wright-Patterson Air Force Base, OH

J. T. Morgan, and L.

45433

E. Matson

ABSTRACT

The success of powder metallurgy (P/M) and processing depends strongly upon the

ability to produce economically a near-net-shape (NNS) part having a required set of

service properties. These properties are generally controlled by the final microstructure

which evolves through billet-preparation schedules and subsequent metal-working operations.

Quality billets must be produced to have optimum aggregate properties which are controlled

by the fundamental characteristics (composition, microstructure, size, hardness, and surface

properties) of the particles and their size distribution in the matrix. These characteristics are controlled to a large extent by the average and discrete quenching rates and,

hence, by the particular RSR powder-making process. Therefore, in this paper the powder

making processes, consolidation, and plastic working are considered as a total system and

discussed with emphasis on the plasticity theory for porous materials, analytical process

modeling of various metal-working processes, and applications of process-modeling techniques

for controlling the metal -flow pattern and achieving required microstructure and shape.

Experimental data on constitutive equations and deformation-process modeling for X7091 A1

2024 Al-20 vol% SiC, and Ti-6242 alloy systems are presented along with the discussions.

INTRODUCTION

Powder Metallurgy (P/M) is a relatively new area of process engineering. Over the past 25

years, most of the information in the literature has been devoted to studies of powder-making

processes for difficult-to-form materials, cold compaction techniques for producing near-net-shape

compacts, and the sintering process. The recent development of P/M alloys and the advent of Rapid

Solidification Technology (RST), however, has demanded hot-consolidation and subsequent plastic

working to insure the required shape and mechanical properties. Therefore, the P/M process —as it is

understood today--can be divided into three general areas: (1) powder making, (2) preform or blank

making (consolidation), and (3) plastic working of blanks or preforms. Although this is a simplistic

breakdown of a complex process (aside from the final operations), it is sufficient to provide an overall perspective. An issue in P/M technology that can be influenced by these processing steps

leading to the finished part is economics. Each step in the P/M process is important because defects

which occur during these steps generally cannot be removed by sintering or heat treatment. Therefore, the objectives of this paper are (1) to examine P/M processing from the vantage point of a vendor (forging or extrusion industry) considering the above areas as a total system as shown in Fig.

1, and (2) to present some recent results of deformation process modeling which was conducted for the

purpose of improving the properties of RST P/M materials in a finished part.

POWDER FLOW RATED

• (^TAP DENSITY^

• ^degassing

)

(f • (COMPRESSIBILITY)

G

AGGREGATE PROPERTIES

G

• (SECONDARY PROCESSING PROPERTIES]

• (SERVICE PROPERTIES$

Fig. 1. P/M process perspective.

519

Fig 2. Technological properties.
POWDER MAKING

In the P/M process, the powder-making step is of the utmost importance because the fundamental

powder characteristics of the individual particle such as composition, size, shape, microstructure,

| and surface structure are controlled by the particular process used to produce the particles. These I powder characteristics, in turn, control the technological properties, especially the aggregate

properties of the consolidated material which are listed in Figs. 2 and 3. Several popular powdermaking processes which are being investigated as possible methods for making RST powders are: air

atomization, centrifugal atomization plus forced convection helium-gas cooling, the rotating electrode | process, and the Osprey process. Although other processes exist, they will not be discussed here.

The powder produced by each of these processes has its own distinct characteristics. Two of the

factors most strongly influencing the subsequent sequences of any P/M process are (1) the particle

size distribution and (2) the environment which controls the average cooling rate and--as shown in

Fig. 4--the particle shape, particle contamination, surface layer, and specific surface area. The

particle size distribution controls the discrete cooling rate of each particle and, hence, determines || which of the particles will be quenched to the microcrystalline and which to the dendritic structure. ' Therefore, the particle size controls the hardness, composition gradient, surface layer, and micro- ' structure. A range of microstructures is usually found for the alloy powders produced by the various ' processes. It has been observed that the centrifugal-atomization process usually produces a more narrow size range, as shown in a study conducted by Hildeman [1] on the Al-Ni-Fe alloy system (see ! Fig. 5). This can be an advantage in cases where it is necessary to have a control led-microstructure |j distribution. 1

SERVICE PROPERTIES

I • YIELD STRENGTH

1 • ULTIMATE TENSILE STRENGTH

I • DUCTILITY

I • TOUGHNESS

I • LOW CYCLE FATIGUE

1 • HIGH CYCLE FATIGUE

; • HIGH TEMPERATURE PROPERTIES

PLASTIC WORKING PROPERTIES

I • INTERPARTICLE BOND STRENGTH ]

I • RELATIVE DENSITY |

I • CONSTITUTIVE RELATIONS |

1 • WORKABILITY '

Fig. 3. Aggregate properties. Fig. 4. Powder process effects.

Alloy systems, designed for high- temperature applications, which contain strong intermetal 1 ic

compounds present some unique difficulties because of the wide range of microstructures which develop

as a result of the discrete quenching rate for each particle associated with a particular particle i size distribution. A large difference in the temperature dependence of the microhardness value is j expected for particles having the microcrystalline structure and those having the dendritic structure. ^ From studies to investigate the LCF behavior of nickel -base P/M materials, it is known that the inter- ^ faces of coarse particles of dendritic structure are sites for fracture initiation. It has been

observed that some critical strain is required to ameliorate the difficulties associated with the ! coarse powders and foreign matter introduced by powder handling. Overcoming these effects is a ; metalworking problem which could be simplified if one were working only with powder having a homogeneous microstructure as well as chemical composition. The coarse dendritic particles retain their | hardness with increasing temperature; therefore, they behave as rigid bodies, moving with the metal

flow pattern. They do not shear and often agglomerate, which acerbates the problem. The metalworking i problem then is to control the distribution and strengthen the matrix to permit the transfer of stress!

and strain into the dendritic particles. The latter is required to condition the particles in such a

' way that they will be energetically favorable for recrystallizing during subsequent processing steps. ,

The surface layers of metal -alloy powders are enriched with impurities; and therefore, the

physical properties may differ considerably from those of the base material. This layer may be

enriched with some solute element, an oxide layer, a hydrated oxide layer, a metallic carbide, an intermetal 1 ic compound, or a combination of these, e.g., an oxide layer plus an intermetal lie compound. The surface composition is controlled by the size and shape of the particle since both affect

the specific surface area. As an example, Hildeman, Lege, and Vasudevan [1] showed the variation of

oxygen content with surface area for angular, flake, and spherical Al-Ni-Fe alloy powder (see Fig. 6).

The powder-making process controls the nature of the surface layer and the physical properties of the

base material. The surface layer is an important part of the particle structure since it accepts the

520
0.5 T T T T

o - ATOMIZED POWDER (ANGULAR) A - SPLAT (FLAKE) 0.4 _

- RSR POWDER (SPHERICAL) o

B.B 0.1 0.2 0.3 0.4 0.5

BET SURFRCE HREfl ISQURRE LETERS/GRRa

Fig. 5. Coulter counter particle size

distribution of Al-3,7 Ni-1.5 Fe powders

contact load when the powder is poured into the mold. The contact pressures on the surface layer,

due to the small surface area or particle contact area, become so great that plastic deformation or local fracture of the particle occurs. The stress state of the particle is always at

some limiting or critical value, and the surface layer in the contact zone will either plastically

deform or fracture. The most general case of a contact load during the compaction process is a combination of normal and tangential loads. The surface layer is always perturbed by the interparticle contact loads. Surface layers with hydrated oxide films are of particular importance to aluminum alloys made by air atomization using RST. These layers must be converted to some stableoxide form before subsequent processing steps can be initiated. During this conversion process, the

volatile components are liberated over different temperature ranges as shown by MSA and TGA studies.

The results are presented in Figs. 7 and 8 for X7091 aluminum powder.

100 200 300 400 500 600 700

TEMPERATURE IN C TEMPERATURE CO

Fig. 7. Plot of ion intensity versus Fig. 8. AM/ AT versus temperature (data

temperature for CT-91 green obtained by TGA).

compact (data obtained by MSA).

Failure to make this conversion will lead to a phenomenon known as blistering which occurs when

the above volatile components are liberated during subsequent thermal exposures. This problem can be

solved for certain aluminum powders by subjecting them to a suitable degassing treatment. In Fig. 8,

degassed and VHP compacts show little liberation of any volatile matter up to 400°C. However, around

500°C, Mg and Zn are volatilized and liberated. Therefore, a proper degassing temperature should be

selected to avoid loss of these elements for this particular alloy. Details on degassing kinetics

and selection of degassing temperature are given in another paper [2] by the present authors.

Nickel-base alloys, in contrast, should not be allowed to form a stable oxide on their surface

because this acts as a site for heterogeneous nucleation of a y' surface for these alloy systems.

Inert powder-making and -handling systems should be used for this class of alloy system.

Fig. 6. Percentage of oxygen versus surface area.

521
CONSOLIDATION

The preform or blank-making process is controlled to a large extent by the powder-making process

and by the physical properties of the individual particles. A powder compact before application of

external pressure consists of many physical particles which are more or less uniform in composition.

The particles range in size from fine to coarse, and their corresponding microstructures range from

microcrystalline to coarse dendritic. At tap density, these powders usually retain their shape, and

they can exert pressure on the surrounding medium. However, they cannot resist tensile stresses.

The compact at tap density occupies a position somewhere between a liquid and a solid. The technological characteristics of this powder, as listed in Fig. 2, depend upon the resistance of the

individual particles to compression, shearing, and bending. The high- temperature RST aluminum and

nickel-base P/M alloys have considerable resistance to these types of loading. The higher the

hardness of the particle and the greater its surface roughness, the lower the tap density of the

powder compact. Correspondingly, the punch must travel a greater distance in the case of low tap

density, which results also in higher tool wear.

The properties of this compact continuously change during pressing. The individual particles

are linked together by different bonds which depend upon the dimensions of the particles. The bonds

are usually between individual particles in contact with other particles and the bonding is due to ' adhesion at the contact surfaces. Particles which are soft and plastic form strong adhesion bonds

during cold compaction, while those which are hard and less ductile do not. A practical comparison

can be made between the cold compaction of X7091 which is a soft, ductile aluminum alloy and the cold

compaction of Al-8Fe-4Ce which is a high-hardness alloy. The latter composition cannot be ejected

from a simple cylindrical mold when cold compacted without severely cracking, while the former

material develops excellent green strength and can be ejected as a cylinder under similar conditions.

The compacting pressure tends to cause the particles to shift and occupy the most stable geometric position; this is accompanied by contact distortion and shearing of the particles. The basic

particle shape tends to remain undistorted. Angular particles, relative to spherical particles, do

not form better green-strength compacts since green strength is principally controlled by the strength

of local adhesion. Vacuum hot pressing or hot isostatic pressing must be used to density a compact

when the particles are hard and resist compression. The aggregate properties of the compacts produced:

by VHP or HIP should be equivalent to the service properties of the components. Otherwise, these

compacts demand subsequent forming operations which are explained in detail in the following section..

In this case, the aggregate properties of the billet should be at least equivalent to the plasticworking properties.

PLASTIC WORKING i

i: Plastic working of P/M preforms is being used increasingly in powder metallurgy as a means of [ shaping and strengthening the finished part. Plastic working has been observed to improve markedly j the properties and microstructures of finished shapes by decreasing the residual porosity, increasing ' the strength of the matrix, and imparting a better surface finish. Additionally, controlled plastic

working of preforms or billets is now being considered as a means of strengthening and increasing

their workability. This additional working is also thought to be essential for controlling the

properties in the finished shape. The results of many empirical observations indicate that a critical

strain is required to overcome the effect of large particles and foreign debris which are introduced

during powder handling. Parts forged directly from an untreated preform or billet often do not

receive sufficient plastic work for production of good second- tier mechanical properties. Failure to achieve sufficient plastic deformation and proper strain distribution often leads to increased

scatter and reduced values of these critical mechanical properties. j

Plastic working is essential for overcoming powder-making defects such as discrete foreign •!

particles and size-distribution-related defects. To perform proper plastic working on P/M material,

it is essential to obtain a good understanding of the material behavior during plastic shaping. This

understanding necessitates construction of a plasticity model for porous bodies which satisfactorily ij describes the experimental results. While the problems involved in the plastic working of P/M }| materials have much in common with those of conventional alloys, a number of differences exist which

| are the direct result of the discrete nature of the individual powder particles. One important

feature of the plastic deformation of a porous body is that it undergoes a permanent change in

volume, which affects its Poisson's ratio [3-4]. Another characteristic is that the mechanical

strengthening experienced by a porous solid during plastic deformation is brought about not only by a change in porosity but also by a strengthening of the matrix. i|

A continuum-mechanics approach is being used for modeling different P/M metalworking processes

using the finite-element method (FEM) for obtaining solutions. This requires a plasticity theory

applicable to porous materials and constitutive equations with density as one of the parameters for i studying densification during deformation, as explained below. I]

522
Plasticity Theory - Porous Material

The design of billet-consolidation processes and of the dies used to forge or extrude porous P/M

materials to full density requires a special yield function for development of the plasticity analysis.

The plastic-flow behavior of porous P/M materials is more complicated than that of ingot materials

because the hydrostatic component of stress influences the onset of plastic flow. The effect of

hydrostatic stress is taken into account by considering a yield function of the form

AJ^ + Bj2 = = 6y2.. (1)

Here J 2 is the second invariant of the deviatoric stress, and Ji is the first invariant or the

hydrostatic component of stress. Yg and Yr are the yield stresses of fully dense and partially

dense materials, respectively. A, B, and 6 are functions of relative density.

Many researchers have determined these constants through heuristic arguments and the use of

experimental results. Green [5] presented an analytical method which considered a uniform cubic

array of spherical voids in a solid under states of stress corresponding to pure shear and hydrostatic compression. The results for these two stress states allowed determination of the two variables A and B. He assumed the stress distribution to be uniform in two directions through the

minimum section of the array and the effect of the voids on the stress in the third direction to be

negligible on the planes midway between the voids. Oyane, Shima, and Kono [6] determined A and B using more stringent assumptions. They found poor agreement between their theoretical and experimental results and, therefore, reported experimentally determined values in a later paper. Shima and

Oyane [7] abandoned this analytical approach entirely and, instead, refined empirical relations

obtained for their experimental values. Kuhn and Downey [8] also presented experimentally obtained

values for these two variables. However, in the present study, the authors derived these variables

taking into account the distortion energy due to the total stress tensor. The final equation obtained

is

(2 + R^)J^ + = (2R^-1) Yq (2)

Theoretical curves along with experimental results for the uni-axial state of stress are shown

in Fig. 9 for X7091 A1 alloy. This figure not only shows good agreement between the theoretical and

experimental results but also explains the influence of hydrostatic stress at various density levels.

When full densification is achieved, the function automatically becomes the von Mises yield function.

Using this yield function, effective stress- and effective strain- relative density relationships

have been derived and compared with the experimental results for the uni -axial state of stress for

the same alloy. Excellent agreement has been observed between theoretical and experimental results,

as shown in Figs. 10 and 11. The theory is also applied to various P/M processes such as frictionless closed-die forging, frictionless plane-strain compression, and hydrostatic compression. These

results are presented in Figs. 12 and 13. The results show that beyond 95 percent of theoretical

density, it is difficult to obtain full densification in hydrostatic and closed-die forging by

application of pressure alone since the pressure required for full densification increases beyond

that obtainable by commercial presses. Therefore, this suggests that pressings should be made at a temperature and strain rate which will accelerate diffusion to permit full densification to be

obtained.

Jl/Yo

Fig. 9. Yield surface in Jl and J2 space. Fig. 10. Relative density versus plastic strain.

523
<

Q.£

UJ

q:K

(n

hiD

a.h

RELATIVE DENSITY Q ^ (\J 0) PLASTIC STRAIN Fig. n. True stress versus relative density

for uni-axial state of stress.

Fig. 12. Relative density versus plastic

strain for different processes.

In order to study the effect of strain rate upon densifi cation during hot working, a constitutive

model has been developed by the present authors for CT 91 (X7091) A1 alloy. Uni-axial compression

tests have been conducted at different temperatures and strain rates, as shown in Fig. 11; and using

these results, the hot-working temperature has been selected. The values of flow stress are plotted

in three-dimensional coordinates as a function of density and strain rate using the computer, as

j shown in Fig. 14. This data base has been used in the model for studying geometric strain hardening ' (densifi cation) during various metal -forming operations for X7091 A1 alloy. ^

Q

O

G

G

d

REL. DENSITY

G

G)G

Fig. 13. True stress versus relative density

for different processes.

Fig. 14. Constitutive behavior model for

X7091 A1 alloy.

I

I

I

I

I

<

t

{

Analytical Modeling

j Several methods are available for the analytical modeling of metal-working processes. However, ‘ the FEM has become popular for the solution of metal-working problems due to the rapid development of i, computers and numerical methods. Many studies have been conducted to solve various metal-working ( problems using this method. These were reviewed and discussed in detail by the present authors in

another paper [9]. In all these studies the von Mises yield criterion has been used since the material

under consideration is always fully dense; however, for P/M applications the influence of hydrostatic'

stress must be taken into account. In the current study the method developed by Kobayashi and Oh

[10] has been used for predicting the metal flow through streamlined extrusion and disk forging since

[

the starting billets are fully dense material. However, a program called ALPID will be modified to : incorporate the yield function presented above for future studies which may consider the use of

| porous billets. From this point of view, the relevant equations required for modification are summarized below. i

524
Equilibrium conditions, neglecting body forces;

i

I a . . . = 0

I

; where a. . is the stress tensor and ,j denotes particle differentiation with respect to j.

I

* ^

I Compatibility conditions:

( 3 )

Mass is constant, # 0

Constitutive relations:

aV, aV,

aV, aV^

^12 ' 1/2

Pi =

e (2 + R^) (1 - R^)

[2(1 - R^) e. + 3R^ ey]

(4)

(5)

Boundary conditions:

“ij "i ' Fj on Sp

Vj = on (6)

and |fs| = frictional stress with the proper sign, where n^- is the unit vector normal to the

surface.

Equations (3) - (6) can be put into the variational principle as

= /

Yp V 3(2 + R^)

2 [(^1 Eo)^ + (eo - En )^ + (e- eJ2] +

1 3(1-R^)

1/2

+ /m;^AV-/Fi V. dS. (7)

Application - Extrusion OF P/M 2024 Al-20 vol% SiC

A new type of P/M aluminum alloy is being developed which incorporates SiC whiskers into the

powder agglomerate for the purpose of increasing the modulus of elasticity and the strength properties

of aluminum alloys. These alloys have many attractive properties, but they are difficult to fabricate

by metalworking processes. A number of applications exist for these materials in aerospace structural

components if they can be extruded or forged into complex geometries without fracturing the whiskers

and changing their aspect ratio. The strength properties are strongly dependent upon this ratio,

while the effective modulus of elasticity is less strongly influenced by it beyond an aspect ratio of

= 12. Figure 15 shows the change in aspect ratio of SiC whiskers as a function of processing steps,

and Fig. 16 shows the dependence of the modulus of elasticity upon the aspect ratio.

The standard shear dies commonly used to extrude aluminum alloys by industry are not satisfactory for extruding this new class of materials. The internal shearing and turbulence of the material

caused by the die geometry fractures the SiC whiskers and forces them to be agglomerated. The

distribution of the SiC whiskers and their fracture can be controlled through the use of streamlined

extrusion dies.

A streamlined extrusion die, by definition, is one whose material velocity profile remains

constant during the reduction from the billet inlet to the die to the material exit. Friction

between the workpiece and the die surface in actual practice prevents "ideal" metal flow from

occurring; therefore, the metalworking problem is to find a practical compromise in die design and in

the process parameters (temperature, strain rate) in order to achieve the goal of approximating

streamlined flow. Through process modeling this goal can be achieved with a minimum number of

trial -and-error iterations. The rigid-viscoplastic FEM discussed above was used to simulate the

i

525
process on the computer and to select the proper die geometry and processing conditions. The die

geometry was modeled by means of a method developed by Gunesekera, ^ a]_. [9]; and this model, in

turn, was used to formulate the FEM model.

Fig. 15. Aspect ratio versus various Fig. 16. Effective modulus versus aspect

processing steps. ratio.

Constitutive relations were developed for the 2024 Al-20 vol% SiC whisker material by compression

testing small cylindrical specimens as a function of temperature and strain rate. A M0 S 2 lubricant

was used to assure homogeneous deformation at a strain rate around 0.50_e and temperatures of

400°C to 550°C. The results are shown in Fig. 17 for a strain rate of 10"2/sec. Information

concerning the billet-processing steps will now be introduced in order to provide a fuller appreciation of the importance of billet conditioning. The 2024 Al-20 vol% SiC whisker material is first

properly degassed to remove the water attached to the powder surface layer and blended to achieve

good mixing of the powder particles and whisker material--an important step. The mixture is then : vacuum hot pressed to a density of 97% or greater. The material workability after vacuum hot '

; pressing is insufficient for subsequent metalworking operations such as forging and complex extrusion, j Therefore, this billet is thermomechanical ly processed by streamlined extrusion to improve the

workability. The constitutive data shown in Fig. 17 were obtained from this material. The SiC

whiskers in the powder compact are essentially random or possess discrete random orientation and i become oriented in the plane normal to the compressive axis in accordance with the laws of classical t mechanics. However, the SiC whiskers after being extruded again develop a discrete random orientation,

somewhat similar to that developed after the blending step. The fiber orientation after this billet-^-

conditioning step is determined by the material point velocity-field. The area-reduction ratio If should be based upon the degree of fiber orientation desired. The random orientation of the SiC ; whiskers is desirable for additional hot working by either extrusion or forging. i

The die length for approximating streamlined flow was determined from the FEM analysis using the ;i constitutive relations developed for the 2024 Al-20 vol% SiC whisker material. A die length ;| was selected, and a die was manufactured. Another, shorter die was produced in order to approximate i the metal-flow pattern of a shear die. The point-velocity field for a typical shear die, a short

conical die, and the streamlined die is shown in Figs. 18 and 19. Analysis also shows that increasing,

the die length beyond the optimum length provides no additional benefit. Isothermal extrusion at 475°C (880°C) was carried out for both the optimum-length streamlined die and the short die.

Microstructures for the two cases are shown in Figs. 20 and 21. The optimum-length streamlined

die produced the desired effect. The SiC whiskers were not fractured and were oriented parallel to the axial direction. The short die caused the SiC whiskers to fracture in the regions where

turbulent flow occurred. This is in agreement with the velocity- field predictions in Fig. 15 for

these dies. It should also be noted in Fig. 15 that the 180-deg. shear die produces a more turbulent:

flow condition.

This case history clearly shows that process analysis is a valuable tool in determining the

proper processing conditions and die geometries for producing a finished shape having controlled

| microstructures and properties. 'f

526
TRUE

STRESS

CMPA>

TRUE PLASTIC STRAIN

Fig. 17. True stress versus true strain curves Fig. 18. Velocity field during extrusion

for 2024 Al-20 vol% SiC. using shear and conical dies.

Fig. 19. (a) Effect of die length upon metal flow during streamlined extrusion (extrusion ratio 9:1)

(b) Effect of extrusion ratio upon metal flow [streamlined die length, 3 in. (76.2 mm)l.

Application - Design of an Advanced Turbine Disk Having Superior LCF Properties

Considerable service experience has been accumulated by the aerospace industry to conclude that

the most probable site for LCF failure initiation is at the interface of large powder

particles which are essentially undeformed. The microstructures for nickel-base alloys (Alloy-X) hot

worked by extrusion [11], high-temperature aluminum (Al-8Fe-4Ce) alloys after forging [12], and REP

Ti 6-4 alloy after HIPing and heat treatment [13] have many particles which remain undeformed, even after severe plastic deformation, as shown in Figs. 22-24. Metallurgical examination clearly

indicates that these particles are much harder than those of the continuous matrix material. Since

both of these alloy systems derive their high- temperature strength from intermetallic compounds, it can be deduced, that these powder particles contain intermetallic compounds whose hot hardness

decreases more slowly with increasing temperature than that of the matrix material. Thus, these

particles behave as rigid bodies and flow in the same pattern as the surrounding material. Figure 18

shows that the metal -flow pattern can cause these rigid bodies to accumulate in certain areas to form

527
Fig. 20. Microstructure of extruded product

through streamlined die of 3-in.

length.

Fig. 21. Microstructure of extruded product

through streamlined die of 2-in.

length.

a larger site for LCF failure initiation. Therefore, it can be conjectured that conventional

extrusion with shear-type dies is not effective in completely breaking up the coarse, hard particles

and that it can, in fact, cause larger failure sites to be created by the metal -flow pattern. This

observation suggests the following approach for improving P/M materials.

Fig. 22. RSR nickel -base alloy after

extrusion.

Fig. 23. Al-Fe-Ce alloy powder after forging.

Hot isostatically pressed material should be thermomechanical ly treated by extruding the

billet material through streamlined dies. This step will improve the workability and strength of

the workpiece material while maintaining a favorable distribution of the larger and harder coarse

particles. The purpose of this step is to change the nature of the flow curve from one which

exhibits flow softening to one which does not. An example of this behavior is shown in Fig. 25

for the case of Ti-6242 alloy. The significance of this material behavior, as exhibited by the

titanium alloy, is that the two curves approach each other at a true strain of 1.0. When the curves touch, both specimens have essentially the same microstructure and improved hot workability.

Once this material behavior has been achieved, the material can be hot worked at a lower temperature.

The streamlined extrusion die required to produce this effect, i.e., the area-ratio reduction, can be

determined by this test. For the case of Ti-6242, the extrusion ratio would be 'v- 9:4, which would

produce an effective strain of -v- 1.0. The actual extrusion ratio is thus based upon the effective

true strain after the completion of work softening.

528
I

!

Fig. 24. REP Ti 6-4 alloy after HIPing and

heat treatment.

Fig. 25. Flow curves for 6 and a + S Ti-6242 alloy.

Fig. 26. Effective strain contour during Fig. 27. Ti disk forging,

streamlined extrusion.

After thermomechanical treatment the material can be used as the preform for the hot die or the

isothermal forging process required to produce the disk. Designing the forging process requires an understanding of the manner in which shear stresses and strains are partitioned from the matrix

material into the harder, coarser particles remaining after billet conditioning. Recalling the

process for extruding the SiC whisker material described above, the metalworking problem with titanium

and nickel -base alloy systems is the reverse. In order to prevent the SiC whiskers from fracturing,

it was necessary to extrude the material at 475®C—just below the solidus temperature where the flow

stress of the matrix material is significantly less than the fracture stress (approximately equal to

its yield strength) of the SiC whiskers. However, it is necessary to increase the flow stress of the

matrix material of the P/M nickel -base alloys such that it is greater than the flow stress of the

coarse, dendritic-like powder particles. In this way, the matrix material can .partition energy into

the coarse powder particles. This situation can possibly be achieved by thermomechanical ly conditioning the as-HIPed billet in the manner previously described. The increased workability which

accompanies billet conditioning allows the forging process to be carried out at a lower temperature

where the flow stress of the matrix phase is close to that of the coarse particles.

529
The effective strain distribution for a streamlined die which was designed for conditioning these i RST P/M materials is shown in Fig. 26. This uniform effective strain distribution assures that: (1)

all material will receive the same amount of plastic work in a practical sense, and (2) the larger

and harder particles will not cluster into a larger failure site. The FEM model allows the extruded

product to be fully characterized with respect to its processing history. This processing history in

terms of its plastic strain can be transferred to the FEM model for forging the disk.

Figure 27 shows a typical simulation of an isothermal disk-forging process. The strain distribution can be easily observed by the grid distortions and also by plotting the effective strain contours.

Other field quantities such as the temperature or velocity can be displayed, since they are always

computed by the model. However, it is necessary to display only the field quantity which is pertinent

to the process-modeling objective.

Different preform geometries can be tested in the computer by a series of iterations, and the

optimum preform geometry which yields the effective strain distribution and microstructure can be

selected. Immarigeon [14], for example, has developed a constitutive-equation form which allows both

the effective strain and grain-size distribution to be predicted for the 713 nickel -base alloy.

Process modeling allows the designer to be innovative and to create new processing possibilities.

For example, he can address the possibility of using process control as another means of guaranteeing | quality assurance, thus permitting him to forge closer to net shape beyond the sonic envelope.

Plasticity analysis also can be used to gain new insight into the goals of powder making. The

current goal of powder making appears to be biased in the direction of producing mostly fine powder

particles having a microcrystalline microstructure. The justification for this objective is based

upon the observation that (1) this type of microstructure responds differently to heat treatment than

the coarser, dendritic microstructures, and (2) 100% microcrystalline powder particles do not contain

any sites (other than possibly foreign ceramic particles) for LCF failure initiation. However,

a study by Buzzanell and Lherbier [15] to investigate specifically the effect of varying structural

conditions and powder size upon the LCF behavior of these materials showed that -60 mesh powders

exhibited significantly better LCF behavior than powders having microcrystalline microstructures.

This consistent observation in the LCF behavior could not be explained. Their materials were con- solidated by CAP-plus-isothermal forging.

The micromechanical plasticity analysis in the present study suggests that one should consider

producing powders which have a constant microhardness from particle to particle in the powder compact.

Practically speaking, it is unlikely that 100% microcrystalline material can be produced, implying

that a certain volume fracture of LCF failure sites will always be present in the form of coarse, dendritic particles. Manipulation of the powder-making process to control the particle-size distribution such that uniform microhardness values are produced should be a goal in addition to that of i producing 100% microcrystalline material. The most probable explanation for the coarse particles ! producing significantly better LCF properties than the microcrystalline material [15] is the manner

in which the load during forging is transferred from particle to particle. A rigid matrix can transfer stress and strain homogeneously, making it possible to reduce the defect size below a critical value. j'

CONCLUSIONS

1. The plastic working of P/M material represents an increasing trend toward improvement of the

service properties of components.

2. A yield function of the form (2 + R^) (1-R2)/3 = (2R2-1)y2 is required for the

development of plasticity theory and analytical process modeling of porous materials.

3. Plasticity theory for porous materials suggests that beyond 95% densification, application of

pressure alone is insignificant in the case of hydrostatic and repressing processes. To achieve

full densification, proper control of temperature and material is required.

4. Analytical process modeling provides the ultimate capability of achieving the full advantages of

materials produced by RST.

5. More attention should be given to understanding the powder-making process in order to permit the

size distribution and the average quenching rate to be adjusted systematically for controlling

the physical properties of the powders.

6. The goal of achieving a uniform chemical distribution should be extended to the control of the

microhardness distribution. The latter can have a profound effect in controlling the LCF

behavior of a P/M material.

530
ACKNOWLEDGEMENT

The authors would like to thank Mr. A. M, Adair, AFWAL/MLLM, WPAFB, OH, for providing many

technical reports which were useful in writing this paper. Drs. Gunasekera and Doraivelu wish to acknowledge the NRC-AFSC Research Associateship Program and Universal Energy Systems, Dayton, OH, for

financial support during the course of this investigation. Thanks are also due to Mrs. M. Whitaker

and Mrs. H. L. Henrich (Systems Research Laboratories, Inc.) for editorial assistance and preparation

of this manuscript.

REFERENCES

[1] G. J. Hildeman, D. J. Lege, and A. K. Vasudevan, "Fundamentals of Compaction Processes for

Rapidly Quenched Prealloyed Metal Powders," Alcoa Technical Center, PA, Oct. 1982.

[2] J. T. Morgan, H. L. Gegel , S. M. Doraivelu, L. E. Matson, I. A. Martorell, and J. F. Thomas,

"Consolidation of Metal Working Preform of X7091 (CT-91) Aluminum Powder," presented at

the TMS-AIME Annual Meeting, Dallas, TX, Feb. 1982.

[3] G. M. Zhdanovich, "Theory of Compacting of Metal Powders," Translated from Teorize

Pressovaniye Metzillchaskikli Poroshkov, pp. 1-262, 1969, by The Foreign Technology Division,

WPAFB, OH, June 29, 1971 (Translation No. FTD-HCC-23-775-70)

[4] H. A. Kuahn, "Deformation Processing of Sintered Powder Metals," Chap. 4 in Powder Metallurgy

Processing: New Techniques and Analysis (Academic Press, NY, 1978), P. 99.

[5] R. J. Green, Int. J. Mech Sci. 14, 215 (1972).

[6] M. Oyane, S. Shima, and Y. Kono, Bull. JSNE 16, 1254 (1973).

[?! S. Shima and M. Oyane, Int. J. Mech. Sci. 18, 285 (1976).

[8] H. A. Kuhn and G. L. Downey, J. Eng. Mat. Tech. 41 (1973).

[9] J. S. Gunasekera, H. L. Gegel, J. C. Mai as, S. M. Doraivelu, and G. Griffin, "Computer Aided

Engineering Approach to Metal Forming," in Proc. 2nd Int. Computers in Engineering Conf.,

Vol. 1 (ASME, NY, Aug. 1982), pp. 1-6.

[10] S. Kobayashi , "Thermoviscoplastic Analysis of Titanium Alloy Forging," AFWAL-TR-81-4130 (Air

Force Wright Aeronautical Laboratories, Wright-Patterson AFB, OH, 1981); S. I. Oh, "Finite

Element Analysis of Metal Forming Processes with Arbitrarily Shaped Dies," Int. J. Mech. Sci.

(to be published).

[11] J. A. Miller, "RSR Powder Processed Blades and Vanes," PWA/GPD FR-16247 (DARPA, Arlington, VA,

30 Apr. 1982); A. Adair (AFWAL/MLLM), WPAFB, OH 45433, Private communication, 1982.

[12] S. D. Kirchoff and W. M. Griffith (AFWAL/MLLS) , WPAFB, OH 45433, Private communication, 1982.

[13] P. R. Smith, (AFWAL/MLLS), WPAFB, OH 45433, Private Communication, 1982.

[14] J. P. Imangeon, D. Murphy, and A. K. Koul, "Plastic Flow Behavior of Powder Processes

Nickel-Base Superalloys During Isothermal Forgings," presented at the Metal Congress

during TMS-AIME Annual Meeting, St. Louis, MO, Oct. 1982.

[15] J. P. Buzzanell and L. W. Lherbier, "Processing Effects on the Properties of P/M Rend Near-Net

Shapes," in Proc. 4th Int. Symp. on Superalloys (ASM, Metals Park, OH, 1980), pp. 149-158.

531
PROCESSING OF P/M ALUMINUM ALLOYS X7090 AND X7091

F. R. Billman, J. C. Kuli, Jr., G. J. Hildeman, J. I. Petit and J. A. Walker

Alcoa Laboratories, Alcoa Center, PA 15069

ABSTRACT

The role of various processing sequences during both the consolidation of P/M

aluminum alloy billets, from rapidly quenched (^10^°C/sec) commercially atomized

powders, and their subsequent fabrication to wrought product shapes is discussed.

Beginning with effective techniques for hot degassing of initial powder particle

surfaces, the judicious usage of common time- temperature-deformation parameters is the

key to development of premium quality aircraft structural materials.

Hot die consolidation of the original powder into the intermediate billet stage

introduces shear-type deformation strains within an evacuated porous metal compact so

as to simultaneously achieve 100% density and intimate particle-to-particle bonding.

Residual porosity and improperly dispersed powder particle surface films act as impediments to metallurgical bonding, causing affected billets to exhibit poor

ductility and fracture toughness. Metal temperatures must be chosen with due

consideration to: powder surface outgassing characteristics, elevated temperature flow

stresses, and each P/M alloy's sensitivity to microstructural coarsening.

Once properly consolidated, P/M billets are readily fabricable by conventional

aluminum alloy extrusion, forging and rolling processes. Exemplary influences of

various extrusion process variables will be described to demonstrate X7090 and X7091

P/M alloy structure/process/property correlations.

INTRODUCTION

Premium Properties from Preferred Structures

The successful development of commercial 7XXX type aluminum alloy powder metallurgy (P/M)

products for high strength, "high technology" applications is based on the metallurgical premise

that premium properties are achieved through the controlled evolution of preferred microstructural

features. Such preferred microstructures are usually finer and more homogeneous than can be

developed in competitive ingot metallurgy (I/M) products. Metallurgical characteristics of P/M

products are developed through control of the initial powder particle solidification process, the

thermal and deformation practices used to consolidate the powder into the final wrought component

shape, and final heat treatments used to achieve the desired balance of strength, toughness and

corrosion resistance.

A schematic overview of Alcoa's commercial aluminum alloy Wrought P/M Process is shown in

Figure 1. Process sequences have been developed over the past 33 years, with careful attention

being given to the interdependence of specific alloy compositions and process parameters. [Refs.

1-16] This presentation will focus primarily on manufacture of heat treatable 7XXX type P/M

extruded products to describe current commercial P/M structure/process/property correlations.

532
Table I compares nominal chemical compositions of current 7XXX type P/M alloys with some of

their I/M counterparts. Alcoa's commercial X7090 and X7091 alloy compositions and three

experimental compositions being studied by Kaiser are shown. In contrast to the I/M alloys, P/M

compositions usually incorporate alloying additions of cobalt to yield Co pAl „ type dispersoid

particles within the aluminum alloy matrix. Chromium, zirconium, manganese, ^iron, nickel and

other dispersoid-forming ingredients can also be readily utilized with or without cobalt in P/M

alloys; often in amounts significantly greater than is feasible for commercial ingot products.

[Refs. 9, 10, 12, 13]

Faster Melt Quenching Provides Finer As-Cast Structures

The ability to effectively utilize novel alloying additions and extended amounts of

conventional I/M alloying ingredients through increasingly rapid melt quench rates has been widely

explored and publicized. [Ref. 17-23] A generalized relationship for aluminum alloys,

illustrated here by Figure 2, predicts the capability for achieving progressively finer

microstructural features with increasingly rapid liquid -> solid melt quenching techniques. The

plot shown here intentionally suggests overlap between coarse and fine powders and between fine

powder and splat-quenched materials (flakes, ribbon, wire, etc.). Specific quenching rate

boundaries of these rapid solidification processes, particularly the various powder atomization

methods, can vary quite significantly as a function of atomized droplet size, the manner in which

such droplets are created and the relative amount of undercooling achieved within the molten

droplets

Convectively quenched droplets (i.e. Alcoa's gas atomized and gas quenched powders or PWA's centrifical ly atomized and gas quenched powders) of a given size will exhibit significantly finer

microstructural features in a given alloy than will similarly-sized powders produced in a closed

batch-type atomization chamber wherein a relatively small volume of gas is used to achieve the

dual function of droplet formation and droplet quenching. For a given alloy composition,

beneficial effects of droplet undercooling on the as-cast powder particles have been reported.

[Refs. 24-27]

Recent investigations have also shown that measured dendrite cell sizes, for a given melt

quenching process, can vary significantly as a function of alloy type and composition. [Refs.

28-32] Also, the relative melt quenching rate actually obtained in both atomized powder and

splat-quenched materials is strongly influenced by conditions at the mel t/quenchant interface

where surface films and static gas layers will inhibit optimum heat transfer. Because of these

alloy/process quench rate/structure complexities, plots like Figure 2 can only serve as a generalized illustration of the metallurgical incentive for using rapidly solidified powder

relative to conventionally cast ingots.

Fine Powder Structures Retained in Wrought Products

Figure 3 contrasts the measured sizes of three metallurgical features (grain size, dendrite

cell size and interparticle spacing) in 7075 alloy extrusions manufactured from two grades of gas

atomized powder and from two sizes of conventionally cast direct chill (DC) ingot. As-cast

structural advantages of the original, more rapidly solidified powders have been retained through

the various billet consolidation and wrought product manufacturing sequences. [Ref. 33] Under

such controlled processing conditions, the relative sizes of other microstructural features in

comparable I/M and P/M products are shown in Table II. Both I/M and P/M materials are fully dense

in their commercial wrought product forms as a consequence of using sufficient hot deformation to close and heal original ingot shrinkage porosity and initial P/M compact or billet porosity. Hot

deformation effectively fragments and disperses prior powder particle surface oxide films in a manner rather analogous to the refinement of as-cast ingot grain boundaries.

RAPIDLY SOLIDIFIED POWDER PRODUCTION

Atomized powders exhibit size distribution ranges according to the specific (and usually

proprietary) atomizing practices used, as illustrated in Figure 4. [Ref. 34] Similar size

distributions in both irregular and "regular" powder shapes can often be achieved by gas

atomization techniques. Less reactive atomizing gas mixtures (i.e. those having significantly

lesser amounts of air, oxygen, moisture, CO 2 . etc.) enable the newly atomized droplet to reach a "regular" and more energetically favorable shape before the outer oxide skin is sufficiently

established to fix the particle morphology, whereas surface oxide formation in more highly

533
SCHEMATIC

FLOW

CHART

OF

ALCOA'S

COMMERCIAL

ALUMINUM

ALLOY

WROUGHT

P/M

PROCESS

TABLE

O

m IT) m Lo LTJ

534

GENERALIZED

RELATIONSHIP

FOR

ALUMINUM

ALLOYS
1

)

reactive atomizing gas mixtures occurs so rapidly as to better preserve the highly irregular

1 molten droplet shapes initially created by the atomization process itself. [Ref. 35]

' All Aluminum Powders Have Surface Oxides

1 The generation of an aluminum powder particle without the presence of a continuous surface

oxide film is extremely improbable in any commercially feasible environment since even the solid

I state o_xidation reaction to form AI2O3 thermodynamically requires an oxygen partial pressure of

only 10"®^ atmospheres. [Ref. 35-37] Aluminum alloys containing Mg or Li are even more reactive.

i For a given type of alloy, and a given type of powder, air-atomized irregular Al-Mg-X powder

for instance, oxygen content is a linear function of the mean particle diameter as shown in

j Figure 5. "Regular" and truly spherical powder shapes will still exhibit this linear

' relationship, but their data bands will be displaced to lower total oxygen contents as indicated.

I The reason for such displacement is suggested by the Figure 6 plot of oxygen content versus powder

I surface area for an Al-Ni-Fe alloy composition where the three significantly different types of

I rapidly solidified particulate are characterized with regard to their respective surface areas.

' [Ref. 38-39] The linear relationship in Figure 6 implies a similar oxide thickness and

composition for all three types of particulate produced in the same alloy composition. Expected

! shifts in this relationship for different alloy compositions having either more or less reactivity

I with the atomization and quenching environments are noted.

i Pertinent characteristics of some aluminum alloy powders are given in Table III. Additional

data on the three Al-Ni-Fe particulates is included here to verify that the oxide film thickness

is, in fact, virtually identical for all three types. Note also the measured differences in

average dendrite cell sizes. In dispersion-hardened nonheat treatable aluminum alloys, the

progressively finer and more homogeneous microstructural features associated with increasingly

rapid melt quenching rates often provide direct strength and ductility improvements in the final

' wrought product. As an example, room temperature tensile properties of forgings identically

fabricated from the three different Al-3.7Ni-l.5Fe powders described in Figure 6 and Table III were reflective of their original melt quenching rates:

Most Rapid ^ Alcoa Splat - 350 MPa UTS, 276 MPa YS, 16.2% EL, 38% RA; Rapid -> Alcoa Irregular Powder - 335 MPa UTS, 244 MPa YS, 15.2% EL, 34% RA; Least Rapid PWA RSR Powder - 295 MPa UTS, 214 MPa YS, 10.2% EL, 16% RA. [Ref. 38-39]

For heat treatable 7XXX aluminum P/M alloys any initial microstructural advantage of

splat-quenched materials over fine, irregular powders can be substantially normalized during the

inherently high temperature billet consolidation sequences and/or during final solution heat

treatments. Further discussion of commercial 7XXX P/M products is therefore directed toward the

use of atomized powder. The remainder of the loose powder data in Table III characterizes 7XXX

type alloy powders produced by Alcoa, Reynolds and Kaiser using a variety of gas atomization

techniques. (Refs. 34-35, 38-40] Average powder sizes and total bulk oxygen contents vary widely

as a consequence of the atomizing nozzle parameters and gas environments used; but, oxide

thickness measurements are relatively comparable. Similar data would be expected for 7XXX alloy

sperical RSR powder.

Outgassing of Aluminum Powder

In addition to their surface oxide films most, if not all, aluminum powder surfaces also

contain hydroxides. For Mg and Li bearing alloy compositions, reaction with even trace amounts of

moisture (i.e. - 90°F dewpoint gas environments) is a virtual certainty. When either pure

aluminum or aluminum alloy powders are heated, gaseous by-products are evolved. Adsorbed surface

moisture can be vaporized and chemically combined hydroxides can decompose. Figure 7 is a generalized illustration of the principal outgassing events possible. [Refs. 8, 15, 40-56] The

exact reaction mechanisms have not been fully defined; but a growing appreciation of the various

possible by-products is being evidenced in the technical literature.

Discrepancies among the various investigations with regard to whether hydrogen or water vapor

is the predominant outgassing species are believed to reflect variations in (1) powder or compact

density and cross-section, (2) heating rate and thermal gradients across the compact cross-section

and (3) differences in the relative degassing efficiencies of the various flowing inert gas and vacuum systems used. The total amount of gas liberated should be a function of the powder alloy,

the powder particle surface area, the powder atomizing and handling environments and the maximum

preheat/degassing temperature used. Zn and Mg vaporization losses do occur at the highest solid

535
100.0

536

Mean

particle

diameter

(microns)
Pertinent

Characteristics

of

Some

Aluminum

Alloy

Powders

^ e — c o «H 42

angstroms

30 35

50-100 50-100

1 1 I 1 87 123 150 101 102

FNA

Oxygen

Content'*^

i

0.020

wt.%

j

0.139 0.366 0.230 0.369

1

0.204 0.295 0.145 0.063

0.056

1

0.55 0.50 0.42 0.16 0.22

Dendrite

Cell

Size

—1

micron

<0.5 -0.5

1 1 1 1 1 1 1

CM « O m

9^ 1 CM CT)

Particle size'”

77

microns

median

-6/+200

mesh

31

microns

median

15

microns

APD

13 26 24

microns

APD

27 115 116 1

1

microns

Linear

APW

13 16 29 61

Powder Type

PWA-RSR

spherical

Alcoa

splat

flakes

Alcoa air

atomized

Flue

gas

atomized'*^

Air

atomized

Air

atomized

Air

atomized

Argon

atomized

Air

atomized

Argon

atomized

Alcoa

X7061

Air

atomized

Reynolds

X7091

Air

atomized

Kaiser

MR64

Flue

gas^*’

atomized

Alloy

Al-Ni-Fe (Refs.

36-36)

Al-Zn-Mg-Cu

X7091 by

Alcoa

(Ref.

35)

Al-Zn-Mg-Cu

By

Alcoa

(Ref.

34)

3o

* §

c

'll ^

< B

> U 3 u. </> O • • • •

537
state degassing temperatures, but result in only very modest losses at the extreme compact

surfaces unless very long, hard vacuum exposures are used.

CONSOLIDATION OF FULLY DENSE BILLETS

For both safety and process efficiency reasons Alcoa preconsolidates its fine aluminum alloy

powders into a 70-80% dense compact using conventional cold isostatic pressing technology. These

compacts are then preheated for degassing prior to vacuum hot pressing into 100% dense billets.

Compact Preheating and Degassing

Table IV outlines the important features of an efficient P/M compact degassing and P/M billet vacuum hot pressing methodology. Figure 8 illustrates how a 7XXX alloy preheat temperature versus

time cycle can vary as P/M compacts increase from <25 cm (10 in.) in diameter to >76 cm (30 in.)

in diameter. For the smaller compacts, typical of Alcoa's commercial 50 kg (110 lb) P/M billets,

a dynamic vacuum heatup and 1 hour soak cycle can be achieved in 7 hours. For the much larger

compacts necessary for consolidation of ^ 1545 kg (3400 lb) P/M plate billets heatup is accomplished more rapidly under Ar atmosphere (static or hot flowing argon) with a 1 hour

evacuation period at the end of the 'X/30 hour cycle. Actual soak times at the o<510°C (950°F)

maximum degassing temperature will vary appreciably from '^24 hours at the outer periphery of a large compact to only 1 hour at the compact's center. [Ref. 16, 57] Optimum degassing

temperatures should be both below a particular alloy's solidus temperature and in a time- temperature regime where microstructural coarsening is minimized. For nonheat treatable

aluminum P/M alloys, degassing temperatures are usually _<399®C (750°F). [Refs. 30, 39, 58-60]

Vacuum Hot Pressing the Billet

Alcoa's standard P/M billet process involves the use of a welded aluminum alloy canister to maintain a relatively gas-free environment within the interconnected pores of the 70-80% dense

compact while it undergoes hot deformation to full density. Canister wall thickness, welded

end-cap designs and compact/canister/cylindrical die geometries must be matched to preclude

rupture of the deforming canister and loss of vacuum.

Recently Alcoa has also developed several canless techniques for vacuum hot-die consolidation

of P/M billets or fully dense forged preforms, depending on the size and shape of the hydraulic

press die cavity used. Simple P/M compacts are first preheated and outgassed in a dry argon

furnace and then quickly loaded into the appropriate hot die cavity. Special vacuum sealing

methods permit the evacuation of the hot die chamber and the hot P/M compact within a 5-15 minute

period, following which the press ram or punch is continuously advanced so as to deform the

canless hot compact into a 100% dense article. Three process variants are possible: Canless P/M

Billets can be consolidated and then ejected from a tapered-diameter cylinder for subsequent

fabrication into extruded or forged components; canless billets can be consolidated in a straight-walled cylinder followed by immediate substitution of an extrusion die for the blind die,

and renewed ram motion to fabricate a Direct P/M Extrusion ; and fully dense forged preform shapes

can be produced from simple P/M compacts in the Direct P/M Forging process. [Refs. 61, 62] As in

Alcoa's standard canned billet process, the introduction of significant amounts of shear-type hot

deformation strains within the evacuated porous metal compact serves to fragment and disperse

original powder particle surface oxide films during the densi fication process.

The necessity for preventing long time, high temperature exposures to either nitrogen or air

has been confirmed in these Direct P/M processes. Both nitrogen and air react with the hot P/M

compacts to form nitrides and oxides along the channels of interconnected porosity. Consequently,

fracture toughness levels in the subsequent wrought products can be degraded by 50-70%. [Ref.

61-62]

Toughness as Billet Quality Indicator

In Alcoa's experience the room temperature notch tensile/yield strength ratio indicative of

fracture toughness is a sensitive evaluator of 7XXX P/M billet quality. Thus P/M billet slices

are routinely taken, heat treated to the appropriate -T7 temper and both smooth and notched

tensile tests are performed. The resultant NTS/YS ratio is used, along with other room temperature tensile properties, as ACCEPT/REJECT criteria for billet release. NTS/YS and other

typical characteristics of Alcoa's X7091 and X7090 alloy billets, in sizes of 27 kg (60 lb) to

538
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I—

CO

a:

<Ji O

O

X ^

00

—

Q- < O

uj

O

cr o ^< o>

CM

.— > q;

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539

Comparative

Small

vs.

Large

Billet

Process

Simulations

(Refs.

16,

57)

Figure 6
1545 kg (3400 1b), are listed in Table V. Bulk oxygen contents of both the original powder and

resultant billet materials are given here to demonstrate the reduction in total oxygen content

achieved through removal of H 2 O vapor during the compact preheat/degassing step. [Refs. 16, 40,

57, 61, 63] Although an oxygen content of <0.5 weight percent has been specified for X7090 and

X7091 wrought products (see Table I), there is no direct correlation between oxygen content and

billet or product properties.

Bulk hydrogen contents also decrease from >30 ppm in fine atomized 7XXX alloy powder to <10

ppm in the billet. With the exception of the extremely low hydrogen contents (<.2ppm) reported by

Voss for his 7075 alloy P/M extrusions fabricated from small laboratory compacts which had been

rigorously degassed in a 90 . hour cycle incorporating relatively hard vacuum levels (10~^ Pa),

other investigators invariably measure hydrogen contents in P/M billets and mill products which

are 10-100 times more than is considered acceptable in a 7XXX ingot. [Refs. 52, 53] Part of this

P/M versus I/M difference is due to the fact that only gaseous hydrogen is believed to exist in

commercial quality ingots whereas high temperature hydroxides (and perhaps other chemically

combined species) do reside in P/M materials.

As an example, three adjacent metal samples were taken from the 1545 kg (3400 lb) billet

material described in Table V. Two samples were analyzed by the standard 800°C ITHAC fusion

extraction technique which resulted in total hydrogen gas evolution from the molten alloy of 5.41

and 5.81 ppm. The third 9 mm diameter x 40 mm long sample was subjected first to hot vacuum solid

state extraction for 4-1/2 hours at 500°C to yield 1.62 ppm of evolved gas, and was then heated to 800°C liberating an additional 3.91 ppm of gaseous by-products. Total gas liberated from this

third sample was 5.53 ppm -- in substantial agreement with the other two fusion determinations.

The conclusion to be drawn from such data is that a large portion ('^^70% here) of the total gas

evolved during fusion extraction of aluminum P/M materials probably does not exist in a detrimental gaseous form in the solid state article. [Ref. 57] This may help to explain why P/M

billet and P/M product properties, including fracture toughness, seem relatively insensitive to fusion extracted hydrogen levels <10 ppm.

FABRICATION OF P/M EXTRUSIONS

Billet Properties Improved by Additional Hot Deformation

Although the vacuum hot pressed P/M billets are 100% dense and relatively well bonded their

strength, ductility and toughness levels are all improved with additional hot deformation as represented schematically in Figure 9. For a given 7XXX alloy, powder type and billet

size/configuration, increasing amounts of hot deformation strain serve to develop relative

"property plateaus". Ultimate tensile strengths seem to stabilize first, then the tensile yield

strength and ductility parameters, and finally the NTS/YS or Kj^ toughness values. [Refs. 57, 64]

Billet to Product Quality Correlation

With controlled thermal and deformation processing, a billet-quality to extruded productquality relationship can be demonstrated using their respective toughness indices as shown in

Figure 10. Excellent extruded P/M product properties have been verified for all three Alcoa

processes (Standard scalped canned P/M billet extrusion. Canless P/M billet extrusion. Direct

P/M billet/extrusion) as shown in Table VI.

Strong Toughness versus Strength and Orientation Relationships

Variations in good quality billet and extruded product toughness levels usually reflect minor

shifts in chemical composition within a given alloy's composition limits. As shown by Figure 11,

a strong relationship exists between room temperature toughness and tensile yield strength values

in 7XXX P/M products. Note also the similar correlation but displaced toughness versus yield

strength trend for the three different test orientations within the same extruded bar shape. Such

reduced toughness levels in the transverse planes are also common to 7XXX I/M products and are reportedly associated with elongated grain shapes, grain boundary orientations and deformation

textures. [Refs. 65-75] Precise resolution of metallurgical structure in P/M products is most

challenging because of the extremely uniform and small features present, but fracture surfaces

within good quality wrought products exhibit ductile-dimple failure in all three planes.

540
Influences of Dispersoid Coarsening and Extrusion Parameters on Toughness

In a recent Alcoa study, billets processed by all three preheat/degas cycles shown in Figure

8 were extruded into three different bar shapes at three different extrusion billet temperatures.

[Ref. 57] Resultant room temperature tensile properties of a given alloy were all measured to be

virtually constant for all extrusion conditions but NTS levels were influenced as shown in Table

VII. NTS in the long- transverse (LT) test orientation were most affected in both alloys, but

longitudinal NTS values in X7090 alloy extrusions could also be significantly influenced.

Extrusion ratio and extrusion billet temperature were the most influential variables.

C 02AI 9 dispersoid coarsening in the richer X7090 alloy (1.5 weight percent Co) was measured

as a consequence of the 24 hour soak at 'v510°C (950°F) cycle and NTS values were decreased 6-7% in

both L and LT failure planes as a consequence of increased amount of C02AI 9 particles in the

0.2-0. 8 micron size range. Earlier Alcoa investigators reported strong toughness degradation with

C 02 AI 9 particle coarsening in the 0.4-1. 4 micron size range. [Refs. 11, 28]

Figure 12 provides a visual illustration of the influences of both extrusion ratio and

extrusion billet temperature on the NTS/YS toughness indices of X7090-T7E71 extruded bars. Yield

strength was essentially constant for all conditions. Extrusion ratios >12:1 and 589K (600°F)

billet temperatures provided the best L and LT toughness characteristics.

Actual Kj compact-tension type fracture toughness tests were performed on 12:1 extrusion

ratio products^ fabricated from P/M billets having been preheated/degassed with the short (1 hour)

and the long (24 hour) soak cycles at 'v510°C (950°F). As shown in Figure 14 the usual NTS/YS

versus K. relationship exists for all processing conditions except where toughness was degraded

by coarsened C 02AI 9 particles present in X7090 alloy. The benefits of lower billet extrusion

temperature are dramatically evidenced for both alloys. And, the relative insensitivity of NTS/YS

to extrusion practice variables in X7091 is explained by the very shallow slope of the NTS/YS

versus K. toughness curve for NTS/YS ratios ^1.3. In such high toughness regimes, the material's

actual fracture toughness performance must be judged on the basis of results.

Although this same study also included the use of several strain rates (ram speeds) for each

set of extrusion conditions, no direct effect of strain rate could be discerned on any of the

mechanical properties measured. In light of the initial billet temperature effect, one might

question why increased adiabatic heating within the more rapidly strained billets did not also

create measureable toughness differences. Because these laboratory tests involved short billet

lengths, and the length of billet extruded at any given ram speed was quite short (L/D 'vl),

adiabatic heating influences might have been inadvertently minimized. Ram speed (strain rate)

influences on toughness of production extrusions (where extrusion billet L/D >4) must still be

considered possible.

The extrusion ratio parameter (R = Area of Cylinder 4- Area of Product) alone does not

adequately quantify the strain history of the wrought shape. Aspect ratio (A = Product Width v Product Thickness) can also be an important descriptor. The longitudinal strength advantage in

highly deformed, axisymetric (low aspect ratio) extrusions is relatively well known. [Refs. 30,

65] Other effects can be more subtle. In 7050 alloy I/M extrusions, for example, product aspect

ratio does have a significant effect on strength, toughness and stress corrosion properties, with

the optimum overall property combinations in all test orientations being developed in extrusions

fabricated with both high extrusion ratios and high aspect ratio. [Ref. 66 ] Experimental studies

are currently underway at Alcoa Laboratories to explore the independent influence of aspect ratio

(at constant extrusion ratios) in 7XXX alloy P/M extrusions. [Ref. 76]

Structural Observations

In addition to the metal lographic finding that long time, high temperature degassing cycles

did measurably coarsen C02 AI 9 dispersoid particles in the Co-enriched X7090 alloy, various other

structural investigations sought to identify microstructural or macrostructural reasons for the

significant influence of extrusion billet temperatures on resultant wrought product fracture

toughness properties. Results are summarized in Table VIII. Subtle evidence of finer grain and

subgrain sizes was found in the higher toughness, low temperature extruded, products. The

predominant failure mode, for all extrusion conditions, was ductile-dimple fracture wherein dimple

size and spacing was in relative agreement with dispersoid size and spacing. C 02AI 9 dispersoid

particles were often seen within individual dimples. [Refs. 57, 77]

541
542
16.5 22 27.5 33 38.5 44 49.5 55

L-T K,c fracture toughness, MPa (ksi VmT)

Correlation of NTS/YS and K|c Fracture Toughness (Ref. 57)

Figure 1 3

TABLE VIII

EFFORTS TO IDENTIFY STRUCTURAL REASONS FOR DEFORMATION TEMPERATURE VERSUS TOUGHNESS EFFECTS IN X7090 AND X7091 P/M ALLOY EXTRUSIONS [Ref. 57, 77]

Structural Evaluation

Optical Metallography

X-Ray Pinhole Patterns

Guinier X-Ray Phase Analysi

SEM-Measurement of

Grain and Subgrain

Boundary Intercepts

in 3 Dimensions

Comparative Results As P/M Billet Extrusion Temperature Reduced

From 700K(800°F) to 589K(600°F)

Both grain and subgrain size was subtly finer.

All interior grain sizes >ASTM10.

Surface recrystallized grain depth increased; but maximum depth ^0.250 mm (0.01 in.) in X7091 . Depths shallower and intermittent in X7090.

Ring Spotliness (Possible Coarsening) Clear.

M' M and C 02AI 9 present in same amounts for all extrusion temperatures.

Calculated Boundary Area per Unit Volume

(S - mm^/mm^). As extrusion temperature decreased,

S '^increased by ^^S% in X7090 and 3-20% in X7091 .

SEM Fractography All tensile and toughness failures by dimple frac- ture with dispersoid particles within dimples. Large dimple size/spacing in the lower-dispersoid containing

X7091

543
Plant trials have confirmed the benefits of reduced extrusion billet temperatures on measured

fracture toughness properties, but the verification of definitive Cause and Effect relationships

among the observed X7090 and X7091 P/M microstructural features is a continuing challenge. The

deleterious effect of large (>1 micron) second phase particles seems firmly established and loose

powder or P/M compact degassing practices must be designed to minimize such dispersoid coarsening

reactions. This is especially important in the consolidation of large P/M billets for plate and

sheet product fabrication. Complimentary studies are underway to better define the interactive

roles of second phase Co 2AI 9 dispersoid particles with deformation texture, recrystallization

characteristics, and fracture toughness and corrosion resistance. [Ref. 78]

SUMMARY

Commercial manufacture of X7090 and X7091 alloy Wrought P/M extrusions and forgings is

already underway using Alcoa's patented alloy/process combinations. Fine, irregular-shaped, and

very rapidly solidified powder is continuously atomized in a 60 million pound per year plant

facility. Degassed and vacuum consolidated billets up to 159 kg (350 lb) weights are then

routinely produced for subsequent fabrication of extrusions or forgings in Alcoa's various

production facilities. 1545 kg (3400 lb) billets have been experimentally produced and ^ 3636 kg

(8,000 lb) billet capability is being pursued for manufacture of full scale P/M alloy plate and

sheet products before 1986. Enroute to achieving these production capabilities, extensive R&D studies have outlined common processing needs for 7XXX and other types of aluminum alloy P/M

products.

Production of high quality powder is followed by outgassing of the original oxide-coated

powder particle surfaces and vacuum consolidation into fully dense billets. Oxide film thickness

of aluminum powders is a function of the alloy composition but is virtually the same for different

atomizing processes, whether they involve air or "inert" gas atomizing/collection environments.

Oxygen content is a direct function of the surface area/volume ratio for the various powder

morphologies. Bulk oxygen contents of the billet and wrought product materials are usually lower

than oxygen contents of the as-atomized powders themselves, in verification of effective hot

degassing techniques for removal of adsorbed and chemically reacted moisture from the surface

oxide films. Hydrogen contents measured in high quality 7XXX alloy P/M billets usually fall

within a 1-10 ppm range when measured by fusion extraction techniques, but a significant portion

of this hydrogen is only released as a gas after melting has occurred. Consequently, the actual

amount of detrimental hydrogen gas in optimally processed solid state P/M products may be

comparable to the £ 0.3 ppm levels of similar I/M materials.

Room temperature tensile and notch tensile testing of P/M billets can be used to develop

ACCEPT/REJECT lot release criteria which enables significant billet-quality to product-quality

correlations. Property levels in high quality, well-bonded billets are still further improved by

subsequent hot deformation enroute to the final wrought component shape. Fracture toughness

levels are particularly sensitive to the relative amount and type of deformation strains incurred.

Strength and toughness properties are sensitive to variations in actual lot to lot chemical

compositions, to test orientation, to the presence of relatively coarse second phase dispersoid

particles, and to other conventional 7XXX process/property relationships.

ACKNOWLEDGEMENTS

The continued contractural R&D support of the U.S. Army Armament R&D Command, the U.S. Air

Force Systems Command and the Defense Advanced Research Projects Agency (Department of Defense) is

gratefully acknowledged. Equally important to the active development and commercial

implementation of Wrought P/M Aluminum Alloys is the strong desire of the Materials Users to embrace this technology, and the dedicated commitment of the Aluminum Company of America's

Laboratory and Corporate Management. Similarly, the interest and contributions by the Operating

and Metallurgical personnel of Alcoa's commercial Powder and Fabricating Plants are most

appreciated.

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544
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12. Ibid. , Phase II Final Report, June 1971.

13. Lyle, J. P., Jr. and Towner, R. J., U.S. Patent 3,637,441, Jan. 1972.

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June 1973.

16. Cebulak, W. S., op. cit. . Phase IVB Final Report FA-TR-76067, Apr. 1977.

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24. Mirth, J. P., Met. Trans. , Vol. 9A, 1978.

25. Mehrabian, R., Rapid Solidification Processing-Principles and Technologies , 1978.

26. Perepezko, J. H. and Rasmussen, D. H., Met. Trans. , Vol. 9A, 1978.

27. Gill, W. N., Chem. Eng. Common. , Vol. 12, 1981.

28. Lyle, J. P., Jr. and Cebulak, W. S., Met. Trans. , Vol. 6A, Apr. 1975.

29. Billman, F. R., Granger, D. A. and Fujioka, T., Unpublished Alcoa Research, 1976-78.

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545
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preparation.

546
POST RSPD PROCESSING EFFECTS ON GRAIN GROWTH

R.A. Amato, M.R. Jackson and P.A. Siemers

General Electric Corporate Research and Development

Schenectady, New York 12301

ABSTRACT

Rapid solidification plasma deposition (RSPD) is a process by which material is built up by

layers onto a substrate to form a near net shape component. Materials fabricated by plasma spray

deposition have structures characteristic of the rapid solidification processes. Recent studies

have indicated that these structures can provide enhanced properties at low and intermediate

temperatures. This study focuses on attempts to increase the grain size of RSPD materials in order to optimize the mechanical properties over a broader temperature range. An evaluation of grain growth resulting from a broad range of post deposition treatments has

been conducted on a number of Ni and Ni-Fe base alloys, along with variation in tensile and stress rupture properties. The techniques studied include isothermal or thermal gradient annealing,

strain and isothermal or thermal gradient annealing, hot rolling, and forging.

Grain coarsening has increased to sizes approaching those of cast alloys following strain and

anneal processing. Other techniques are, in general, unsuccessful in producing large grains. As

expected, hi^ temperature mechanical properties increase with the grain size in RSPD material.

INTRODUCTION

The fabrication of near net shape structures by the rapid solidification plasma deposition (RSPD)

process has evolved from coatings technology and early flame spraying techniques(l,2). The process involves

the injection of fine powders into a high energy plasma jet where they melt and are directed towards a

substrate. Upon impacting the substrate, the material is rapidly solidified. The ability to build up layers of metal on a substrate in low pressure has lead to sound structures requiring little subsequent working and/or

machiningO). The technique is especially attractive when using materials for which casting and further

processing are difficult. Rapidly solidified materials (4,5) have ultra-fine grain sizes with excellent strength

at low and intermediate temperatures and chemical homogeneity manifested by an increase in solidus

temperatures over cast or wrought material.

While nearly full densification results from high temperature isothermal annealing, grain sizes have

been limited to 25-50 ym in some Ni and Ni-Fe based alloys. Poor high temperature strength and stress rupture life are due to the abundant grain boundary area. Difficulty in achieving grain coarsening is not

exclusive to the RSPD process. For superalloy components made by conventional powder metallugy, limited

coarsening is attributed to MC and ^23^6 oxides, and even y' (Ni^Al) particles impeding grain

boundary motion(6,7). Superalloy material “typically contains significant quantities of these phases, and the

thick RSPD structures deposited in low pressure have oxygen levels in the 300-500 wppm range.

Nucleation and growth considerations lead to two possible grain growth techniques. By raising the total

free energy of a specimen sufficiently, for example as strain energy, recrystallization and subsequent grain

growth upon annealing can be promoted. Alternatively, eliminating the phases pinning boundaries will increase mobility and promote larger grains. These basic concepts, reviewed by Burgers and Aust (8,9), suggest that grain growth can be achieved in RSPD matericils by prior deformation or by controlling

composition.

This investigation sought to increase the grain size of RSPD materials through techniques employing

both deformation and heat treatments. Specifically these techniques included: strain followed by isothermal

or thermal gradient annealing, hot rolling and forging, and chemistry studies using Rene' 80 and experimental

Ni-Fe base alloys. In addition, the effect of thermal gradient annealing on material in the as-sprayed

condition was tested.

These attempts to increase grain sizes were intended to improve high temperature mechanical

properties. To measure this effect in RSPD alloys, mechanical properties were determined at various

temperatures on materials exhibiting increased grain size.

547
EXPERIMENTAL PROCEDURE

Deposition Processing and Alloy Selection

All deposits were made in a low pressure plasma spray facility. The vacuum chamber pressure was

reduced to below 0.33 kPa (2.5 torr) pri^r to back-filling with plasma gases to 8 kPa (60 torr). An EPl Model

03CA gun, with an anode having a 20 powder feed port angle, was operated at 55 volts and 1300 amps.

Argon and helium plasma gases had inlet flow rates of 1.40 liters/sec (179 scfh) and 0.36 liters/sec (46 scfh),

respectively. All gases were of commercial purity (99.99% pure, <10 ppm moisture, <20 ppm O

2

).

Metal powder was fed into the gun using two Plasmadyne Roto-Feed Hoppers at a combined mass flow

rate of 10.5 kg/hr (4.7 rpm at feeder). The Ar carrier gas had a volume flow rate of 0.12 liters/sec (15 scfh).

Powders sieved to -400 mesh and having the composition listed in Table 1 were supplied by either Alloy

Metals, Inc. of Troy, MI, or Homogenous Metals, Inc. of Herkimer, NY. Rene' 80 is a Y ' (Ni-Al) precipitation

hardened Ni-base alloy containing the titanium carbide (MC) which breaks down into ^nd ^^ 03^6

boundary c^bides during subsequent heat treatment(lO). The 47 v/o y' phase is in soluuon at temperatures

above 1150 C. The six experimental alloys are strengthened by either a Ti-rich y ' or a Cb-rich y

" precipitation phased 1). A sequence was chosen to gradually vary the volume fractions of the matrix and

strengthening phases. An 1150 C anneal was found to be sufficient to drive the precipitates into solution for

all six alloys. The Rene' 80 and Hastelloy X are the standard compositions available commercially. The

experimental alloys, designated C1-C6’, have compositions similar to IN 718, Hastelloy X and U- 212. Powder was deposited onto 15cm x 15cm (6 in x 6 in) steel plates that were grit blasted and degreased

in Freon prior to coating. The gun traveled at a velocity of 15 cm/sec (6 in/sec), 30.5 cm (12 in) above the

substrate. The gun was perpendicular to the substrate for a 90° impingement angle.

TABLE I

Alloy Ni Fe Cr

Nominal Compositions of Alloys Atomic Percent Mo Cb Ti A1 Si Mn C B W

Cl 62 10 14 2 6 2 2.5 0.6 0.3 0.4 0.2 C2 55 19 14 3 4.5 1.5 1.5 0.6 0.3 0.4 0.2 - C3 51.5 19 19 4 7.5 1 1.5 0.6 0.3 0.4 0.2 - C<f 48 19 24 5 0.5 0.5 1.5 0.6 0.3 0.4 0.2 - C5 38.5 34 19 3 0.5 2.5 1.0 0.6 0.3 0.4 0.2 - C6 29 49 14 1 1.5 4.5 0.5 0.6 0.3 0.4 0.2 - Rene' 80 66.8 - 16 2.5 - 6 6.5 - - 0.8 0.1 1.3 Hastelloy X 46.9 19.3 24.6 5.5 1.0 0.5 0.5 0.2

Post Deposit Processing

Plate deposits were sectioned to provide 3.5 cm x 1.5 cm x 0.15 cm (1.4 in x 0.6 in x 0.06 in) coupons

whose faces were perpendicular to the plane of deposit and whose longitudinal direction was parallel to the

direction of gun motion. Each slice was subsequently surface ground to make parallel coupon faces.

Thermal gradient treatments, similar to those used to align grains in dispersion strengthened

superalloys, were performed by passing Rene' 80 specimens through a water cooled Cu chill inserted into a

vertical tube furnace. The specimens were driven through the system at a rate of 2.5 cm/hr (1 in/hr).

Temperature gradients of 100 C/cm at

y

' solvus temperatures were measured by a series of thermocouples

mounted on a test specimen. Thermal gradients decreased as the specimen moved away frogi the Cu chill. A

thermal gradient of 80^C/cm was measured at the maximum furnace temperature of 1245 C located 4.5 cm

from the Cu chill.

Specimens that were to be strained prior to annealing were first given a high temperature solution

anneal in order to achieve their full^ density and maximize the grain size. Rene' 80 specimens received ^

2 hour solutioning treatment at 1220 C which in some cases was followed by an additional hour at 1240 C.

This increased the starting grain size to 25 y m. The Ni-Fe base alloys were given solution anneals for 1 hour

at 1220°C with the exception of Alloy Cl, w^ich showed evidence of incipient melting at temperatures as low

as \^170 C. This alloy was annealed at 1150 C. All isothermal anneals were done in a 1.33 x 10 kPa (1 x

10 torr) vacuum furnace.

The solution annealed coupons were further sectioned into specimens 1.0 cm x 0.5 cm x 0.15 cm (0.40 in X 0.20 in X 0.06 in) and cold rolled on a two-high 6 x 6 rolling mill. The specimens were subsequently

recrystallized for 1 hour, at the same temperatures used for solutioning, and the resulting grain size

determined.

548
Specimens were selected for the strain and thermal gradient anneal study based on response to strain

and isothermal anneal treatments. Alloys C2 and C4 were cold rolled 4% and 22%, respectively, following

the solution annealing procedure previously described. The specimens were driven into the thermal gradient

furnace at 2 cm/hr (0.8 in/hr) and 10 cm/hr (4 in/hr) and the resulting grain sizes determined.

Hot rolling specimens, 2.5 cm x 1.5 cm x 0.3 cm (1 in x 0.6 in x 0.12 in) were sliced from plate deposits

of alloys C2, C3, C5, Hastelloy X, and Rene' 80 as previously described. The coupons were packed in 304L

stainless steel after being coated with a thin layer of MgO to prevent bonding. The packs were heated in a Muffle furnace and reduced in a single pass on a two-high Birdsboro Mill at the conditions listed in Table 2.

Hot forging was done on a hammer forge on 2 cm x 3 cm x 3 cm specimens heated to 1200°C. A single

stroke deformation of 27% reduction in thickness was followed by reheating and a second stroke of 30%

reduction.

TABLE 2 SINGLE PASS HOT ROLLING CONDITIONS

Alloy Temp. (°C) % Reduction

C2 1220 52 C2 1220 26 C3 1220 51 C3 1220 27 C5 1220 50 C5 1220 27

Hastelloy-X 1230 53 Hastelloy-X 1230 26 Rene 80 1230 50 Rene 80 1230 23

Material Evaluation

The effect of grain size on the mechanical properties of RSPD alloys was determined for critically

strained and isothermally annealed 3.5 cm x 1.5 cm x 0.15 cm (1.4 in x 0.6 in x 0.06 in) coupons prepared as described in a previous section. The large grained specimens were tensile and stress rupture tested^at various

conditions. The flat test specimens had a 0.63 cjn (0.25 in) gage length. A strain rate of 0.08 min”^ was used on Rene' 80 specimens while a rate of 0.04 min was used on the other specimens. Tensile tests were done in vacuum while stress rupture tests were done in argon and air, for Rene' 80 and^the experimental alloys,

respectively. The Ni-Fe base experimental alloys were heat treated 4 hrs at 900 C and 12 hrs at 760 C to precipitation strengthen the alloys.

Metallographic investigations were done to determine the grain size that resulted from the different

processing techniques. The specimens were sectioned and mounted to reveal the three orthogonal planes of

rolling. Grain sizes were determined utilizing a line intercept method with 12-30 measurements averaged to give grain sizes and error bands indicating the 90% confidence limits.

RESULTS

Microstructures

Strain and Isothermal Anneal; The effects of strain and anneal processing on RSPD materials was first investigated using Rene' 80. The extent of grain coarsening due to deformation and annealing, 1 hour at 1250'^C, is shown in Figure 1. These results indicate that at very low reductions in thickness, i.e. less than

3%, no increase in grain growth is achieved. In the 3-6% deformation range large grains on the order of 100-

150 y m are produced. The typical microstructure which developed after the 3-6% deformation and

subsequent anneal is shown in Figure 2 along with the microstructure developed after a solution anneal. The

primary recrystallized grain size decreased at deformations greater than the critical amount. Grain sizes were uniform through the thickness of the specimens in all cases.

The six experimental alloys along with Hastelloy X were subjected to a similar series of deformation

and anneal heat treatments. The results of these tests are presented in Figures 3 and 4. The alloys show a

similar behavior to Rene' 80, having a critical deformation range below which no increase in grain growth is observed but within this range large grain sizes can be found. Alloy Cl was an exception which showed only a

slight increase in grain size. 549
GRAIN

SIZE,

mi

crons

The microstructures developed in alloys C2, C4, and Hastelloy X after critical strain and annealing,

along with the microstructures characteristic of the materials given an isothermal anneal treatment only, are

given in Figures 5-7. Alloy C2 had the largest increase in grain size of the experimental alloys and after just a 3% deformation and anneal grew grains to over 400 pm, see Figure 5. While grain growth proceeds through

the individual pass layers, there is a tendency for the grain boundaries to become aligned with these layers.

The increase in grain size of alloy C4, Figure 6, after the critical strain and anneal process compared to the

isothermally annealed RSPD material is characteristic of alloys C3-C6 and commensurate with the increase

found for Rene' 80. Hastelloy X responded most favorably to the strain and anneal treatments with

recrystallization and grain growth resulting in 500 pm grains after just 2% reduction in thickness.

Beyond the critical deformation range the primary recrystallized grains decrease in their ability to grow beyond the 25 pm levels which developed without pre-straining. The recrystallized grain sizes found

after the 14% reductions are characteristic of the results following larger deformations. Large secondary

recrystallized grains, however, were found in alloys C4, C5, and Hastelloy X after 18-25% deformation and

subsequent anneals, see Figure 8.

Grain boundary phases, which can cause inclusion-inhibited growth and can lead to secondary

recrystallization, were extracted from alloy C4 and inspected by transmission microscopy. They have been

determined to be rich in Cr and Mo with some Fe and Ni additions. The FCC structure has a lattice

parameter of 10. 8A which suggests either a ^^23^6 carbide.

Fig. 1. Grain growth behavior of plasmasprayed Rene' 80 as a function of the

amount of cold rolling.

Fig. 2. Typical grain sizes developed in RSPD

Rene' 80 following (a) an isothermal

anneal were 25 pm while those

resulting from (b) critical strain (3-6%

reduction in thickness) and anneal

processing were 100-150 pm.

550
Grain

Size,

(nm)

500

% Reduction in Thickness

O

in

o

o

o

o

m

-I— —I—I—I—I—I—I—I—I—I—

1—I—rGRAIN GROWTH

BY CRITICAL

STRAIN ANNEALING

LUN

I—

CO

<

cr

LD

a

a

(\j

o

o

\

HASTELLOY X

4 6 8 10 12

% REDUCTION

14 16

Fig. 3. Variation in grain size with deformation following recrystaUization anneal

heat treatment. Experimental alloys

C2-C6 were annealed 1 hr. at 1220*X

while alloy Cl was annealed 1 hr. at 1150°C. Error bars indicate 90% confidence limits.

Fig. 4. Grain growth as a function of reduction in thickness following recrystallization anneal at 1220°C for 1 hr.

Strain and Thermal Gradient Anneal; For thermal gradient annealing studies, the furnace temperature

profiles could be varied. In the present work, the profiles were controlled to achieve a lOO^C/cm thermal

gradient at the location where the y' solvus temperature, 1150 C, was first exceeded. For Rene' 80 samples

in the as-deposited condition, no increase in grain size was observed compared to isothermal annealing.

Alloys C2 and C4 showed a propensity to grow large primary or secondary recrystallized grains after

reductions in thickness of 4% and 22%, respectively. The alloys were tested in these conditions to determine

the effect that a 100°C/cm thermal gradient anneal would have on grain growth. Neither alloy C2 nor C4

grew elongated grains due to their exposure to a thermal gradient. Microstructures were characteristic of those resulting from an isothermal anneal rather than from a directional recrystallization process.

Hot Deformation; As an alternative to cold rolling with subsequent anneal treatments, a study was conducted

to determine the feasibility of hot rolling RSPD superalloys to control grain size. Typical microstructures

developed during hot deformation are shown in Figures 9 and 10. Similarities to the as-sprayed material exist

with evidence of unmelted particles and a fine grain size. Rene' 80 seen in Figures 9(a) and (b) may have

undergone dynamic recrystallization as the grain morphology remained equiaxed after 50% reduction in thickness. The Ni-Fe base alloy C2 may have undergone a similar effect as these grains are also equiaxed

while unmelted particles have become elongated in the rolling direction, see Figure 10(a). A subsequent

anneal again produced grains 25 ym in size as seen in Figures 10(b) and (d). The observations from hot

forging confirmed these results. 551
Fig. 5. Typical grain sizes developed in RSPD experimental alloy C2 following (a) 1 hr. at 1220°C were 25

ym while those resulting from (b) critically straining 3% (reduction in thickness) before annealing

averaged 230 ym.

Fig. 6. Typical grain sizes developed in RSPD experimental alloy C4 following (a) 1 hr. at 1220°C were 25

ym while those resulting from (b) critically straining 4% (reduction in thickness) before annealing

averaged 150 y m.

1

Fig. 7. Typical grain sizes developed in RSPD Hastelloy X following a) 1 hr. at 1220°C were 50 ym while

those resulting from (b) critically straining 3% (reduction in thickness) before annealing averaged

240 ym.

552
Fig. 8. Secondary recrystallization is apparent in the RSPD aUoys C4, C5,

and Hastelloy X. This micrograph

shows the rolling plane of experimental alloy C4 after 1 hr. at 1220°C,

cold rolling 22%, and 1 hr. at 1220°C.

Fig. 9. Hot-rolled RSPD Rene' 80, reduced 50% in thickness at 1230°C, appeared to have undergone

dynamic recrystallization as equiaxed grains 10 ym in diameter resulted. Micrographs taken (a) of the rolling plane and (b) perpendicular to the rolling plane.

An interesting observation regarding the microstructures of hot rolled and annealed RSPD material is the evidence of grain boundary motion beyond unmelted prior particle boundaries. It can be seen in Figure 1

that the grain boundary has advanced beyond the unmelted particle boundary which is decorated by fine

carbides.

Material Properties

To determine the effect of grain coarsening on the high temperature properties of the RSPD alloys,

critically strained and isothermally annealed specimens were tensile and stress rupture tested at various

conditions. The results of these tests, along with data for the isothermally annealed RSPD material, are presented in Table 3 for Rene' 80 and Table U for the experimental alloys. Hastelloy X was not tested after

strain and isothermal anneal treatments.

The major trends which are found when comparing isothermally annealed RSPD material to critically

strained and annealed RSPD material are: an increase in high temperature strength, a decrease in low

temperature strength and ductility, and an increase in stress rupture life. Rene' 80 tensile strengths at 1000 C increased from 213 MPa to 432 MPa as the grain size changed from 20 y m to 80 ym. The

experimental alloys (alloy Cl is omitted from the comparison since it did not show a marked increase in grain

size cifter processing, and the anomalous room temperature tensile data of isothermally annealed alloy C3 is not used in the comparison) did not show as significant an effect at high temperature as was found for Rene'

553
if

^ V

;-'1 ' ' «.

• >-:

r n) . vt, r r -

' f'>^' V

^ ^ S * /

if,

'v /

Fig. 10. Typical structures of RSPD alloys following hot deformation. AUoy C2 hot rolled 20% reduction in thickness at 1220°C. (a) rolling plane as deformed, (b) rolling plane deformed and post annealed, (c) longitudinal section as deformed, (d) longitudinal section as deformed and post annealed. Etchant:

glycergia.

Fig. 11. Example of grain growth beyond

unmelted particle boundary. RoUing

plane of AUoy C2 rolled 20% reduction

in thickness at 1220°C. Etchant:

glycergia.

554
80. The exception to this was alloy C2 which increased from 64 MPa to 180 MPa, at 1060 C, when the grain

size went from 25 pm to over 400 pm. The increases in grain size of alloys C5 and C6 did not result in an

increase in the high temperature tensile strengths. The room temperature tensile strength and reduction in area of Rene' 80 decreased from 1330 MPa and 17% to 1280 MPa and 8.3% as the grain size increased. The

experimental alloys were in general decreased by 10% in tensile strength and twice that in ductility at low

temperatures. The results of the stress rupture tests indicate that the time to failure for Rene' 80 increased

by 100% and for the experimental alloys increased by nearly one Larson-Miller parameter with larger grains.

Critically strained and annealed RSPD Rene' 80 material was compared to cast and fully heat treated

mat^ial. The plasma sprayed material had strengths (1280 MPa at 20°C, 772 MPa at 850°C, and 432 MPa at 1000 C) that^were equal to or greater than the cast material (1030 MPa at 20 C, 750 MPa at 850°C, and 345

MPa at 1000 C). The low temperature elongation of the plasma sprayed material, 8.3%, was slightly greater

than that found in the cast material, 7.0%. At high temperatures, the ductility of the plasma sprayed

material wa^ significantly less than was found for the cast material: 3.9% compared to 23.0% at 850^C. The RSPD material was 2-3 Larson-Miller parameters below the cast and fully heat treated

material.

DISCUSSION

There is a lack of normal grain coarsening behavior in certain RSPD materials following isothermal

anneal heat treatments. This indicates that the energy necessary for grain boundary motion is greater than is generated by a reduction in grain boundary area. For similar reasons, the application of a temperature

gradient to plasma sprayed Rene' 80 was unsuccessful in promoting grain boundary motion beyond that found

after isothermal anneals.

Grain growth of RSPD materials can be achieved, however, by straining and annealing. The largest

amount of grain growth occurs following a small but critical amount of deformation. The strain and anneal

process has previously been successful in promoting grain growth in powder metals consolidated by extrusion

(7,12). The small amount of deformation provides enough energy to nucleate a few recrystallized grains, and

their subsequent growth results in large sizes.

Grains in Rene' 80, which grow on an average to 25 ym after an isothermal anneal, grow to 100-150

ym after a 3-6% critical strain and anneal. The six experimental alloys had grain sizes 25 ym or less after an

isothermal anneal yet developed grains to 150-200 y m by the critical strain and anneal process. Alloy C2

grew the largest grains, averaging over 400 ym in diameter. This may be a result of a eutectic reaction,

involving the carbide particles, that this alloy undergoes at temperatures only slightly above the anneal

temperature. A reaction of this type could significantly affect the diffusion characteristics of the carbides

and increase the grain boundary mobility. Alloy Cl after being critically strained and annealed grew grains

that were only slightly larger than those developed 1^ an isothermal anneal. The occurrence of a low

temperature melting reaction, however, required a 70 C decrease in this alloys heat treating temperature.

Hastelloy X had grain sizes that increased by an order of magnitude over the 50 ym size from an isothermal

anneal following 2% deformation and a subsequent anneal. TABLE 3

GRAIN GROWTH AND MECHANICAL PROPERTIES OF RENE 80

Additional 0.5 h at 1250*C during alloy preparation. ***M.R. Jackson

Aerospace Siruaural Metals Handbook, for conventionally cast Rene 80

w>2 h/1220"C, 1 h/1240"C, 5% cold rolling, 1.5 h/1250*C

TENSILE TEST RESULTS

Ultimate

Alloy Preparation Temperature (X)

Grain Size (mhi)

0.2 %

Yield Stress (MPa)

Tensile Strength (MPa)

Reduction

In Area

(%)

RSPD, 20 70 910 1280 8.3 2 h at 1220”C 850 100 705 772 3.9

1 h at 1240'C (nonuniform) Cold Roll 5.0% 1000 80 425 432 2.9

1.5 h at 1250X 1000"> 70 296 381 10.7

RSPD, 20 20 827 1330 17

2 hat 1250”Ci2i 850 20 524 723 12 1000 20 186 213 9

As-Cast, 20 400-800 860 1030 7.0 Fully Heat Treated^^' 850 400-800 560 750 23.0 1000 400-800 <275 345 20.8 STRESS RUPTURE RESULTS

Conditions Time to Failure (h) Temperature Stress (RSPD, Isothermal Anneal (As-Cast, Fully, (”C) (MPa) RSPD'*' 2 h at 1145°C, 4 h at 870°C)'" Heat Treated)'^’

871 345 4.3 2.33 100

1.91 (at 262 MPa)

982 138 6.1 3.17 (at 117 MPa) 100

555
TABLE 4

TENSILE BEHAVIOR OF EXPERIMENTAL ALLOYS

(Strain rate: 0.044/min)

T(*C)

Full Heal Treaiment*'* YS(MPa) UTS(MPa) V

RA

Strain !l Heat Treatment*^’ YS(MPa) UTS(MPa) e^^ «i RA

Cl 20 1103 1203 1.8 1.8 7.0 900 1390 18.5 18.7 16.6 610 974 1176 2.8 2.8 5.4 800 1139 10.8 16.8 14.0 760 647 688 1.2 4.0 3.1 600 645 2.8 6.8 6.0 910 214 232 1.2 5.4 5.6 259 274 1.2 7.2 4.0 1060 28 39 2.8 29.0 21.0 70 73 2.2 13 8 6.0

C2 20 596 963 28.0 36.0 45.0 273 714 50.5 54.8 34.4 610 383 696 24.0 38.0 39.0 219 536 50.1 54.0 28.0 760 369 379 7.2 22.0 22.0 232 376 13.1 53.7 34.7 910 134 135 0.5 25.0 30.0 128 146 7.5 65.3 65.3 1060 63 64 0.4 12.3 12.1 167 180 2.8 53.4 32.0

C3 20 272 352 3.0 3.0 8.6 299 735 55.4 64.1 41.3 610 304 647 41.0 49.0 36.0 266 579 47.2 52.1 30.0 760 389 444 4.0 26.0 22.0 297 403 7.1 48.4 33.3 910 183 188 0.6 20.0 17.8 152 160 4.5 44.9 26.0 1060 58 62 1.0 17.8 13.5 63 69 3.5 11.3 7.8

C4 20 446 906 28.0 34.0 36.0 347 808 40.2 48.2 37.9 610 359 676 26.0 43.0 42.0 397 561 10.2 14.0 13.7 760 314 382 7.8 58.0 45.0 No Test - - - -

910 155 172 4.8 23.0 25.0 159 172 3.7 80.0 40.0 1060 59 62 1.5 8.0 8.3 66 73 1.7 11.8 5.9

CS 20 787 1179 18.8 24.0 28.0 542 1006 22.3 23.8 20.7 610 734 1016 14.2 22.0 26.0 541 856 20.1 32.1 29.4 760 620 638 0.8 17.8 16.3 466 512 1 0.0 27.7 23.5 910 130 132 0.4 30.0 30.0 124 134 15.7 69.2 39.2 1060 54 55 4.5 31.0 28.0 46 51 16.9 28.2 18.0

C6 20 967 1265 9.7 9.7 10.7 829 946 2.1 2.1 2.7 610 880 992 5.2 15.8 27.0 750 878 9.2 9.9 12.2 760 592 615 1.2 19.2 23.0 562 589 5.6 15.1 14.0 910 128 143 4.0 30.0 38.0 120 138 6.0 48.7 31.9 1060 41 48 3.7 49.0 43.0 35 36 4.2 24.3 28.6

T —lest temperature YS —yield strength UTS —ultimate tensile strength —strain at maximum load (%)

ej^ —strain at failure (%) RA —reduction in area (%)

1 h/l220“C*, 4 h/900°C, 12 h/760°C, M.R. Jackson

I h/I220”C", cold roll 2-6%, 1 h/1220°C’, 4 h/900”C. 12 h/760”C

'except Alloy Cl which used IISO'C anneals

Note: All alloys deposited in EPI Ar-He plasma

Stress Rupture Properties in Air for Experimental Alloys

Alloy Condition*

Test Temp.( C) Stress (MPa) Life(h)

Cl 1 610 827 3.1 2 _ - -

1 760 310 1.0

2 760 310 3.3 C2 3 610 483 14.0

4 610 483 17.5

3 760 193 4.1

4 760 193 190.9 C3 3 760 193 26.9

4 760 193 110.2

3 910 48 11.5

4 910 48 52.9

C4 3 910 48 31.5

4 910 48 95.9

3 1060 14 4.8

4 1060 14 54.2 C5 3 760 193 58.3

4 760 193 237.2

3 910 48 41.4

4 910 48 229.8 C6 3 610 483 572.6

4 610 483 1143.9

3 760 193 81.6

4 760 193 180.1

P(K,h) Elong.(%) R.in A. ("6)

15.1 3.4 5.9

20.7 3.0 3.2 21.2 0.8 0.4

18.7 10 12 18.8 29 41 21.3 15 14 23.0 18 24

22.1 13 12 22.8 19 23 24.9 51 32 25.7 31 21

25.4 27 27 26.0 29 19 27.6 32 19 59.2 29 6.0

22.5 4.1 6.6

23.1 6.8 6.5 25.6 35 28 26.5 16 8.7

•Condition: 20.1 4.5 7.4 20.4 (unloaded)

22.6 14 13 23.0 4.4 0

1 - 1 h/1150°C solution, M.R. Jackson

2-

1 h/1150°C, cold roll to critical strain, 1 h/1150 C

3-

1 h/1220°C, solution, M.R. Jackson

4 - 1 h/1220°C. cold roll to critical strain, 1 h/1220 C

(All had 4 h/900 C & 12 h/760 C final anneals)

556
The difference in the grain sizes of alloy and Hastelloy X both before and after critical strain

and annecd treatments is surprising since the alloys have basically the same chemistries. Alloy C4 (along with

all the other alloys tested in this study except Hastelloy X) contains minor additions of Cb, Ti and Al. It appears that the lack of these major y oxide and carbide forming elements in Hastelloy X resulted in much

more mobile grain boundaries than are in the experimental alloys and Rene' 80. In addition, the evidence of grain growth through the individual pass layers after critically straining and annealing implies that the

interfaces have good bonding and are not responsible for impeding grain boundary motion. It is more likely

caused by solute atoms or phase(s) which can be abundant at these locations.

The smaller grain sizes developed as deformation increased beyond the critical strain regime are a

result of the large number of nuclei forming throughout the lattice. This limits the extent each new grain can coarsen. The primary recrystallized grain size behavior of the alloy investigated in this study is typical

of the effect of prior cold work on recrystallized grain size. Examining the microstructure produced after

these large deformations and subsequent anneals revealed some large grains (over 100 pm) in Rene' 80,

experimental alloys C4 and C5, and Hastelloy X. Secondary recrystallization resulted in these grains being

significantly larger than the surrounding primary recyrstallized grains.

Directional recrystallization of RSPD alloys, by a thermal gradient anneal, was not found to be

successful in materials that demonstrated exaggerated grain growth during isothermal anneals. The

conditions for critical strain annealing and secondary recrystallization, as demonstrated in this report,

involve relatively small deformations. Post deposit treatments of these types are of interest since RSPD is meant to be a near net shape process.

Hot deformation processes produced microstructures which retained characteristics of the as- sprayed material treated isothermally at the working temperatures. Equiaxed grainy limited to 5-10 ym in diameter, developed in Rene' 80 specimens after 50% reduction in thickness at 1230°C. Similar sizes were

observed in the experimental alloys and Hastelloy X. The 25 ym grain sizes produced by a subsequent anneal were typical of the structure developed following an isothermal anneal of unworked material. The results from forging operations coincide with the observations that grain growth was not enhanced by hot

deformation.

The major mechanical property effects observed after an increase in grain size on RSPD material were

a decrease in low temperature strength and elongation, an increase in high temperature strength, and an

increase in stress rupture life. These effects were observed in both Rene' 80 and the experimental alloys,

when critically strained and annealed materieil was compared to isothermally annealed material. The

decrease in low temperature strength as grain size increases is predicted by the Hall-Petch relationship. The

decreasing elongation is a result of less uniform slip as the grain size gets larger. The increase in high

temperature strength and stress rupture life is expected with the decrease in grain boundary area. The small decrease in strength at low temperatures compared to a large increase at high temperatures

in Rene' 80 indicates that the amount of strengthening from grain boundaries is relatively small in this alloy.

The experimental alloys, however, which show larger decreases in low temperature strength would appear to be more dependent on grain boundary strengthening. The increase in high temperature strength of alloy C2

after the critical strain and anneal process confirms the grain size effect in these alloys. The lack of overall

agreement of the strength data is probably a result of the alloys having relatively low high temperture

strength. The alloys do, however, concur with the trend based on the stress rupture results, increasing by

nearly one Larson-Miller parameter with larger grain sizes.

The results clearly show that developing larger grains in RSPD Rene' 80 was sufficient to give the

plasma sprayed material equal high temperature strength to that of the cast and fully heat treated material.

The stress rupture life, however, appears to be more heavily dependent on grain size, and the RSPD material

(80 y m grain diameter) is 2-3 Larson Miller parameters below the cast material (400-800 y m grain

diameters).

CONCLUSIONS

An evaluation of grain growth resulting from a broad range of post deposition treatments has been

conducted on a number of Ni and Ni-Fe base alloys, along with the variation in tensile and stress rupture

properties. The principle findings and conclusions are:

(1) The most successful method for achieving grain growth in RSPD material was straining with a subsequent recrystallization anneal. Grains were increased, to sizes approaching those typically found

in cast materials, after a small but critical amount of cold rolling followed by an anneal treatment.

(2) Grain growth was not increased by a thermal gradient anneal beyond the 25-50 ym size developed

during an isothermal anneal.

557
(3) Directional recrystallized of RSPD alloys, by a thermcil gradient anneal, was not achieved in materials

that demonstrated exaggerated grain growth during isothermal anneals.

(4) Hot deformation processing was ineffective in promoting grain growth. Hot rolled specimens have

microstructures characteristic of the as-sprayed material: unmelted prior particle boundaries and a

fine grain size. Subsequent heat treatments increase these grains to the size found after isothermally

annealing unworked material.

(5) Coarsening the grain sizes of RSPD materials by critically straining and annealing resulted in improved

high temperature mechanical properties. The extent to which this affected the tensile properties of an

alloy was dependent upon the materials inherent high temperature strength and the relative importance

of grain boundary strengthening.

(6) Rene' 80 material that was plasma sprayed before critically straining and annealing had strengths that were equal to or greater than the cast and fully heat treated material.

(7) Carbides or oxides formed after additions of certain elements, appear to be possible impediments to grain growth. Thus, the alloy compositions of interest for RSPD are such that some resist grain growth

during annealing.

(8) Some alloys show a tendency for secondary recrystallization during annealing following cold rolling.

Secondary grains grew an order of magnitude larger than the uncoarsened matrix grains, but secondary

recrystallization was never complete.

ACKNOWLEDGEMENTS

This program has been funded by DARPA (Contract No. F33615-81-C-5156), through the offices of Lt. Col. L.A. Jacobson, with H.L. Gegel of AFML serving as technical monitor. The authors would like to thank

S. Rutkowski who prepared the plasma sprayed deposits along with T. Douglas, R. Laing, and P. Martiniano

who did the rolling operations. In addition we would like to thank A. Ritter who identified the carbide

particles by transmission electron microscopy.

REFERENCES

(1) T.A. Roseberry and F.W. Boulger, "A Plasma Flame Spray Handbook", Report No. MT-043, Naval Sea

Systems Command, Dept, of the Navy, 1977, March.

(2) W.E. Ballard, Metal Spraying and Sprayed Metal , Charles Griffin and Company, London, 1948.

(3) M.R. Jackson, J.R. Rairden, J.S. Smith, and R.W. Smith, "Production of Metallurgical Structures by

Rapid Solidification Plasma Deposition", J. of Metals, 1981, November, 23.

(4) D.J. Looft and E.C. van Reuth, "Rapid Solidification Processing: An Overview", Proc. of the Int. Conf. on Rapid Solidification Processing, Reston, Va., 1977, 1.

(5) R.J. Patterson, II, A.R. Cox, and E.C. van Reuth, "Rapid Solidification Rate Processing and Application

to Turbine Engine Materials", J. of Metals, 1980, September, 34.

(6) L. Winberg and M. Dahlen, "Recrystallization in a Powder Metallurgy Nickel-Base Superalloy", J. of Mat. Sci., 1978, 13, 2365.

(7) M. Dahlen and L. Winberg, "Grain Coarsening of PM Nickel-Base Superalloy by Critical Strain

Annealing", Metal Sci., 1979, 13, 3 and 4, 163.

(8) W.G. Burgers, "Principles of Recrystallization", in The Art and Science of Growing Crystals, J.J. Gilman ed., J. Wiley and Sons, NY, 1963, 416.

(9) K.T. Aust, "Large Crystals Grown by Recrystallization", in The Art and Science of Growing Crystals, J.J. Gilman ed., J. Wiley and Sons, NY, 1963, 452.

(10) E.W. Ross, "Rene' 80: A Cast Turbine Blade Alloy", Metal Progress, 1971, March, 93.

(11) M.R. Jackson, "RSPD for Fabrication of Advanced Aircraft Gas Turbine Components", P.A. Siemers

ed., DARPA Contract F33615-81-C-5156, Interim Report 1982, January.

(12) M. Dahlen and H. Fischmeister, "Grain Size Control in PM Superalloys", European Symp. on PM VI, Stockholm, Sweden, 1978, June 149.

558
Ni-Mo-B ALLOYS: METALLIC GLASS TO DUCTILE CRYSTALLINE SOLID

S. K. Das, L. A. Davis, J. R. Y. Wang* and D. Kapoor*

Allied Corporation

Materials Laboratory

Corporate R & D P.O. Box 1021R

Morristown, NJ 07960

ABSTRACT

The microstructures and mechanical properties of a series of rapidly solidified

Ni-Mo-B alloys, containing from 20 to 30 at%Mo and 10 to 20 at%B, have been investigated

using analytical electron microscopy and hot hardness testing. The as-cast metallic

glasses are quite brittle and can be readily pulverized to powder and consolidated into

bulk crystalline solids by hot isostatic pressing (HIP) or other means. The consolidated

bulk samples show fine dispersions of borides (M02 NiB 2 ) which contribute to the strengths

and hardnesses of the materials. Mechanical behavior is strongly influenced, also, however,

by the matrix composition; when the Mo/Ni ratio in the matrix exceeds 0.25 the hot hardness

is enhanced, as the matrix is strengthened through the presence of ordered phases (Ni 4 Mo,

Ni 3 Mo and Ni 2 Mo depending on composition and heat treatment). When Mo/Ni< 0.25 hot hardness

is reduced, even when the B content is as high as 20 at%. Concurrently, room temperature

toughness is enhanced in the absence of ordered phases.

INTRODUCTION

The formation of metallic glasses by rapid solidification processing and their application as magnetic materials is now well known [1]. To date, however, metallic glasses have found little

application as structural materials [2]. In contrast rapid quenching of crystalline alloys for

structural applications has become a field of considerable activity [3]. The goal, usually, is to produce consolidated bodies from powder/particulate, which, itself is microcrystalline in the as quenched state. Relatively little attention has been given to structural applications of crystalline alloys obtained by devitrification from the glassy state. Recent work at Allied Corporation

has shown that devitrified metallic glasses of certain compositions, when consolidated into bulk

compacts, exhibit outstanding properties [4]. As an example, in the present paper we report certain unique mi crostructural features of a class of Ni-Mo-B alloys which have attractive strength,

hot hardness and wear resistance properties.

EXPERIMENTAL

The compositions of the alloys investigated in the present study are listed in Table 1. The

alloys were rapidly quenched from the melt into continuous ribbons using either jet casting or the

Planar Flow Casting (PFC) process [2]. The recrystallization temperatures and the incipient melting points of the alloys were determined from differential scanning calorimetry and differential

thermal analysis of the as cast ribbons and are also listed in Table 1. The ribbons were consolidated into bulk compacts by pulverizing to -35 mesh powder and subsequent HIPping at 1100°C and a pressure of 15,000 psi. Thin foils suitable for transmission electron microscopy were prepared by a combination of electrolytic jet thinning and ion milling. The foils were examined in a Phillips

EM400T scanning transmission electron microscope (STEM) equipped with an energy dispersive x-ray

analytical system.

* Consolidated Metal Products Department, Allied Technologies Company, Allied Corporation.

559
RESULTS and DISCUSSION

The as cast ribbons were examined by x-ray diffraction analysis and appeared to be amorphous

for all the compositions listed in Table 1. However, transmission electron microscopy revealed the

presence of extremely small crystallites in jet cast ribbons of alloys containing 10% B (Fig. 1).

Table 1. Nominal compositions of the Ni-Mo-B alloys.

ALLOY

NO. COMPOSITION (atomic%) RECRYSTALLIZATIONTEMPERATURE (°C)

INCIPIENT MELTING

Ni Mo Fe B POINT (°C)

1 54 26 20

2 49 31 - 20 561 1235

3 62 23 - 15 - 1245

4 60 30 - 10 574 1238

5 56.5 23.5 10 10 521 1250

6 66.5 23.5 ” 10 535 1235

Fig. 1 Transmission electron micrographs of as cast ribbons of Ni5oMo3oBio alloy using

(a) - (c) jet casting and (d) - (f) the Planar Flow Casting technique; (a) and

(d) are bright field micrographs, (b) and (e) are corresponding microdiffraction

patterns and (c) and (f) are corresponding dark field micrographs.

560
The dark patches seen in Fig. 1(a) are indeed crystallites as evidenced by the microdiffraction pattern (Fig. lb) and the dark field micrograph (Fig. Ic). The same alloy when cast using the RFC process does not show any detectable crystallites (Fig. Id). No descrete spots can be seen in the

microdi ff raction pattern (Fig. le). The spotty contrast seen in the dark field micrograph (Fig. If)

taken using the amorphous ring, can result from statistical density fluctuations even in a completely amorphous matrix [5]. These results show that a somewhat higher cooling rate can be obtained

in the RFC process as compared to jet casting. The microstructures of the as cast ribbons of the

alloys containing 15 and 20% B are similar to those shown in Fig. 1(d) - 1(f). Whether small crystallites are present or not the as cast ribbons are brittle and readily pulverized into powder and

consolidated to full density by HIRping using the conditions specified above.

A typical microstructure of as HIRped Ni54Mo26B20 alloy is shown in Fig. 2. It consists of a fine grained nickel -base matrix together with a dispersion of boride platelets (some are marked B in Fig. 2a). The boride phase was identified by electron diffraction (Fig. 2b) and x-ray diffraction to

be orthorhombic Mo 2 NiB 2 , which is isostructural with M02C0 B 2

. The chemistry was further confirmed

by quantitative x-ray microanalysis (Fig. 2c) which gave a Ni :Mo ratio of ~ 0.6:1 in the boride particle. This ratio was found to hold for most of the alloys studied (see Table 2); the alloy containing 10% Fe was an exception in that a small amount of Fe substituted for Ni . It was found, in

addition, that the matrix composition varied with the alloy-composition (Table 2). Also alloys that

had a Mo:Ni ratio > 0.25 in the matrix , exhibited ordering reactions in the matrix, as shown in

Fig. 3 for Ni5oMo3oBio and Ni4gMo3iB20 alloys. In the sample quenched in iced brine from 1100°C

the matrix exhibits short range order (s.r.o.l as seen in the [100] diffraction pattern (Fig. 3b)

where diffuse scattering is observed at { IV2 O} positions. The presence of { IV2 O} s.r.o. diffuse scattering has been observed in binary Ni-Mo alloys containing 20 - 25at% Mo [6,7]. Upon annealing at 700°C fine ordered domains (50-100A) can be identified in the matrix as shown in the dark field

micrograph in Fig. 3(c). From the electron diffraction pattern (Fig. 3d) the domains can be identified as Ni 4 Mo. In addition to the reflections from two variants of Ni 4Mo there are some additional spots (marked by the arrow in the diffraction pattern in Fig. 3d), which are due to the

presence of a small volume fraction of Ni 2 Mo phase; this phase has been observed as a metastable

phase during early stages of ordering of binary Ni-Mo alloys [8]. On prolonged annealing metastable

Ni 2Mo phase disappears and Ni 4 Mo domains grow as shown in Figs. 3(e) and 3(f) for Ni4gMo3iB2o alloy.

The same samples annealed at 800 and 850°C for 5-50 hours show fully ordered equilibrium Ni 3 Mo phase

(not shown here).

Fig. 2 Bulk microstructure of as HIRped Ni54Mo26B20 alloy, (a) Bright field TEM,

(b) selected area diffraction pattern and (c) x-ray microanalysis of one of

the boride particles.

561
Table 2. X-ray Microanalysis of Ni-Mo-B Alloys

ALLOY

NO. NOMINAL BULK COMPOSITION(Atomic%)

MICROANALYSIS OF ORDERED PHASES PRESENT mm IN MATRIX

Ni :Mo :Fe BGITTMNi :Mo:Fe

1 Ni 54^026^20 1:0.2:- 0.6:1:- No

2 Ni 49^031^20 1:0.27:- 0.6:1:- Yes

3 Ni62M023Bl5 1:0.2:- 0.6:1:- No

4 Ni 60^030^10 1:0.3:- 0.6:1:- Yes

5 Ni 56. 5^023. sFeioBio 1:0.28:0.2 0.55:1:0.08 No

Fig. 3 Effect of heat

and (e) - (f)

in iced brine

using a Ni 4 Mo SAD of matrix

aged at 700°C

treatment on microstructure of (a) - (d) Ni5oMo3oBio slloy

Ni4gMo3iB20 alloy, (a) Bright field TEM of sample quenched

from 1100°C, (b) [100] SAD of matrix in (a), (c) Dark field

spot in (d) from a sample aged at 700°C for 20h, (d) [100]

in (c), (e) Dark field of Ni 4Mo spot in (f) from a sample

for 50h, (f) [100] SAD of matrix in (e).

I

(

562
Fig. 6 Tensile properties of as HIPped Ni 56.5FeioMo23.5

-Bio alloy as a function of temperature.

from these alloys are quite ductile.

Figure 6 shows hot tensile properties

for the alloy Ni 55,51^023. sFeigBio* At higher temperature (> 800°C) the material

becomes superplastic due to its extremely

fine grain size. This alloy has a room temperature transverse rupture strength of 490 Ksi and

charpy impact (unnotched) strength of 50 ft. - lbs. The alloys that exhibit ordered phases have

somewhat lower ductility than the ones that do not. However, by proper choice of heat treatment the

degree of ordering in the matrix and the domain size can be varied, which in turn affects the toughness. The effect of such heat treatments on impact toughness is shown in Table 3. The as HIPped

material is fully ordered due to the slow cooling from the HIPping temperature and exhibits an impact strength of only 4.0 ft-lbs. On quenching from 1100°C the matrix is short range ordered

(Fig. 3a) and the impact strength increases to 21.5 ft-lbs. (Table 3).

563

Elongation,

%
Table 3. Effect of Heat Treatment on Room Temperature Impact Strength (unnotched) of Ni5oMo3oBio

A1 1 oy

No. HEAT TREATMENT IMPACT STRENGTH

(ft-lb.)

ROOM TEMPERATURE

HARDNESS, Rc

1 As HIPped 4.0 62

2 Annealed 1100°C/2h> quenched 21.5 53

in iced brine

3 As in 2) + aged at 600° C for Ih 11.0 57

4 As in 2) + aged at 650°C for Ih 10.0 56

5 As in 2) + aged at 700°C for 50h 4.0 61

In addition to the good mechanical strength the fine dispersion of borides in a fine grained

matrix gives excellent wear resistance properties, a discussion of which is beyond the scope of this

paper.

SUMMARY

The results presented here show that rapidly solidified Ni-Mo-B metallic glasses can be fabricated into bulk crystalline solids with useful engineering properties, e.g.. high tensile strength,

and high hot hardness. These properties are derived from their unique microstructural features such

as fine boride dispersions with good high temperature stability and the presence of ordered phases

in a fine grained matrix. This class of alloys has great promise for use as hot working tools and

dies and in other wear resistance applications.

ACKNOWLEDGEMENTS

The authors are greatful to Ms. E. Musso for her experimental assistance and Mr. R. St. Amand

for generating some of the mechanical property data.

REFERENCES

1. L. A. Davis and R. Hasegawa, Metal 1 urgical Treatises , p. 301, Conference Proceedings, eds.

J. K. Tien and J. F. Elliot, The Metallurgical Soci ety of AIME, Warrendale, Pennsylvania,

U981).

2. L. A. Davis, N. J. DeCri stofaro , C. H. Smith, "Metallic Glasses: Science and Technology", p. 1, Vol . 1, Conference Proceedings, eds. C. Hargitai, I. Bakyoni and T. Kemeny, Central Res. Inst,

for Physics - Budapest (1981).

3. e.g. Rapid Solidification Processing, Principles and Technologies, II, eds. R. Mehrabian,

B. H. Kear and M. Cohen, Claitor's Publishing Division, Baton Rouge, Louisiana (1980).

4. R. Ray, Proceedings of 4th International Conference on Rapidly Quenched Metals, p. 1515, eds.

T. Masumoto and K. Suzuki, The Japan Institute of Metals, Sendai, Japan (1982).

5. J. M. Cowley, Proc. 13th Annual EMSA Meeting, p. 560, ed. C. J. Areeanaux, Claitor's Publishing

Division, Baton Rouge, Louisiana (1972).

6. P. R. Okamoto and G. Thomas, Acta Met. 825 (1971).

7. S. K. Das, P. R. Okamoto, P. M. J. Fisher and G. Thomas, Acta Met. 913 (1973).

8. S. K. Das and G. Thomas, Phys. Stat. Sol (a) 177 (1974).

564
MICROSTRUCTURAL MODIFICATION OF ALUMINUM-SILICON ALLOY

POWDERS DURING DYNAMIC CONSOLIDATION*

W. H. Gourdin

Lawrence Livermore National Laboratory

Livermore, CA 94550

J. E. Smugeresky

Sandia National Laboratories, Livermore

Livermore, CA 94550

ABSTRACT

We describe the preliminary results of a study of the dynamic consolidation of a rapidly solidified aluminum-6% silicon alloy powder. The distinctive microstructure of

this alloy, which consists of a dispersion of silicon precipitates at the boundaries of

essentially pure aluminum grains, provides an excellent grid against which to observe the

Changes induced by shock compaction. Examination of recovered specimens, shocked with

either explosives or high velocity projectiles, shows that at low stress, high density is

achieved without substantial interparticle bonding, and the microstructure is essentially

that of tne starting power. For stresses between 1.7 and 2.6 GPa we observe increasing

compact integrity, and the appearance of regions of altered microstructure at the

interfaces between powder particles. Above 4.0 GPa, such modified material is a distinctive feature of the microstructure, and compacts are well bonded. Increasing the

stress serves not so much to increase the density as to increase the amount of energy

deposited in the specimen. Bonding is the result of the local deposition of energy at particle interfaces during the passage of the shock. This energy may be calculated from

the measured Hugoniot (shock compression) curve of the powder. We suggest that through

proper choice of stress level, structurally sound compacts can be formed with limited

modification of the powder microstructure.

Introduction

Recent studies have shown that shock waves can be used to density metal [1-4] and ceramic [5-8]

powders at high stress in times of the order of microseconds. Such consolidation avoids the

thermally activated changes which accompany conventional thermo-mechanical processing. As such, the

technique is ideally suited to the processing of rapidly solidified metal powders which possess

unique or metastable microstructures. [9]

In this paper, we report the preliminary results of a study of the dynamic consolidation of a rapidly solidified aluminum-6% silicon alloy powder. Our observations suggest that the bonding of

consolidated material is a result of energy deposition at powder particle interfaces, which produces

localized melting and rapid resolidification. The amount of energy deposited can be determined from

the measured Hugoniot of the powder. Through proper choice of stress level, structurally sound

compacts can be formed with limited modification of the initial powder microstructure.

Experimental

The as-received, ultrasonical ly gas atomized A1 + 6% Si powder^ was put through a 60 mesh sieve to remove very coarse material. The properties of the resulting -250 ijm powder are given in Table 1.

*This work was supported in part by the Army Research Office and performed under the auspices of

the U.S. Department of Energy by the Lawrence Livermore National Laboratory under contract No. W-7405-Eng-48.

•prepared by P. Domalavage and N. J. Grant of MIT.

565
Table 1

Al + 6% Si Ultrasonical ly Gas Atomized Powder: Characteristics

Packing density: 1.78 gm/cc

Theoretical density: 2.691 gm/cc

Size distribution: % Less than size

Specific area: 1450 cm^/gm

Eutectic temperature, T,|,: 577°C

Size ( u)

100

7550

25

250

70

34

23

The size distribution was determined using standard sieve analysis and the surface area was measured by the BET method. The microstructure, shown in Figure 1, is cellular with silicon

precipitates at cell boundaries. A eutectic structure is also found in some particles. Lattice

parameter analysis shows the aluminum matrix to be essentially pure. We have found no evidence for

supersaturation within individual particles. However, splat caps (Figures 1, 5) at the surface of some particles show no precipitation, even though electron microprobe analysis shows approximately

the same average Si content as the phase-segregated particle. This suggests that these regions may

be supersaturated with respect to silicon, presumably because they cooled more rapidly than the

particles themselves. As discussed below, thermally modified regions observed following dynamic

consolidation have a microstructure similar to that of these splat caps.

Specimens were compacted using both high explosives and high velocity projectiles. In all cases, the geometry was planar, ensuring that passage of at least the first shock occurred under

conditions of one-dimensional strain. The compaction assembly used in the gas gun is shown

schematically in Figure 2. Impact of the aluminum projectile at speeds of 200-750 meters/sec

generates a shock which propagates through the powder. Carbon piezo-resistive sensors, placed at the front and rear of the powder, were used in the determination of the Hugoniot. A massive

stainless compaction chamber was used to contain the sample for later recovery.

Figure 1. Microstructure of the

as-received Al + 6% Si ultrasonical ly gas atomized powder. Note the "splat cap" on the larger particle.

Figure 2. Schematic illustration of the assembly

used in gas gun experiments.

566
The fixture used in explosive experiments is shown in Figure 3. A plane wave lens, 5.6 cm in

diameter, was used to generate a plane detonation wave in the explosive. This impinged on a copper

cover plate, which transmitted a shock into the powder, contained in a massive mild steel block.

The stress delivered to the copper cover was measured for each of three different explosives, PETN 0.88 qm/cc, PETN 0.95 gm/cc and HMX 1.1-1. 2 gm/cc. These stresses are in good agreement with those

calculated from the equations of state for these explosives. [10, 11] The stress delivered to the

powder is calculated from the known stress in the copper.

Results

In this paper, we consider several specific experiments, spanning 1. 3-4.0 GPa, which are representative of a range of powder response and compact structure. The details of these

experiments are summarized in Table 2, and their position on the powder Hugoniot is indicated in Figure 4.

Low Stresses

The microstructure of experiment 71 (1.25 GPa) is shown in Figure 5. The overall

microstructure is essentially that of the starting powder and the compact has little strength

despite its nigh density (Table 2). We find no evidence of structural modification from the shock

consolidation, other than material flow.

"Plane Wave" Compaction

block

Consolidated powder

Figure 3. Plane wave compaction using explosives.

Table 2

A1 + 6% Si Dynamic Compaction Experiments

Expt. Type Stress(GPa)* Final Density^ E(10^ erg/gm) Comments

71 Gun 1.25 98% 0.95 90 No bonding

A1 proj., 527

m/sec

95 Gun 1.65 99.6 1.51 144 A1 proj., 660 m/s

7P Expl. 1.8 97 1.56 148 PETN, 0.88 gm/cc

94 Gun 2.0 100 1.92 183 A1 proj., 754 m/s

IIP Expl

2.6 98 2.47 235 PETN, 0.95 gm/cc

15P Expl. 4.0 98 3.96 377 HMX 1.1-1. 2 m/cc

Well bonded compact

*Stress of the first shock in the powder.

'''Mercury porosimeter.

567
SPEC. VOLUME V/V0

Figure 4. Stress-volume Hugoniot measured for the

A1 + 6% Si powder. Stresses for the explosive

experiments were determined from the Known stresses

in the copper cover plate.

Figure 5. Microstructure of gas gun

experiment 71. Note the residual splat cap

surrounding a particle in the center of the

photograph.

The featureless regioh surrouhding one grain in Figure 5 is a splat cap, found in the starting

powder. Numerous cracks in the sample result from tensile waves generated during release to ambient

pressure.

Intermediate Stress

Between 1.7-1. 8 GPa (Figure 6a) and 2.6 GPa (Figure 6b) regions of modified microstructure

appear at particle interfaces in increasing numbers. Although most of specimens 94, 95 and 7P resemble Figure 5, some particle interfaces have the structure shown in Figure 6. The featureless

areas bear a strong resemblance to the splat caps mentioned above, but their turbulent appearance is

in marked contrast to the structure of Figure 5. Some fine precipitates are visible in the

interfacial region, but the general appearance and the presence of unaltered surrounding grains

suggests that the material has undergone local thermal modification. Smaller powder particles are

no longer distinct. The compacts themselves become increasingly well bonded in this stress range,

and can oe handled without difficulty.

High Stress

For stresses >4.0 GPa (15P) considerable modification of the microstructure is evident, as Shown in Figure 7. Altered material now forms a nearly continuous phase between and around powder

| particles and only the larger particles retain a semblance of their original microstructure. ( Modified regions show a dispersion of fine precipitates and a generally coarser structure than at intermediate stresses (Figure o). Some areas show severe deformation. Despite large ruptures

resulting from late-time distortion of the compaction fixture, sample 15P could be broken only with i difficulty. \

Discussion

The contrast with the original powder microstructure as well as the similarity to the splat

caps suggests that the modified microstructures are a result of thermal alteration which occurs

during the passage of the shock wave. Such modification is apparently local, occurring primarily at

568
Figure 6. Microstructures of compacts formed in the intermediate stress region, 1.7-2. 6 GPa:

(a) 7P, formed with PETN explosive, 0.88 gm/cc; (o) IIP, formed with PETN, 0.95 gm/cc.

Figure 7. Microstructure of compact 15P,

formed at 4.0 GPa using HMX explosive,

1.1 gm/cc.

powder particle interfaces. If the observed effects are indeed thermal, the energy deposited into

the powder by the shock is of particular interest. From the Hugoniot of Figure 4, we can estimate

that the density behind the first shock is close to solid in all experiments. Subsequent shocks

deposit substantially less energy into the compact. Under such circumstances the energy deposited

in the powder is just:

E = Eg - Er (1)

Where Eg is the energy of the initial shock,

Es = n/2) P (Vq-V) (2)

569
and is the energy regained upon release. Because the powder is close to solid density, Ep is generally small when compared with E

5 , Out may be estimated from the Hugoniot of solid aluminum.

The approximate energies E, for each of the experiments shown in Figure 4, are collected in Table 2. Apparently energies in excess of ~2xl0^ ergs/gm are required to produce the observed microstructural

Changes. This is significantly less than the energy required to heat and melt the entire mass, ~9xl0^ ergs/gm, and the average temperature rise (Table 2) is substantially less than the

eutectic, Tn, (Table 1). However, if we imagine that this energy is deposited primarily through

the particle surfaces, as a result, for example, of frictional heating as the particles deform

around eacn other, the surface temperature can easily exceed T,^, producing a layer of molten

material at the interface. A simple model for calculating the surface temperature of a typical

powder particle has been recently developed and has been found to provide a satisfactory explanation

of local microstructural changes observed in 4330V steel compacts. [12] Application of this model to the A1 + 6% Si powder studied here shows that surface temperatures are expected to exceed Tm at a stress of 2.25 GPa, for which E = 1.96 x 10^ ergs/gm. These values define a critical state, above

which interfacial melting begins. The correspondence between our observations and the results of

the model calculation suggest that microstructural changes and bonding in the compacts are due

primarily to localized melting followed by rapid resolidification. If the molten layer is thin, we

can expect cooling to be rapid, much like a splat cap. As the energy input increases, the average

temperature rises (Table 2), more material will melt and the cooling rate will decrease. The fine

precipitates and generally coarser structure of sample 15P relative to IIP may reflect this.

The surface temperature achieved during the passage of the shock wave depends not only on the

thermal properties of the material, but upon the characteristics of the powder as well.

[

12 ] Different powders, even of the same material, will, in general, begin to show interparticle melting

at different stress levels. Similarly, at a given stress level, the final density of different

powders may differ. Density and bonding are thus separate, though related, issues in dynamic

consolidation. One can achieve high density in the absence of bonding, as demonstrated by the

microstructure of shot 71 (Figure 5). However, above a critical value, a stress can be chosen so

as to Obtain both high density and adequate bonding in a compact while minimizing alteration of the

initial powder microstructure.

Summary/Conclusions

We have found that at low shock stresses, 4,1.3 GPa, dynamically consolidated A1 + 6% Si powder

has a high density but very little strength. The microstructure is essentially the same as that of

the powder. As the stress is increased to 1.7-2 .

GPa, interparticle bonding begins to occur, and

for stresses in excess of 4.0 GPa the compact has substantial strength. Bonding of the compacts

apparently results from melting and rapid resolidification at the surfaces of powder particles.

Such modified material begins to appear at intermediate stresses, and for high stresses ^4.0 GPa,

forms essentially a continuous feature. Tne energy deposited at the interfaces can be calculated

from the measured Hugoniot of the powder.

From this work, we conclude the following:

1) The bonding in a dynamically consolidated specimen results from rapid energy deposition at powder particle interfaces. It is distinct from, but related to, the compact final density.

2) For the A1 alloy powder studied here, bonding begins to occur between 1.7 and 2.6 GPa. Stresses in this range or in excess of 2.6 GPa are required to produce structurally sound,

nigh density material.

3) A critical stress level appears to be necessary to achieve both high density and adequate

cohesion.

Acknowledgments

The authors wish to thank P. Curtis for preparing the metal lographic specimens and S. L. Weinland and T. Sullivan for performing the gas gun experiments. F. Helm performed the explosive

experiments

570
References

[Ij D. RayDould, D. G. Morris, and G. A. Cooper, J. Mat. Sci ., 14:2523 (1979).

L2J D. Rayoould in Shock Waves and Hign-Strain-Rate Phenomena in Metals , M. A. Meyers and L. E. Murr, Editors, Plenum Press, New York, 1981.

[3j 0. Rayoould, J. Mat. Sci ., 16:580 (1981).

[4J D. G. Morris, Metal Sci ., March 1981, p 116-124.

[5j R. A. Prummer, Ber. Dt. Keram. Ges ., 50:75 (1973).

[bj C. L. Hoenig and C. S. Yust, Amer. Cer. Soc. Bull. , 60:1175 (1981).

[7J W. H. Gourdin, C. J. Echer, C. F. Cline, and L. E. Tanner in Proc . 7th International Conf. on High Energy Rate Fabrication, University of Leeds, 1981, p 233-241.

[8J W. H. Gourdin, S. L. Weinland, C. J. Echer, and S. L. Huffsmith, to be published.

[9J 0. G. Morris, Metal Sci ., June 1980, p 215-220.

[lOj B. M. Dobratz, LLNL Explosives Handbook: Properties of Chemical Explosives and Explosive

Simulants , Lawrence Livermore National Laboratory UCRL-52997.

[llj E. L. Lee, private communication.

[12J W. H. Gourdin, to be published.

571
MICROSTRUCTURE OF RAPIDLY SOLIDIFIED STAINLESS STEEL AND SUPERALLOY POWDERS:

A COMPARISON OF POWDER-MAKING TECHNIQUES

John E. Smugeresky

Sandia National Laboratories

Livermore, CA 94550

ABSTRACT

Several complex ferrous alloys (1-903, JBK-75, HP-9-4-20, and managing steel) in the

form of either Inert Gas Atomized (IGA) or Centri fugal ly Atomized (CA) powder have been

characterized and compared. Powder particle size, microstructure, morphology and packing

density were used for the characterization. Definite differences in particle size and

microstructure are generally observed, while particle morphology and packing densities are similar. Particle sizes can be separated and/or normalized by mechanical separation and

differences in microstructure are also observed for differences in powder particle size for

each technique. The microstructures are those of rapid solidification and are compared to those of slower conventional ingot metallurgy solidification. Because of the rapid solidification of atomization, the microstructure of the powder particles are metastable and

consequently are easily transformed by the thermomechanical process of consolidation. The

relation between powder-making techniques, powder characteristics, and properties in the

consolidated condition are summarized.

Introduction

Complex alloys containing many elements are often prone to compositional variation caused by

solute segregation during solidification, especially in large commercial ingots. This phenomenon,

commonly known as "coring," occurs by solute rejection, which leaves solute-rich regions between dendrite arms. Subsequent homogenization can in some cases minimize the segregation but it ususally is accompanied by undesirable excessive grain growth. Consequently, additional thermomechanical processing to achieve recrystallization is required to obtain smaller grain sizes.

Alternatively, reducing the size of the ingot or providing chilled molds to enhance heat extaction can reduce solute segregation at dendrite interstices by increasing the cooling rate. Atomization techniques for making metal powders can in effect reduce the "ingot" size to less than 250 uii by

producing molten droplets that solidify in the absence of a mold, and are thought to provide cooling

rates in excess of 10*^ K/sec. Thus, rapid solidification by atomization can provide significant

improvements in compositional homogeneity of complex alloys (see Fig. 1).

Inert gas atomization (IGA) and centrifugal atomization (CA) are two rapid solidification powder

making techniques which are believed to provide sufficiently different cooling rates and/or powder

characterisics such that differences in properties of the consolidated powder should be observed. In this investigation, four iron-base complex alloys were atomized by IGA and CA to compare the

techniques.

Experimental

The compositions of the alloys are given in Table I. IGA was done either by Udimet Powder

Division of Special Metals or by Alloy Metals using argon gas. These powders were subsequently

screened to remove particles larger than 250 pm. CA was done by Pratt & Whitney. These powders were screened to remove particles larger than 177 pm. Particle size distributions were determined by a screen analysis while apparent and tap densities were determined using ASTM techniques B212-76 and

B527-70. Powders were consolidated by hot isostatic pressing (HIP) in evacuated and sealed stainless steel containers at 900 and 1200°C and 103 MPa for three hours. Since these alloys are precipitation hardened, post-HIP treatments were utilized prior to mechanical property measurements.

572
Table I. Nominal Alloy Composition

JBK-75 1-903 HP 9-4-20 250 Maragi

Cr 15 - <1 -

Ni 30 38 9 18

Ti 2 1.4 - 0.3

A1 0.2 0.7 - 0.04

C 0.02 - 0.2 0.02

Co - 15 4 8.4

Mo 1.3 - 1 4.5

Fe Bal Bal Bal Bal

Nb 3.0 _

Results

In the space provided here, it is not possible to include all of the results to date. However,

I will highlight the data to summarize the general characteristics of the powders so that a comparison can be made between IGA and CA. Powder characteristics are listed in Table II. The IGA powders

have a wider particle size distribution containing significantly more particles less than 37 pm than

the CA powders. Although the apparent densities depend on the technique used, the tap densities are

not. The strengths of compacted powders HIP at 1200°C are generally similar to those of ingot

metallurgy (I/M) processed material and independent of powder technique except for 1-903, while ductility, as measured by RA, is similar to I/M material except for 1-903 and compacts of 250 managing

steel made from CA powder. The as-solidified microstructure of the powders (Figs. 2-5) is dendritic

for CA powder but is cellular for IGA powder. The smaller the powder particle is the finer its

microstructure. For the IGA JBK-75 powder the finer particles show evidence of dendrites.

In all cases, the as-solidified structure was transformed during HIP at 1200°C and post-HIP

heat treatment to that similar to conventional I/M products except for a slightly finer but more uniform grain size and prior powder particle boundary (PPPB) precipitate decoration (Figs. 6 - 8 ). These

PPPB precipitates have been identified as TiC in JBK-75,^ M3 C in HP 9-4-20,^ and Ti , Ni , Mo rich

precipitates in 250 grade maraging steel. ^ However, for JBK-75 the use of lower HIP temperatures has

resulted in microstuctures and mechanical properties that depend on powder technique. Figure 9 summarizes the yield strengths, ductility, and density while Figure 10 summarizes the microstructure.

Discussion

The microstructure of rapidly solidified powders is affected by cooling rate (aT), thermal gradient (G), and solidification rate (R) in a manner described previously > » . Decreases in particle size generally result in finer microstructure but not necessarily a change from cellular to dendritic, suggesting that the G to R ratio remains constant for all particle sizes. However, the G to R ratio is lower for CA than for IGA as evidenced by the dendritic structures of the CA powder

versus the celluar structure of the IGA powder. Dendrite arm spacing of the CA powder is significantly smaller than that of an ingot (see fig. 1). This suggests that the G to R ratios of CA and

conventional castings are similar, but the aT's are dramatically different. Although extrapolation

of dendrite arm spacing versus cooling rates for conventional castings to that for powders is a tempting approach for estimating the cooling rate in the powder, the extrapolation range is too large

to expect the estimate of cooling rate to be any better than that using a random number generator.

The metastable character of the as-solidified structures is demonstrated by the transformed

microstructures of HIP samples. Temperatures in the 1200°C range have routinely been used to ensure

adequate interparticle bonding during densification.^ In these cases, all evidence of the

as-solidified structure has been wiped out, and microstuctures of HIP samples cannot be used to distinguish powder-making techniques. Fractography of HIP samples confirms that true metallurgical

573
Figure 1. Comparison of As-Solidified Structure in (a) Conventional Ingot and (b) Centri fugal ly

Atomized Powder Particle "Ingot"

Figure 2. Microstructure of JBK-75 (a) IGA Powder (b) CA Powder

,t i_.

\

' ' '' ?

tM- • • ' -< • -r.' .-Li c'-\ ( r , ^ - ' • -lyi

Figure 3. Microstructure of 1-903 (a) IGA Powder (b) CA Powder

574
Figure 4. Microstructure of HP-9-4-20 (a) IGA Powder (b) CA Powder

Figure 5. Microstructure of 250 Grade Maraging Steel: (a) IGA Powder (b) CA Powder

Figure 6 (Left). JBK-75 After Consolidation at 1200°C. Microstructure is similar to that of ingot

metallurgy material except for prior powder particle boundary precipitates (arrows). Figure 7 (Middle). 1-903 After HIP at 1200°C. Figure 8 (Right). 250 Grade Maraging Steel After HIP at 1200°C.
900

“I 1 1

1 I I 1 aoo 700 800 900 1000 noo 1200 1300 HP T (°a

g

&(/)

Figure 9. Yield Strength (a). Density (b, top curves), and Density (b, bottom curves) of Compacts of JBK-75 in the Solution Treated and Aged Condition as a Function of HIP Temperature.

Figure 10. Microstructure of JBK-75 After HIP at 900°C. (a) IGA Material. Note evidence of

both the as-solidified structure and transformed fine grain structure of this fully dense

compact, (b) CA Material. Note the presence of porosity and the as-solidified structure as well as the absence of movement of prior powder particle boundaries.

576
TABLE II

Powder Characteristics

250 Grade

JBK- 75 I- 903 HP-9- 1 ro o Maraqe

IGA CA IGA CA IGA CA IGA CA

Max. Particle Size (vm) 250 177 250 180 250 212 250 212

Avg. Particle Size (ym) 90 75 70 97 100 85 96 96 Amt. Powder (<37 ym (pet) 15 3 21 1 18 1 12 1 Morphology* SPH+SAT SPH SPH+SAT SPH SPH+SAT SPH SPH+SAT SPH

Apparent Density (% T.D.) 49 59 53 57 41 58 49 60

Tap Density (% T.D.

64 65 66 65 55 65 58 68

*SPH = Spherical

SPH+SAT = Spherical with satellites

TABLE III

Powder Compact Properties*

JBK-75 1-903 HP-9-4-20 250 Grade Maraqe

IGA I/M CA IGA I/M CA IGA I/M CA IGA I/M CA

UTS (MPa) 1130 1135 1116 1185 1410 1315 1440 1720 1750 1655

YTS (MPa) 758 820 744 955 1200 1220 1245 1710 1680 1550

^T (Pet) 24 22 25 7 14 13 16 N/A 14 4

RA (Pet) 50 45 49 12 38 55 63 52 62 14

*After HIP at 1200°C and heat treatment

bonding occurs when transparticle-transgranular fracture is observed. After HIP at lower temperatures, interparticle fracture is observed in some cases. Although interparticle fracture is at first glance faceted like intergranular fracture, the facets usually consist of "planar" ductile dimpled fracture characterized by very small, very closely spaced dimples. In several cases the dimples

are associated with fine precipitates which form either on the powder surface prior to interparticle

contact or at the interface during consolidation.^ In cases where the prior powder particle boundary

becomes a migrating grain boundary, interparticle failure is generally not observed. Consequently,

control of precipitation at PPPB's during consolidation and control of grain boundary migration are important in obtaining desirable properties and/or fracture mode in rapidly solidified and powder

processed material.

In contrast, using lower HIP temperatures for JBK-75 demonstrates that powder-making techniques

which provide quite different microstructures can provide consolidated material with quite different

properties and microstructure. For example, HIP at 900°C provides good strength and ductility in

compacts made from IGA powder, but provides reduced strength and negligible ductility in compacts

made from CA powder. The difference is not only due to differences in density (98% theoretical density for CA and 100% theoretical density for IGA) but also to a lack of metallurgical bonding between

CA powder particles. The dendritic structure of the CA powder is apparently stronger and more stable

than the cellular structure of the IGA powder. Being stronger, the CA powder does not compress to

full density during HIP at 900°C, and less than fully dense compacts have lower properties. However,

such a drastic drop in ductility (from 32 to 3% RA) with only a 2% drop in density indicates the lack

577
of metallurgical bonding between powder particles. As seen in Fig. 10a, the PPPB's are still very

evident in CA powder compacts, but quite difficult to detect in IGA powder compacts (Fig. 10b). In the latter case, the PPPB's are no longer individual grain boundaries, but have probably transformed

into a series of small grains such that a trace of the PPPB passes through the interior of the grains

rather than at the grain boundaries. Thus, regions that were formerly boundaries are actually interiors of single grains and thus represent truly metallurgically bonded material. Since this is not the case for CA powder compacts, they lack evidence of true metallurgical bonding.

Despite the transformation of the metastable as-solidified cellular structure, we see in

Fig. 10b evidence of how to take advantage of rapidly solidified powders. In general the transformed structure contains very fine grains (approximately 5 to 10 times smaller than for powder HIP

at 1200°C). By stopping the transformation at this time instead of allowing grain growth to occur,

grain size strengthening is possible. Quenching from the HIP temperature can provide additional metastability of a supersaturated solid solution in precipitation strengthened alloys such as this so that subsequent aging using nonconventional heat treatments can provide increased strength. This

hypothesis for exploiting rapidly solidified powders is presently under investigation.

I

! I Summary/Conclusions j

Despite differences in microstructure of several ferrous alloy powders dependent upon rapid

solidification powder making techniques, use of high temperatures during HIP wipes out all remnants of the as-solidified structures, and results in compacts with similar properties.

Use of lower HIP temperatures, compacts are obtained with microstructures and properties

dependent upon powder-making technique. IGA powders provided higher strength, ductility, and density

j after HIP at 900°C than CA powders. Although the microstructure in all cases transforms during HIP, , the resulting structures for low temperature HIP are much finer than those of I/M material.

'

It is hypothesized that by combining HIP and post-consolidation heat treatments to control ' grain size and precipitate characteristics, rapidly solidified powder structures can be ! exploited to take advantage of differences in cooling rate between powder-making techniques. |

Acknowledgements i

This work was sponsored by DOE contract #DE-AC04-76DP00789; the CA powders were provided | through DARPA.

^

References I

1. J. E. Smugeresky, Met. Trans A, 13A (1982), 1535.

| 2. T. F. Klimowicz, SAND81-8229, June, 1981. ^

fi 3. R. M. German and J. E. Smugeresky, Met. Trans A, 9A (1978), 405.

j1 4. M. Cohen, B. H. Kear and R. Merabian, in "Rapid Solidification Processing, Principals and

| Techniques, II," Claitor's Publishing Division, 1980, pg. 1. !

5. S. R. Coriell and R. F. Sekerka, ibid, pg. 35. f

6. J. E. Smugeresky and R. M. German, Mat. Sci. & Eng. 36 (1978), 222. f

7. J. E. Smugeresky and R. M. German, Met. Trans A, 9A (1978), 253. |

8. J. L. Bartos and P. S. Mathur, in "Superalloys: Metallurgy and Manufacture," Proceedings ^

of Third International Conference, September, 1976. |

9.

R. H. Van Stone and R. D. Gangloff, Reference 4, pg. 317

10.

C. C. Law and M. J. Blackburn, Modern Developments in Powder Metallurgy, Vol . 14,

APMI, Princeton, NJ, 1980, pg. 93.

1

i

1

i

578
CONSOLIDATION, THERMOMECHANICAL PROCESSING, AND MECHANICAL PROPERTIES OF

RAPIDLY SOLIDIFIED, DISPERSION-STRENGTHENED TITANIUM ALLOYS

S. M. L. Sastry, T. C. Peng, and J. E. O'Neal

McDonnell Douglas Research Laboratories

St. Louis, Missouri 63166

and

L. P. Beckerman

Air Force Wright Aeronautical Laboratories

Dayton, Ohio 45433

ABSTRACT

The feasibility of using rapid solidification processing techniques to produce

dispersion-strengthened titanium alloys for use at temperatures up to 700°C was investigated. Titanium alloys containing large volume-fractions of fine (< 100 nm diam)

incoherent dispersoids were produced by* splat quenching electron-beam-melted Ti-Er and TiNd alloys. The splat-quenched flakes were consolidated by vacuum hot pressing at 800-

850°C at stresses of 56 MPa (8 ksi) and 70 MPa (10 ksi). The consolidated alloy disks were thermomechanically processed by rolling at 25-700°C and annealing at 500-700°C. The

microstructures, hardness values, and tensile properties of thermomechanically processed

alloys were determined. Significant increases in tensile and creep strength were observed

in dispersoid-containing alloys. The effects of dispersoids on recrystallization, grain

growth, and tensile and creep properties were determined.

Introduction

Dispersion strengthening is an effective method for increasing the useful service temperature of an alloy system. Production of dispersion-strengthened Ti alloys by conventional ingot casting and

processing methods has been only partially successful [1-3] because additions of large amounts of

dispersoid-forming elements result in small volume-fractions of coarse, equilibrium, constituent particles, which are ineffective for dispersion strengthening.

The solubility of potential dispersoid-forming elements in Ti can be increased significantly by

increasing solidification rates to above 10^ K/s, and the solid solutions can be annealed subsequently

at suitable temperatures to obtain incoherent dispersoids [4-6]. Binary Ti-Er and Ti-Nd alloys were

evaluated for the production of fine, incoherent dispersoids by rapid solidification. The thermal

stability of dispersoids at 700-1000°C ( 1292-1832°F) and the effects of dispersoids on deformation,

recovery, and recrystallization, and strengthening of Ti were determined.

Experimental Procedure

5-kg ingots of Ti-1.5Nd and Ti-l.OEr and 5 kg of 12.5-mm (0.5 in.) diam rods of Ti were procured

from RMI Co. The Ti-Nd and Ti-Er alloy ingots were made from high-purity titanium (maximum oxygen

concentrations of 0.08 wt%) and erbium and neodymium chips supplied by Research Chemicals, Inc. Each

alloy was doubly melted by consumable-electrode arc melting, first in vacuum and then under argon gas

pressure. The ingots, 7.5-cm (3.0 in.) diam and 40-cm (16 in.) long, were extruded to 12.5-mm

(0.5 in.) diam rods by RMI Co. The extruded rods were used as feed stock for the production of

rapidly solidified flakes by electron-beam melting and splat quenching [7]. The rapidly solidified

flakes, which were 75-150 pm thick and had cooling rates of « 10^ K/s, were reduced by a rotary sheetmetal cutter to 2 X 2 mm particulates. The particulates were cleaned with acetone and a chem-mill

solution consisting of hydrofluoric acid, nitric acid, and water, and consolidated into 50-mm (2.375

in.) diam and 90-120 mm (0.36-0.5 in.) thick disks by vacuum hot pressing at 800-850°C at stresses of

579
70-84 MPa (10-12 ksi) for 2-16 h. The consolidated disks were examined for complete densif ication by

density measurements and optical metallography and scanning electron microscopy of through-thickness

sections

The consolidated Ti-alloy disks were isothermally forged at 800°C to obtain 4.0-mm (0.16-in.)

thick plates suitable for further rolling. The high temperature and slow strain rates of deformation

used in the isothermal forging operation produced fully recrystallized microstructures. The 4.0-mm

(0.16-in.) thick plates were rolled at 25, 500, and 700®C with 10, 20, 30, 40, and 50% thickness

reductions. The thickness reductions and temperatures selected were Intended to establish the dependences on strain and temperature of the degree of recrystallization and the recrystallized grain size

Tensile properties at 25-700°C (77-1292°F) and creep properties at 482®C (900°F), 600°C (1112°F)

and 700°C (1292°F) were determined using test specimens 62.4-mm (2.5-in.) long having a 1.5 x 4.1 mm

(0.06 X 0.16 in.) reduced-gauge section.

Results and Discussion

Dispersoids in Ti-Er and Ti-Nd Alloys

Typical dispersoid distributions in rapidly solidified Ti-l.OEr and Ti-1.5Nd alloys annealed at 700°C for 2 h and 100 h are shown in Figures 1 and 2. Whereas the ingot-processed alloys contained

coarse (> 0.5 pm) particles, the rapidly solidified and annealed alloys contained 40-100 nm diam

dispersoids with interparticle spacings of 90-200 nm. The Ti-1.5Nd alloy also had a small volumefraction of coarse (> 0.1 pm diam) dispesoids with interparticle spacings greater than 1 pm. Rapid

solidification refines the dispersoids by at least an order of magnitude in the alloys.

Figure 1. Microstructures of rapidly solidified Ti-l.OEr annealed at (a) 700°C for 2 h and (b) 700°C

for 100 h.

1 jLtm (b)

1 (a)

Figure 2. Microstructures of rapidly solidified Ti-1.5Nd annealed at (a) 700°C for 2 h and (b) 700°C

for 100 h.

580
Effects of Dlspersoids on Recrystallization and Graia Growth of Ti

The Ti-l.OEr and Ti-1.5Nd alloys rolled at 25°C with a 10% thickness reduction were partially

recrystallized upon annealing; all other processing and annealing conditions resulted in complete

recrystallization. The partially recrystallized Ti-l.OEr alloy had a dislocation cell structure with

a cell size of « 1 pm. The 700°C annealing of differently processed alloys resulted in complete

recrystallization with smaller recrystallized grain sizes in Ti-Er and Ti-Nd alloys than in Ti.

To produce partially recrystallized or unrecrystallized microstructures by rolling and annealing

Ti-l.OEr and Ti-1.5Nd, the amount of deformation should be « 10% and the annealing temperature should

be « 600°C. Higher reductions and annealing temperatures result in complete recrystallization.

Because of the success of retaining the partially recrystallized microstructure in Ti-l.OEr by rolling

at 25°C with 10% strain and annealing at 600°C, the Ti and Ti-l.OEr specimens were subjected to four

cycles of cold rolling at 25°C and annealing at 600°C and 650°C. Ti-l.OEr specimens subjected to the

cyclic cold-rolling and annealing have elongated dislocation cells (Figures 3) with the short dimension of the cells » 1 pm. The unrecrystallized fibrous-grain structure typically found in dispersionstrengthened nickel and aluminum alloys could not be produced in rolled and annealed dispersionstrengthened Ti-l.OEr and Ti-1.5Nd alloys.

Effects of Dispersoids on Tensile and Creep Properties

Tensile properties at 25°C of differently processed Ti, Ti-l.OEr, and Ti-1.5Nd alloys are shown

in Figures 4a-4c.

Specimens cold rolled with 10% thickness reduction have a substantial amount of internal energy

in the form of dislocation tangles and cells. Specimens annealed at 600°C have recovered dislocation

substructures with finer cell sizes in Er- and Nd-containing alloys than in Ti. Specimens annealed at

650 and 700°C have fully recrystallized grains. Annealing at 700°C for 100 h was performed to determine the effect on tensile properties of dispersoid coarsening that may occur at temperatures and

times encountered in service use of this class of alloys. Under all processing and annealing conditions, the Ti-Er and Ti-Nd alloys have significantly higher yield strengths and ultimate tensile

strengths than in Ti. The strength increments are greatest in the fully recrystallized specimens.

The magnitudes of the strength increments depend on concentration of the rare-earth alloying element and the thermomechanical processing schedule. The numbers above each bar in Figure 4a are the

strength increments observed in Ti-Er and Ti-Nd alloys expressed as percentages of the unalloyed Ti

values. The Ti-Er and Ti-Nd alloys annealed at 700°C for 100 h have the highest strength increments,

but the ductilities in this case are significantly lower (Figure 5c). Annealing at 700°C for 2 h produces the best combination of strength and ductility in Ti-Er and Ti-Nd alloys. Annealing at 700°C

for 2 h produces 402 MPa (66.1 ksi) yield stress and 25.1% total elongation in Ti-2.0Er (a 76% in- crease in strength and 7.7% decrease in ductility compared with Ti) and 553 MPa (79.4 ksi) yield

stress and 18.4% total elongation in Ti-3.0Nd (111% increase in strength and 43% decrease in ductility

compared with Ti).

1 um

(a) (b)

Figure 3. Dislocation cell structure in rapidly solidified (a) Ti and (b) Ti-l.OEr subjected to repeated 10% cold rolling and 650°C annealing cycles.

581
Figure 4

40

35

30

5i :o o

Ti ^ Ti-l.OEr Ti-1.5Nd

(c)

ll

^ ]

i

' *• mh-i

1

1

pr.. Pi

PP

Cold rolled

\. iih

U'^O : Clili-. ' i..i

Cold rolled

IQt^o

n h

Cold rolled 10%

650^C/2 h

Annealed at 700=C/2 h

i

i:

Pl?

i!

Annealed at 700=C/100 h

Effect of thermotnechanical processing on (a) yield stress, (b) ultimate tensile strength,^

and (c) ductility of rapidly solidified Ti, Ti-l.OEr, and Ti-1.5Nd. t

582
The steady-state creep rate and stress rupture lifetimes of Ti-l.OEr and Ti-1.5Nd alloys are

shovm in Figures 5 and 6. The Ti-l.OEr and Ti-1.5Nd alloys have significantly lower creep rates than

titanium, with the effect being more pronounced at 700°C than at the lower temperatures. Of the three

alloys, Ti-l.OEr has the lowest creep rate at all temperatures and stresses investigated. The

dispersoid-containing alloys have order-of-magnitude longer stress-rupture lifetimes than Ti (Figure

6). For example, upon creep deformation at 600°C and 19 MPa (28 ksi), dispersoid-f ree titanium failed

in 12.5 h whereas Ti-1.5Nd sustained 68 h and Ti-l.OEr sustained 288 h before fracture. Similar

improvements were observed at 482 and 700°C. The proximity of plots of stress-dependence of stressrupture time of Ti-l.OEr at 600°C (1112°F) and Ti at 700°C (1292°F) in Figure 6 clearly demonstrates

the 100°C (200°F) increase in service temperature capability obtained by dispersions in Ti

S^::SS (k<;i) 12 5 10

S;.e^5 (MPa)

Stress (ksi)

Stress (MPa)

(a) (b)

Figure 5. Effect of dispersoids on stready-state creep rate of Ti at (a) 600°C and (b) 700°C.

Figure 6. Effect of dispersoids on stress-rupture lifetimes of rapidly solidified Ti, Ti-l.OEr, and

Ti-1.5Nd.

583
1. Conclusions

Rapid solidification at cooling rates > 10^ K/s produces nonequilibrium, supersaturated, solid

solutions and fine (< 100 nm) dispersoids in Ti-l.OEr and Ti-1.5Nd alloys.

2.

The equilibrium, coarse, constituent particles, which are generally observed in conventionally

cast Ti alloys containing 1-2 wt% Er and Nd, are totally absent in rapidly solidified alloys.

3.

Compacts with > 99% theoretical densities are produced by vacuum hot pressing rapidly solidified

particulates at 820-840°C and stresses > 70 MPa (10 ksi).

4.

Thermomechanical processing by rolling at 25-500°C and annealing at 600-700°C does not produce

fibrous-grain microstructures typically found in dispersion-strengthened nickel-base and aluminumbase alloys

5.

Annealing at 600°C of rolled Ti-l.OEr and Ti-1.5Nd alloys results in partially recrystallized

microstructures with 1-pm dislocation cells. Upon annealing above 650°C, the alloys readily

recrystallize to equiaxed-grain microstructures with finer grain sizes in Ti-Er and Ti-Nd than in

titanium.

6.

The incoherent dispersoids in Ti-Er and Ti-Nd alloys increase by 40-110% the yield strength and

ultimate tensile strengths of Ti with no significant loss in ductility.

7.

At 700°C (1292°F), the target service temperature for dispersion-strengthened Ti, the dispersoidcontaining alloys have 40% higher yield strength and ultimate tensile strength than dispersoidfree titanium.

Acknowledgements

This research was conducted under Air Force Contract No. F33615-81-C-501 1

References

[1]

F. R. Morral, Dispersion Strengthening of Metals, Batelle Columbus Laboratories MCIC Report MCIC77-30 (1977).

[2]

C. F. Dixon and H. M. Skelly, Dispersion Strengthening Titanium with Refractory Oxides, Canadian

Met. Quart. 11, 491 (1972).

[3]

M. J. Buczek, G. S. Hall, S. R. Seagle, and H. B. Bomberger, Grain Refinement of Titanium Alloys,

AFML-TR-74-2557 (Nov 1974).

[4]

J. E. O'Neal, T. C. Peng, and S. M. L. Sastry, Microstructural Studies of Laser Melted TitaniumRare Earth and Titanium-Boron Alloys, in Proc. Electron Microscopy Soc. Amer., ed. by G. W. Bailey (Claitor's Publ. Div., Baton Rouge, LA, 1981), p. 66.

[5]

S. M. L. Sastry, T. C. Peng, J. E. O'Neal, and L. P. Beckerman, Dispersion Strengthening of

Titanium Alloys, TMS-AIME Fall Meeting, St. Louis, MO, 24-28 October 1982.

[6]

P. J. Meschter, J. E. O'Neal, and S. M. L. Sastry, Characteristics and Thermal Stability of

Dispersed Phases in Rapidly Solidified Titanium-Rare Earth Alloys, TMS-AIME Fall Meeting,

St. Louis, MO, 24-28 October 1982.

[7]

T. C. Peng, S. M. L. Sastry, and J. E. O'Neal, Rapid Solidification Processing of Titanium

Alloys, Third Conf. on Rapid Solidification Processing, Gaithersberg, MD, 6-8 December 1982.

il

584
EFFECT OF AGING ON THE ^11 CROSTRUCTURE OF RAPIDLY COOLED

Ti-5AI-2.5Sn WITH Si ADDITIONS

A. C. Jackson*, T. F. Broderick**, F. H. Froes** and J. Moteff***

* Systems Research Laboratories, Inc. 2800 Indian Ripple Road

Dayton, OH, 45440

** Air Force Wright Aeronautical Laboratories

Materials Laboratory (AFWAL/MLLS)

Wright-Patterson Air Force Base, OH, 45433

*** University of Cincinnati

Cincinnati, OH, 45221

ABSTRACT

The structure of splat-cooled Ti-5AI-2.5Sn with Si additions was characterized

as-produced and in various heat treated conditions. Additions of 0.5 or 1.0

wt.pct. Si produce a martensitic structure, while a cellular structure is produced

with 5 wt.pct. Si. After heating at 535°C (1000°F) for 5, 10, 20, and 40h, the

microstructural changes observed in the 0.5 wt.pct. Si alloy were a transformation

of the martensite to alpha, grain growth, and a small amount of precipitation. In the 5 wt.pct. Si alloy similar heat treatments produced a decomposition of the

cellular structure to a dispersion of small precipitates. Electron diffraction and

EDS analysis indicated that the precipitates were probably Ti^Si^. Similar alloys

prepared by pendant drop melt extraction show a more dendriuc microstructure in the as-produced condition. The results are discussed in terms of the solidification

and cooling rate characteristics of the two processes used.

I. INTRODUCTION

In contrast to aluminum-based materials and superalloys, titanium powder metallurgy (PM) has

been driven by cost considerations rather than improved performance (1). This situation is now

changing, however, with PM/Rapid Solidification Technology (RST) of Ti alloys receiving attention

because of the potential improvements in mechanical properties due to characteristics such as increased solubility and/or improved homogeneity. This should be particularly true for

strengtheners such as Si, for which the equilibrium solubility is low. In the present paper,

microstructures of splat-cooled and pendant drop melt extraction (PDME) (2-5) Ti-5AI-2.5Sn with

additions of Si will be presented. The Ti-5AI-2.5Sn base alloy was chosen since it is a simple alpha

alloy; the Si additions, since this element is a potent alloy strengthener particularly for high

temperature applications.

II. EXPERIMENTAL PROCEDURES

The Ti-5AI-2.5Sn alloy containing 0.5% silicon is very close to the eutectoid composition and,

consequently, evidence of a eutectoid decomposition would be expected. The structure of this alloy

in the as-received condition is illustrated in Fig. la. The structure contains a very fine

martensite, characterized by a high dislocation density within the grains. The corresponding

diffraction pattern is shown in Fig. 1b, the zone being identified as [I2l6]. No precipitates are

evident, indicating that the silicon is all in solution. On aging the material in vacuum for periods

of up to 40h, the microstructure changes considerably. Precipitates present after 40h are shown in Fig. 2. No definitive diffraction pattern has been obtained yet, but Energy Dispersive

Spectrometry (EDS) data indicates that the precipitates are Si-rich. Thus, the precipitates are

tentatively identified as Ti^Si^.

Fig. 3a shows the general morphology observed in the alloy containing 1 wt.pct. Si in the

585
as-received condition. The microstructure appears to consist of an extremely fine martensite with S no specific orientation and a high dislocation density. No precipitates were observed even though I the 1% Si level is hypereutectoid. The diffraction pattern, using a relatively large selected area

^ aperature, shown in Fig. 3b, consists of partially formed ring patterns indicating that a wide range fl of orientations are present. This fine microstructure is typical of Ti alloys which have been 3 rapidly cooigd. Aging of this alloy initially yields a Si-rich eutectoid structure, which after 40h at ^ 535°C (1000°F) has decomposed to a dispersion of Ti^Si^ precipitates (Fig. 4).

j The microstructure present in the 5 wt.pct. Si produced by splat-cooling, is cellular with \ evidence of a very fine martensite within the cells. The cell boundaries are poorly formed and |

appear to contain many dislocations. This structure is interpreted to be the cellular structure I expected as a result of a eutectic decomposition. A eutectic structure is not observed, indicating ^ that solidification was sufficiently rapid to suppress its formation, but not sufficiently rapid to ; suppress the formation of the cell structure. In addition to the poorly formed cell structure, very | small precipitates can be seen in the matrix (<10nm). Attempts to obtain diffraction patterns from

j the cell boundary structure were not successful, but the diffraction patterns from the general area

| indicated an HCP structure (Fig. 5a), with a [0111] zone axis. Examination of a portion of the i boundary near the edge of the foil revealed a fine structure present within the boundaries. At 40h ,1 aging, the cell structure has broken up into irregularly-shaped precipitates, as shown in Fig. 6, ' which is a Si x-ray area scan of a portion of a STEM foil. I

The PDME material was expected to solidify at a slower rate than the splat-cooled allovs. ' Microstructural differences due to this cooling rate variation, would be large if the variation covers ^ the critical rate for formation of metastable phases, dendrites, etc. Evidently, the cooling rates in the PDME were above the critical rate, since the microstructures observed are essentially the same

j as observed in splat-cooled alloys. Fig. 7 shows the microstructure for the 0.5 wt.pct. Si alloy.

Again, martensite is evident, but the lath size is slightly larger than in the base material. The : absence of any precipitate particles is as expected, since at this concentration, the Si is in solid

solution. A major change in microstructure is apparent in Fig. 8, which shows the 5 wt.pct. Si I alloy produced by the PDME method. Examination of the microstructures suggests that the cooling

of this fiber was not uniform. The fine structure around the outer edges is associated with a rapid

quench, while the larger, more well developed structure in the interior indicates a slower quench

rate. Some martensite is also present in the matrix, but it is not well developed. EDS x-ray

analysis of the cells and cell boundaries indicated that the concentration of Si was significantly

| higher in the boundaries than in the matrix.

III. DISCUSSION OF RESULTS I

j The splat-cooled material exhibits a fine grain size as-cooled. An apparent increase in solid

j solubility in the 0.5 wt.pct. Si alloy was sufficient to suppress formation of the eutectoid,

producing a metastabie condition. The 1 wt.pct. alloy showed similar behavior. Such extensions in 1 solubility have been seen in Al alloys and superalloys and Ti alloys appear to follow the same

behavior. The cellular structure seen in the 5 wt.pct. Si alloy indicates that the solubility

extension is less than this level of solute. Since the eutectic is at about 8 wt.pct., a strong j eutectic formation is expected in this alloy, but this is not seen. The cellular structure indicates a high interface velocity and a high thermal gradient sufficient to produce a metastabie structure in which the eutectic is suppressed (6). | The PDME material shows similar effects (5). l.ow-solute solid solubility is extended, and the #

eutectic is suppressed in the high solute case (5 wt.pct.). In this process, the liquid/soiid 1

interface velocity and related cooling rate (liquid temperature to solid temperature) is slower than .1 in the splat-cooled method. The effect of the interface velocities on the microstructure appears to S

be the same for both processes. That is, the only effect is on the size of the structures 1' produced; the splat-cooled structures being finer. During both processes, solidification in alloys t containing less than 5 wt.pct. Si occurs fast enough for solute rejection to be minimized,

suppressing the eutectoid. For alloys containing 5 wt.pct. Si, rejection of solute occurs during ' cooling, giving a cellular structure. Subsequent cooling after initial solidification results in further ' modification of the microstructure by diffusion mechanisms and metastabie phase transformations.

For each process, there are at least two distinct temperature regimes which can significantly , influence the final microstructure* ; one spanning the transformation from liquid to solid, the other j

^ it Ti recognized that other temperature ranges may also contribute to the microstructures

developed. For example, super-heating can completely destroy all nearest neighbor relationships | resulting in a more random material than if quenching is carried-out from a temperature which is

| just above the liquidus. !'

586
after complete solidification has occurred. In the former temperature regime, the structure

produced is strongly influenced by interface velocity and thermal heat transfer coefficient. The

second temperature range, where the alloy has solidified, is characterized by a slower coding rate.

In titanium, the martensitic transformation occurs on rapid quenching, and such structures are observed in these alloys. The microstructure observed thus appears to develop in two stages.

That is, at temperatures just above the liquidus, the short-range order present is frozen by the

very high solid/liquid (S/L) interface velocity and the rapid removal of heat. This results in suppression of the eutectic and eutectoid decomposition, except in the high solute alloy (5 wt.pct.),

where cells are formed. During the slower cooling after solidification, martensitic transformations

take place, changing the initial metastable solidification structure, and producing the final martensite and cell structures observed.

IV. CONCLUSIONS

Production of rapidly-cooled Ti-5AI-2 . 5Sn , containing various levels of Si by splat-cooling and

PDME methods, results in very fine grained martensitic material for low solute levels (up to at least

1 wt.pct.) and cellular structures for high levels of solute (5 wt.pct.). Aging of the splat-cooled

alloys results in TLSi^ precipitation in the low solute alloys either as a dispersion or as a weakly

formed eutectoid. The cellular structure present in the as-produced high solute alloys decomposes

to a dispersion of irregularly-shaped Ti^Si^. For both processes, there are two distinct

temperature regimes (that span solidification and a lower temperature range in which the material is completely in the solid state) which together contribute to the final microstructures developed.

V. ACKNOWLEDGEMENTS

The authors gratefully acknowledge the assistance of R. Omior and F. Deutscher in the

portions of the TEM work and R. Bacon for his contributions to the SEM work. Helpful discussions

with Dr. L. Beckerman are appreciated. This work was partially supported by Air Force Contract

F33615-80-C-5079.

VI. REFERENCES

1. F. H. Froes, D. Eyion, C. E. Eichelman and H. M. Burte, JOM, Vol. 32, No. 2, Feb. (1980),

p. 47.

2. A. C. Jackson, Ph.D. Thesis, University of Cincinnati, 1983.

3. A. C. Jackson, F. H. Froes, A. Clauer and R. Carbonara, TMS-AIME Fall Meeting, Louisville,

October, 1981.

4. S. Sastry, T. Peng and J. O'Neal, TMS-AIME Fall Meeting, St. Louis, October, 1982.

5. A. C. Jackson, F. H. Froes, A. Clauer, and R. Carbonara, to be published.

6. R. Mehrabian, Int. Met. Rev., 27 (4), 185 (1982).

587
FIGURE 2:

FIGURE 3(a): Microstructure of alloy with

1 wt.pct. Si,

FIGURE 3(b): Diffraction pattern from area

shown in Fig. 3(a).

Same alloy as Fig. 1(a), aged 40h

at 535°C , showing evidence of

precipitation.

FIGURE 4: Same alloy as Fig. 3(a) aged 40h

at 535®C.

588
FIGURE 5(b): Diffraction pattern from area

shovm in Fig. 5(a), B = [0111].

FIGURE 6: X-ray area scan (enhanced) showing

the broken-up cell structure after

aging 40h at 535°C. Same alloy as Fig. 5(a).

FIGURE 7: SEM micrograph of Ti-5Al-2.5Sn^0,5Si

PDME material. Microstructure is

similar to Fig. 1(a).

FIGURE 8; SEM micrograph of Ti-5Al-2.5Sn with

5Si. Microstructure exhibits a cellular structure.

589
ON THE TEXTURE DEVELOPMENT - DIRECTIONAL RECRYSTALLIZATION INTERACTION

OF RAPIDLY SOLIDIFIED Ni-Mo-Al-W AND Ni-Cr-Al-W ALLOYS

H. A. Chin, S.E. Hughes and J. S. Andrus

PRATT AND WHITNEY AIRCRAFT/GPD

WEST PALM BEACH, FLA 33402

J. L. Larson

UNIVERSITY OF ILLINOIS, URBANA-CHAMPAIGN , ILL 61801

ABSTRACT

Producing wrought single crystals from consolidated rapidly solidified powders will significantly increase the high temperature capability of nickel base superalloys. Through

rapid solidification, Ni-Mo-Al-W and Ni-Cr-Al-W alloys were developed at PWA/GPD (under

Government contract) which exhibited improvements in creep/stress rupture, oxidiation re- sistance and fatigue over available cast alloys. Such improvements are due to a combination of (i) alloy design, (ii) RST processing to a metallurgically homogeneous condition,

(iii) solid state directional recrystallization (DR) to single crystals, and (iv) thermomechanical process development to specific crystallographic textures for controlled

orientation DR crystal growth.

This paper discusses the development of preferred orientation in rolled Ni-Mo-Al-W and

Ni-Cr-Al-W alloys as a function of processing parameters. Rolled textures of both alloys

are observed to be similar, exhibiting preferred orientations typical of intermediate to high stacking fault energy face-centered cubic alloys. Cold straight rolled product ex- hibited a texture centered at the [

1 1 3 }<332>with a spread in ideal orientations to {110}

<112>. By cross rolling, the texture spread was effectively converted to a singular

texture centered at ideal orientation [110KH2> Intermediate annealing below the solvus

of the alloy caused no change in the rolling texture though the ro 1-1 ability of the alloy

is restored. Textures thus developed are observed to control the initiation and growth

of specific crystal orientations on directional recrystal 1 ization . These texture-DR

crystal orientation results are discussed and rationalized as a directional secondary re- crystallization phenomenon.

INTRODUCTION

It is common knowledge that the most efficient means for increasing gas turbine engine performance is

by increasing the turbine inlet gas temperature. However, since the engine environment of the Pratt

and Whitney Aircraft/Government Products Division (PWA/GPD) FIDO gas turbine already taxes current

design and materials capabilities, further improvements require combinations of better materials and

designs. For several years PWA/GPD has investigated the utility of wafer constructed airfoils having

high cooling effectiveness. Concurrently, PWA/GPD pursued development of better materials for use in

these wafer constructed airfoils. Envisioned was a stronger powder alloy composition which, when

processed to single crystalline form, would represent an excellent material condition for high tem- perature creep loaded airfoils. Under joint funding from the Defense Advanced Research Projects

Agency (DARPA) and the Air Force Materials Laboratory (AFML) two nickel base superalloys, 185 and 116 were developed for turbine blade and vane applications, respectively. Processing methods for generating single crystalline wafer stock from both al loys were developed. This paper presents the re- sults of an investigation into the development of preferred orientation and how that preferred orientation affects crystallographically oriented directional recrystallization. Discussions on the development of deformation and annealing textures in f.c.c. metals are. available in the literature

[1-5].

The studies described in this paper examined the relationships between texture formation and hot isostatic pressing, extrusion, hot rolling, cold rolling and both hot and cold cross rolling for these

alloys. Material processed by these methods was subjected to directional recrystal 1 ization (DR) heat

590
treatments to establish the relationships between deformation textures and both DR crystal orientation and growth rates.

The alloys used were developed at PWA/GPD under the combined sponsorship of DARPA and AFWAL. The

alloy compositions are listed below:

Composition, Wt%

Ni Mo Cr A1 W C Alloy 185 bal 14.4 - 6.8 6.25 0.04

Alloy 116 bal - 9.2 8.3 9.4 0.04

Both alloys utilized a high volume fraction of an ordered f.c.c. ' - NigAl phase, imparting high

temperature strength through precipitation hardening. In addition, molybdenum was added to alloy

185 for solid solution strengthening, and chromium to alloy 116 for environmental resistance.

Both alloys were produced by a rapid solidification rate (RSR) powder process developed at PWA/GPD

utilizing rotary atomization and forced convective cooling with helium gas under inert conditions

[6 - 12]. Cooling rates attained were in excess of 1 ,000,000°c/second, which virtually eliminated

macrosegregation. The powder particles produced were mostly in the 10 to 100 micron range and were

spherical

EXPERIMENTAL PROCEDURE

Alloy 185 and alloy 116 powders, produced by the PWA/GPD atomization rig, were placed in stainless

steel consolidation cans immediately after hot dynamic outgassing at 1000°F and 10~4 torr pressure.

Rectangular cans were used for hot isostatic pressing (HIP) consolidation and cylindrical cans for

extrusion consolidation.

HIP was performed at 2250°F for 2 hours at 15 ksi pressure. The extrusion parameters investigated

are listed in Table I. Both HIP and extrusion parameters resulted from earlier unpublished studies.

Hot rolling studies were conducted on canned alloys 185 and 116 consolidated by HIP. They were

rolled using the parameters listed in Table II on a 12" diameter mill. Reductions per pass started

at 5 % and gradually increased to 20% as rolling progressed.

Cold straight rolling studies were conducted on extruded alloys 185 and 116. Extrusions 5E9, 5E19,

5E19, 5E21 , 5E15, 6E7, 6E10 and 6E18 were all cold rolled to 15, 25, 35, 45 and 55% reductions.

Annealing was performed every 5 to 10% in reduction for 5 minutes at 2200°F. Reductions per pass were on the order of 1%.

Cold cross-rolling studies were conducted on HIP plus hot rolled alloy 116 and 185 stock, and hot

cross rolling studies were conducted on HIP alloy 116 and both HIP and extruded alloy 185. Parameters

were as for above, except for total reductions and the inclusion of cross rolling. Cross-rol 1 ing was performed as either the final portion or an intermediate portion of the rolling. The cross-rol 1 ing

and total reduction parameters are listed in Table III.

X-ray diffraction methods were used to determine texture orientation of the extrusion, HIP product

and rolled strip. Plotting of the pole figures was accomplished using a computer and software

written by Love [13].

TABLE I EXTRUSION PARAMETERS AND RESULTING TEXTURES

I.D. # ALLOY

EXTRUSION

TEMPERATURE, °F

EXTRUSION

RATIO

FIBER TEXTURE

COMPONENTS

5E9 185 2000 15:1 <100>, <111> weak

5E19 185 2100 25:1 <100>, <11 0> , <11 1> weak

5E21 185 2250 15:1 <110>, <11 1> weak

5E15 185 2300 26:1 <11 1> weak

6E7 116 2100 15:1 <100> strong , <1 1 1 > weak

6E10 116 2250 8:1 <111> strong

6E18 116 1900 8:1 <100>, <11 1> weak

591
TABLE II TABLE III HOT ROLLING PARAMETERS CROSS ROLLING PARAMETERS

ROLLING TOTAL ROLLING TOTAL SR:XR

TEMPERATURE REDUCTION TEMPERATURE REDUCTION (:SR)

ALLOY °F % ALLOY °F % RATIO

185 2275 90 185 70 85,78,73 60:40

185 2250 78 185 70 73 70:30

185 2100 75 185 70 85,78,67 75:25

185 2000 84 185 70 78,73 80:20

116 2250 75 185 70 71 34:48:18

116 2100 75 185 70 71 42:42:16

185 70 71 60:25:15

116 70 67 65:35

116 70 67 75:25

185 2200 80 60:40

185 2100 91 70:30

185 2000 84 50:50

116 2300 91 65:35

Directional recrystallization heat treatments were performed on the processed alloys under vacuum

(typically 10“^ torr) in a zone annealing furnace. A graphite susceptor, heated by induction,

radiated heat onto the sample passing through. The gradient at the region of insertion into the

susceptor provided the driving force for columnar grain growth. The measured thermal gradients at 2300°F ranged between 200°F/in and 300°F/in and insertion rates for smooth continuous grain growth

varied from 1/8 inch/hour to 2 inch/hr. Furnace control was achieved with W/W-Re thermocouples imbedded in the susceptor such that the maximum specimen temperature peaked between the 7' solvus and

incipient melting temperatures. Wafer stock from ten different rolling schedules was cut into 1/2

inch by 2% inch strips with a wedge shaped leading edge for DR (the wedge increased the tendency

toward single crystalline growth). Usually, columnar grains were grown to within 1/2 inch of the

trailing edge and by back reflection Laue technique the crystallographic orientations of all major

grains were recorded.

RESULTS AND DISCUSSIONS

CONSOLIDATION

The subsolvus HIP parameters were found to effect full consolidation of the powder billet. Pole

figure analysis of the HIP product revealed a random distribution of crystal orientations or non- textured condition in both alloys.

For the extrusion parameters used, full consolidation of the powder billet resulted. Pole figure

analysis of the extruded products revealed that the two classes of alloys, Ni-Mo-Al-W (185) and

Ni-Cr-Al-W (116) developed extrusion textures typical of moderately high and low stacking fault

energy fee metals, respectively.

Alloy 185 developed four different textural conditions, as listed in Table I and predicted in Figure

1. Alloy 116 developed the three textural conditions also listed in Table I and predicted by Figure

1. The range of textures produced in both classes of alloys followed the changes expected in f.c.c.

alloys from varying Stacking Fault Energy (SFE) and processing temperatures. The high temperature

extrusions of both alloys appeared to occur above the 7' solvus (as single phase 7 ) due to adiabatic heating during deformation. Under this condition, alloy 185 exhibited the singular<l 1 1>

fiber texture typical of higher SFE f.c.c. alloys undergoing a temperature related texture transition.

Alloy 116 exhibited a singular<100> fiber texture typical of a low SFE f.c.c. alloy. Lowering the

processing temperature resulted in a texture transition to the commonly observed duplex<lll> + <100>

fiber texture for the moderate-to-low SFE f.c.c. alloys. The existence of the <1 10> fiber component

found in alloy 185 processed in the intermediate temperature range has not been reported in the

literature for f.c.c. metals. The development of that texture component is not understood at this

time and requires further investigation.

592
Extrusion Temperature, °F (°C) Extrusion Temperature, °F(°C)

Figure 1: Fiber Textures in Alloys 185 and 116 as Influenced by Extrusion Ratio and Temperature,

ROLLING

Hot straight rolling was found in impart similar texture components to both alloys 185 and 116. Both

exhibited a texture spread about the end orientations { 113}<332>and [110}<112>with the most

prominent component at {113}<332> The texture strengths were weak, typically at 2 times random.

The texture strength was not significantly affected by the deformation temperature, although the

alloy 185 strip rolled at 2000°F (the extreme lower hot rolling workability limit) exhibited a 3 to 4 times random texture strength.

The absence of a temperature-related texture transition in the hot rolled 116 and 185 was thought to

be a result of heat losses during rolling. Thus, actual rolling temperatures were lower than for

extrusion and texture transitions were insignificant, as with the lower temperature ( 2250°F) extrusions.

The low texture strengths of the hot rolled 185 and 116 are not fully understood. However, it is

thought that at these high temperatures and low deformation rates these fine grained materials deform in part by grain boundary sliding and consequently lose texture strength.

Four different input textures of extruded alloy 185 and three of extruded alloy 116 were cold

straight rolled. All responded in a similar fashion. Regardless of input texture, both alloys developed a weakly defined f.c.c. type sheet texture with as little as 9% reduction. With increased

reduction, the sheet texture increased in definition (i.e. reduced spread) and strength. Textures

could be described as a spread about the end orientations {1 1 3}<332>and {110]<112>. The sheet

textures were 2.5 x random at 25% reduction and 4 - 5x random at 36% reduction. Reductions greater

than 35% resulted in even stronger and better-defined textures.

The deformation texture developed for both alloys was found to be unchanged by the many subsolvus

(2200°F/5 min/air cool) intermediate annealing treatments used through the rolling schedule. Such

behavior has been rationalized as crystallization in situ , whereby restoration occurs by subgrain re- arrangement in the absence of movable recrystallization fronts. This was caused by the strong inhibitive effect of the ordered 7' phase on the growth of the classic f.c.c. [100]<001>annealing

texture nuclei [5] .

The texture development observed for both alloys appears to follow the trends reported in the

literature for other f.c.c. metals and alloys. In the case of alloys 185 and 116, the end orientation {110]<112> is always present for the range of reductions studied. The major texture component, {113}<332>, lies 10° off the {1123<lll>by rotation about a<110>and is thought to be a product of the low to intermediate SFE alloys for the specific processing conditions. This is consistent with Dillamore's hypothesis [24] which characterizes the high SFE rolling texture as centered

at the {4,4,11]<11 ,11 ,8>, this component being a product of cross slip ( [1 1 3}<332>1 ies 2 ® from the [4, 4, 11 } <11, 11, 8>). Theories for the characteristic plane strain rolling textures

have been discussed extensively in the literature [2, 14 - 19 ] .

593
Through cross rolling, both alloys 185 and 116 were observed to develop similar textures. The major

effect of cross rolling was to change the straight rolling texture spread into a symmetric singular

texture, centered at the {110} <112> component. The degree of textural singularity was observed to

be dependent on the ratio of the straight to cross roll reductions, i.e. SR:XR. For alloys 185 and

116, a fair degree of singularity was achieved using SR;XR ratios within the bounds 80:20 and 60:40.

Above and below these bounds the textures more closely matched those of straight rolled product,

parallel to and 90° off the original (straight) rolling direction, respectively. Textural singularity also was observed in hot cross rolled product, centered at the {110}<112>, although the tex- ture strength was significantly lower. Figure 2 shows a pole figure of cold cross rolled Alloy 185.

Figure 2: (111) Pole figure for cold cross rolled Alloy 185. 73% total

reduction 70:30 SR:XR Ratio

1 in. /hr

in. /hr

2 in. /hr

2y, in. /hr

2 cm

Figure 3: DR Rate Response of Cold Rolled

Alloy 185.

Figure 4: Alloy 185 DR Strips Showing Wider

Columnar Grains in the Cold-Cross

Rolled Material (Left) Than the

Straight Rolled Material (Right)

DIRECTIONAL RECRYSTALLIZATION

Hot rolled, cold straight rolled and cold cross rolled strips of each alloy were directionally re crystallized. The resulting grain growth patterns and orientations were documented.

594
In the as-rolled condition, only the cold rolled material responded to DR heat treatments. The cold

rolled material exhibited continuous columnar grain growth from a maximum rate of 2 inch/hr down to

1/8 inch/hr, shown in Figure 2, whereas the hot rolled material exhibited only equiaxed grain

coarsening. The as-rolled materials differed in texture strength (cold stronger than hot) and microstructure. Interrupted DR trials revealed that the cold rolled material, microstructure transformed

during DR, producing grain boundary 7' necklaces. This transition, previously linked to positive

DR response, was not found with the hot rolled material. However, a subsolvus conditioning heat

treatment was developed for fine grain alloy 185 that established the necklace structure in previously non-DR responsive material. Hot rolled alloy 185 was condition heat treated to a necklace

structure and subsequently exhibited columnar grain growth at insertion rates of 1/2 in/hr and slower.

So it appears that both texture and microstructure significantly affected DR response. The desirable

features of cold rolling versus hot rolling can be summarized as:

1. Stronger deformation texture

2. Finer and more homogeneous as rolled microstructure

3. DR response in the as-rolled condition

4. Continuous DR at insertion rates 4x faster

5. Greater propensity for single crystalline growth

The straight rolled material repeatedly grew narrower columnar grains than the cross rolled as shown

in Figure 3. Cross rolling, with its singular texture, allowed fewer orientation combinations than

straight rolled sheet. As the leading edge of a straight rolled sample underwent abnormal grain

growth a larger number of allowed orientations grew resulting in narrower columnar grain widths.

60

O

c

o

3?w

U.c

o

50

40

30

20

10 I i

(001 ) < 110> m( 112) < 111 >

inch/hr, ^

inch/hr.

n

1

i

1i

1

p

1

( 111 ) < 110> Q

inch/hr,

I M %

inch/hr

(Wafer Plane) <DR Insertion Axis>

ilii

( 111 ) < 112 >

I

III IIP ( 111 ) < 111 >

Figure 5: DR Grain Orientation Frequencies for Different Strip Insertion Rates Using Results From

All Ten Rolling Schedules.

The columnar grain inspection and Laue data indicated that a strong cold cross rolling texture en- hanced both single crystalline growth and orientation control. However, the effects of the cross

rolling ratio variations were not discernable.

Repeated growth observations showed that the polycrystalline texture, though a necessary condition

for allowed DR crystal orientations was not sufficient to guarantee a specific orientation (Frequencies of DR orientations for various growth rates are shown in Figure 4). However, by use of appropriate selection techniques (seeding, grain isolation, growth rates) any of the reported low index

orientations could be produced as single crystals. In this respect, single crystal parts can be

fabricated with selected mechanical properties compatible with service requirements.

The results of the study are consistent with previous studies on secondary recrystallization by Rosi

et al [20] , Aust [21] and Beck [22]. They reported that a few selected crystal orientations are

promoted by the presence of a strong texture. The resulting crystals had a preference for specific

orientation relationships to the matrix, the number of possible growth centers being limited by the

existing texture such that other orientations are generally scarce. Further, the growth rate or mobility of a DR grain was strongly dependent on orientation. Previous studies showed similar re- sults [ 20 - 23] where the<lll> orientation exhibited the lowest mobility of all orientations,

whereas crystals in the [111] plane or zone exhibited the highest mobilities eg.<110> and<112>.

595
The observations and the data presented represent an extensive evaluation of alloy 185. Admittedly

the number of DR trials with similarly processed alloy 116 are fewer. However, the results of these

trials with hot, cold straight, and cold cross rolled material duplicate those of alloy 185.

CONCLUSIONS

This study shows that, in general, deformation textures produced in alloys 116 and 185 follow the

trends found for f.c.c. metals. Hot extrusion textures are dependent on alloy type, extrusion ratio

and temperature, while hot rolling textures are virtually independent of rolling temperature, alloy

tyoe and seldom exceed 3 times random. Cold rolling texture tyoes are relatively indeoendent

of alloy type, and texture strengths are higher and highly dependent upon percent deformation, with

35% reduction required to completely erase previous textures. Cross rolling imparts textural

singularity upon both the hot and cold rolled material.

There is a strong correlation between crystallographic texture and both DR rate response and orientation control. Cold cross rolling produces the strongest, most singular textures, which promote

the most rapid and most easily controlled DR. The DR crystal orientation control provided by the

strongest textured material is not sufficient for production of DR single crystals, however, but

must be supplemented with grain isolation techniques, furnace design, growth rate manipulation, and

microstructural manipulation for growth of large single crystals. Studies investigating these

parameters are underway.

ACKNOWLEDGEMENTS

The authors wish to thank 0. B. Cavin (ORNL) and J. Umbach (PWA/GPD) for Pole figure generation,

J. Zindel for the extrusion studies and R. E. Anderson and J. A. Miller for their consturctive

discussions and support.

This work is jointly sponsored and monitored by the U.S. Defense Advanced Research Projects Agency

and the Air Force Wright Aeronautical Laboratories.

The views and conclusions contained in this paper are those of the authors and should not be ‘'[Tter preted as necessarily representing the official policies, either expressed or implied, of the U.S.

Defense Advanced Research Projects Agency, U.S. Air Force, or the U.S. Government.

REFERENCES

(1) H. Hu, Texture, 1974, Vol . 1, p. 233.

(2) H. Mueller, Osterr. Akad. Wiss. Math Naturw. 1958, Vol. 7, p. 117

(3) H. Hu and S.R. Goodman, Trans. Met. Sco. AIME., 1963 Vol. 227, p. 627

(4) H. Hu and S.R. Goodman, ibid, 1963 Vol. 227 p. 1454

(5) E. Hornbogen and H. Kreye, Proc. Inti. Symposium on Textures in Research and Practice,

Clausthal - Zellerfeld, 1968, p. 274

(6) A. R. Cox and E. C. VanReuth, Proc. 3rd Inti. Conf. Rapidly Quenches Metals, Univ. Surrey.

U.K., 1978, p. 225

(7) M.R. Glickstein, R.J. Patterson II, and N.E. Shockley, Proc. Inti. Conf. on Rapid

Solidification Processing, Baton Rouge, LA, 1977, p. 46

(8) P.R. Holiday, A.R. Cox and R.J. Patterson II, ibid, p. 246

(9) R.G. Bourdeau and J.B. Moore, ibid, p. 334

(10) D.B. George, J.B. Moore, and E.C. VanReuth AIAA/SAE/ASME Joint Propulsion Specialist Conf.,

Las Vegas, Nevada, 1979, p. 1 (11) A.R. Cox, J.B. Moore and E.C. VanReuth, Proc. AGARD Conf #256, Adv. Fabrication Processes,

Neuilly Sur Seine, France, 1978

(12) R.E. Anderson, A.R. Cox, T.D. Tillman, and E.C. VanReuth, Proc. 2nd Inti. Conf. on Rapid

Solidification Processing, Reston, VA, 1980, p. 416

(13) G.R. Love, Oak Ridge National Lab. Tech. Memo, ORNL-TM-2018, 1968

(14) I.L. Dillamore and W.T. Roberts, ACTA Met., 1964. Vol. 12, p. 281 (15) I.L. Dillamore and W.T. Roberts, Met. Review, 1965, Vol. 10, No. 39, p. 271 (16) I.L. Dillamore and H. Katoh, Brit. Iron and Steel Res. Assoc. Report # MG/39/71, 1971

(17) I.L. Dillamore and H. Katoh, Metal Science, 1974, Vol. 8, p. 21 (18) T. Kamijo, Trans. JIM, 1969, Vol. 10, p. 238

(19) M. Hatherly, and I.L. Dillamore, J. Australian Inst. Metals, 1975, Vol. 20, p. 71 (20) F.D. Rosi, B.H. Alexander and C.A. Duke, Tran AIME, 1952, p. 189

(21) K.T. Aust, The Art and Science of Growing Crystals, J.J. Gilman, ed., John Wiley and Sons, 1963

(22) P.A. Beck, Trans. AIME, 1951, Vol. 191, p. 475

(23) B.B. Rath, and H. Hu, Trans. Met. Soc. AIME, 1966, Vol. 236, p. 1193

(24) I.L. Dillamore, Met. Trans. 1970, Vol. 1, p. 2463

596
THERMAL STABILITY OF HIGH TEMPERATURE P/M ALUMINUM ALLOYS

Young-Won Kim*, M. M. Cook* and W. M. Griffith**

* Metcut-Materials Research Group

P.O. Box 3351 1 Wright-Patterson Air Force Base, OH, 45433

** Air Force Wright Aeronautical Laboratories

Materials Laboratory (AFWAL/MLLS)

Wright-Patterson Air Force Base, OH, 45433

ABSTRACT

Emerging high temperature aluminum P/M alloys have been characterized for

thermal stability. The alloys selected are Alcoa aluminum powder alloys containing

Fe (5.7 to 7.4 wt%) and a second alloying element (Ni, Ce, Co). These alloys are strengthened by finely dispersed intermetallic compounds that are formed during the

melt atomization process. The materials' properties, especially strength, depend on

the thermal stability of these compounds. A series of isochronal and isothermal

experiments on vacuum hot-pressed material were conducted at temperatures

between 316^C and 552^C. The annealing response was monitored by hardness

measurements and microstructural changes were observed by SEM. The isochronal

annealing experiments show that all the alloys have good stability up to 350°C and

that the stability is rapidly lost at higher temperatures. The deterioration of

stability differs in degree depending on the second alloying element, being the

worst with Ce addition. The difference in the instability among alloys is further

substantiated when the material is isothermally annealed. The coarsening of the

intermetallic compounds is responsible for the decrease in stability and strength.

Microstructural observations show the coarsening rates of the compounds vary from

alloy to alloy.

I. INTRODUCTION

The need for stronger, lighter and more economical materials for aircraft structure and engine

components has led to the development of new powder metallurgy (P/M) aluminum alloys (1,2,3).

Aluminum P/M alloys for use at elevated temperatures are one such class of alloys. Characterized

by a fine microstructure and a high volume fraction of thermally-stable intermetallic compounds,

these alloys will compete with heavier steel and titanium alloys that experience service temperatures

up to 33o’^C. The elevated temperature strength of the alloys depends on the thermal stability of

these compounds. Optimum processing conditions have not been established for high temperature

alloys in part because of the limited understanding of the thermal behavior of the strengthening

compounds.

Three of the most promising alloys (Al-Fe-Co, Al-Fe-Ni and Al-Fe-Ce) , currently under

development by Alcoa, were selected for study. The thermal behavior for vacuum hot-pressed

material was studied following isochronal and isothermal annealing. The coarsening of

microstructure and strength change was monitored by SEM observations and hardness measurements.

II. EXPERIMENTAL

Several ternary alloys, based on Al-Fe, were supplied by Alcoa as alloy billets (1) that had

been cold-compacted and then vacuum hot-pressed at 343°C. The compositions of three alloys

selected for the present experiment are listed in Table I.

597
Table I. Alloy Compositions (wt%)(1)

Alloy Alcoa No. Fe Ni Co Ce Al

A 513671 5.70 — 6.20 — Bal

B 513670 5.70 6.10 — — Bal

C 513672 7.40 — — 3.40 Bal

Small coupon specimens were prepared by sectioning the billets and polishing to 600 grit

finish. Isochronal annealing experiments were conducted in air circulation furnaces, at temperatures between room temperature (RT) and 538°C and the annealing time was 1h. Two series

of isothermal experiments were conducted in a salt bath at several temperatures selected from the

isochronal annealing results. Changes of hardness values, after every annealing, were monitored

by Rockwell B hardness measurements. Microstructural changes were observed by SEM of selected

annealing conditions and for the starting material condition (as-hot-pressed).

III. RESULTS AND DISCUSSION |

1 . Isochronal Annealing 1

The three alloys were isochronally (1h) annealed at temperatures between RT and 538°C. The j results are shown in Fig. 1 where RT hardness values, measured by Rockwell R (Rg). plotted

against annealing temperatures. The isochronal annealing responses of the alloys” appear to be < similar. Hardness values are relatively stable up to 350°C and they rapidly drop at higher ^ temperatures. Further scrutiny, however, reveals observable differences in the annealing ,• responses among the alloys, in addition to the differences in starting material's hardness. Alloy A a (AI-5.7Fe-6.2Co) and C (AI-7.4Fe-3.4Ce) are very stable up to 350°C, showing hardness values J

virtually unchanged. Alloy B (AI-5.7Fe-6.1Ni), having a higher initial hardness value, exhibits J some instability in this temperature range; however, this does not necessarily mean that Alloy B is unstable compared with the other two alloys. When Alloy B is heat-treated, for example, for 1 h at ? 350°C to have R_ 92 and then isochronally annealed, the material exhibits the same stability as that

of Alloy A or C T6,7).

At around 350°C, all the alloys begin to soften and at 538°C, a considerable fraction of

hardness (from R„ 90 to 65) is lost. Coarsening of the strengthening intermetallic compounds is j responsible for the loss. The rate of decrease with increasing temperature differs between alloys,

Alloy A exhibits the slowest coarsening and Alloy C, the fastest, although the difference in the '> rate does not appear to be substantial. As will be shown, however, isothermal annealing « substantiates the difference. i

2 . Isothermal Annealing

Using the results of the isochronal annealing experiments (Fig. 1), two series of isothermal | annealing experiments were conducted over a wide range of elevated temperatures (316^C to 552^C). '

a) In the first series of experiments, three temperatures (316, 349, 382°C) were selected j around the critical temperature (350°C, Fig. 1), at which the stability to instability transition takes ; place for all the alloys. These temperatures bracket the maximum use temperature (goal) of the

j developing elevated temperature P/M Aluminum alloys. Accordingly, it is of engineering importance | to characterize the thermal response of the alloys in the region. Three coupon- specimen groups j were prepared from as-hot-pressed billets for the three-stage isothermal annealing experiments.

The specimens of the first group were immersed in a salt bath at 316°C simultaneously. After , a certain period of elapsed time, one specimen was removed and tested for hardness (Rockwell B

scale, Rp). The longest annealing time at this temperature was 50h, The second and third groups

of specimens were annealed at this temperature (316^C) for 50h. The second group of specimens was then isothermally-annealed at 349®C up to 33h. The third group of specimens were then given

an additional annealing at 340°C for 33h and isothermal annealing experiments were conducted at ; 382°C up to 27h. The above annealing experiments thus provide continuous isothermal annealing J curves at three intermittent temperatures. The results are shown in Fig. 2.

598
I At 316 C, Alloys A and C exhibit an initial small hardness drop followed by fairly constant

I hardness values, whereas. Alloy B reveals constantiv decreasing hardness values. This is

; consistent with the isochronal curve (Fig. 1) where Alloy B exhibits a similar trend. At 349°C, the

j isothermal annealing responses of the alloys are all similar; the hardness values constantly

I decrease. Up to this stage, it may not be unreasonable to say that the alloys have quite similar

i stability. When the isothermal annealing temperature is increased to 382°C, the annealing response

I becomes markedly different between alloys: Alloy C is the most unstable and Alloy A exhibits the

highest stability. Alloy B, which was initially unstable at 316°C, now demonstrates a greater

j stability than Alloy C. It is worthwhile to point out, however, that the total decreases in hardness

' values are likely to be acceptable from the engineering viewpoint. When hot-pressed. Alloy R is

! extruded at a 20:1 ratio at 400°C, for example, the hardness value decreased to R_ = 82, but the

1 tensile strength was 470MPa (67Ksi) (6,7). The hardness, R„ = 82, is compared with the final

j hardness (R_ 81.5) of Alloy B, Fig. 2. This means that the material still retains the engineering

strength after rather vigorous thermal treatments (Fig. 2).

b) In the second series of annealing experiments , five temperatures (380, 423, 470, 507, and

552°C) were selected from the later stage iTi isochronal annealing curves (Fig. 1). Isothermal

annealing experiments at these temperatures have double purpose. Firstly, in the case that the

material has to experience such high temperatures, for example, for degassing, it is important to

I know how the material loses the stability and thus, strength. Secondly, understanding the

j coarsening behavior at higher temperatures is of scientific importance. The experimental procedure was similar to that in the first series in that the isothermal annealing was continuously conducted at

j successively higher temperatures after an isothermal anneal. The results are plotted in Fig. 3.

At 380°C, the initial rapid decrease in hardness, shown by all the alloys, slows gradually.

When the temperature is increased to 423^C after 12h at 380^C, the drop continues, however, no

dramatic change is shown. A significant change occurs when the temperature is raised to 470°C.

Alloy C exhibits a drastic decrease in hardness with time. While alloys A and B follow a similar

I trend. Alloy A is more stable. At higher temperatures, the difference in hardness changes between

* Alloy A or B and Alloy C is more obvious, indicating the coarsening of intermetallic compounds in Alloy C is occurring at a much faster rate. This will be shown, nicrostructurally , in the next

I section. One interesting feature in Alloy C is that the rate of decrease is not necessarily

accelerated with increasing temperature. For example, the hardness decrease in Alloy C at 470*^C is drastic and the decrease continues with time (see dotted line). At the next higher temperature

(507^0, the initial decrease is significant but stability is quickly reached. At 552^0, however, the

rapidly dropping rate is again observed to be continuous for a long period of time, although that

trend is not shown to its conclusion in Fig. 3.

3. Microstructure

Microstructural observations were made for the as-hot-pressed condition and selected annealing

conditions. All the observations were done by SEM after etching the specimens with Keller's

etchant,

a) As-Hot-Pressed Material : There are distinguishable differences in microstructures between

alloys, as shown in Fig. 4 where typical regions of the hot-pressed alloys are shown. In Alloys A

and B, the typical microstructure is eutectic, with the eutectic nature more pronounced in alloy B,

Figs. 4a and b. Both microstructures were taken from relatively large powder particles. The

eutectic structure in Alloy A appears to be disturbed which is not understood clearly: either heat

fluctuation during solidification or deformation during hot-pressing or both may be responsible. It is mentioned, however, that the disturbance is found infrequently. The eutectic is coupled in many

cases, as shown in Fig. 4b, which also exhibits grain boundaries inside a powder particle. In contrast to Alloys A and B, Alloy C shows virtually no eutectic; instead, dendrite or cellular

structure is predominant. Fig. 4c. It is likely that the addition of Ce in Al-Fe suppresses the

formation of eutectic which is detrimental for material's ductility. In fact. Alloy C has shown the

best ductility among the alloys after forging (2). It has been reported that improvement in the

ductility of Alloys A and B can be achieved by destroying the eutectic structure using a severe

deformation means such as extrusion (6,7).

It was also observed that Alloys A and B exhibit fairly uniform microstructure within a powder

particle, while Alloy C exhibits a mixture of fine and coarse structures. Fig. 4c, However, overall

microstructure in bulk material appears to be more uniform in Alloy C. Although, not shown here.

Alloys A and B contain large powder particles having coarse intermetallic compounds in a dendritic

form.

b) Heat-Treated Material: The SEM micrographs in Fig. 5 show the coarsening of

599
intermetallic compounds to which the hardness changes may be attributed. Overall, Alloy C

exhibits coarse microstructure that probably is responsible for the low hardness values compared

with other alloys. Fig. 1. When heat-treated at 398^C for 1h, Alloys B and C show some sign of

slight coarsening. Figs. 5d, e and f. This is consistent with the corresponding hardness changes

in Fig. 1. When heat-treated at a higher temperature (476°C) for 1h, there is a significant

coarsening regardless of the alloy. Careful observation shows, however, that there is a considerable difference between alloys. Alloys A and B have fine microstructure regions that

usually are within small powder particles. In contrast. Alloy C exhibits overall uniform coarse microstructure which makes the identification of powder particles difficult. This is, again,

consistent with the observations that the hardness of Alloy C is decreased at a faster rate over the

other alloys. It is noted that the intermetallic compounds generally become

spherical after higher temperatures.

IV. CONCLUSIONS

1. All of the alloys investigated are isochronally (1h) stable up to 350°C. At higher temperatures,

the hardness decreases linearly with increasing temperature, Al-Fe-Ce showing a slightly faster

decrease.

2. Isothermal annealing experiments show that at 316°C, Al-Fe-Co and Al-Fe-Ce are quickly

stabilized in hardness with time after initial small drops. Al-Fe-Ni shows constantly decreasing

hardness although the drop is small. At higher temperatures, however, Al-Fe-Ce exhibits the least stability with time. Al-Fe-Co and Al-Fe-Ni behave similarly, but overall, Al-Fe-Co is more stable.

3. Al-Fe-Co and Al-Fe-Ni still maintain acceptable hardness after continuous annealing consisting

of 316°C/50h, 349*^C/33h and 382°C/27h. Al-Fe-Ce falls below acceptable hardness after a few

hours at 382°C following the first two heat treatments.

4. The initial morphology of compounds differs between alloys. Al-Fe-Ce contains a dendritic or

cellular type structure. In Al-Fe-Co and Al-Fe-Ni, eutectic structure predominates. The eutectic

structure is detrimental to ductility. The rod-type compounds in eutectic become spherical after a high temperature annealing.

REFERENCES

1. W. L. Otto, Jr., "Metallurgical Factors Controlling Structure in High Strength P/M Products",

AFML Technical Report, AFML-TR-76-60 , May 1976.

2. R. E. Sanders, Jr. and C. J. Hildeman, "Elevated Temperature Aluminum Alloy Development",

AFWAL Technical Report, AFWAL-TR-81 -4076 , Sept. 1981.

3. S. L. Langenbeck et.al, "Elevated Temperature Aluminum Alloy Development",

Lockheed-California Co., AF Contract #F3361 5-81 -C-5096 , Interim Report, LR29977, Oct. 1981.

4. W. M. Griffith, R. E. Sanders, Jr,, and C. J. Hildeman, "Elevated Temperature Aluminum

Alloys for Aerospace Applications", Proceedings of High Performance Aluminum Powder Metallurgy

(in press), Feb. 14-1 8, 1982 , Dallas, TX.

5. S. D.Kirchoff, R. H. Young, W. M. Griffith and Y-W. Kim, "Microstructure/Strength/Fatigue

Grack Growth Relations in High Temperature P/M Aluminum Alloys", Proceedings of High

Performance Aluminum Powder Metallurgy (in press), Feb. 14-18, 1982, Dallas, TX.

6. Y-W. Kim and L, R. Bidwell, "The Effects of Thermomechanical Processing on Al-Fe-Ni P/M

Alloy", TMS-AIME Annual Meeting, Feb. 22-26, 1 981 , Ghicago, IL.

7. Y-W. Kim, L. R. Bidwell and W. M. Griffith, "The Effects of Microstructure on Tensile

Properties of P/M AI-5.7Fe-6.1Ni", in preparation to publish in Met. Trans.

600
ROCKWELL

B

HARDNESS

(Rb)

ROCKWELL

B

HARDNESS

(Rd)

FIGURE 1

ANNEALING TIME (hr)

FIGURE 3

ANNEALING TIME (hr)

FIGURE 2

FIGURE 1 : Isochronal annealing curves for

alloys A (Al-5.7Fe-6.2Co), B (Al5.7Fe-6.1Ni) , and C (Al-7.4Fe-3.4Ce).

FIGURE 2: Isothermal annealing curves for

alloys A (Al-5.7Fe-6.2Co), B (Al5.7Fe-6.1Ni) , and C (Al-7.4Fe-3.4Ce) near the critical temperature.

FIGURE 3: Isothermal annealing curves for

alloys A (Al-5 . 7Fe-6 . 2Co) , B (Al5.7Fe-6.1Ni) , and C (Al-7.4Fe-3.4Ce)

at high temperatures.

601
FIGURE 4: Typical microstructures of as-hot-pressed alloys: a = Al-5.7Fe-6.2Co;

b = Al-5.7Fe-6.lNi; c = Al-7.4Fe-3.4Ce.

FIGURE 5: The effect of heat treatment on the microstructures of alloys A, B and C: a-c = as-hot-pressed; d-f = annealed, 398°C/lh; g-i = annealed, 476 C/lh.

602
EVALUATION OF VOIDS IN POWDER CONSOLIDATES USING LONG WAVELENGTH NEUTRONS

M. Faterai and B. B. Rath

Physical Metallurgy Branch, Naval Research Laboratory, Washington, DC 20375

ABSTRACT

The capabilities of small angle neutron scattering method (SANS) have been shown to be

unique in evaluating small defects in alloys fabricated by consolidating fine particulates.

Potential applications in detecting small voids and precipitates either left in the alloy

during fabrication or subsequently formed by nucleating and growth have been realized. The

availability of high-flux research reacters, high speed, large capacity computers, and

efficient position sensitive neutron detectors has made quantitative defect characterization

possible. An analysis of radially averaged SANS profiles, obtained from area detector

plots, provides information on the range and relative number of defects, such as ralcrovoids

in powder of relevant processing parameters. Results of a study on the detection of voids

in powder compacted Ti-alloys and steels are described. Correspondence is established

between SANS profiles and relative void size distributions with consolidation parameters

such as pressure and temperature. It is shown that hydropressed Ti-6A1-4V powder exhibits

a reduction in the number and size of voids with an increase in consolidation pressure from

60 ksi to 100 ksi. Similarly, hot-pressed stainless-steel powder shows a systematic

decrease in porosity upon Increase in compaction temperature from 800° to 1200°C. Results

of the non-destructive evaluation by SANS, with supporting evidence from transmission

electron microscopy, is presented.

Introduction

The application of non-destructive test methods to the problem of evaluating microstructural

defects in materials has been a subject of ever increasing importance in materials science. To a large extent, this trend is due to the development of engineering materials of sophisticated

composition aimed at higher durability and performance. It has been indicated (1-3) that, most

non-destructive evaluation methods suffer from several limitations, particularly when the defects of

interest are very small. Small-angle neutron scattering (SANS) is an appropriate method for flaw

detection, where defects of interest are in a size range of approximately lOO-lOOOA. The neutron

intensity profile scattered at small angles which generally decreases monotonically as a function of

the scattering angle, contains information on the size and number density of the scattering centers in

materials. The shape of the net scattering curve, obtained from the difference between the reference

standard and the sample under investigation can be analyzed to evaluate a relative size distribution

of the defects.

Methods for paritlce size distribution have been devised to obtain information, on the average

size of defects and their distribution as a function of processing variables (4-7). These evaluations

are based upon the assumption that the observed intensity is the integral of the product of the

size-distribution function and the single-particle intensity function. Since there are no analytical

forms that can be used in evaluating size distributions from various types of defects, numerical

evaluation methods are utilized. Different algorithms have been used to derive size distributions of

defects from SANS profiles (4-7), by numerically solving an integral equation of the type:

I(q) = K / D(R)m2(R)Io(qR)dR -( 1 )

603
which relates the number density, D(R), of particles of radius R to the measured Integrated Intensity,

I(q). In the above equation q = 4tt sin 0/ A, m is the excess nuclear scattering length density in the

particle and IgCqR) is the single-particle scattering function which depends on the particle shape and

dimensions.

However, if the sample contains a uniform disperson of nearly similar particle sizes or the

Intensity data is limited in the q range, the Guinler method (8) may be employed which suggests that

the natural logarithm of intensity, Inl, proportional to the square of the wave-=vector , q^, such that

the

slope

Alnl

A(q^)

-( 2 )

where R is the radius of gyration of the spherical particles. A straightline behavior Inl vs. q^ over

a significant portion of the scattered intensity is an indication of nearly uniform size. A trend of

varying slopes indicates either particles of varying size and distribution or other complex ' scattering interactions. If the presence of a widely varying particle or void size distribution is

recognized a more elabrate size distribution calculation is warranted. In the present investigation

the analysis of the void sizes have been made using the Guinier method. ^

Powder consolidates of different metals and alloys, are known to contain a small volume fraction >' of voids, particularly if they have not been subjected to a post-consolidation hot working process, w

The source of thse voids are due to: (1) a lack of complete densif ication which leaves unsintered

regions in the final product, (2) shrinkage during solidification and (3) trapped gases within the |

powder particles. The characteristics of these cavities present in titanium alloy powders and powder 4

consolidates have been described by Eylon et. al. (9). They have shown that the gas pores, present 2

inside the powder particles do not close during the consolidation process. These cavities 2

subsequently serve as the initiation sites for cracks (10). SANS provides a unique non-destructive || method to quantitatively determine the void sizes in powder consolidates. The method can indeed be

|| used to optimize processing variables to reduce the number density and size of voids in sintered or hot-lsostatically pressed products. |

Results and Discussion Ii

A systematic study of porosity in powder consolidates of stainless steels HIPed at a constant 15 | ksi pressure and varying temperatures and a T1-6A1-4V alloy hydropressed at 1315°C at varying

pressures has been conducted using SANS method. Electron microscopic examination of the samples | provided supportive evidence of voids in these alloys systems.

Porosity in Powder Compacted Stainless Steels

i Disks of stainless steel specimens prepared from gas-atomized powder (Fe-21Cr-6Ni-9Mn-.02C-.25N2,

wt%) from cylindrically extended rods were obtained from Sandia Corporation. Specimen designations

with corresponding HIPing temperatures and final densities are listed in Table I.

I A nominal thickness of 2.5mm was chosen for all samples, with surfaces polished on 400 and 600 i paper, and finished with Linde A powder. A standardized sample preparation procedure was adopted to

| ensure uniformity and a minimization of surface strain, since minor variations in surface conditions

have been shown to affect small-angle scattering profile (11).

Neutron scattering measurements were performed at the SANS facility of the Oak Ridge National

Laboratory. After preliminary measurements, it was established that the significant part of the

scattering was confined to very small angles (q = .005 to 0.015A“^). Therefore, the

sample-to-detector distance was set at 16 meters for improved resolution on the 64 x 64 element .

604
position-sensitive detector. The measurements were made in a vacuum of 10“^ mm of Hg. The scattering

characteristics of a standard sample (wrought stainless steel) were measured and substracted from each

of the test samples to obtain a true measure of the scattering that is attributable to defects. Data

processing included correction for background measurements and normalization of all data by the

transmission and thickness factors.

TABLE I

Sample # Temp. , °C Density*

B26 800 98.7

B28 800 98.9

B27 900 99.5

B33 1000 99.5

B34 1000 99.2

B36 1100 99.3

A5 1200 99.6

A6 1200 99.8

STD 100.0

*Nominal density is measured relative to wrought material as 100.0.

The Guinier plots from SANS measurements are shown in Fig. 1. For clarity in the display, two different scales are used; the lower set of curves for varying from 0 to 8x10”^, A“^, and the upper

set of curves for varying from 0 to 2x10“^, A~^. The upper set shows the detailed variation in the

maximum pore size, (corresponding to low-q values), while the lower set indicates the relative

intensity of the smaller pores (related to higher-q values).

Figure 1 - Guinier plot of SANS profiles from (21-6-9) stainless steel, showing the relationship

between intensities and compaction temperatures.

605
It is seen that the scattering decreases with increasing temperature. Similarly, the relative |: density quoted in Table I appears to be directly related to the temperature, the lowest temperature

(800°C) corresponding to the lowest density (98.9%) and the highest temperature (1200°) to the highest

density (99.8). The difference in scattering between these two extreme cases, larger than one order

j of magnitude, is consistent with the effect of voids as the major source of scattering. Furthermore, i; the slope of each curve provides information about relative pore sizes. Thus, specimen # B36 shows a f larger maximum pore size compared to other samples, with pore number density somewhat less than

^ # B33-34, while both have nearly the same measured density of 99.3%.

fi Despite some uncertainty in the measured densities there is a strong evidence of a systematic I variation of density with extrusion temperature. The pore size in all samples, calculated from the

Guinier slope, ranges from 300A to 800A and the relatively low scattering Intensity Indicates u correspondingly low void densities. A simple calculation of void density on the basis of the given j weight density yields a number of about 10^^ voids/cm^, assuming an average void size of 500A. This |i number may be converted to a surface density distribution appropriate for observation by SEM or f similar techniques. For the material of density 99.8%, the calculated void spacing corresponding to jl the above number density will be about 0.5 ym (5000A) . Attempts to observe these voids by scanning f electron microscopy proved difficult because of interference from surface irregularities in the |i specimen. I

I Transmission electron microscopy was performed for direct observation of voids. This procedure

is not without problems since the voids which are exposed to the surface are enlarged during the ^ electrothinning processes and contribute to an error in the size distribution analysis. However, the i; interior voids remain unaffected and can provide a good estimate of number density and size. :

Figures 2 and 3 demonstrate the use of two TEM methods used in the observation of voids for

sample B28. In Fig. 2, the specimen is set at S=0 (exact Bragg condition), where S is the deviation

for Bragg angle. Interference fringes from variation of sample thickness are observed. Here, the i intensity contrast from voids changes alternately from white to black to white. The method is thus ' suitable for observation of small voids. In this micrograph voids in a range of sizes varying between

100 and 400 A along with a very large hole (3000A) are visible. In Fig. 3 a multibeam condition with

S^, larger voids (approximately 700A) are visible. In order to minimize the adverse effects of !j electrothinning, a procedure consisting of electropolishing followed by ion-milling (a minimum of Ih : to remove first few hundred anstrom layer from the surface) has been developed. .

I 5-H

5000 A

Figure 2 - TEM micrograph showing the interference thickness fringes

and change of contrast for voids in stainless steel

specimen # B28 processed at 800° C. Figure 3 - Electron Micrographs of Voids in Powder Compacted

Stainless Steel (Fe-21Cr-6Ni-9Mn-0.02C0.25N), Processed at 800°C.

1000 A

606
Porosity In HIPed Ti-Al-4V

Three samples of hydropressed T1-6A1-4V were supplied by Gould Corp. Titanium alloy powder was consolidated in autoclave at a temperature of 1275°C and at pressures of 60, 80, and 100 ksi,

respectively. The powder compacted samples and a standard Ti-6A1-4V were ground to a thickness of

approx. 3.3 mm, then polished on No. 400 and No. 600 Emery paper. As in the preceeding case, the

sample-to-detector distance of 16 m was better suited due to the absence of significant data at larger

scattering angles. Figure 4 shows SANS profiles of the three specimens.

An analysis of pore sizes gives a range of values for pore size ranging from 300A at one extreme

to 1150, 1020 and 880A at the other extreme corresponding to 60, 80, and 100 ksi, processing

pressures, respectively. It may be seen that at higher scattering angles the intensity associated

with the 80 and 100 ksi specimens is higher than the corresponding intensity for the 60 ksi specimen.

However, at very small angles the intensity from the 60 ksi sample exceeds those from the other two. This behavior indicates a refinement of pores with increasing pressure, and is consistent with the

effect of compaction pressure on pore size. Transmission electron micrographs of two samples in Fig.

5 show typical pore sizes observed.

Results of this investigation clearly shows that intensity profiles of long wavelength neutrons

scattered from powder compacts containing very small volume fractions of residual voids can be

effectively analyzed to yield quantitative information on void size and distrubutlons. Furthermore,

the method demonstrates the potential for non-destructive evaluation of defects in rapidly solidified

powder consolidates.

COZ

UJ

LU>

IJJ ir

Figure 4 - Guinier plot from SANS profiles of HIPed Ti-6A1-4V alloys prepared at 60, 80, and 100 ksi.

607
I

Figure 5 - Electron micrographs of voids in Ti-6A1-4V, Compacted at 2400°C (A) 60 ksi, (B) 80 ksi.

Acknowledgments ! This work was funded in part by the Office of Naval Research and has been performed in

cooperation with the National Science Foundation (Oak Ridge National Laboratory SANS Facility). The

authors would like to thank Drs. J. E. Sumgeresky and R. Stoltz of Sandia Corporation, and Dr. M. Dustur of Imperial Clevite Corporaton for supplying the alloy samples and Dr. C. S. Pande for the

| electron microscopic measurements. I References !

1. M. Fatemi and B. B. Rath, "Non-Destructive Evaluation of Defects in Structural Materials Using

Long-Wavelength Neutrons," NRL >femorandum Report No. 4226, May 1980.

2. M. Fatemi, C. S. Pande, M. A. Imam, and B. B. Rath, AIME Fall 1981 Meeting, Louisville, Ky. j

3. M. Fatemi and B. B. Rath, "Non-Destructive Evaluation of Defects in Structural Materials," Report ! to Office of Naval Research, September 20, 1980.

^ 4. G. C. Vonk, J. Appl. Cryst. , 433 (1976). !

5. H. A. Mook, J. Appl. Phys. , 43 (1974). |

!' 6. 0. Clatter, J. Appl. Cryst., JJ, 7 (1980).

j 7. I. S. Fedorova and P. W. Schmidt, J. Appl. Cryst., J_l, 405 (1978). !

8. A. Guinier and G. Fournet, Small-Angle Scattering of X-Rays , J. Wiley and Sons, NY (1955). '

1 9. D. Eylon, R. E. Omlor, R. J. Bacon and F. H. Froes, Powder Metallurgy of Titanium Alloys , Ed. F. j H. Froes and J. E. Smugeresky, The Met. Soc. of AIME, Warrendale, Pa., (1980). j

10. H. H. Hausner, Jr. of the Japanese Soc. of Powder Met. 7 (1960). :

11. M. Roth, J. Appl. Cryst., 10, 172 (1977). J

608
RAPID SOLIDIFICATION - PAST PROMISES AND CURRENT REALITY

A. M. Adair,* L. A. Jacobson,** E. C. van Reuth,** and J. Dickson***

*AFWAL Materials Laboratory, Wright-Patterson AFB OH

**Defense Advanced Research Projects Agency (DARPA), Arlington VA

***Systems Planning Corporation, Arlington VA

ABSTRACT

A number of ongoing efforts are being conducted to develop higher performance structural

alloys by means of rapid solidification processing. In almost every program, some difficulties have been encountered, and additional work has been necessary to address the fundamentals of the specific areas in which problems were not anticipated. The historical background

of rapid solidification is briefly reviewed. The undertaking of development efforts to define the exploitation route for applications requiring improved materials performance is

given. These efforts are described next, with emphasis on those factors which differ from

our original high expectations. Recent studies which have addressed both the general

aspects of current techniques for achieving and consolidating rapidly solidified materials,

and the specific factors which will bear on the future application of such materials to high performance defense systems are summarized. Finally, some suggestions for future,

rapid solidification efforts are made. These are based on lessons learned from experience

and should accelerate the application of rapidly solidified structural alloys.

Introduction

There can be little doubt that we are on the threshold of a new era in metallurgy and more

broadly, in materials science, brought about by the application of Rapid Solidification Technology

(RST) to the production of new and improved materials. As the introductory paper for a session

concerned with applications of materials made by RST, we will treat somewhat selectively the past

work that has brought us to this threshold, and outline some of the expected future advances that

will come from application of RST. Concerning our choice of title for this paper, it is our feeling

that over the past several years, there has been a commendable spirit of enthusiasm growing for the

potential benefits of RST. Also, there has been a body of experience accumulating from efforts to apply RST to the task of developing useful new materials. We have been keenly aware of the many

difficulties that have been encountered in these latter efforts, and these difficulties are what has

prompted us to select our title for this review. The implication is, of course, that our current

experience is not fully bearing out the past promises of RSI. We do not intend to convey a pessimistic note; rather, we hope to stimulate objective examination of this new technology, and a clearer

perspective as to the advantages it offers.

We will limit our discussion to those areas in which RST has been applied to structural materials.

There is, of course, an additional large body of information concerning the successful production of

amorphous alloys by RST. The application of such materials in a variety of useful devices also

depends on the unusual physical properties of the material which play a crucial role in the achievement of significant performance improvements.

It is in the structural alloys arena that RST must face head-on the challenges posed by competitive conventional processing methods, new alloys not dependent on RST, and alternative materials such

as composites that cover a wide range of temperature capability. This is a very competitive arena,

one in which there is now a very keen interest in those mechanical properties which fal I in the

"second tier." First tier properties are those which have the most direct bearing on the performance

of a component--strength , stiffness, and density. Second tier properties are those which influence

the useful lifetime of the component—fatigue with its crack initiation and crack growth components,

corrosion and the effects of corrosive environments on crack growth, creep and stress rupture,

thermal fatigue, oxidation and hot corrosion. Somewhat further in the background, but increasing in

importance, is the area of fabricabil ity and inspectabil ity of structural components. In the efforts

to apply RST materials, we must consider how they can best be made into useful shapes and how they

may differ in their ability to be inspected, based on demanding requirements to guarantee that they

609
are free from defects that could lead to the development of critical flaws in service.

An appropriate conclusion to this introductory section is a discussion of cost. RST involves

additional costs compared to conventional structural materials processing, and we must be aware constantly of the necessity to justify these higher costs in realistic terms which balance them

against achievable benefits. A recent study by Kearns has clearly put materials cost considerations

into perspective [1]. When a materials choice must be made, it will invariably be the lowest cost

material which will accomplish the needed performance. On the other hand, it is very difficult to resolve cost-benefit trades in a generalized context. It is almost impossible to compute the system

life cycle costs saved by incorporating a higher cost, but more durable material. Judging from recent

decisions in the Department of Defense, it appears that the initial acquisition cost of a system is

of greater concern than the future costs of maintenance, inspection and repair. These latter costs

do not clearly show up in the materials selection decision process. We expect this to be more important in the future, but we also must recognize that in the advocacy of new RST materials, we must recognize that RST is a higher cost option to conventionally processed materials. This higher

cost will often have to be justified on bases over which the materials developer has little control.

We suggest that this is a current reality against which the past promises of Rapid Solidification

Technology need to be evaluated. Fortunately, we are faced with many moving targets of opportunity,

and so we proceed with an examination of RST promises as related to some present materials challenges.

Historical Background of RST

The historical development of RST can be viewed from two perspecti ves--the processes by which

rapid quenching from the melt (as the general field has been known until recently) can be achieved,

and the application of such processes to produce useful new alloys. Examining the specific processes,

one finds that in 1871 a patent was issued to E. M. Lang for a method of casting solder wire on the

periphery of a rotating drum [2]. After a rather long passage of time, the next major development

was a chiH mold technique described by Falkenhagen and Hofmann in 1952 [3]. One year later, R. B. Pond, Sr. was awarded a patent on the melt spin method for making metallic filaments [4]. It is of

interest to note that in 1958, Soviet investigators published work on the piston/anvil technique and

the structure of alloys produced by this rapid solidification method [5].

In 1960, Duwez et. al . published a description of yet another rapid quenching method involving a gun or shock tube as the means of projecting molten alloy against various substrates [6]. This work

is generally cited as having first recognized that unusual new solids could be achieved through rapid

solidification. Duwez reported that complete (metastable) extension of solid solubility in Ag-Cu

alloys could be achieved. This alloy system was examined by other investigators who were concerned i with new techniques for achieving rapid quenching from the melt. A few of the more significant are: Mutsuzaki et. al . in 1963, electrical discharge machining [7]; Moss et. al . in 1964, plasma jet

spraying [8]; and Pond and Maddin in 1969, rapid solidification melt spinning [9]. Aluminum alloys were the objects of investigations concerning self-substrate quenching using either an electron beam

(Hiller, 1968) [10] or a laser (Jones, 1969) [11].

Amorphous metallic alloys were first reported by Duwez et. al . in 1960 [12], and research on these materials continued to increase, culminating with the recognition given to Duwez in his selection

to present the American Society for Metals Campbell Memorial Lecture in 1967 [13]. The early 1970's ; were marked by a number of significant process and composition patents issued to investigators working at the Allied Corporation and at the Battelle Memorial Institute. In 1972, Jones and Suryanarayana

published an annotated bibliography of rapid solidification work that covered the period 1958-1972

and included a total of over 400 references, many of which were from work done at universities [14].

A key link between RST and the established field of powder metallurgy was made in the mid 1970's

by Cox et. al. who combined centrifugal atomization with forced convection cooling (inert gas quenching), and achieved rapid solidification rates for powders of nickel, aluminum, and iron base alloys

'

[15]. Other powder process developments occurred subsequent to this time period, among them ultrasonic'

gas atomization as reported by Grant et. al . [16]. A wide variety of processes for making alloy

powders are now being explored as part of the effort to advance rapid solidification technology. Of

particular importance are the nickel base alloys for which increased homogeneity has demonstrable

benefits (especially oxidation resistance) and aluminum alloys with high iron content and either

molybdenum or cerium additions which retain higher temperature strength properties and cannot be

processed by conventional ingot methods.

It has been suggested that the close ties that are developing between powder metallurgy and RST

tend to detract from the stature of the latter as a separate field of endeavor. For structural 1; alloys, however, apart from the capability of rapid solidification powder making methods, it is clear

that other advancements in powder metallurgy (such as isothermal forging and hot isostatic pressing)

provide the most acceptable routes to the fabrication of useful structural shapes. Thus, the many

jf) 610
powder consolidation methods which have been developed constitute an important part of the historical

background. Without this background, much of the rapid progress in RST could not have been achieved.

In fact, many would agree that the production of FlOO turbine engine components from IN 100 powder

beginning in 1972 was the first significant application of RST.

The promise of RST is founded on a sound basis of process techniques and exploratory applications

to a wide variety of alloy systems. You have heard, and you will hear in this conference, many

examples of how the promise of RST is being realized. We turn now to a few points that bear on current realities which we expect will stimulate discussion, thought, and further experimental

efforts to build on the momentum already established.

Current Reality

We intend in this section to place emphasis on several lessons learned, drawing from recent

experience gained on both completed efforts, and programs that have yet to be completed.

Aluminum Alloys

While there has been keen interest for some time on major programs sponsored by the DoD, there

is as yet no volume application of RST aluminum alloys. Most of the effort has been on the high

strength alloys of the Al-Zn-Mg compositions, and they are now undergoing a round-robin evaluation by

industry. The concerns are that their overall property improvements are not sufficient to justify

the additional cost, and, for specific applications, that improved fatigue crack growth resistance is

needed. If these alloys are processed under properly controlled conditions, they offer a narrower

variation in properties, improved fatigue strength, and improved stress corrosion resistance. We should emphasize that the final product quality is equally dependent on the initial powder quality,

its careful handling, method of consolidation, and various subsequent processing steps.

Application of RST to chemical compositions which are normally used in ingot metallurgy can even

have a negative effect. Most high strength aluminum alloys are used in a partially recrystallized

condition, i.e., a wrought-elongated grain structure. However, rapid solidification can result in a fine dispersion of constituents that inhibit recrystallization [17]. The net result of deformation

processing plus heat treatment on RST products can then result in a fully recrystal 1 ized product.

We may learn that high modulus, low density alloys based on lithium additions, processed initially

by RST, can also be processable by ingot methods, depending on lithium content. Work in progress

should provide the answer, as properties of materials made by each method will be compared. The high

transition element content alloys for either higher modulus or higher temperature requirements will

probably require all-inert processing or novel low temperature consolidation techniques, or both.

Much of the benefit of RST processing has been lost due to the need for high temperature outgassing

treatments that invariably result in coarsening of the strengthening particles. High temperature

consolidation methods can lead to the same result. It is not yet clear whether particle surface

oxide can be tolerated, given the current emphasis on both strength and toughness. Recent evidence

presented by Hildeman et. al. [18] suggests that variation in oxide content which is due primarily to particulate (powder) morphology affects the alloy behavior in slow bend toughness testing; the lower

oxide content giving clear indications of improved toughness. Further, they report that even for

spherical products having the lowest total oxide content which were compacted by controlled vacuum

hot pressing, subsequent hot deformation in excess of a 60% reduction in height would probably be

required to achieve optimized mechanical properties.

Nickel Base Alloys

Production of IN 100 turbine engine components continues using a combination of extrusion and

isothermal forging by a patented process. Consolidation of powder to near net shaped components by

hot isostatic pressing (HIP) has been used to produce turbine engine disks for production and experimental studies; however, the use of components after HIP without any hot deformation operations ("AS

HIP") has been brought into disfavor for reasons having most to do with the undesirable consequences

of contamination. Most of the activity on disk alloys has been to improve the overall process or

improve existing alloy compositions rather than investigate a broad range of possible approaches to utilize the advantages of RST. Such a program is planned in the near future. HIP consolidation of

melt-spun ribbon has been attempted, but results indicate that more effort would be needed to achieve

uniform winding, and consequent uniform consolidation shrinkage.

Improved oxidation resistant alloy compositions intended for use in turbine airfoils have

required that elements such as chromium and tantalum be added (which originally had not been expected

611
to be needed). However, the amounts added are much lower than current casting alloy compositions,

and the alloys still do not contain cobalt. Turbine airfoils produced from wafer blocks that were directionally recrystallized to elongated grain structures (similar to directionally solidified cast

airfoils) had to be abandoned due to unacceptable transverse properties. While some of the compositions made by RST have also been successfully cast as single crystal test bars, the best compositions,

containing small amounts of hafnium and yttrium, are not amenable to current practice for casting

single crystal airfoils having complex geometry cores for cooling air distribution.

Rapid solidification plasma deposition of nickel base alloys is an emerging process, but it is too soon to give a fair assessment of its potential. Major programs on this technology are currently

underway.

Iron Base Alloys

While we in the DoD may tend to be concerned with the "more exciting" alloy systems, the ferrous

alloys may well be the "sleeper" in this technology. One of the first successful commercial applications of RST technology was the PM tool steels produced by Crucible. They demonstrated several years

ago that the more uniform chemistry and the more evenly distributed, smaller carbide size led to more efficient cutting and longer tool life. Crucible combined this product with hot isostatic pressing

to produce a variety of tool shapes. With current efforts underway to decrease machining costs by

various approaches, the demand for such improved tooling materials is expected to increase significantly.

Bearing materials have always been a "very interesting" product, at times almost in an atmosphere

of intrigue; due no doubt to the efficiency of the failure process associated with large bearings.

AISI 52100 and M50 steels have long been industry standards; anything new never could find anyone

wanting to change or willing to re-evaluate the relative cost. Most bearing problem areas center

around corrosion and rolling contact fatigue failures, undoubtedly affected by the presence of

corrosion product. Inherent fracture toughness is also a concern for modern turbine engine bearings.

RST bearing steel compositions [19] have been produced and evaluated, whose properties suggest that

attitudes may very well change, especially as the new turbine engine designs contain higher rotor

speeds and fewer stages. Highly corrosion resistant bearing steel compositions have been produced

which have equivalent rolling contact fatigue resistance to today's M50 steel; other compositions

have shown very significant increases in rolling contact fatigue resistance. There is little doubt

that these materials will be very much in demand in the future. Applications are also expected to be

made where high levels of wear resistance is required, such as cam surfaces in diesel engines. As we try to get increased durability for all products we manufacture, specialty applications for these

uniquely microstructured steels will evolve.

By no means should we neglect moderate to high strength steels that are quenched and tempered to achieve their properties. It has been shown that the impurity elements present in relatively low

levels in cast and wrought steels take on new roles in RST produced alloys [20]. Their distribution

is very uniform, they are finely divided, and they effectively prevent grain growth of the austenite

phase during heating. This characteristic opens up a much larger processing and heat treatment

window and offers an opportunity for more favorable combinations of strength and toughness in these

steel s

Iron-aluminum alloys are not new, but achieving grain size control to effect useful room temperature ductility is new and has been a consistent result of RST technology. These materials have

excellent oxidation resistance over a wide temperature range [20], their density is significantly

lower than nickel base alloys, and their opportunities for the application of thermomechanical

processing are many. Burner cans in turbine engines and several static structures are already under

consideration for early implementation. These products must compete with the high chromium containing

stainless steels and the "leaner" composition superalloys, both nickel and cobalt base.

Future Projections

An important suggestion was made at a recent Government-Industry workshop on RST. One participant

expressed concern that the preponderance of Government sponsored RST efforts were addressing the

application of these new materials in some of the most structurally critical components. He suggested

that application efforts might proceed more smoothly if some less demanding components were selected

for initial demonstration. We admit that this suggestion is eminently sensible, but is not very

likely to have much effect, given the circumstances in which new RST structural alloys must compete

with other materials for limited development and demonstration resources. This competitive situation

is unlikely to significantly change in the future, so the efforts to achieve the highest possible ' payoff in performance will continue to emphasize the use of RST materials for the most challenging

612
components. The generation of a knowledge base for RST materials based on first-hand experience will

consequently be delayed.

A more basic factor which continues to cloud the RST picture for the future is the lack of

information on the unequivocal effects of higher solidification rates on the properties of consolidated

shapes representative of a particular application made from a single alloy composition. Until

controlled studies are performed, in which solidification rate is the single variable, other factors

being held as constant as possible, we will not have a clear perspective on the real benefits of

rapid solidification. For example, the recently completed work by Hildeman et. al. [18] used the same aluminum alloy compositions, and when cooling rate was varied among splat, gas atomization, and

centrifugal atomization, the morphology changes associated with these processes had a greater effect

than differences in cooling rate.

Another fundamental problem has to do with the effects of different powder production and

handling environments on the properties of consolidated RST alloys. We expect that different alloy

systems will differ in their sensitivity to environmental inertness, and might also anticipate that

different alloys within a given system would also respond differently. Controlled studies are also

needed to elucidate the effects of different environments on mechanical properties, especially second

tier. These last two points serve to remind us that our efforts to date have been concerned with

determining sufficient conditions for achieving benefits through RST, but we have not really sought

out the necessary conditions. If the necessary conditions turn out to involve lower solidification

rates or less stringent requirements for inert atmospheres, strict adherence to a set of sufficient

conditions may impede RST materials application potential.

A final point bearing on RST structural applications is our perception of the need for more understanding of the influence of various consolidation processes on the properties of the RST

material. In addition to the more well established methods of consolidation such as HIP, hot pressing

and hot extrusion, other methods such as plasma spray or melt-spray deposition, dynamic compaction

and ultrasonic compaction should receive greater emphasis. Work is already underway on a number of

these consolidation techniques, but results are not sufficient to predict the possible advantages or disadvantages of each. Again, it can be expected that different RST alloys will respond differently

to the various compaction methods.

Concl usions

Perhaps the most important lesson learned to date that applies to the field of Rapid Solidification Technology for structural alloys is that an atmosphere of optimistic over-advocacy should be

carefully avoided. The principles that have guided the development of current widely used high

performance structural materials are still valid ones, even for RST. However, RST processing has

distinctive capabilities that have tempted us to place it outside the mainstream of past alloy

developments. True, RST does offer unique new possibilities for control of alloy microstructures,

but the question remains: Given this new control, what sort of microstructures should we try to produce? We do not yet seem to be capable of establishing alloy microstructure goals for the simultaneous optimization of both strength and toughness, based on clearly demonstrated fundamental

principles. Perhaps the next steps needed are the devising of programs and the conducting of critical

experiments, using RST materials, to shed some light on the critical microstructural parameters

affecting mechanical properties of interest, both singly and in combination.

In the context of recent lessons learned, which have combined to add an air of caution to some

of our enthusiasm, what can we now predict for the future of RST? It is useful to remind ourselves

that actual directions which some lines of inquiry take do not always follow those that are predicted

for them at their early stages. In his 1967 Campbell Memorial Lecture, Professor Duwez made the

following remarks in his conclusion: "As Professor Gilman told you (in his lecture given the previous

year) . . . alloy monocrystals are a reality and may be the next generation of turbine blades in jet

engines. I would not dare to predict such a spectacular future for alloys quenched from the liquid

state, and I am sure they will never be used in jet engines." [13]

Ten years after that prediction, a braze alloy made by rapid quenching from the melt in ribbon

form had clearly shown superiority over its powder counterpart. This alloy is now used to join jet

engine components. And fourteen years later, the potential of rapidly quenched nickel base alloys

for turbine blades was firmly established through engine testing of two blades fabricated from

Ni-Al-Mo-W alloy powder (RSR 185). Variations of this alloy are now showing the clear benefits that

rapid quenching from the melt can give to high temperature strength and oxidation resistance, improvements that are due to the high degree of chemical homogeneity that results from rapid solidification.

So one past prediction has not, in fact, been borne out by later experience. There are also

examples for which the actual experience has exceeded expectations. Despite our high expectations

613
for structural alloy performance improvements, we know that the higher cost of RST will have to be

offset by some other form of economic benefit, and we suggest that the best context in which to make

such cost/benefit determinations is the specific use environment. The nickel base alloy disks are

the only structural alloy RST applications for which the benefits have been balanced against costs.

We must have other examples of actual experience to build a broad-base economic acceptance of RST

applications. Competition from other quarters will also cause the perspective and expected benefits

to change.

References

[1] Kearns, T. , Private Communication.

[2] Lang, E. M. , "Machines for Casting Solder Wire," U.S. Patent No. 5,862, Reissued 1874.

[3] Falkenhagen, G. and Hofmann, W. , Z. Metal Ikunde 43 (1952) 69.

[4] Pond, R. B., "Metallic Filaments and Method of Making Same," U.S. Patent No. 2,825,108; 1953.

[5] Salli, I. V., Liteinoe Proizvodstvo , 1958, No. 7, 22-25.

[6] Duwez, P., Willens, R. H. and Klement, W. , J. Appl . Phys. , 31 (1960), 1136-7.

[7] Mutsuzaki , K. , Suzuki, M. , and Kawai, E., J. Jap. Inst. Metals , 27 (1963) 424-7.

[8] Moss, M. , Smith, D. L., and Lefever, R. A., Appl. Phys. Letters , 5 (1964) 120-1.

[9] Pond, R. and Maddin, R. , Trans-Met. Soc. AIME , 245 (1969) 2475-6.

[10] Hiller, W. , Met. Contr. Mecan (La Metallurgie) 100, No. 1 (1968) 35-6.

[11] Jones, H. , Mater. Sci. Eng. 5 (1969) 1-18.

[12] Klement, W. , Willens, R. H., and Duwez, P., Nature 187 (1960) 869-79.

[13] Duwez, P. Trans. Am. Soc. Metals , 60 (1967) 607-33.

[14] Jones H. and Suryanarayana , C., J. Materials Science , 8 (1973) 705-53.

[15] Cox, A. et. al.. Reports on Contract F3361 5-76-C-51 36.

[16] Anand, V., Kaufman, A. J., and Grant, N. J., "Rapid Solidification of a Modified 7075 Aluminum

Alloy by Ultrasonic Gas Atomization," Rapid Solidification Processing Principles and Technologies II , Edited by R. Mehrabian, B. H. Kear and M. Cohen, Claitors Publishing Division, Baton Rouge, LA (1980)

273-286.

[17] Jacobson, L. A., Pierce, C. M. , and Cook, M. M. , "Microstructures of Powder and Conventionally

Processed 7075 Aluminum Alloy," AFML-TR-71 -240, W-PAFB, Ohio, December 1971.

[18] Hildeman, G. J., Lege, D. J., and Vasudevan, A. K. , "Fundamentals of Compaction Processes for

Rapidly Quenched Prealloyed Metal Powders," AFWAL-TR-82-41 56, W-PAFB, Ohio, October 1982.

[19] Adam, C. M. , Bourdeau, R. G., and Brock, J. W., "Application of Rapidly Solidified Alloys,"

AFWAL-TR-81-4188, W-PAFB, Ohio, February 1982.

[20] Proceedings of Briefing/Workshop on RApid Solidification Technology , R. Mehrabiati & Organizing

Committee, COMAT; National Bureau of Standards, 1-2 July 1981.

614
AEROSPACE STRUCTURAL APPLICATIONS OF RAPIDLY SOLIDIFIED

ALUMINUM-LITHIUM ALLOYS

. * * + + **

R. E. Lewis , I. G. Palmer , J. C. Ekvall , I. F. Sakata , and W. E. Quist

* Lockheed Palo Alto Research Laboratory, Palo Alto, CA, 94304

+ Lockheed-Cal ifornia Company, Burbank, CA, 91520

** Boeing Commercial Airplane Company, Seattle, WA, 98124

ABSTRACT

Advanced aluminum-lithium alloys are being developed to save a significant amount of

weight (10 to 20 percent) in new aircraft airframe structures, by replacing the currently

used 2xxx and 7xxx high strength aluminum alloys. The weight saving of aluminum-lithium

alloys is directly related to their lower density, higher modulus, and high strength.

Both powder metallurgy and ingot metallurgy techniques are being used in alloy development programs. Rapid solidification processing of the alloys is currently being evaluated because improved ductility and toughness may be achieved through fine grain size,

freedom from massive segregated phases, and an increased amount of elements that form

fine, stable dispersoids. A number of property goals established for the development of

aluminum-lithium alloys is reviewed. A comparison with these goals of the best properties of alloys in current development is presented. Specific, analyses of applications of

aluminum-lithium alloys to different types of aircraft are made, with particular attention

to the advantage of using a variety of low density alloys exhibiting different properties.

Finally, a number of technical issues regarding the development and acceptance of the ma- terials by the supplier and user communities are reviewed.

INTRODUCTION

In the development of new aerospace structures, both increased structural efficiency and reliability are goals to be achieved, and at the lowest cost. Improvements in structural efficiency are achieved by applying more accurate or effective structural designs or by using new structural materials that will reduce weight, or a combination of these. For a majority of aerospace structures, designers believe that with currently available materials, redesign alone will not result in a significant further improvement in structural efficiency. Any substantial additional gains in weight savings

will have to be achieved by applying advanced materials that are lighter, stiffer, stronger, more

fatigue resistant, or more damage tolerant. The successful development of such materials will benefit

all aerospace applications including aircraft, missiles, and space vehicles.

Improvements in aluminum and titanium alloys as well as composite materials are all topics of

high current interest in the aerospace community. Of these, improved aluminum alloys, such as aluminum-lithium, will have obvious advantages because of their comparatively low cost and easy assimilation using the extensive existing manufacturing facilities. Lithium additions to aluminum are particularly effective as lithium reduces the density and at the same time increases the stiffness of aluminum. Both property changes are important contributions to weight savings in aerospace structures,

with density decrease being the most important [1]. Only one other element, beryllium, when added to aluminum has a similar effect on the properties, but it is not an attractive choice for most applications because of its high cost and toxicity.

Rapid solidification processing (RSP) of new aluminum-lithium alloys is currently being evaluated

with the objective of improving ductility, toughness, and strength. These improvements may be

achieved using the attributes of RSP including finer grain size and solidification structure, extended

solid solubility, and suppression of coarse second phase particles.

The status of current activities in this important alloy development and processing field is re- viewed briefly below, along with appropriate aluminum-lithium ingot metallurgy developments which

share the same general objectives. The major emphasis of this paper, however, is to review the airframe structural designer's assessment of the potential contributions of aluminum-lithium alloys, and

why there is increasing interest on the part of both prospective suppliers and users. In the final

615
section, a review is presented of the technical issues that are involved in the successful development

and acceptance of these new materials.

i

PROPERTY GOALS FOR ALLOY AND PROCESS DEVELOPMENT

Current property goals of a number of aluminum-lithium alloy development programs are shown in | Table 1. Included are two powder metallurgy (PM) alloy development programs being performed by Lockheed Missiles & Space Company, Inc. (LMSC) and the Boeing Commercial Airplane Company using rapidly ' solidified particulate, a study by Lockheed-Cal ifornia Company (CALAC) in which the availability of ] both (PM) and ingot metallurgy (IM) alloys was assumed, and IM alloy development programs being per- ' formed by Alcoa, Alcan and British Alcan Aluminium Company (formerly British Aluminium Company). In all cases the baseline alloys are those currently being used on existing aircraft, and the property i goals represent for the most part improvements in density, modulus and specific strength with other

j properties remaining the same, so that the new alloys would be used on a direct substitution basis for

j existing alloys, thereby minimizing the requirement for redesign. A number of smaller alloy develop- i ment programs and programs involving mechanical alloying are not included in Table 1. Details of many

j of these programs are given in the proceedings of the First International Aluminum-Lithium Conference ' [2] and will be supplemented in the upcoming Second International Aluminum-Lithium Conference [3]. I

i The LMSC/DARPA program started in 1978. The objectives are to develop alloys meeting the property

' goals shown in Table 1, and to establish a metallurgical basis for manufacturing scaleup and application to new aerospace systems. Phase 1 (Alloy Development) has been completed and reported [4,5,6]. ^ In this phase three iterations of alloys were evaluated, involving twenty-one compositional variations

made from both splat and fine atomized powders. In the first iteration of alloys Al-Li-Cu and Al-Li- ' dispersoid alloys were examined, made from splat particulate. In the second iteration Al-Li-Cu-Mg al- | loys were evaluated, also made from splat particulate. In the third iteration the most promising com- ] positions from the first and second iterations were evaluated, made from fine atomized powders.

In Phase 2 (Scaleup) of the program two Al -Li-Cu-Mg-Zr alloys were prepared in the form of both

^ PM and IM alloys. The target compositions were Al-3Li-l.5Cu-lMg-0.2Zr and Al-3Cu-2Li-lMg-0.2Zr. The ] PM alloys were prepared by Pratt and Whitney Aircraft from chill cast remelt stock using centrifugal 4 atomization in helium, followed by screening, degassing, canning, and extrusion. The weight of the ex- trusion billets was approximately 30 lb. The IM alloys were prepared by British Alcan Aluminium by

i direct chill (DC) casting of 140 lb billets, followed by homogenization and extrusion. Preliminary

j tensile properties of the alloys have been measured, and a full spectrum of engineering property data I is currently being obtained. Plans for further work on this program involve the optimization of com- | position in the Al -Li-Cu-Mg-Zr alloy system, including in particular the evaluation of higher zirconium I levels (in the range 0.5 to 2.0 wt percent) in order to utilize the potential advantages of rapid sol idij

ficationmore fully than were used in the Phase 2 alloys.

The Boeing/AFWAL program started in 1982. The objective is to develop a family of aluminum alloys [ with significantly lower density and higher specific modulus and strength than existing high strength j| ingot alloys. The primary thrusts of the program are goal assessment, alloy development, and applica- ij tion payoff studies. Phase 1 (Team Organization and Goal Assessment) has been completed and Phase 2 (Overall Alloy and Process Development) is on-going [7]. The property goals shown in Table 1 are the

| preliminary goals defined at the start of the program; the actual goals have not yet been published. i Three different classes of alloys are being examined, representing three different production methods.

Class I comprises rapidly solidified powder alloys produced by Pratt and Whitney Aircraft by centrifuj gal atomization in helium. Class II comprises mechanically alloyed powder alloys produced by International Nickel Co., and Class III comprises ingot alloys produced by DC casting by British Alcan J

Aluminium. The primary alloy systems include Al -Li-Cu-Mg-Zr and Al-Li-Zn-Mg-Cu-Zr in Class I, Al-Li-Mg ij and Al-Li-Cu-Mg in Class II, and Al-Li-Mg-Zr in Class III. Powder has been prepared and consolidated ; for alloys in Classes I and II and mechanical property and microstructural evaluation of the alloys is

j currently being performed.

j The CALAC/NASA program was started in 1980 and completed in 1981; the final report of the program

j has recently been issued [8]. The study identified the weight savings and economic benefits that

v/ould result from the incorporation of advanced aluminum alloys in future commercial aircraft. In ad- | dition an estimate was made of the future demand for new PM and IM aluminum alloy products. Realistic 1 property goals were established for advanced aluminum alloys, in coordination with ALCOA, to maximize ,,j aircraft system effectiveness. Five alloy classes were identified to meet the requirements for future

| airframe design, and were classified according to the major property requirements as follows: A,

’ strength; B, strength and corrosion resistance; C, strength, stiffness and corrosion resistance; D, durability and damage tolerance; E, low density, high stiffness. Two of these alloy classes, C and E, j would contain lithium. Detailed properties were defined for all five classes of alloys for sheet, r plate, extrusion and forging products, and are described in the final report [8]. )

616
TABLE 1. Abbreviated Property Goals of Current

Development Programs for AluminumLithium Alloys

COMPANY/SPONSOR MFC TYPELOW STRENGTHMEDIUM STRENGTH

HIGH STRENGTH OADTA

LMSC/DARPA PM 7075-T76 7075-T76

+30% E/P + 20% e/p

+20% ay/p BOEINC/AFWAL * PM 2024-T3,T4 7075-T73 707S-T73 -104P +30% E/p +20% E/p

+ 20% (ry/P CALAC/NASA PM,1M 2024-T3 7075-T76 -10%P 10-20%<rv

84 t TICUE + 10% FA- TICUE ALCOA W/BOEINC * IM 2024-T3 7075-T6 2324-T39

-10% -7%P -9%p BACO. RAE 1 2024-T3 2014-T6 7075-T73 ALCAN, RAE 1 IM -10IP -10%P -10%P BRIT, AEROSPACE j ALCOA/NAVAIR IM 2024-T3 7075-T6

-8%/> 8%P

•PRELIMINARY PROPERTY COALS.

TABLE 2. Tensile Properties of LMSC/DARPA

Phase 2 Al-Li-Cu-Mg-Zr PM Alloys

7075-T76 BASELINE GOAL A GOAL B ORIGINAL RtVlSED TARGET* ACTUAL TARGET* ACTUAL

F. , ksi ty 67 70 67 82.4 76 82.2

F, , ksi tu

73 80 73 90.4 87 86. 1

ELONG., %

10 10 10 4.1 10 6.7

p, lb In.

0. 101 0.101 0.091 0.090 0.095 0.093

E, 10® psi 10.3 10.4 12.

11.8

E/p, 10® Ibf in. lb"' 102 103 133 124

CTy /p , 10^ Ibf in. lb

^ 663 693 736 916 796 384

‘TARGET PROPERTIES SHOWN WERE DEFINED AT THE START OF THE PROGRAM AND

WERE BASED ON THE ORIGINAL BASELINE VALUES FOR THE 7075-T76 ALLOY.

The ALCOA IM alloy development program started in about 1981, and represents a continuation of

previous in-house work and work funded by NADC. Very little information on the program has been re- leased, apart from the three property goals shown in Table 1, which were defined in collaboration

with Boeing [9]. The intention is to have alloys commercially available by 1985. Alcoa is also performing a smaller program, sponsored by NADC, with similar property goals, as shown in Table 1.

The British Alcan Aluminium (BAACo) work started in about 1978, and was a continuation of earl ier

work performed by Fulmer Research Institute and High Duty Alloys. BAACo has concentrated on developing the casting and processing technology to make sheet and extrusions. A DC casting technique is being used, with some modifications to improve the safety aspects. BAACo is currently casting 800 lb

ingots of dimensions 20 x 7 in, and the intention is to scale up to 400O-8000 lb rolling slabs by

1985 [10]. Initial work was performed on the Al-Mg-Li system, but most of the work is now being performed on an Al-Li-Cu-I^-Zr alloy developed by the Royal Aircraft Establishment (RAE), Farnborough, UK. The work at ALCAN International started in 1980, and the intention is to develop the casting technology to make large ingots for plate. In order to avoid the hazards associated with DC casting an alternative casting system is currently being used. ALCAN is presently casting 310 lb ingots of dimensions 5 X 12 in, and the intention is to scale up to 6000 lb ingots by 1984 ([11]. Most of the

work is currently being performed on the RAE alloy. The alloy developmeht goals shown in Table 1 for

both ALCAN and BAACo were defined in collaboration with British Aerospace.

For the alloy development programs shown in Table 1 mechanical property data is only available

at the present time for the LMSC/DARPA PM program and the BAACo IM program. Mechanical property data

on Phase 1 of the LMSC/DARPA program have been reported [4,5,6]. Data on Phase 2 of the program consists of preliminary tensile data generated on small scale extrusions of cross-section dimensions

0.28 X 0.56 in, made at LMSC using the same fine atomized powder as for the larger Phase 2 extrusions.

Data have been obtained for alloys solution heat treated at 1000°F, water quenched, stretched 2 percent and aged at 375°F to the peak aged condition. The data are shown in Table 2 where they are also

compared with the program property goals A and B. Both alloys show higher yield strengths than the

program goals but lower ductility. The high yield strengths are attributed partially to unrecrystallized microstructures . The effect of various underaging treatments on strength and ductility is cur- rently being studied, in order to define aging conditions that will produce lower strength levels

that are closer to the program goals, and at the same time higher tensile ductility. Mechanical property data on the BAACo program consists of tensile and other data on an alloy of the RAE composition.

Recent tensile data are shown in Table 3 for sheet, plate and extrusion products, where they are com- pared with data for a draft specification for sheet, and target data for plate [lOJ. The data for the

three different property goals were obtained by different aging treatments on a single alloy composition. Strength levels are equal to or greater than the target values; however, elongations are slightly lower, Th<= sheet data are for 0.063 in thickness, the plate data are for 1.0 in thickness equivalent in degree of reduction to 3.0 in. plate, and the extrusion data are for a cross-section 2.5 x 0.625 in. All materials were solution heat treated, water quenched, stretched and artifically aged.

APPLICATION STUDIES FOR SMALL MILITARY AIRCRAFT

The aircraft selected for the weight savings analysis was the S-3A patrol aircraft. The S-3A is

a subsonic, carrier-based, anti-submarine warfare aircraft that is fully operational with the U.S.

617
TABLE 3. Longitudinal Tensile Properties of

BAACo Al-Li-Cu-Mg Alloys Prepared

from DC Cast Ingots

TABLE 4. Simplified Formulas for Weight Ratio

& Margin of Safety Correction Factors

CATEGORY NO FAILURE MODE

W2

WEIGHT RATIO MARGIN OF SAFETY CORRECTION FACTOR

1 TENSILE STRENGTH

"1 f'"2

1

1 « M.S

2 COMPRESSIVE STRENGTH Us. !!l!

"1 f'V2

1

1 M.S

3 CRIPPLING

^2

^1

6,

•1 -2S

Fcy2

2S

i 1

1 1 M.S.

L J

.5

4 COMPRESSION SURFACE COLUMN & CRIPPLING

"2 £1£2

4

Fcv2

.2

1

1 « M.S.

5 BUCKLING

COMPRESSION OR SHEAR

^2 p .33

1

1 * M.S.

33

6 AEROELASTIC STIFFNESS

“2 E,

"1 £2

1

1 * M.S

7 DURABILITY AND DAMAGE

TOLERANCE ASSESSMENT CUTOFF (DAOTAJ

“2 fl

'“l £2

1

i * MS.

8 GENERAL INSTABILITY. COMPRESSION OR SHEAR a £1£2

.5

1 MS.

.5

9 MINIMUM GAGE

P2 <2 1

1 MS.

Navy aircraft carrier fleet. It was designed and manufactured by the Lockheed-Cal ifornia Company and

all necessary weight and structural analysis data are readily available. The S-3A structure is con- sidered to be representative of other small military aircraft such as trainers and fighters. It is also representative of a future subsonic Vertical/Short Takeoff and Landing (V/STOL) aircraft of interest to the Navy as illustrated in Fig. 1. The S-3A performance and structure have many similarities with leading versions of the subsonic V/STOL aircraft that are currently in the preliminary design stage at several aircraft companies.

Weight Evaluation Methodology . The numerical method developed by Lockheed [1,12J for accurately predicting weight savings of an aircraft structure from properties of the advanced aluminum alloys was used for this investigation. The flow chart for this structural weight evaluation is shown in Fig. 2. The structural weight of each of the components to be evaluated is distributed according to the primary failure modes that size each component. Nine failure modes are considered. Equations that re- late changes in weight to the. basic material properties are then used to calculate the weight savings.

The maximum potential weight savings are determined for the improved material properties assuming

that the primary failure modes do not change. The relationship between the structural weight, W2 , for

the advanced material and the structural weight, W] , for the baseline material is obtained by substituting the appropriate material properties into the weight ratio equations of Table 4. The potential

weight savings are then determined from the following equation:

Weight savings in pet of W-] = 100 (1 - W2 /W] ) 0 )

Fig. 1. An artist's rendition of one version

of the subsonic V/STOL airplane

Fig. 2. Structural weight evaluation for a new material

618
Fig. 3. Examples of critical failure modes in

S-3A patrol aircraft

TABLE 5. S-3A Patrol Aircraft Weight Statement

ITEM WEIGHT (lb)

PERCENT

OF CROSS

WEIGHT STRUCTURE 13,336 27.9

(WING)* (3,415) (7,2) (FUSELAGE)* (3, 388) (7,1) (EMPENNAGE)* ( 729) (1.5) (CONTROL SURFACES)* (1,393) (2.9) (OTHER) (4,411) (9.2)

PROPULSION 3,394 7.1 FIXED EQPT AND SYSTEMS 9,656 20.2 PAYLOAD, CREW, OPER. EQPT 10,420 21.8 FUEL 10,994 23.0

CROSS 47,800

^ALUMINUM COMPONENTS CONSIDERED FOR WEIGHT SAVINGS ANALYSIS

If some primary properties are improved considerably more than others, or if some properties are improved while others remain unchanged or are degraded, other failure modes may become critical for portions of the structure. To determine the real potential weight savings, the effects of alternate

failure modes must be considered. To do this, the ratio W2/W] for the alternate failure mode (or

modes, if more than one alternate mode is critical) is multiplied by the factor given in Table 4. The

margin of safety used in these correction factors is the margin of safety of the alternate mode relative to the primary failure mode. For any given portion of structure, the lowest value of W2 /W] , including the effects of alternate failure modes and margins of safety, is substituted in equation (1) to obtain the real potential weight savings.

Allocation of Component Weight by Design Failure Modes . The structural arrangement and examples of

the critical failure modes in the S-3A patrol ai rcraft are shown in Fig. 3. The weight breakdown used

in the analysis is shown in Table 5. Only the primary wing, fuselage, empennage structure and the

control surface structure (e.g., flap, aileron, rudder) were considered. The landing gear, nacelle

and air induction system are primarily made of steel or titanium alloys. The primary and secondary

structural components considered weigh 8925 lb, 67 percent of the total structure weight or 19 percent of the gross weight.

Allocation of weights of individual components into the nine design failure mode categories was made by reviewing the S-3A structural analysis reports and applying prior applicable experience. This

weight breakdown in shown in Table 6. For example, thirty percent of the structure is designed by tensile strength. The primary wing structure weight consists of structure designed equally by 6 categories: (1) tensile strength, (2) crippling, (3) compression surface, (4) buckling, (5) aeroelastic

stiffness and minimum gage, and (6) durability and damage tolerance. Tensile strength is dominant for

the fuselage basic structure design or 55 percent of the fuselage weight.

Alternate failure modes and margins of safety for the S-3A structure were determined from stress

analyses and are presented in Table 7. Note that category 7, durability and damage tolerance, is the

next critical failure mode for the tensile strength critical structure.

TABLE 6. Allocation of Component Weight by

Design Failure Mode for S-3A Structure

FAILURE CATEGORY DESCRIPTION WING

(lb) FUSELAGE(lb) EMPENNAGE(lb) CONTROL

SURFACES

(lb) TOTAL(lb) (%)

TENSILE STRENGTH562 1873 75 178 2688 30.

2 COMPRESSION STRENGTH

3 CRIPPLING 581 247 452 1280 14.4

4 COMPRESSION SURFACE

583 138 721 8.1

5 BUCKLING 583 740 54 321 1698 19.0

6 AEROELASTIC

STIFFNESS

445 184 80 275 984 11.0

7 DADTA 527 476 60 47 1110 12.4

8 GENERAL

INSTABILITY

91 91 1.0

9 MIN GAGE 134 24 75 120 353 4.0 TOTAL 3415 3388 729 1393 8925 100

TABLE 7. Alternate Failure Modes and Margins

of Safety for S-3A Structure

PRIMARY

FAILURE MODE

CATEGORY

ALTERNATE FAILURE MODES

NEXT CRITICAL SECOND NEXT CRITICAL CATEGORY MS CATEGORY MS

1 7 0.15 6 0.30

2 NOT APPLIC)0iBLE NOT APPLIC/(BLE

3 6 0.30 7 0. 30

4 6 0.30 7 0.30

5 3 0.30 6 0. 30

6 5 0.30 7 0. 30

7 1 0.20 6 0.25

8 6 0. 15 5 0.20

9 6 0. 10 5 0.20

619
Advanced Aluminum Alloy Applications . The aforementioned weight savings evaluation methodology was used to determine the benefits from the application of advanced aluminum alloys to the baseline S-3A

aircraft. The results of this analysis are shown in Fig. 4. These weight savings versus density

ratio are based on the following premises.

0 Baseline material properties are for 7075-T76 extrusion,

0 Goal A properties are the same as baseline except for 30 percent increase in E/p,

0 Goal B properties are the same as baseline except for 20 percent increase in E/p and

20 percent increase in specific (density compensated) tensile, compression, and shear

strength ,

0 Optimization of weight savings by combined selective use of different alloys was not attempted. However, if the alloy being evaluated resulted in a negative saving in a category, the baseline material was not replaced and the weight savings assumed zero |

for that category,

*

0 The allocation of weight by design failure mode and margins of safety was in accor- | dance with Tables 6 and 7, respectively. J

The resulting weight savings for Goals A and B are presented versus density ratio. As indicated ? in Fig. 4, the effects of the alternate failure modes are significant for the higher density ratio ^ materials and must be included to determine the actual weight savings.

^ Example Analysis of Weight Savings . An advanced aluminum-lithium alloy with an assumed material pro-

| perty set representing Goal B with reduced density of po/p-] = 0.94 was selected for this analysis. j Since Goal B is a 20 percent increase in E/p, F^^,/p » ^na Fj.y/p; tbe ratios of E 2 /E,,

i Fcyo/Fcyi equal to 1.2 x 0.94 = 1.128. The^values of F^'^/\[)j/\ and t^iin assumed unchanged. The

' maximum'^potential weight savings for this aluminum-lithium alloy system is summarized in Table 8 . The | weight saving by category is shown for the S-3A structure. The use of an aluminum-lithium alloy havj ing these assumed properties results in a calculated weight savings of 1022 lb or 11.5 percent for the | S-3A aircraft. The actual weight savings and the maximum potential weight savings are equal for this

| example, indicating that the margins of safety are sufficiently high to preclude any influence of the

[ next critical failure mode on weight savings. ;

Effect of Property Trends on Weight Savings . Since the solute content essentially fixes both the denity and modulus in the candidate alloy systems, the weight savings for each alloy system can be esti- | mated as a function of strength level. Shown in Fig. 5 is the change in modulus and weight saving ! plotted against density for four strength categories for Al-Li-Cu-Mg alloys and with copper content I varying with lithium content, not exceeding the solid solubility limit.

Fig. 4. Structural weight savings versus density ratio for the S-3A aircraft

TABLE 8 . Weight Savings Example Analysis for

S-3A Patrol Aircraft

STRUCTURE MAXIMUM POTENTIAL CATEGORY WEIGHT WEIGHT SAVINGS

(lb) (lb) {%)

1 2888 448 16.6

2 0 0 0

3 1280 147 11.5

4 721 91 12.6

5 1898 185 9.7

S 984 183 18.6

7 1110 0 0

8 91 10 11.5

9 353 0 0 TOTAL 8925 1022* 11.5*

BASELINE ALLOY 707S-T7t: F^75k»l; CS kai; E.^10.« M»i; 101 Ib/in.^ ADVANCED AJ-ICu-2U-1Ma-.2Zr: F„ M kal; F^ 72 kal; E^ll.6 Mal; P.0S*lb/ln.^

*ALSO EQUALS ACTUAL WEIGHT SAVINGS

620
Effect of Variation of Individual Properties on Weight Savings . The effect of variation of individual

property improvements on weight savings for the S-3A aircraft is illustrated in Fig. 6. Density affects the weight savings in all failure mode categories: the structural weight savings in each category is directly proportional to the density ratio. Improvements in either strength or stiffness affect structural weight by about 2.5 to 3.5 percent per 10 percent property improvement. Fig. 6 indicates that DADTA property improvements have the least impact on overall structural weight because of

the small amount (10-12 percent by weight) of structure in Category 7. Flowever, since Category 7 is often one of the next critical failure modes in Categories 1 through 6, its impact on weight savings

can be considerably more significant than indicated. The results shown in Fig. 4 also indicate the

importance of considering secondary failure modes on projected weight savings for new alloys.

Additional Weight Savings Due to Selective Use of Several Alloys . Maximum weight savings for aircraft

structures can best be achieved by selective use of several advanced aluminum alloys. A high specific

strength alloy is best for failure modes 1 and 2. A high-modulus, low-density is best for failure

modes 5 and 8. A combination of strength and stiffness, both with low-density is best for failure

modes 3 and 4. Finally, high DADTA properties are best for failure mode 7. A combination of advanced

alloys with different improvements in specific properties would show a higher weight savings potential

than the weight savings resulting from application of a single new alloy. However, the properties for

new alloys should be achieved without significant reductions in the other mechanical properties or alternate failure modes with low margins of safety would become critical and limit the weight savings.

APPLICATION STUDIES FOR TRANSPORT AIRCRAFT

Several government and industry studies have recently been conducted on the use of low-density

alloys in transport and other aircraft [1,4,7,8,13,14]. Each of these investigations has shown that

substantial reductions in structural and total aircraft weight can be achieved through the use of lowdensity aluminum alloys and that aggressive alloy development on this system is appropriate. Of

course, other types of large subsonic and supersonic aircraft (to Mach 2), including cargo and bomber

types, can also make use of these alloys and to approximately the same degree of effectiveness as with transport aircraft. At speeds above Mach 2 aerodynamic heating becomes an increasing factor and

alloys designed specifically for elevated temperatures are required.

^tu Fty 655 (95) 586 (85) 586 (85) 517 (75)

517 (75) 448 (65) 448 (65) 379 (55)

12

(/) „ IO

10

DENSITY -J 1_

3.6 3.2 2.8 2.4 J L _L

2.0 1.6 _J 1_

I I AI-3CU1.2 0.8 1 Mg X Li _J Al - 1 .5 Cu - 3.6 3.2 2.8 2.6 2.0 1.6 1.2 0.8 1 Mg - X Li LITHIUM CONTENT - WT %

10 15 PROPERTY IMPROVEMENT (%)

20 25

Fig. 5. Relation of density to lithium content,

modulus, and S-3A aircraft weight

savings for Al-XLi-YCu-lMg alloys

Fig. 6. Relationship of individual property

improvement on S-3A aircraft weight

savings

621
Approach , ihe approaches now being considered for the implementation of low-density aluminum-lithium

alloys on transport aircraft are as follows:

1. Substitution of low-density aluminum alloys into current aluminum structure with little

or no redesign.

2. Redesign and resizing of structure and airframes to best utilize low-density aluminum

alloys. Non-aluminum components may also be incorporated into this redesign.

Both of these approaches offer certain advantages in economics, alloy design and utilization,

and aircraft structural efficiency. For example, the first approach allows a more rapid introduction

of new alloys into current and near production aircraft, and will yield a substantial weight saving

to the user and minimize complications for the manufacturer. For alloy producers the attractive features of aluminum-lithium alloys are the immediate broad market potential and the opportunity to quickly amortize alloy development and new facility costs. The utilization method based on alloy substitution, however, does not allow optimum use of the attributes of the aluminum-lithium alloys as the new alloys must meet all of the in-place design parameters. This, in effect, dictates that density decrease is the only property that will make a substantial contribution to improved structural

efficiency albeit the effect will be an important one.

The second approach can bring all the unique properties of the aluminum-lithium alloy system into

play, including enhanced fatigue crack-growth-rate resistance and substantially increased modulus.

With these and other alloy property attributes considered in conjunction with aircraft resizing,

studies have shown that an increase of about 50% over the original weight savings may often be obtained for large transport aircraft [8]. The substitution of low-density aluminum alloys for other

aircraft structural materials (principally titanium and steel) may also contribute additional weight

savings but because the majority of the current transport structure is aluminum, as shown in Fig. 7,

the degree of possible substitution for these materials is limited.

Fig. 7. Distribution and type of structural

materials and forms in a typical

advanced transport aircraft

TABLE 9. Airframe Applications for

Lithium Alloys

AluminumPRODUCT

FORM TYPICAL AIRFRAME APPLICATIONS

BILLET SIZE REQUIRED

(lb)

FORCING PIN JOINTS, FITTINGS, STRUTS, FRAMES

100-1,000 SOME 5,000

EXTRUSION STIFFENERS, SUPPORTS, RAILS 100-2,000

SHEET FUSELAGE AND EMPENNAGE SKINS AND CONTROL SURFACES

300-2,000

PLATE UPPER AND LOWER WING COVERS 2,000-20,000

TABLE 10. Transport Aircraft Weight Savings

Studies Use of Aluminum-Lithium

Alloys

WEtCHT SAVED

BASELINE ALLOY(S) ADVANCED Al-Li ALLOYS*

^ LOCKHEED STUDY L1011 AIRCRAFT; ELEMENTS OF WING, BODY, EMPENNAGE

7075-T76 DENSITY -10% STRENGTH 0, +20% FATIGUE AND +20% TOUGHNESS

MODULUS +20%

12.2 16,784

BOEING STUDY, 757-200 AIRCRAFT; ELEMENTS OF WING, EMPENNAGE

2024-T3

7075-T6 2224-T3

7150-T6

DENSITY -9 to -11% STRENGTH 0 to + 7% FATIGUE AND -8 to +20% TOUGHNESS MODULUS +13 to +23%

10.9 1,129

BOEING STUDY, 747-F AIRCRAFT; FLOOR BEAMS, SUPPORT STRUCTURE

7075-T6

7075-T73

ALLOY 1 ALLOY 2 DENSITY -10% -10% STRENGTH +10% 0% FATIGUE AND 0% 0% TOUGHNESS MODULUS +10% 0%

(1) 13.2 (1) 686

(2) 10.0 (2) 519

*PROPERTY DIFFERENCES FROM BASELINE ALLOY(s)

622
! The major thrust of weight saving studies with rapidly solidified PM low-density alloys has been

i for primary structural applications, as this area offers a large and straightforward weight savings

: benefit. All structural subgroups as well as the entire wing, body, and empennage sections have been

considered. This is possible because it is anticipated that PM aluminum-lithium alloys can be made

in all common structural forms required for large transport aircraft with the possible exception of

^ the largest plate products. These plates may require billet sizes in the 10,000 to 20,000 pound

, range (see Table 9). However, non-structural items or secondary structure can also provide an impor-

! tant increment of weight savings if constructed with low-density alloys, as a substantial increment

I of aircraft weight comes from small items such as clips and brackets, as well as large components

such as floors, cargo handling equipment, aerodynamic smoothing surfaces, and parts of the cabin interior. Alloy development studies to date have indicated that new aluminum-lithium alloys will achieve

the engineering property levels of 2024-T3, 7075-T73, and 7075-T6 with commensurate density reductions

of about 10 percent.

Methodology . The methodologies used for assessing structural efficiency (weight savings) in transport aircraft obtainable with RSP or IM aluminum-lithium alloys follow procedures comparable to that

described above.

Example Analyses . Three recent industry/government studies dealing with the use of aluminum-lithium

alloys on transport type aircraft are the Lockheed/DARPA [13], Lockheed/NASA [8J, and Boeing/AFWAL

[7,14] programs. The aircraft used in these studies were the L-1011, B-737, F-27, B-747, and the

B-757-200, which has received the most detailed analysis. The latter aircraft is shown in Fig. 8.

Elements of the results of three of these studies are shown in Table 10, where it will be noted

that weight savings of from 10 to over 13 percent are achieved in the affected structure. This is particularly noteworthy because of the diversity of baseline alloys, and the projected properties for

the aluminum-lithium alloys between the various studies. The validity of these weight saving estimates are further substantiated by the Lockheed/NASA study [8] and the most recent findings of the

Boeing/AFWAL study [7]. It should be emphasized that these calculations are primarily based on simple

alloy substitution and relatively small changes in design (i ,e. , airplane dimensions were not changed).

Structural weight savings of 10 to 13 percent have a major effect in payload and/or range improvements. For example, improvements calculated for the L-1011 aircraft are 21.3 percent (or 17,040

lbs.) in payload and 23.7 percent in range [13] (see Table 11). These impressive improvements appear

to be indicative of what could be expected for other transports and/or similar aircraft, and thus

underline the attractiveness of replacing currently used aluminum alloys with aluminum-lithium alloys

as they become available.

If a total aircraft redesign can be considered such as downsizing of the aircraft dimensions, additional structural weight can be saved than when direct substitution only is utilized. This is be- cause as the basic structure is made lighter, the landing gear, engines, fuel tanks, and so on, can also be made both smaller and significantly lighter. One of the iterative weight saving programs described in the literature is Lockheed's ASSET program [8,13] which projects for aluminum-lithium alloy

applications an additional 1.9 to 5.3 percent reduction in total structural weight by downsizing various transport aircraft (see Table 12). For transport aircraft of the 100-passenger size and larger,

total structural weight reductions are about 15 percent. For the L-1011 aircraft, a 15.9 percent re757-200

Fig. 8. The Boeing 757-200 Transport Aircraft

TABLE 11. The Effect of Aluminum-Lithium Alloys

on the Payload or Range of an L-1011

Transport

CHANCE CAUSED

BY APPLICATION OF Al-Li ALLOYS BASELINE

(nmi)

CHANCE CAUSED

BY APPLICATION OF AI-LI ALLOYS

• BASELINE ALLOY: 707S-T76

• AI-LI ALLOYS:

STRENCTH

FATICUE/ FRACTURE MODULUS

DENSITY

CHANCE FROM BASELINE

(1) (2) (3) (4)

•*20 SAME SAME SAME

SAME -*-20 SAME SAME

SAME SAME -^20 SAME

-10 -10 -10 -10

623
TABLE 12. The Effect of Advanced AluminumLithium Alloys on the Weight, Cost,

and Performance of Transport Aircraft

AIRCRAFT

REDUCTION IN STRUCTURAL WEIGHT REDUCTIONS IN FUEL AND COSTS

RESULTING FROM RESIZING AIRCRAFT NO REF. RESIZING

(%)

WITH

RESIZING

(%)

PRODUCTION

COSTS

$M (t)

FUEL

(%)

DIRECT

OPERA. COSTS

(%) OPERA. COSTS

$M (%)

L-1011* -10.2 -16.1 — -6.7% - - 13

L-1011* -10.6 -15.9 -1.70 (-2.5%) -6.4% -5.9% -1.68 (-3.94%) 8

B-737 -9.8 -14.6 -0.27 (-2.1%) -5.5% -3.7% -0.24 (-2.17%) 8

F-27 -7.8 -9.6 -0. 10 (-1.6%) •2.6% -1.9% -0.05 (-0.91%) 8

•DIFFERENT BASELINE AND ADVANCED ALLOYS WERE CONSIDERED IN THESE STUDIES.

TABLE 13. Air Vehicle Structure Components

Where Weight is Affected by DADTA

Properties

DADTA PROPERTY COMMERCIALPASSENGER AIRCRAFT

HIGH PERFORMANCE

MILITARY AIRCRAFT

FATIGUE ALL HIGHLY TENSION LOADED STRUCTURE

ALL HIGHLY TENSION LOADED STRUCTURE

PLANE STRESS FRACTURE TOUGHNESS.FUSELAGE AND WING LOWER SURFACE SKINS DOES NOT GENERALLY

IMPACT WEIGHT

PLANE STRAIN FRACTURE TOUGHNESS.

*«lc

MOST SUBSTRUCTURE

ELEMENTS NEED MINI- MUM K|^ PROPERTIES

MUST MEET MINIMUM K,^ REQUIREMENTS IN HIGHLY TENSION LOADED STRUCTURE CRACK GROWTH.

da/dN

SKIN MATERIAL AND SOME MAJOR SUBSTRUCTURE

ELEMENTS

MOST AREAS WITH MEDIUM TO

HIGH TENSION LOADINGS

CORROSION. STRESS CORROSIONNEW ALUMINUM ALLOYS AND PROTECTIVE SYSTEMS GENERALLY ADEQUATE

duction in structural weight translates into a 5.9 percent lower direct-operating cost, a 2.5 percent

lower production cost, and a 4 percent lower total operating cost [8J. For the B-737 aircraft, a 14.6

percent lower structural weight translates into a 3.7 percent lower direct operating cost, a 2.1 percent lower production cost, and a 2.2 percent lower total operating cost [8].

CONSIDERATIONS FOR DURABILITY AND DAMAGE TOLERANCE ASSESSMENT (DADTA)

Evaluation of durability (fatigue) and damage tolerance characteristics, of a structure requires

materials property data on fatigue, fatigue crack growth, fracture toughness, and subcritical crack

growth in the applicable corrosive environments. Various types of test data are required to properly

evaluate these properties. The test types required include constant-amplitude fatigue, spectrum

fatigue, fatigue crack growth, fracture toughness (Kj^ and Kq), and stress corrosion cracking (Kigcc^

The properties obtained from these tests are required to properly evaluate potential weight savings

for aircraft structural applications.

Some of the aircraft structural components sized by the DADTA properties of materials are given

in Table 13. The areas of the structure where the tensile stresses are high for the normal operating

service loads are particularly affected by the DADTA material properties. Based on these data, the

ultimate design allowable tension stresses are reduced as required to provide adequate fatigue life

and damage tolerance capability. The structural areas generally affected include wing lower surfaces

and the fuselage shell structure. Since the maximum allowable stresses in these skin structures are limited, the allowable stresses in substructure elements attached to these skins are also limited

even though they may not be DADTA critical. Therefore, a considerable portion of the structural

weight can be affected by the DADTA material properties.

For lower wing surfaces of passenger aircraft, the plane stress (K^) fracture toughness and crack

growth properties have the largest effect on the damage tolerance characteristics of the structure.

The critical damage size (crack length) is of the order of the wing stiffener spacing or larger and

the stiffeners are intended to provide some crack arrest capability.. For high performance fighter

type aircraft the critical crack sizes in the wing surfaces are smal ler and flight safety is provided

by inspections at intervals based on the crack growth properties of the material. Fatigue is also

important for both types of aircraft: long service life requirements for passenger aircraft and se- vere high amplitude fatigue loadings for high performance (fighter) aircraft.

For the fuselage shell structure of passenger aircraft, the fatigue, plane stress (K^) fracture

toughness, and crack growth properties all affect the durability and damage tolerance characteristics

of the structure. The critical crack sizes in the skin are of the order of 20 to 40 inches. Fail

safe design is provided by the frames, stringers, and in some cases reinforcements attached to the

skin. The limitations in stress are based on both the fatigue and residual strength requirements.

Fuselage structure for high performance military aircraft are sized to meet crack growth and fatigue

requirements. Therefore, the plane stress fracture toughness properties generally do not have a weight impact in the fuselage structure of high performance military aircraft.

Protection from corrosion must be provided for aluminum alloys on aircraft structure. The corrosion protection systems in current use include anodizing or cladding of the basic material. These

systems, including the use of sealants and paints, have generally provided adequate corrosion protection for most areas of the airframe structure of aircraft in current use. Cladding and anodizing,

however, have a detrimental effect on the fatigue properties of the aerospace aluminum alloys which

must be accounted for in establishing the allowable stresses for fatigue. The effect of corrosion

624
protection systems on the fatigue properties of the new advanced aluminum alloys will require careful

evaluation.

TECHNICAL ISSUES RELATED TO SUCCESSFUL ALLOY DEVELOPMENT AND APPLICATION

There are a number of technical factors or issues which are related to successful development and

application of the aluminum-lithium alloys. Some of these issues are peculiar to either the RSP powder metallurgy or ingot metallurgy products; many of the issues pertain to both products. It is hoped

that the discussion which follows will increase awareness of the factors affecting eventual utilization and thereby contribute to the orderly, efficient establishment of this important new family of

aluminum alloys.

Durability and Damage Tolerance Assessment . The durability and damage tolerance assessment of new aluminum-lithium alloys requires that fatigue, fatigue crack growth, fracture toughness, and other re- lated properties be established. Some results have recently been reported for the ZOZO aluminum alloy

developed over 25 years ago by Alcoa L15J. The lithium content of this alloy is low, only 1.1 weight

percent, well below the 2 to 3 weight percent in the most promising new alloys. By mid-1983, some fracture and fatigue data for the new alloys will be published by various aerospace contractors to AFWAL, DARPA and NASC.

In addition to the DADTA related data, verification of the applicability of conventional fatigue

and fracture mechanics methodology will be required. This will include accuracy of predicting actual

crack growth rates in structural elements from laboratory test data of the new alloys. The effects

of variables such as load ratio, overloads, and chemical environment on fatigue crack growth behavior

will also need to be assessed. Inspection techniques used for detecting initial crack sizes in airframe components will need to be demonstrated for the new alloys, particularly the PM alloys because

of their fine grain size. Crack arrest characteristics of sheet need to be demonstrated, including

the effectiveness of stringers, straps, and the frames used to support the fuselage skin. The behavior of mechanically fastened joints typically used in airframe construction will need to be evaluated.

Finally, the suitability of current corrosion protection systems including paints, coatings, and sur- face conversion layers for the new aluminum-lithium alloys will need to be assessed.

Availability of RSP Facilities . Atomization of aluminum-lithium alloys to produce fine powder will

require inert gas protection. These facilities are at present quite limited in both capacity and a- vailability [16] and undoubtedly new capital facilities will have to be developed to support the potentially large demand. New production facilities for melting and casting will also be required.

Ingot melting of aluminum-lithium alloys requires special protection of the liquid metal to avoid ex- cessive hydrogen pickup and also special care in the casting procedure to avoid explosions due to liquid metal/water reactions.

Control of Particulate Quality . Some of the facilities used to date in the development of fine

atomized aluminum-lithium powders are used for a variety of other metal alloy systems. Consequently,

in some cases foreign particles 30 to 50 pm or larger in size have been found in each batch of aluminum-lithium powders. Nickel-base contamination particles in the aluminum alloy powder lot, for ex- ample, promote the formation of a brittle intermetal 1 ic shell of nickel aluminides around the particle.

This feature can then act as a local defect in the final consolidated product and lower the ductility

[4]. Other particulate^qual ity factors of concern include large variations in the oxide film thickness (100 to over 1000 A in the splat particulate)[4,5J, large variations in fineness of solidification structure (particularly in splat)[4,6], and hydrogen contamination from water vapor in the inert

shielding or atomizing gas and possibly from other sources.

Consolidation of Large Powder Metallurgy Billets . Although the requirement is not unique to PM aluminum-lithium alloys, some of the component parts in aircraft airframes will need large billets for mill

processing. For most extrusions, forgings, and sheet, billets weighing less than 1000 lb will satisfy

a large part of the requirements. Plates for upper and lower wing covers in transport and bomber aircraft, however, will require rolling slabs weighing up to 20,000 lb. It is unlikely that powder metallurgy aluminum billets of that maximum size will be made in the forseeable future, at least without

an enormous investment.

Cost. The cost of both PM and IM aluminum-lithium alloys is an important issue. The raw material

cost of an aluminum alloy containing 3 weight percent lithium is twice that of the alloy without lithium. This is due to the high current cost of 1 ithium--$30-$60 per pound depending on purity level.

Alcoa is hopeful that the IM aluminum-lithium alloys they are now developing will not cost more than

3 times the commercial aerospace aluminum alloys [17]. Note that Alcoa sold the 2020 aluminum-lithium

alloy in plate form about 25 years ago for $2.00 to $2.50 per lb, at which time 7075 alloy plate sold

for about $.75 per lb [17]. The PM aluminum-lithium alloys will likely cost somewhat more than the

IM aluminum-lithium alloys because of the powder metallurgy processing.

625
Another perspective of cost is that of the user. In general, airplane manufacturers value a pound of weight saved in an airframe at $200 to $600. As described in previous sections, 9 lb of

aluminum-lithium alloy would typically be required to replace 10 lb of current aerospace aluminum alloy in the airframe in order to save 1 lb of weight. At the above value of a pound saved the substituion would be valued at least $200 and possibly up to $600. Thus, the substitution of the new aluminum-lithium alloys would be cost effective if they cost $11 to $33 per pound, assuming a buy-to-fly

ratio of 2. The above costs must be appropriately factored to account for the user's particular efficiency of material utilization (for example, specific "buy-to-fly ratios").

A third perspective of cost is the investment capital required for the producers to develop large

scale facilities and the associated production technology. One major aluminum producer has estimated

the usage of aerospace grade aluminum alloys between the years 1990 and 2000 at 350-500 million lb per

year [17]. Of this amount perhaps half will be alloys already in current production; the other half

perhaps new materials such as PM aluminum high strength and high temperature alloys, aluminum-lithium

alloys, and some materials replacing aluminum such as composites. This company suggests that the suc- cessful development of aluminum-lithium alloys which are cost effective may be expected to result in

the use of between 30 and 100 million lb per year [17]. This anticipated usage is sufficient to stimU'

late significant private sector investment by prospective producers.

Ingot Versus Powder . Since both ingot metallurgy and RSP powder metallurgy approaches can be used to develop new aluminum-lithium alloys, some comments regarding known and prospective advantages and disadvantages of each are appropriate. Two of the IM product advantages expected in this case are lower

cost and larger billet size, within some upper lithium content. Advantages of the RSP PM product include a higher lithium content potential. For example, one complex aluminum-lithium alloy designed to exhibit a yield strength of at least 65 ksi tended to crack extensively when DC cast in 8 inch dia.

billet size if the lithium content was in excess of 2. 7-2.8 wt percent. Such limitation in lithium

content due to cracking does not apply to the PM process. Another advantage of the PM process is that

higher levels of transition elements can be added for special reasons, whereas in the IM alloys, the same elements must be restricted to low levels to avoid formation of coarse, embrittling intermetallics. Such advantage for the RSP approach makes it feasible to use the technique to develop a low

density, high temperature aluminum alloy by combining the attributes of lithium addition for decreased

density and selective transition element addition for increased modulus and to form thermally stable

dispersoids.

Ingot metallurgy disadvantages may include increased composition heterogeneity, coarser grain

size and coarser solidification structure. These may in turn affect the consistency and reliability

of the mechanical properties.

In addition to the higher cost of the RSP powder metallurgy product, current suppliers of aluminum-lithium powder are generally not historic producers of aluminum mill products and thus are less

familiar with the requirements of producing and distributing aluminum aerospace quality material.

Finally, another disadvantage is the increased effort required to establish and maintain consolidation

and processing procedures to achieve consistent and reliable interparticle bonding and disruption of

the preexisting oxide films covering the surface of every particle.

Obtaining Designer Confidence . The aerospace structural designer's confidence will be gained only

after the following steps are accomplished.

• Verify all the property goals,

• Establish the consistency of the product and associated properties,

• Establish scaleup in product size to meet application requirements,

• Verify performance in prototype flight hardware,

I Establish supply base and acceptable cost.

These accomplishments will require close coordination between suppliers, users, and sponsors. Major

additional investments will be required, particularly to establish the production base and to verify

prototype hardware performance.

626
CONCLUSIONS

1. The application of aluminum-lithium alloys to aircraft airframe structures will result in significant (10 to 20 percent) weight savings. These weight savings are primarily due to the decreased

density of the alloys.

2. Both RSP powder metallurgy and ingot metallurgy aluminum-lithium alloys have identifiable contributions, with each process having distinct advantages and disadvantages.

3. A significant interest in aluminum-lithium alloys has been established by both major users and

suppliers. This is based on the promising cost effectiveness of the alloys as the result of

weight savings, and the significant future market for the alloys.

4. A major effort including significant additional investment is still required to gain the aerospace structural designer's acceptance of the new aluminum-lithium alloys.

ACKNOWLEDGMENT

The authors thank Drs. T. M. F. Ronald, AFWAL/MLLS, WPAFB, OFI, L. A. Jacobson and E. C. van Reuth,

Materials Science Division, DARPA, Arl ington, VA, for their encouragement and support for an important

part of the work on which this paper is based: AFWAL Contracts F3361 5-78-C-5203 (DARPA Order 3417)

and F33615-81-C-5053.

REFERENCES

[1] J. C. Ekvall, J. E. Rhodes and G. G. Wald, "Methodology for Evaluating Weight Savings from Basic

Material Properties," Design of Fatigue and Fracture Resistant Structures , ASTM Special Tech.

Publ. 761, 1982, pp 328-341.

[2] Aluminum-Lithium Alloys , Proc. First International Aluminum-Lithium Conference, T. H. Sanders, Jr.,

and E. A. Starke, Jr., eds., AIME, Warrendale, PA, 1981.

[3] Second International Aluminum-Lithium Conference, T. H. Sanders, Jr., E. A. Starke, Jr., and

R. Schmidt, organizers. Scheduled for April 12-14, 1983, Monterey, CA.

[4] R. E. Lewis, I. G. Palmer, D. D. Crooks, H. G. Paris, F. R. Billman, L. Otto, E. A. Starke, Jr.,

T. H. Sanders, Jr., A. Gysler, R. Crooks, E. E. Underwood, F. S. Lin, G. Wald, E. Himmel,

E. Rhodes and R. Simenz, "Development of Advanced Aluminum Alloys from Rapidly Solidified Powders

for Aerospace Structural Application," Interim Technical Report for Period September 1978 - March 1981, Technical Report No. AFML-TR-81 -XX , March 1981, AFWAL Contract F3361 5-78-C-5203,

DARPA Order 3417, to be issued.

[5] I. G. Palmer, R. E. Lewis, and D. D. Crooks, "The Design and Mechanical Properties of Rapidly

Solidified Al-Li-X Alloys," Aluminum-Lithium Alloys , T. H. Sanders, Jr. and E. A. Starke, Jr.,

eds., AIME, Warrendale, PA, 1981, pp 241-262.

[6] I. G. Palmer, R. E. Lewis and D. D. Crooks, "Development of Al-Li-X Alloys Using Rapidly Solidified Powders," Proc. 1982 AIME Conf. on Aluminum Powder Metallurgy, Dallas, Texas, Feb. 1982,

AIME, Warrendale, PA, 1983.

[7] G. H. Narayanan, W. E. Quist, B. L. Wilson and A. L. Wingert, "Low Density Aluminum Alloy Development," Interim Technical Report for Period 4 Jan. 1982 to 1 July 1982, AFWAL Contract No. F33615-81-C-5053, Aug. 1982.

[8] I. F. Sakata, "Systems Study of Transport Aircraft Incorporating Advanced Aluminum Alloys - Final Report," NASA CR-165820, WASA Contract NASl-16434, Jan. 1982.

[9] W. C. Pohlmann, Private Communication, Sept. 1982.

[10] C. Baker, Private Communication, Dec. 1982.

[11] R. Jordan, Private Communication, Oct. 1982.

[12] R. E. Lewis, et al, "Development of Advanced Aluminum Alloys from Rapidly Solidified Powders for

Aerospace Structural Applications." Lockheed Palo Alto Research Laboratory Interim Technical

Report for period March 1980 to December 1980, Air Force Contract F3361 5-78-C-5203, DARPA Order

3417, March 1980.

627
[13] R. E. Lewis, D. Webster, I. G. Palmer, "A Feasibility Study for Development of Structural Aluminum Alloys from Rapidly Solidified Powders for Aerospace Structural Applications," Lockheed

Palo Alto Research Laboratory, Interim Technical Report for Period June 1977-May 1978; Air Force

Number AFML-TR-78-1 02, DARPA Order No. 3417, July 1978.

[14] Boeing Document D6-49955-1, "Low Density Aluminum Alloy Development," Boeing Military Airplane

Company, Page 4-21, May 1981.

[15] R. C. Malcolm, A. K. Vasudevan, R. J. Bucci, P. E. Bretz, "Evaluation of the Engineering Properties of a Commercially Produced Aluminum Alloy 2020-T651 Plate," Alcoa Report for Amendment to NASC Contract NOOOl 9-79-C-0258, Dec. 18, 1981.

[16] National Materials Advisory Board Committee on Rapid Solidification Processing, "Rapid Solidifi

cation Processing: Status and Facilities," Report NMAB-401 , National Academy of Sciences,

Nov. 1982.

[17] E. Spuhler, Private Communication, Nov. 1982.

628
APPLICATIONS OF RAPIDLY SOLIDIFIED ALUMINUM ALLOYS

AT ELEVATED TEMPERATURES

C. M. Adam*, J. W. Simon and S. Langenbeck"*"

Pratt & Whitney Aircraft/Government Products Division

West Palm Beach, Florida

*"Lockheed California Company, Sunnyvale, California

ABSTRACT

During the past five years rapid solidification technology has been successfully

applied to a number of novel aluminum alloys. By using conventional powder metallurgy

techniques these alloys have been fabricated into components which appear able to compete

on a specific strength basis with Ti-6A1-4V components at temperatures up to 450°F (500K)

In this paper payoff studies for both airframe and gas turbine engine applications are identified and discussed.

INTRODUCTION

The use of the PWA/GPD rapid solidification technology to produce new aluminum engineering alloys

has been recently reviewed in three detailed publications (References 1-3). These alloys rely on a relatively thermally stable 300-800 a dispersion of metastable intermetallic compounds of the modified

Al^Fe type which are produced by a series of thermomechanical processing steps during consolidation of

rapidly solidified aluminum alloy powder of appropriate composition. As a result of the relatively

slow thermal differsivity of the major alloying elements in the new alloys of Al-8Fe-2Mo and Al-8Fe-4Ce

compositions these alloys have an extended operating temperature range beyond conventional aluminum

alloys containing copper, for example alloy 2219.

It is appropriate therefore to consider whether such alloys will be viable competition on a comparative (specific) strength basis for titanium alloys, and what payoff in weight savings might

result from use of the properties of these new alloys in airframe and engine applications. Early

studies on gas turbine engine components in the fan and low temperature compressor, and on airframe

components, indicated weight-saving payoffs ranging from 5 to 25% for the properties generated during

1980 and 1981. A second iteration of these payoff studies during 1982 (Reference 4) has identified

property goals which should lead to greater weight savings if optimal strength and toughness levels

can be met. In the following sections of this paper it will be clear that significantly greater market

volume opportunities for rapidly solidified aluminum alloys lie in Mach 3+ airframe applications than

in gas turbine engine applications, however the impact on the delicate performance/weight/cost balance

for future military gas turbine engines may be such to allow superior advanced engine designs which are

not feasible with conventional titanium alloys.

GAS TURBINE ENGINE APPLICATION STUDIES

Modern gas turbine engines have a complicated variety of component temperature/stress regimes.

Extensive metallurgical research and development over the last three decades, coupled with increasingly

sophisticated design techniques, has improved the performance of military engines from a three-to-one

thrust-to-weight ratio to eight to one (figure 1). In seeking further improvement in these thrust-toweight ratios, designers have sought to reduce engine weight with no concomitant loss of thrust. This

can be achieved by development of lighter alloys with higher temperature capabilities. For rotating

components, specific strength and specific stiffness are both important materials properties. Specific

strength determines the maximum rotational speed for which rotating hardware can be designed; specific

stiffness determines the vibrational frequency response of a rotating component, and the high-cycle

fatigue damage that will be accumulated by aeroelastically induced vibration of that component. Static

structures are invariably limited by Euler compressive buckling stresses which are a function of the

*Now with Allied Corporation, Corporate Technology, Morristown, NJ.

629
elastic modulus of the alloy selected. Development of higher specific strength and specific stiffness

aluminum alloys will therefore allow engine designers to reduce engine weight in those gas-turbine

areas where aluminum alloys have traditionally been used. Table 1 outlines such applications by major

engine manufacturers.

Figure 1. Military Engine Performance from 1940 to 1980 I

I

Engine induced temperature effects in nacelle and structural airframe components could be > alleviated by development of higher temperature capability, and higher specific strength alloys. In general terms, several hundred pounds of aluminum alloys will be used in future advanced transport

engines, and therefore up to one thousand pounds of engine related weight could conceivably be affected

by alloys developed during programs related to elevated temperature strength aluminum alloys.

TABLE 1. Current Use of Aluminum Alloys in Medium-Performance Gas-Turbine Engines

Engine Aircraft Current Aluminum Alloy Application

PWA JT8D-15

PWA JT8D-209

PWA JT9D-7A/7F

PWA JT9D-59A/70A

PW 2037-136

Boeing 727-200

MCD DC

Boeing 747

MCD DC- 10

Boeing 757-200

Second-stage fan vane, outer fan duct,

third-stage compressor

Compressor entry guide vane, fan case, fan duct, second- and third-stage

compressor vanes

Aft fan case

Aft fan case

First-stage fan vane

]

I

•i

1

I GE CFM-56

GE CFM-50

GE FlOl-X

R-R RB-211

Boeing KC-135

MCD DC-10; A300

Boeing 747 ; A300

First-stage fan vane, fan case |

First-stage fan vane, fan inlet case j

First-stage fan vane

j

1 First-, second-, third-, and fourth-

| stage intermediate compressor vanes, intermediate case I'

The replacement of aluminum alloys by titanium in the early 1960's resulted from higher ij temperature environments encountered in high-performance gas-turbine engines for military applications ,j and from lack of strong aluminum alloys for use in the temperature range of 250 to 650 F, which were

competitive with titanium alloys (References 2 and 3) on a strength-to-weight ratio basis. Figure 2 i shows an engine cross section intended for development from 1985 to 1990 which was used as a baseline | application engine for this payoff study. This engine is one of a generic series of demonstrator |
engines, designated JT69 by Pratt & Whitney Aircraft, and has provided a series of design/analytical

studies with a full engine test demonstration scheduled for late 1984. A fundamental difference

between the early 1960 's JT3D engine and the JT69 engine Is In the thrust produced; early JT3D engines

produced 19,000 lb static thrust with an engine weight of 4300 lb, while engines Incorporating the

advanced JT69 technology should produce some 28,000 lb thrust from an engine weight of less than 3000113^

These performance differences arise from a substantial reduction In the number of engine parts, which

combined with more efficient materials and computer-optimized design, yield thrust-to-welght ratios

exceeding ten. A subsidiary benefit of lighter and less mechanically complicated engines comes from

decreased manufacturing costs; current buy/fly ratios of six to ten are common for most engine components. Critical and strategic element considerations frequently Increase the costs of engine alloys

to prohibitive levels, and It Is anticipated that a significant payoff will accrue from lower material

costs and fabrication costs If advanced aluminum alloys produced by rapid solidification are utilized,

rather than currently available titanium alloys.

Development of advanced powder metallurgy aluminum alloys with Increased specific stiffness for

the applications outlined above will require specific alloy properties substantially better than 2219

properties up to 350 F, In order to be competitive with alternate structural titanium alloys. Advanced

aluminum alloys will find applications In gas turbine engines and Mach 3 airframes If they are structurally competitive with titanium on a specific strength basis at temperatures up to 650 F. In

the following sections of this payoff analysis both static and rotating engine components will be

examined as well as airframe applications for- a super-cruiser version of a Mach 3 + airframe to determine weight-savings from application of elevated temperature aluminum alloys.

Figure 2. Gas Turbine Engine Cross-Section for Applications Study

In subsonic engines potential applications for aluminum alloys are more numerous than In high

performance fighter engines which operate up to Mach 2.5. The applications listed In Table 1 are those where prolonged temperatures rarely exceed 200 F In engines with modest thrust-to-welght ratios.

In higher performance engines, the fan and compressor stages produce higher aerodynamic pressures,

and adiabatic heating In these stages raises component temperatures to some 450 F. In military

fighter engines there Is a further effect of ram airspeed at high Mach numbers which results In further aerodynamic adiabatic heating of engine components by some 250 F. In general^ the steadystate temperature Increment per stage In the fan and compressor section Is 100 to 150 F, and these

steady state operating temperatures are Increased for short operational cycle times by the Imposed

ram temperature frequency. In military transport engines, such as that envisaged for the Long Range

Cruise Aircraft (LRCA) , airspeed Is limited to subsonic velocities and engine components experience

lower temperatures. It Is therefore anticipated that a Mach 2.5 engine will provide a realistic but

somewhat conservative data base for • Investigating the applicability of elevated temperature aluminum

alloys, and that such data can be readily transferred to subsonic engines studes where applicability

Is anticipated to be somewhat greater.

In this application study the JT69 engine Is used to compare the performance of both rotating and

static aluminum hardware, using typical mechanical properties of the best alloy developed under a previous Pratt & Whitney Aircraft - Air Force Wright Aeronautical Laboratories Contract F33615-76-C5136 "Application of Rapidly Solidified Alloys." The properties of this alloy, Al-8Fe-2Mo, are listed

In Table 2, and represent a starting point for further alloy development during the course of this

program. These mechanical properties are relatively modest, and fall short of specific strength

631
equivalency with Ti-6A1-4V.

Table 2. Typical Mechanical Properties of Al-8Fe-2Mo Used for

Investigating "Base-Level" Payoff for Elevated

Temperature Aluminum Alloys

75°F 300°F 450°F 650°F

Ultimate Tensile Strength

Yield Strength

Elongation, %

E, Msi

0.2% Creep Stress (450°F)

1% Creep Stress (450°F)

71.3 ksi 61.8 ksi 49.1 ksi 32.2 ksi

68.4 50.4 41.5 29.0

6.7 7 8 10

12.2 11.5 11.0 10.0

- 25 ksi, 100 hr

35 ksi, 400 hr; 30 ksi. 1000 hr

450°F/36 ksi/1000 hr without failure - 3% strain

- 20 ksi -

Creep Rupture

Fatigue, 10^ Cycles

R= -1.0, = 1

In the following analyses of both fan and low-pressure compressor blades and vanes, a temperature!

limitation of 450 F has been arbitrarily imposed in determining weight savings which would result fromj

a direct replacement of T1-6A1-4V with Al-8Fe-2Mo. Further temperature extensions to 650 F will

require substantial alloy development, some of which are currently being developed during the AFWAL

program "Elevated Temperature Aluminum Alloys" by Lockheed* Pratt & Whitney and Alcoa.

Fan and Low -Pressure Compressor Blade Analysis

In the baseline engine chosen for this application study, the 2nd- and 3rd-stage fan blades

fulfill some of the compression requirements of the low-pressure compressor of engines or more modest

thrust-to-weight ratios. Table 3 shows the maximum blade root stresses for Ti-6A1-4V alloys at these

environmental temperatures.

Table 3. The 2nd- and 3rd Stage T1-6A1-4V Blade Temperatures and Stresses

Used as Baseline Data for Calculating Aluminum Blade Alloy Weight

Savings

i

Blade

Second Stage

Third Stage

Temperature

300°F

400°F

Maximum Radial

Stress

80.5 ksi

72.4 ksi

Peak-to-Peak

Vibratory Stress

Capability

29 ksi

27 ksi

Life*

3500 Cycles

7000 Cycles

*Equivalent Type 1 cycles, which roughly approximates to engine start-up, a series of engine

maneuvering accelerations and decelerations, and engine shutdown.

The radial stresses in Table 3 result from centrifugal loadings and are dependent on the density,

P, of the blade alloy and the blade geometry. If the blade has mass, m, length, 1, and root area A^,

then the blade root stress at an engine speed, co, and radius of gyration, r, is given by

2 Blade radial force, F = mco r ; Blade radial stress, = co^plr

^ Blade specific stress, CJ/p= (o ir

It is important therefore to compare the specific strength properties of an aluminum candidate .j blade alloy with T1-6A1-4V in defining weight savings. i

Vibratory blade stresses arise from two sources, both of which are considered in Table 5. [ Fluctuation in gas-flow aerodynamic loadings of approximately 1 Hz frequency lead to high-cycle

632
fatigue damage, and are independent of the blade material of construction. These stresses arise

principally from the interaction of the rotating blades and static vanes and have been shown to be up

to 15 ksi in amplitude (30 ksi peak-to-peak) in FlOO engine studies. In determining applicability of

advanced aluminum alloys, experience should therefore require a minimum high-cycle fatigue capability

of 15 ksi amplitude at the operational steady stress level. The second vibratory stresses considered

in Table 3 are basically low-cycle fatigue stresses in nature arising f^om engine run-up and shutdown

These low-cycle fatigue stresses have a lifetime requirement between 10 and 10“^ cycles, and have a maximum value equal to the blade specific stress. Comparison of Ti-6A1-4V and an advanced aluminum

alloy therefore requires two sets of data. The first data set consists of a modified Goodman diagram

(figure 3) showing the relationship between actual high-cycle fatigue endurance (10^ cycles) stress

amplitude, and blade specific steady-stress. The second data set consists of comparative LCF values.

Figure 3. Modified Goodman Diagram for T1-6A1-4V and Al-8Fe-2Mo Fan Blade Alloys

Table 4 shows calculated JT69 steady-state levels for typical engine fan blades, and comparative

Al-8Fe-2Mo stresses based on the experimental fatigue data presented in Table 2. This data shows that

the present alloy properties do not possess sufficient high-cycle fatigue stress capability to meet

the arbitrary P&WA/GPD design stress requirement of 15 ksi single amplitude oscillatory stress at the

operational steady-stress level of 47.5 ksi. Although specific steady-strength equivalency with

Ti-6A1-4V at 400 F has been demonstrated (Reference 5), the required high-cycle fatigue stress amplitude endurance limit of »30 ksi has not been demonstrated and must therefore represent a desirable

alloy property goal for the program. In achieving this material fatigue goal a direct substitution

of Al-8Fe-2Mo for Ti-6A1-4V would be possible, leading to a hypothetical weight savings of 35%. An

alternative design approach would be to adjust the steady stresses for Al-8Fe-2Mo to % 32 ksi by

redesigning the blade. The allowable oscillatory stress of 15 ksi would then be available, and this

would result in an approximate weight saving of 7%. This hypothetical blade redesign may then result

in unacceptable aerodynamic performance for this stage of the engine, and further detailed analyses

of the interaction of aerodynamic losses and hypothetical weight savings would be required. The

modified Goodman diagrams of figure 3 thus represent the extremes of weight saving; approximately 7%

for the lower fatigue curve and approximately 35% for the upper proposed curve. It is anticipated

from this analysis that fatigue curves falling between these two extremes will result in weight

savings given by simple interpolation, i.e., a high-cycle fatigue endurance limit of 25 ksi single

amplitude would lead to s; 20% blade weight savings. The low-cycle fatigue performance of T1-6A1-4V

and Al-8Fe-2Mo should be identical when an endurance limit stress of ^28 ksi single amplitude for

Al-8Fe-2Mo has been achieved. At current strength levels, Al-8Fe-2Mo has approximately 75% of the

specific low-cycle fatigue strength of T1-6A1-4V, which is adequate for these blade applications.

Fan and Low-Pressure Compressor Vane Analysis

Table 5 shows a comparison of vane stress indicating that the steady aerodynamic loading on T1-6A1-4V vanes is 39.2 ksi. To provide adequate high-cycle fatigue (10 cycles) vibratory capability,

the existing properties of Al-8Fe-2Mo are inadequate, and the vane cross-sectional area would need to be increased 22% to reduce the steady-stresses to a value where a 15 ksi single amplitude oxcillatory

stress is available. As was discussed for the blade analysis, this cross section increase may incur

aerodynamic penalties; however, improved processing of Al-8Fe-2Mo should produce adequate vibratory

capability at increased steady stress levels. The weight savings under the Table 5 stress conditions

633
Table 4. Comparison of T1-6A1-4V and Al-8Fe-2Mo Fan Blade Stresses at 400°F for 2nd- and 3rd-Stage JT69 Demonstrator Engine Blades

Critical Parameter Ti-6A1-4V

1st Generation

Alloy

2nd Generation

Alloy

Centrifugal loading 72.4 ksi 47.5 ksi 47.5 ksi

Steady Stress 94% yield stress 114% yield stress* 95% yield stress

Vibratory capability

10 cycles

14 ksi amplitude 6 ksi amplitude 15 ksi amplitude

Airfoil dimensions A.l A.l A.l

1 1 1 Weight savings - 7% 35%

*Yield stress is unacceptable for current blade genometry. By redesigning the blade

and lowering the steady stress to 32 ksi, a 7% weight saving is hypothetically possible.

vary from 24 to 35% with intermediate weight savings possible at intermediate modiifed Goodman diagram

stress levels. A further design consideration for vanes involves calculation of the vane deflection

during aerodynamic transient pressure fluctuations of tw^ce the stead-state aerodynamic pressure.

Such surge deflections, 6, are given by W1 /KEI where Wl"^ is the vane loading, K is a dimensionless

paramter depending on the method of vane support, E is the alloy elastic modulus, and I is the

moment of inertia. For the two allgys under comparison, Al-8Fe-2Mo and Ti-6A1-4V, the respective

elastic moduli are 11.5 and 15 X 10 psi, respectively, at 400 F, and equivalent deflections would

therefore require an increase in the Al-8Fe-2Mo vane thickness of approximately 9.3%. A set of alloy

development goals for vanes thus requires two improvements: viz an increase in elastic modulus to provide enhanced surge deflection capability, and an increased high-cycle fatigue endurance limit to provide adequate fatigue strength.

Table 5. Comparison of Ti-6A1-4V and Al-8Fe-2Mo Vane Stresses at 400°F

for 2nd- and 3rd-Stage JT69 Demonstrator Engine Vanes

1st Generation 2nd Generation

Critical Parameter T1-6A1-4V Alloy Alloy

Air loading, steady stress 39.2 ksi 39.2 ksi 39.2 ksi

Vibratory capability

single amplitude

27 ksi 12 ksi* 23 ksi

Airfoil dimensions, txb nominal, tb 1.22 tb nominal, tb

Weight reduction nominal 24% 35%

Equivalent surge

deflection

nominal 10% increase in t/b nominal

Weight reduction for

equivalent surge

nominal 25% 35%

deflection

*Direct material substitution will not allow 15 ksi single amplitude HCF vibratory

capability. To increase this capability the steady stress has to be reduced to- 32 ksi,

requiring a 22% increase in vane cross-sectional area.

The foregoing analyses show that specific strength equivalency with Ti-6A1-4V is an important

requirement for achieving direct replacement by advanced aluminum alloys. These analyses show that a high-cycle fatigue single amplitude endurance limit of =5;30 ksi (figure 3) will be required for blades,

and a limit of 25 ksi will be required for vanes. Vane alloys should also have an improved^elastic

modulus to allow for deflection parity with Ti-6A1-4V, and an elastic modulus goal of 13 X 10 psi at 400 F could be artibrarily set for vane applications.

634
AIRFRAME APPLICATION STUDIES

High strength aluminum alloys have been the primary structural materials in high performance

fighter aircraft for many years. The relatively low material and fabrication costs combined with high

strength-to-weight ratios have made the aluminum alloys most cost effective in airframe structural

applications. However, degradation of strength and stiffness of aluminum alloys at elevated temperatures has been a major factor in limiting fighter speeds to the Mach 2 level. The leveling off of

fighter maximum design speeds is shown in Figure 4 (Reference 5). The lone exception, which is not

shown on the figure, is the YF-12 - which has a predominantly titanium structure.

Figure 4 also shows the temperatures due to aerodynamic heating over most of the skin surface of an aircraft designed for the noted maximum speeds. From this figure it can be seen that if aluminum

alloys can be developed with improved high temperature strength and stiffness, the maximum operating

speeds of fighters with aluminum alloy structures can be extended to beyond Mach 3. Such a development would be especially advantageous for application to a future long range supersonic-type fighter

designed for a cruise speed in the range of Mach 2.7 to 2.8 with overspeed capability of Mach 3 plus.

In addition, there appear to be other significant applications for elevated temperature high strength

aluminum alloys in lower Mach Number aircraft. Such structures as engine inlet ducts, nacelles, and

other engine heated components will benefit from aluminum alloys that maintain strength and stiffness

at elevated temperatures.

A significant characteristic of a supersonic aircraft is that, for the major portion (75-90%) of

the airframe structure, the subsonic and transonic loades ate greater than the supersonic loads.

Consequently, much of the structure is designed by room temperature properties. Design loads at speeds at which elevated structural temperatures become a significant factor may be 40 percent or less

than the design loads at the lower temperatures. The actual reduction depends on a number of factors;

including location (such as wing or fuselage), design maneuver load factor, gross weight reduction at high speed, and mission profile. Similarly the continuous load levels contributing to creep at elevated temperature are normally less than 25% of room temperature loads.

Thus, for advanced powder metallurgy (P/M) aluminum alloys to compete with titanium in airframe

structures subject to elevated temperatures, the room temperatCire properties must be competitive on a density basis and must be unaffected by elevated temperature exposure over the life of the aircraft.

In addition the alloys should possess elevated temperature properties in excess of 40% of room

temperature properties and creep strength greater than 25% of room temperature strength. In contrast,

parts designed by elevated temperature loads require elevated temperature properties equivalent to titanium while room temperature properties are not as critical.

Analyses have been conducted to Identify the effects on structural weight of incorporating

elevated temperature resistant advanced aluminum P/M alloys into the airframes of high performance

fighter aircraft and selected components of military transport aircraft. In order to explore the

temperature range of interest of 394 K (250 F) to 616 K (650 F) and to include a varied range of

structural components, three baseline structures have been analyzed; viz. Mach 2.75-3.0 Super635
cruiser Airframe; Mach 2.2 -2.5 Advanced Tactical Fighter Aft Fuselage Structure; and Military

Transport Components.

Supercruiser Airframe Weight Payoff Analysis

The Air Force has for some years been interested in the design of military aircraft with long

range supersonic cruise capability. Two examples of activities in this area are: (1) the design

conference titled "Technology for Supersonic Cruise Military Aircraft" held by the Air Force Flight

Dynamics Laboratory in February, 1976 (Reference 7); and (2) the study program performed by Boeing

for the Air Force Flight Dynamics Laboratory in 1978 (Reference 8). Industry studies were initiated

recently in response to an Air Force Request for Information. As stated in the RFI the intent is "to provide interface and dialogue with industry on the projected threat, needs/deficiencies, operational concepts, and advanced technology fighter options." The information requested is intended to contribute to a basis for an advanced tactical fighter design for full scale development in 1987 with

an I.O.C. of early to mid-1990's.

An important mission to be considered in this program is the strategic air defense of the USA

with a mission radius of 1000, or more, nautical miles. A potential design combat profile to meet

the above mission requirements has been used for the supercruiser airframe weight payoff analysis.

The combat mission speed and altitude versus time diagram resulting from the mission profile consists

of a cruise speed which is assumed to be Mach 2.75 with an overspeed up to Mach 3.0 for approximately

80 minutes at 60,000 feet. The design life of the aircraft under these mission profile conditions

is estimated to consist of 500 flights of the basic combat mission and 4500 flights of an alternate

mission. The resulting structural design skin temperature versus time spectra are shown in Figure 5.

The supercruiser configuration selected for analysis is one of a family of configurations proposed

by the Boeing Company in Reference 8. It is a long range fighter with a design gross weight of

47,627 kg (105,000 lb.) and a cruise speed in the Mach 3 range. The airframe primary structural

material is titanium. The wing, body and strake, and empennage structures have been investigated to determine the effect of substituting P/M aluminum alloy for the existing titanium. Allocation of

weights of individual components into the design failure mode categories was made by reviewing the

data in Reference 7 and applying prior applicable experience. This experience includes the design

and service experience of the YF-12 and F-104, the preliminary design of a number of advanced tactical

fighter type aircraft, and years of research and development of commercial supersonic cruise vehicles.

Experience with aircraft designs similar to the supercruiser fighter has shown that for the major

portion of the wing, fuselage, and empennage structure the subsonic and transonic loads are greater

than the supersonic loads. Consequently, most of the structure is designed by room temperature loads.

Consequently, most of the structure is designed by room temperature material properties after the

applicable temperature/time spectra exposure rather than the elevated temperature properties. The

weight savings evaluation methodology was used to determine the weight effects of substituting P/M

aluminum alloy for titanium. In actual practice, selective substitution would result in a weight

saving of approximately 8%. Optimization of weight savings by combined selective use of several P/M

aluminum alloys optimized for strength, stiffness, or durability and damage tolerance would increase

the weight savings to an estimated 10% or 12%.

Figure 5. Supercruiser Structural Design Temperature (K) Versus Time Spectra

636
Advanced Tactical Fighter Aft Fuselage Structure

The Air Force Flight Dynamics Laboratory has for several years sponsored industry studies under

the project "Built-Up Low-Cost Advanced Titanium Structures (BLATS)". One of these programs was a preliminary design activity conducted by Grumman Aerospace Corporation directed toward establishing

advanced structural concepts for an aft fuselage of an emerging, high performance, Mach 2 class

fighter airframe (Reference 9)

The conventional titanium aft fuselage developed in the BLATS program has been used as the

baseline structure to evaluate the effects of using elevated temperature resistant P/M aluminum

alloys in a Mach 2 plus Advanced Tactical Fighter. Temperatures of the structural elements under

certain ambient temperature, speed and altitude conditions were estimated using the data of Reference

9. From these data, three temperature-time spectra have been postulated which represent the temperature-time histories in the 8000 hour life of a short to medium range advanced tactical fighter for:

(1) structure with little or no engine heating; (2) structure with moderate engine heating; and (3)

structure with maximum engine heating. The design speed/ temperature range was assumed to be: cruise

at Mach 2.2 on a standard day with overshoot to a maximum of Mach 2.5 on a hot day. The weight

savings methodology previously discussed was used to determine the effects of substituting P/M

aluminum alloy for titanium.

In the BLATS report (Reference 9) the structure is analyzed by assuming that critical load and

temperature conditions are encountered at maximum structural temperatures. In this aft fuselage much

of the structure is subjected to heat and sonic noise radiation from the engine. The engine inlet

ducts are designed by air loads and elevated temperatures due to ram air. Therefore, much of the

structure would be critical for maximum temperature conditions. However, the fuselage shell structure

forward of the engines is likely to be critical for room temperature conditions after exposure to the

higher temperatures. In this limited analysis it has not been possible to segregate the structure

into the above two categories. Therefore, the structure has been evaluated for the two boundary

conditions: (1) critical for operations at Mach 1.8; and (2) critical for operations at maximum

temperatures.

If P/M aluminum could be selectively substituted only where it yields a weight saving, a weight

saving of 13% would result. If the aluminum is substituted completely, lower net savings of 6%

is shown in Table 6. In actual practice, unrestricted selected usage is not practical. A realistic,

limited selective usage is estimated to result in a net weight saving of approximately 8%.

Military Transport Components

Certain structural components, currently made of titanium, stainless steel, and aluminum, are subjected to critical load-temperature design conditions in the 394 K (250 F) to 616 K (650 F) range.

Several of these components, identified in Table 7 have been investigated to determine the weight

impact of the use of high temperature resistant P/M aluminum alloys. Shown in this table are: the

currently used materials, the structural weights; the design temperatures; and the primary structural

design conditions.

Figure 6. Estimated Potential P/M Aluminum Weight Benefits Variation With Cruise Mach Number

637
The results of utilizing the DARPA weight ratio equations and dominant design failure mode (or

modes) have been used to determine the effect of changing the titanium structure to P/M aluminum. In the case of the C-5A Engine Nacelle, which is made of titanium and aluminum, it is estimated that half

of the total weight would be converted from titanium to P/M aluminum. Weight savings of approximately

one to eleven percent are possible for the remaining structures.

TABLE 6. SUMMARY OF PERCENTAGE WEIGHT SAVINGS FOR ADVANCED TACTICAL FIGHTER AFT FUSELAGE

PERCENTAGE WEIGHT SAVINGS

DEGREE OF

ENGINE

HEATING

CRITICAL LOAD

AND TEMPERATURE

CONDITIONS

UNRESTRICTED

SELECTIVE

USAGE

COMPLETE

SUBSTITUTION

NONE M ^ 1.8 12.4 5.0

MAX. TEMP. 13.2 5.2

MODERATE

M ^ 1.8 13.8 7.0

MAX. TEMP. 13.4 5.8

MAXIMUM

M ^ 1.8 12.8 4.5

MAX. TEMP. 11.7 -2.0

TABLE 7. HIGH TEMPERATURE ALUMINUM APPLICATION CANDIDATES FOR MILITARY TRANSPORT AIRCRAFT

EXISTING

MATERIAL

WEIGHT

KG (LB)

OPERATING

TEMPERATURE

K (°F)

PRIMARY

DESIGN

CONDITION

PYLONS - C-5A Ti 908 (2002) 422 (300) STIFFNESS

FWD LONGERONS /BULKHEAD Ti/S.S. 500 (1102) 616 (650) STIFFNESS

AFT SKINS/LONGERONS/BULKHEADS

ENGINE NACELLE - C-5A

DOORS Ti/Al 381 (841) 394-644

(250-700)

AIR LOADS

FLAP SURFACES

C-141A (T.E. ASSEMBLY) Ti 596 (1313) 505 (450) SONIC FATIGUE

C-130 (LOWER SURFACE) Ti 80 (17) 422 (300) SONIC FATIGUE

Airframe Payoff Study Conclusions

Analyses performed in this study have shown elevated temperature resistant high strength - high

stiffness aluminum alloys to be viable candidates to replace titanium in aircraft structures subjected

to elevated temperature environments. Weight savings in the range of 5% to 10% havei consistently been

calculated for a varied number of structural components ranging from the engine heated structures of

subsonic transports to the entire airframe of a Mach 3 class Supercruiser/Advanced Tactical Fighter.

From these studies, combined with experience from previous studies of high performance military

aircraft accumulated over the past two decades. Figure 6 has been constructed to show estimated weight

versus Mach Number variations for an advanced tactical fighter airframe. Potential P/M aluminum usage

is shown compared to all conventional 7075-T6 aluminum alloy and to a least weight combination of

titanium and 7075-T6 aluminum. While these curves are estimates and would vary somewhat with the

aircraft configuration and service life, the trends are valid. The elevated temperature resistant

aluminum alloys have the potential to be the primary structural material for high performance military

aircraft up to the Mach 3 cruise speed range.

638
In a recent publication (Reference 10) the material mixes of past and current fighter aircraft were shown with a projection to a future tactical fighter was postulated to be a Mach 2 plus supersonic

cruise aircraft having an initial operating capability (IOC) between 1990 and 1995. Contrary to that

prognostication, it is forecast that with the development of the P/M aluminum alloys, the material

mix in the future tactical fighter will be modified. The high temperature aluminum alloys will replace

much of the projected use of titanium and high strength-highstiffness aluminum alloys will reduce the use of advanced composites.

CONCLUSIONS

Application of rapid solidification technology to advanced aluminum alloys has lead to development

of new alloy compositions which have significant potential weight saving payoffs for a variety of new

applications

o For gas turbine engine vane applications weight savings of 24-35% appear feasible. Blade applications will require improved fatigue capability to achieve weight savings of 35%. Current fatigue

properties will limit weight savings to approximately 7%.

o Airframe application of elevated temperature P/M aluminum alloys will lead to weight savings of

8 to 12% depending on the degree of alloy optimization required for an aircraft of the supercruiser

configuration. The advanced tactical fighter weight savings were computed to be between 6 and 13%

depending on the mission profiles and alloy optimization chosen, with an expected weight saving of

approximately 8%. Weight savings for various military transport components ranged from 1.2 to 10%

depending on the components selected.

ACKNOWLEDGEMENTS

The initial work on elevated temperature aluminum alloys was performed under a Defense Advanced

Research Projects Agency contract monitored by the Air Force Wright Aeronautical Laboratories,

Contract F33615-76-C-5136, "Application of Rapidly Solidified Alloys". Subsequent work has been

performed under a sub-contract from Lockheed California Company under funding from Air Force Wright

Aeronautical Laboratories on contract F33615-81-C-5096, "Elevated Temperature Aluminum Alloys". The

continuing encouragement and support of Dr. E. van Reuth and Dr. L. Jacobsen (DARPA) , and of

Mr. A. Adair and Mr. W. Griffith (AFWAL) as program monitors for these programs is gratefully

acknowledged. The views and conclusions contained in this document are those of the authors and

should not be interpreted as necessarily representing the official policies, either expressed or implied, of the Defense Advanced Research Projects Agency, U.S. Air Force, or the U.S. Government.

REFERENCES

1. Adam, C. M. , R. G. Bourdeau and J. W. Broch, "Application of Rapidly Solidified Alloys," Final

Contract Report, September 1981, Contract F33615-76-C-5136.

2. Adam, C. M. , "Structure/Property Relationships and Applications of Rapidly Solidified Aluminum

Alloys," Proc. 1981 Mats. Res. Soc. Symposia, North-Holland Elsevier (1982) p. 411.

3. Adam, C. M. and J. W. Simon, "Influence of Microstructure on Texture Development in an Aluminum

Powder Metallurgy Alloy," A.S.M. Conference on Textures/Microstructures/Mechanlcal Properties,

1982, to be published.

4. Langenbeck, S., "Elevated Temperature Aluminum Alloys," Air Force Wright Aeronautical Laboratories

Contract F33615-81-C-5096, Semi-Annual Technical Report, 1982.

5. Simenz, R. F., G. G. Wald, S. L. Langenbeck and E. J. Himmel, Lockheed California Company

Technical Proposal, RFP F33615-80-R-5174 , 23 September 1980.

6. Lewis, R. D. , et al. Phase I Final Report, DARPA/AFML Contract F33615-78-C-5203, to be published.

7. Design Conference Proceedings, Technology for Supersonic Cruise Military Aircraft, Air Force

Flight Dynamics Laboratory, February 17-20, 1976.

8. Hahn, Darll W. , Technical Report, AFFDL-TR-78-80, Volume I, July 1978.

9. Paez, C. Technical Report, AFFDL-TR-79-3093, August 1979.

10. Hadcock, Richard N., Technical Paper, NASA Conference Publication 2162 - Part 2, Tactical Aircraft

Research and Technology Conference Proceedings, October 21-23, 1980.

639
GAS TURBINE APPLICATIONS OF RST ALLOYS

D.B. George, R.G. Bourdeau, J.A. Miller

Pratt & Whitney Aircraft

Government Products Division

West Palm Beach, Florida 33402

ABSTRACT

Technological advances in the rapid solidification of alloys useful to the gas turbine

engine show benefits related to performance, durability, and weight reduction. Implementation of this technology has been achieved successfully for superalloy turbine

disks. Near-term plans for implementation have been prepared for RST high temperature

aluminum alloys as a substitute for titanium and RST corrosion resistant bearings to replace conventional M50. Mid-range plans are directed toward implementation of RST

superalloys for airfoils and RST intermetallics for burners and liners. The status of

the RST technology is addressed in this paper and engine design benefits are described

INTRODUCTION

A new era of supremacy for allied tactical air forces began in the mid-1970's when the F-15 and later

the F-16 entered the U.S. Air Force inventory. As of this writing, there are 727 McDonnel Douglas

F-15's and 784 General Dynamics F-16's in service with free-world air forces. By the mid-1980's an estimated 925 F-15's and 1,550 F-16's will be deployed worldwide. Both of these aircraft are powered

by the Pratt & Whitney Aircraft FlOO engine, which featured the first production application of

powder metallurgy turbine and compressor disks.

The qualitative advantage in tactical airpower provided by these and other types of aircraft in

service today is not guaranteed to last indefinitely. The typical life of a high-performance fighter aircraft system is about eighteen to twenty years, so we will need a replacement for the F-15 in

1992 if we are to retain our technological lead. That may seem a long way off, but to achieve a 1992 introduction via a competitive flyoff of prototype aircraft, the program to develop the engine

for the new fighter will have to start in 1983.

Improved fighter maneuverability expressed in terms of turn rate, acceleration, or rate of climb is

directly related to the ratio of available engine thrust to the total aircraft weight. This ratio

has nearly tripled since the 1940 's as a result of the corresponding increase in the thrust-to-weight

ratio of the engine itself. We believe this trend will continue due in large measure to the expanded

use of powder metallurgy materials which feature higher specific strength and temperature capabilities. The revolution begun in the mid-1970's, with the introduction of powder metallurgy disks into

the FlOO engine, has evolved to encompass a broad spectrum of anticipated powder metallurgy applications for the next tactical engine system. Prior to discussing the potential impact of powder

metallurgy on future engine systems, it is worthwhile reviewing the highly successful FlOO engine

experience which serves as the foundation for the further exploitation of this technology.

HISTORICAL BACKGROUND

During the early 1960's, in the development of the J58 duct-burning turbofan engine by Pratt and

Whitney Aircraft, it became apparent that Waspaloy and Rene 41—the most advanced nickel-base alloys

then used for forgings—did not possess the high temperature strength required for the turbine disks

A new wrought alloy, at that time known as Astroloy, did exhibit the required strength and elevated

temperature operating capability. The evaluation of Astroloy forgings, however, revealed that while

the desired strength potential was present, these forgings exhibited excessive scatter in mechanical

properties due to gross structural segregation and lack of homogeneity as shown in Figure 1.

640
FIGURE 1: CROSS SECTION OF AN ASTROLOY DISK FROM 20 IN. DIA. INGOT

The segregation noted in early Astroloy forgings could be traced back to the large columnar grain

structure of the ingots being produced at that time. Metallurgical evidence was compiled, showing

that the segregation existing in conventional ingots was not eliminated by subsequent forging and

thermal treatments.

Based on these findings, a full-scale Astroloy ingot evaluation program was initiated by Pratt and

Whitney Aircraft. The results of the evaluation showed that casting parameters yielding relatively

fine, equiaxed grain structure minimized macrosegregation in the ingot. Production processes were developed which produced such structures in ingots which were forged into full-scale turbine disks

for the J58 engine. These disks have compiled an excellent record in the field.

Subsequent to this, in an effort to develop a stronger high temperature disk composition than Astroloy, ultrafine equiaxed grain ingots of INIOO were cast. Eight-inch diameter disks were forged from

the fine-grained INIOO ingots, however, the macro- and microstructure of these disks, produced from

what was considered an optimum fine-grained ingot, contained a prohibitive amount of segregation.

This barrier to further development of stronger disk materials lead to consideration of a powder

metallurgy approach to the production of large engine disks with a minimum of macrosegregation. The

degree of macrosegregation in a small powder particle would be expected to be less by several orders

of magnitude than that present in the structure of a large ingot conventionally melted and cast

under optimum conditions.

A method of powder production, collection and densif ication was developed by Pratt and Whitney Aircraft to eliminate the problems associated with particle oxidation, scatter in properties, and lack

of reproducibility which had plagued efforts up to that time to utilize powder metallurgy for superalloy parts. To avoid oxidation of individual powder particles, an "all inert" method of powder

production, collection, and billet densif ication was developed wherein the metal powders would not

be exposed to the atmosphere prior to densif ication (Figure 2). Astroloy, the most advanced production nickel-base superalloy disk material at that time, was selected for the investigation.

All-Inert FIGURE 2: ALL-INERT POWDER METALLURGY PROCESSING PREVENTS THE FORMATION OF OXIDE

FILMS ON POWDER PARTICLES

641
This initial program demonstrated that small Astroloy forgings (Figure 3) made by the "all inert"

technique could be compacted and forged over a wide temperature range and showed uniformity of

structure and mechanical properties. The macro- and microstructural segregation of the small powder

disks was significantly reduced from that of even the most optimum structure seen in a disk made

from a cast ingot.

FIGURE 3: POWDER METAL PANCAKE FORGING ILLUSTRATES DRAMATIC REDUCTION IN SEGREGATION

While this Astroloy powder program was underway, it was discovered that iron, nickel, and titanium

base superalloys could be made to behave superplastically . Billet and bar stock of A-286, INCO 901,

Astroloy, Waspaloy, B1900, INIOO, and titanium alloys 8-1-1 and 6-4 were produced by controlled

processing so that they behaved in a superplastic manner over a fairly wide temperature range if the

strain rates were properly controlled (Figure 4).

1 600*F - 46%EL

1650*F • 240%EL

nOO*F - 217%EL

1800* F - 888%EL

1900* F • 1390% EL

2000* F - 1220% EL

2100*F • 230% EL

Note; Temperatures Shown Under Each Specimen

Are Test Temperatures

FIGURE 4: INIOO SPECIMENS AFTER STRESS RUPTURE TESTING

642
The next logical step was to forge these alloys isothermally in the temperature range and at the

strain rate where they demonstrated superplastic behavior. It was found that billet stock previously made superplastic could be forged isothermally into complex configurations using very low unit

pressures. This controlled isothermal forging practice, illustrated in Figure 5, ^ich utilizes

the superplastic behavior of the alloy being forged was given the name GATORIZING ®by Pratt and

Whitney Aircraft and is covered by U.S. Patent 3,519,503.

Containing Chamber

for Argon, Ng,

FlOO ENGINE POWDER METALLURGY EXPERIENCE

The first application of powder metallurgy to production parts was for the high thrust-to-weight

ratio FlOO engine. In the early stages of the development of this engine, the materials selection

was predicated on achieving the highest strength-to-weight ratio that it was possible to obtain.

For the latter stage compressor disks and turbine disks, we were in need of a stronger alloy than

Astroloy, then the most advanced alloy operating in a production jet engine. INIOO, a nickel base

superalloy developed for cast turbine blades, had the best strength-to-weight ratio at operating

temperatures of available alloys. However, this highly segregation prone alloy was unusable for

large engine disk forgings using conventional state-of-the-art melting and forging practices. By

utilizing the "all inert" powder metallurgy concept developed in conjunction with the GATORIZING

forging process, we were able to produce compressor and turbine disks and spacers from the INIOO

alloy for the FlOO engine with an unprecedented level and uniformity of mechanical properties.

Figure 6 is a cutaway of the FlOO engine with a listing of the nine INIOO parts produced by the

GATORIZING process.

The first FlOO engine disk from INIOO powder billet, forged by the GATORIZING process, was produced

in April 1970. Since that time, 3,029 production FlOO engines have been delivered to allied Air

Forces worldwide each containing the nine powder metallurgy components illustrated in Figure 6 for

a total of 27,261 production components. The F-15/F-16 fleet had accumulated a total of 1,385,000

engine flight hours as of November 1982, which equates to 12,465,000 hours of flight experience with

powder metallurgy components. To date, there has only been one incident where a material processing

related disk failure was suspected. The cause of the incident could not be established despite an extensive investigation. To say the least, the operational experience with FlOO powder metallurgy

components to date has been phenominal; particularly when one remembers that this is the first production application of powder technology in a gas turbine engine. Based on the excellent experience

with powder metallurgy components in the FlOO engine, the stage is set to further exploit this technology in the next generation fighter engine, the development of which begins next year.

643
91h-S1age Compressor Disk llth-Stage Compressor Disk 13th-Stage Compressor Disk 13th-Stage Compressor Cone Seal

Ist-Stage Turbine Disk 2nd-Stage Turbine Disk 3rd-Stage Turbine Disk 4th-Stage Turbine Disk 2nd-Stage Turbine Spacer

FIGURE 6: FlOO ENGINE USED GATORIZED © INIOO COMPONENTS PRODUCED FROM INIOO POWDER

METALLURGY BILLET

ADVANCED RST ALLOYS

Having made the revolutionary technology step from ingot to all inert powder in the FlOO engine,

Pratt and Whitney Aircraft with government support has embarked on a Rapid Solidification Rate (RSR)

technology program aimed at further capitalizing on the potential of powder metallurgy alloys. As

illustrated in Figure 7, the development of RSR alloys represents an extension of the increased

cooling rate technology revolution which began with the now state-of-the-art processing of all inert

powder

A device for producing rapidly solidified alloy powders was constructed by Pratt & Whitney Aircraft

which uses forced convective cooling of molten particles accelerated from a central source into a high conductivity gas quench medium. The rig is shown in Figure 8. An alloy charge is melted in

the crucible at the top of the rig. The liquid metal is poured into a transfer system, which in turn

supplies a continuous flow of molten metal to a disk atomizer spinning at approximately 24,000 rpm. The metal is centrifuged into particles ranging from 10 to 100 microns and is subsequently solidified

by an array of helium gas-quenching jets.

Ingot

0.01°C/Sec

Cooling Rate

Argon

All Inert Powder

10,000°C/sec

Cooling Rate

1,000,000°C/Sec

Cooiing Rate

200Mm

1 I 2^m

FIGURE 7: RSR POWDER PROCESSING EVOLVED FROM ALL-INERT POWDER EXPERIENCE

644
FIGURE 8: RSR POWDER PRODUCTION FACILITY

Several hundred experimental aluminum, iron, and nickel base powder metal alloy compositions have

been atomized and evaluated to date under the sponsorship of the Defense Advanced Research Projects

Agency (DARPA) , the Air Force Materials Laboratory (AFML) , the Naval Research Laboratory (NRL), the

Office of Naval Research (ONR) , and the Naval Air Systems Command (NAVAIR) . In addition to further

advances in compressor and turbine disk material capabilities, potential powder metal applications

in the next generation fighter engine include lightweight aluminum fan stators, corrosion resistant

bearings, oxidation resistant/lightweight burner liners, and improved durability turbine airfoils.

A brief description of the alloys being pursued for these applications and the associated benefits

is described in what follows.

Advanced Turbine Disks Materials

An interdivisional alloy development program has been underway at Pratt and Whitney Aircraft for

several years aimed at further refining the INIOO disk material chemistry. Figure 9 illustrates

one dramatic improvement which was achieved in crack growth rate reduction through the use of revised

processing. Studies are proceeding to select the best combination of properties to satisfy the re- quirements of a wide variety of future engine disk requirements, including several derivatives of the

FlOO engine and the 5000 SHP Modern Technology Development Engines (MTDE) for the Army and Na-yy.

FIGURE 9; CRACK GROWTH RATE CHARACTERISTICS OF AN ADVANCED VARIANT OF INIOO (PWA 1074)

645
In addition to improved versions of INIOO, a new chemistry disk material is being developed in

support of the next generation fighter engine program to achieve the properties illustrated in

Figure 10. As shown, a +50°F creep rupture capability in conjunction with a 10% improvement in

tensile properties, and a crack growth rate reduction are desired. The final selection of the

chemistry of this new alloy will be made in 1984.

The improved turbine disk material will permit an increase in the rotational speed of the high com- pressor rotor enabling elimination of one compressor stage. Alternatively, use of this material

could result in extension of the disk inspection interval and/or reduced weight with an attendant

reduction in operational and support costs.

Fatigue Rate

FIGURE 10: RELATIVE PROPERTY GOALS FOR NEW TURBINE DISK ALLOY

RSR Aluminum Compressor Airfoil s

The development of a new high temperature aluminum alloy (Al-8Fe-2Mo) with useful strength capabilities in the 500°F range is proceeding based on the application of rapid solidification rate powder

metallurgy technology. As illustrated in Figure 11, RSR aluminum alloys have been produced which

provide on the order of a 150°F improvement in alloy capability relative to conventional aluminum

alloys such as AA2219. Ultimately, it is believed the useful temperature range could be extended

to the 650°F regime. In view of the improved temperature capability of Al-8Fe-2Mo (developed under

FIGURE 11: COMPARISON OF RAPID-SOLIDIFICATION-RATE ALUMINUM TO CONVENTIONAL

TITANIUM AND ALUMINUM ALLOYS

646
AFML contract F33615-79-C-5136) , the potential exists to replace heavier and more costly titanium

components with aluminum in the first two fan stages.

The density of this new aluminum alloy is on the order of .104 lbs. /in. whereas the density of

Ti-6A1-4V is .160 lbs. /in. . Analyses using the strength properties achieved to date indicate a potential of up to a 20% weight savings for vane applications. During 1982, a set of 2nd stage

fan stator vanes will be fabricated from the current RSR aluminum alloy for an FlOO derivative engine

test. If the goal alternating stress capability of Figure 11 is realized, applications to compressor

blades are also envisioned with an attendant weight reduction for the supporting disk.

Corrosion Resistant Bearings

Recent Air Force Aero Propulsion Laboratory testing, (F33615-76-C-5136) of a series of RSR processed

bearing alloys have demonstrated the potential of significantly improving the rolling contact fatigue

(RCF) and corrosion lives of M50-like bearing steels. As illustrated in Figure 12, the alloy designated RSR 561 displayed a five-fold improvement in the B-10 rolling contact fatigue life relative to the baseline VIM-VAR M50 material. Also noteworthy is an RSR alloy which displayed a modest 20% improvement in rolling contact fatigue but which contains 19% chromium (versus 4% in M50) . Due to its

high chromium content, the RSR alloy should be substantially more corrosion resistant than M50,

prompting the inclusion of an RSR alloy in the government-sponsored Improved Corrosion Resistant Turbine

Engine Bearing Program (F33615-81-C-2023) . In addition to being pursued under that program, the

trial fabrication of prototype RSR bearings will also be undertaken in cooperation with the Fafnir

Bearing Company.

FIGURE 12. ROLLING CONTACT FATIGUE DATA FOR BEARING ALLOYS

Iron-Aluminum Alloys

Iron-aluminum alloys have been recognized for a long time as having desirable properties for high

temperature applications. They have excellent oxidation and sulfidation resistance, significantly

lower density than nickel or cobalt alloys (.22-. 24 lbs. /in. ^ as opposed to .28-. 33 lbs. /in. 3), and

they contain no strategic elements. Until recently, however, iron aluminum alloys containing a sufficient amount of aluminum for good high temperature oxidation resistance were too brittle for

engineering applications and their strength, especially at elevated temperature, was mediocre.

Recently, it has been shown that rapid solidification and thermomechanical processing can produce

fine dispersions of titanium diboride in Fe-Al alloys. The dispersions greatly improve both the

strength, ductility, and temperature capability of the iron-aluminum alloys. While the data are

limited, it is believed that the dispersion strengthened iron-aluminum alloys can be developed into

two classes of alloys; one class has the oxidation resistance (Figure 13) and high temperature creep strength required for use at 1800°F to 2000°F while the other class has a lower density with somewhat

reduced temperature capabilities.

The lower density alloy is viewed as a candidate to replace Inconel 718 in the compressor. The high

temperature oxidation resistant material will be considered for application from the combustor

rearward.

647
2000° F Cyclic Oxidation Resistance Stress-Rupture Strength

FIGURE 13. PROPERTIES OF IRON-ALUMINUM RELATIVE TO TRADITIONAL BURNER ALLOYS

Turbine Airfoil Materials

Currently both RSR vane and blade alloys are being developed under sponsorship of the Defense Advanced i Research Projects Agency (DARPA) and the Air Force Materials Laboratory (AFML) . The RSR vane alloy

features a significant improvement in oxidation resistance (Figure 14) and a +65°F increase in incipient melting point relative to the advanced single crystal alloy PWA 1480 which has been employed

in recent versions of the JT9D commercial engine and the FlOO derivative engine family. The RSR

blade alloy will have a -l-65°F advantage in creep rupture (Figure 15) and provides a significant improvement in oxidation resistance relative to PWA 1480. Both RSR processed alloys are currently

undergoing detailed material property characterization scheduled for completion by the end of this

year. Prototype vane and blade airfoils are scheduled to be fabricated next year for an advanced

engine durability test commencing in early 1984.

The advanced RSR turbine airfoil alloys will provide cost of ownership benefits through either improved airfoil life (reduced maintenance cost) or a reduced cooling air requirement (improved SFC)

FIGURE 14. 2100°F CYCLIC OXIDATION RESISTANCE OF RSR SUPERALLOYS

648
100

FIGURE 15. CREEP RUPTURE STRENGTH COMPARISON OF ADVANCED TURBINE BLADE MATERIALS

SUMMARY

The successful implementation of powder metallurgy disks into the production FlOO engine and the

excellent operational experience of those critical components has set the stage for the incorporation of a wide range of powder metal systems via rapid solidification processing into the next

generation fighter engine. The timely development of advanced aluminum, iron, and nickel base

powder metal superalloys under DARPA/AFML/NRL/ONR/NAVAIR sponsorship will enable the next generation

fighter engine to provide a 20% improvement in engine thrust-to-weight ratio over the current FlOO

engine. A performance improvement of that magnitude will ensure continuance of the era of tactical

fighter aircraft superiority which began with the introduction of the FlOO powered F-15 aircraft.

ACKNOWLEDGEMENTS

The materials developments described herein are being supported by the Defense Advanced Research

Projects Agency (under ARPA Order 3152), the Air Force Materials Laboratory (AFML) , the Naval Research Laboratory (NRL) , the Office of Naval Research (ONR) , and the Naval Air Systems Command

(NAVAIR)

The views and conclusions contained in this document are those of the authors and should not be

interpreted as necessarily representing the official policies, either expressed or implied, of the

Defense Advanced Research Projects Agency, U.S. Air Force, U.S. Navy or the U.S. Government.

REFERENCES

1. R.L. Athey and J.B, Moore, "Development of INIOO Powder Metallurgy Disks for Advanced Jet

Engine Application," Powder Metallurgy, Syracuse University Press, 1972.

2. R.L. Athey and J.B. Moore, "Progress Report on the GATORIZING ® Forging Process," SAE Paper

751047 presented at the National Aerospace and Manufacturing Meeting, Culver City, Los Angeles,

November 17-20, 1975.

3. F.W. McAbee, Jr., "Turbine Technologies of Tomorrow," Air Force Magazine, January 1982.

649
AIRCRAFT ENGINE GAS TURBINE COMPONENT

FABRICATION CONCEPTS USING RSPD

A.M. Johnson, J.S. Kelm, R.W. Smashey

D.V. Rigney & T.G. Wakeman

General Electric Co., Evendale, OH. 45215

ABSTRACT

The Rapid Solidification Plasma Deposition (RSPD) process offers unique opportunities for the

fabrication of structural components. Injection of powder into a high velocity, low pressure plasma jet

with subsequent melting and deposition of the powder onto a removable mandrel results in fine

grained, homogeneous deposits. Structures of varying complexity can be made to near net shape from an extremely wide range of material compositions. Construction materials can be tailored for specific

service environments within a component. Fabrication concepts using RSPD as a flexible

manufacturing process for advanced aircraft gas turbine engine components are discussed as well as some of the expected benefits to be derived. Material and design aspects of RSPD are considered.

INTRODUCTION

The demands of advanced aircraft gas turbine engine operating requirements for increased performance

with minimal influence on component costs present a major challenge to design, materials and

manufacturing engineers. Conventional casting, forging and powder metallurgy (PM) processes have been

under development for many years, and further gains in capability in these areas are difficult to obtain.

Furthermore, final part costs have continued to increase at a steady rate due to accelerating material and

labor costs.

Rapid Solidification (RS) technology offers improvements in alloy design and material homogeneity. A

wide range of RS processes are currently under investigation throughout the world with many promising

results. In most cases, rapidly solidified material is initially produced, followed by consolidation and metal

working steps to reach a desired configuration from which a part is machined. Frequently, the most

challenging aspect of making actual parts from rapidly solidified material is the careful handling and

consolidation processing required.

The Rapid Solidification Plasma Deposition (RSPD) process under development by the General Electric

Company provides the capability for flexible manufacturing of simple or complex shapes wherein the

solidification and consolidation steps are combined. The essence of this technology derives from a coupling

of a high energy plasma gun operated in a low pressure chamber to deposit powder material of a desired

composition onto a substrate. As-deposited densities are sufficiently high that post-deposition heat

treatments are normally adequate to achieve full density. Near net shapes may be formed by proper selection

of mandrel configurations and deposition process control. Structures having a wide range of geometric

configurations, including multiple alloy compositions within a single part, are made possible by this

technique.

The development of RSPD for the fabrication of advanced aircraft gas turbine engine components is actively being pursued. The primary focus of the current effort is to provide a basic understanding of the

process to establish a firm technology base for subsequent implementation. This paper addresses the key

material, component design and fabrication considerations for RSPD technology.

650
BACKGROUND

The extremely high temperatures generated in plasma arcs have been used for a variety of purposes

including; a) purification of high melting temperature materials, b) spherodization of refractory particles, c)

high energy/high temperature testing, and d) deposition of specialized materials to provide unique

characteristics.

Gianinni”’, described the atmospheric plasma jet and projected material development applications. By

the early 1960's, primitive plasma generators manufactured in the United States were used to apply coatings

for repair of used or mis-machined components'^’. At the same time, tungsten rocket nozzles were plasma

formed. Mash and Brown‘d’ describe improvements in bonding and strength of stainless steels, aluminum

and other alloys in inert environments. Densities of the deposits were substantially improved from the usual

range of 75-85% to as high as 92% due to minimal oxidation of the sprayed powder. Further treatment, by

vacuum sintering or hot isostatic pressing (HIP), has served to increase densities of some deposits to near

theoretical*'*’.

In conventional plasma processing, the environment is maintained at atmospheric pressure. In experiments performed by Muehiberger'®’, 80 KW plasma jets produced at ambient pressures of —0.1

atmosphere (76 torr) indicated a jet velocity of Mach 3, whereas previous 80 KW atmospheric plasma jet

velocities seemed limited to Mach 2. More importantly, Muehiberger demonstrated the capability to produce

deposits nearing 100% density with relative ease, whereas Okada and Moruo'®’, using inert gas shielding,

reported 98-99% cast alloy density only under carefully controlled conditions. Rairden et al.'^’, conducted an

evaluation of low pressure plasma deposition for use as a method to deposit reactive MCrAlY coating alloys at the Muehiberger facility. The deposits were near full density with little if any deposit contamination. The low

pressure inert environment was acceptable for preheating components to sufficient temperature to allow

deposition of MCrAlY coatings onto large industrial gas turbine blades. In 1977, simple plasma deposits of

the superalloys IN738 and Rene' 80 were made'®’. These initial plasma deposits, after heat treatment,

showed significant improvements for RSPD IN738 over conventionally cast IN738 strength and ductility at low temperature (<1200° F) as well as improved thermal fatigue resistance of RSPD Rene' 80 over cast Rene'

80.

Plasma-forming has also been used for specialty constructions, in which the base material is sprayed at atmospheric pressure onto a mandrel material to match the desired internal configuration of the component.

After spraying, the mandrel material is removed by machining, leaching, or mechanical separation. The

deposits are then typically finished to final dimensions. Since the plasma process relies on the deposition of

molten or semi-molten material, conformation to the shape of the mandrel can be excellent. As such, the

thermal spray process has been profitably used to construct dies and cast forms for specialty items such as

precision plastic furniture components, decorations, and shoe soles, each of which require dimensional

stability and faithful reproduction of the surface*®’.

PROCESSING FOR STRUCTURAL SHAPES

Atypical low pressure plasma arc chamber used for RSPD processing is shown in Figure 1. The overall

system includes a plasma generating gun, devices to manipulate the gun and target substrate, pumps, filters,

power supplies, powder feeding equipment, plasma forming and powder carrier gas equipment, and controls

to regulate the process and manipulations.

Figure 2 is a schematic of the plasma gun showing an internal injection port for the powder. External

feeding of powder is a feature of other gun designs. They are generally capable of continuous operation in the

50 - 80 KW range, with short excursions to 120 KW. The position of the plasma torch is computer controlled

in four axis; pitch, yaw, horizontal and vertical movement. The target mandrel can be moved laterally along

the horizontal axis and rotated. Such a manipulation system is shown schematically in Figure 3.

651
Inert Gas

Figure 1 RSPD Development Chamber

Gun Power

Supply

Transferred Arc Power Supply

Figure 2 Schematic of a Plasma Spray System.

For Reversed Transferred Arc Cleaning, the Polarity

of the Transferred Arc Power Supply is Reversed

652

Figure 3 Schematic of Six Axis Plasma Gun and

Mandrel Manipulation System
The plasma gun uses modified gas mass flow controllers to regulate the primary and secondary plasmaforming gases. Positive transferred arc, which is used to provide supplemental heating to the mandrel or component mounted on the horizontal axis, can be computer sequenced in time and position with the target

mandrel. Similarly, the powder feed, vacuum pump and valve systems are also computer controlled. Thus, once a specific process program is written, an automatic sequence can be initiated for a) plasma arc

initiation, b) reverse transferred arc (RTA) cleaning of the mandrel surface, c) positive transferred arc and/or

plasma torch impingement preheating of the mandrel, d) spraying and e) shut-down. Pressure within the

chamber is automatically maintained, and the temperature of the substrate is monitored using an optical

pyrometer or contact thermocouple. Process data are acquisitioned and stored using a data-logger connected

to the computer terminal.

The fabrication of RSPD structural components involves several major steps: 1) design and construction of

mandrels, 2) selection of powder size and composition, 3) deposition processing and 4) post-deposition

processing.

The mandrel is an important part of the RSPD process. In production modelling, the cost of mandrel

fabrication and removal has a significant influence on the economic feasibility of the process. Therefore, a mandrel system must either be very simple and inexpensive so that is can be sacrificed after deposition, or durable enough to withstand many deposition cycles. The successful application of either method appears to be dependent largely on the size of the component being constructed. For example, disposable mandrels

may be more suited for small components, whereas more complex reusable tooling may be required for

larger components.

Near net shape deposition requires close dimensional control of both the mandrel substrate and the

deposit. Conventional ambient pressure plasma forming uses mandrels which are not heated, or which are heated to a maximum of only a few hundred degrees Farenheit. The deposit thickness is generally much less

than 0.060 inches - and the densities of the deposits are usually less than 90% of theoretical. The RSPD

process, on the other hand, requires that mandrels be heated to ~ 1600°F to avoid spalling of the deposit as

a result of accumulation of residual tensile cooling stresses and to facilitate adequate bonds. Delivery of

powder to the plasma torch in a uniform, predictable manner requires handling methods which keep the

powder free-flowing and uncontaminated with extraneous materials. Further contamination of the powder

during deposition is avoided by control of interstitial impurities in plasma forming gases and the chamber

environment. Typically, the superalloy powders obtained commercially in the size range of present use (<37

/im) are available with total interstitial (O2, N2) contamination levels of — 250 ppm. In a properly operating

system, superalloys are deposited with no increase in contamination level. Large quantities (>10%) of very

fine, (<10 /ym) diameter powder have been found to impair the smooth delivery of powder to the plasma

generator, and tend to evaporate in the plasma plume

When the molten particles impact on the substrate they "splat”, forming a lenticular deposit 1-3 //m thick,

and up to 100 ywm in diameter. These splats are rapidly solidified at rates estimated to be — 10^ °F/sec,

forming a chemically homogeneous and fine grained (0.2-0.5yum) structure. The as-deposited microstructure

is dependent on a large number of process variables, such as traverse rate, gun power, powder introduction

method and position, chamber pressure, plasma torch orientation to the workpiece, etc. The effects of some

of the variables on RSPD microstructures are discussed by Smith et ar°’ and Jackson et al”^’.

After deposition, removal or separation of the deposit from the mandrel is required. Mandrels made from

materials such as copper or low alloy steel can be dissolved in acid. Machining or other metal removal

techniques can also be used. Since the as-deposited RSPD structures are relatively brittle, a heat treatment is given to ductilize the deposits and to reduce porosity. RSPD deposits made with substantial relative

movement of the gun and substrate are generally over 98% of cast or wrought densities.

After heat treatments of 2150-2300°F, deposit densities generally increase to >99% of theoretical. Grain

size of /'strengthened superalloys increases to ~ 10-30 //m. Solution strengthened alloys, such as

Hastelloy X, show much larger grain growth after heat treatment to — 75-300 /ym.

653
COMPONENT APPLICATION CONSIDERATIONS

The RSPD process offers unique opportunities to be considered in the design of gas turbine components.

These opportunities include increased flexibility of material utilization, near net shape capability, and graded

multi-material configurations. The design, material and fabrication aspects of two widely differing gas turbine

engine components are described below. Special consideration for RSPD processing and the potential

benefits obtainable relative to conventional procedures are discussed.

Airfoil cooling inserts have traditionally been made from sheet metal. Due to the necessity to conform to

rigid dimensional requirements, sheet metal inserts are difficult and expensive to fabricate. Advanced design

blade inserts with complex shapes have dimensional requirements that are typically ± 0.002 inch and have

the additional complexity of requiring variable, localized wall thickness control. To make such designs from

fabricated sheet metal requires extensive machining. To reduce this machining cost, explosive forming of

numerically controlled (NC) taper ground tubing has been attempted. Early in the refinement of this

approach it became apparent that the amount of plastic deformation which is required to produce an insert

configuration would produce a high percentage of localized tubing failures.

Figure 4 RSPD Blade Cooling Insert Fabrication Sequence

Tubing losses during explosive forming have been significantly reduced through the use of ultrafine grain

size RSPD Waspalloy. The RSPD Waspalloy has demonstrated the capability to accommodate more high

speed deformation without failure than its coarser grain, wrought counterpart. Since localized wall thickness can be controlled during the RSPD fabrication process, the expensive NC grinding operation used to produce

tapers in wrought tubing was eliminated. The process used to make inserts using RSPD is schematically

illustrated in Figure 4. In this process, a controlled thickness of Waspalloy was deposited onto a pre-tapered

soft iron mandrel. Deposit wall thickness control was accomplished by varying the rotation speed of the

mandrel and the translation speed of the plasma gun, thereby producing a barber pole type spray path about

the mandrel. Fully heat-treated RSPD Waspalloy tubing used to make the inserts had an oxygen content of

about 260 ppm and was between 98 to 100% dense after sintering. Typical grain size before explosive

forming was about 12 /wm. Utilization of the RSPD Waspalloy tubing reduced manufacturing losses during

explosive forming from over 75% for the wrought material to less than 5%.

654
A finished RSPD Waspalloy insert and a typical cross section showing the fine grained structure after all processing steps are shown in Figure 5. Finishing included the attachment of a welded tip cap and a brazed- on foot assembly. After the brazing operation, holes were lazer drilled through the leading edge and two sides

of the insert.

The tensile strength of RSPD Waspalloy was determined to be about equivalent to that of forged material as shown in Figure 6. In stress rupture, the RSPD Waspalloy was about two Larson-Miller parameters weaker

than forged Waspalloy. This is typical of fine grain material and is considered satisfactory for this insert

application. In tensile ductility the RSPD Waspalloy behaved much like forged material to about 1000°F.

Above 1000°F, the fine-grained RSPD material began to show evidence of reduced ductility,reaching a minimum at about 1600°F. Above 1600°F the tensile ductility increased. This behavior is also typical of cast

superalloys exposed to elevated temperature and is believed to be caused by the presence of oxygen in grain

boundaries.'®’ Component bench testing of the finished RSPD blade inserts showed them to be acceptable for

Figure 5 Finished RSPD Blade Cooling Insert

Stress 1000 PSI

200

160

120

80

40

0

1 D Oq RSpd

^ ^^Waspalloy

Waspalloy

Forging

A. Ultimate

Tensile Strength

—I I i—j i_i I i_

0 400 800 1200 1600

Temperature °F

lOoOOk

Stress 1000 PSI

40

20

10

Waspalloy

Forging

. RSPD

Waspalloy

C. Stress Rupture

‘ J—L

34

P = 38 42 46 48

T (20 + Log t) X

10’^

Stress 1000 PSI

160

120

Percent Elongation

40 -

\ RSPD

Waspalloy

80 - Waspalloy

_ Forging

B. 0.2% Yield

Strength

—I I I I I t i 400 800 1200 1600

Temperature °F

30

20

10

RSPD ^

Waspalloy

Waspalloy

Forging

D. Tensile

Ductility

\

0^1

400 800 1200 1600

Temperature ° F Figure 6 Comparison of RSPD Waspalloy and Forged Waspalloy Mechanical Properties

655
The design flexibility afforded by the RSPD process can be illustrated by considering a gas turbine

combustor liner. A typical annular combustor is shown in Figure 7. At the upstream end, it consists of a dome

through which fuel is introduced and a portion of the airflow passes. The downstream combustion region is contained within an annulus formed by an inner and outer combustor liner. The liners must be designed

utilizing a limited amount of cooling flow for protection against the hot gases. The cooling flow is introduced

through cooling rings spaced along the liner shell. Large axial thermal gradients between the hot liner shell and cooling ring are generated by the geometry and boundary conditions as illustrated in Figure 7C. Thermal

gradients also exist circumferentially and are generally harmonic in nature, being a function of the number of

fuel injection points in the dome. Thus, combustor liner life is limited by cracking due to hoop stress induced

by axial thermal gradients and a combination of cracking, distortion, and oxidation due to circumferential

thermal gradients (hot streaks).

Combustor liner materials have generally been limited to those readily available in sheet or forged form,

such as Hastelloy X and HS188. The RSPD process permits designs utilizing graded, multi-material

configurations. Opportunities are available to combine materials of differing strength, modulus, and thermal

coefficient of expansion in unique fashions to address design requirements. For example, the designer may

consider different coefficient of expansion materials in thermal gradient fields to reduce thermal stresses and thereby increase life. An example of this concept would be to fabricate the hot liner panel from Hastelloy

X, while the cooling ring would be fabricated from A286 which has a higher coefficient of thermal expansion

and greater low temperature (<1200° F) strength.

550 °F

C. Liner Cooling Ring Hoop Stress

l^roblem Due to Thermal Gradient

Figure 7 Combustor Liner Cooling Ring

Two concepts for manufacturing a bi-metal combustor liner by the RSPD process are shown in Figure 8. Method A is applicable to multiple material configurations, such as the Hastelloy X and A286 combination

noted above. Hastelloy X sheet would be formed on a reusable mandrel and subsequently sprayed locally

with a circumferential ring of A286. These local regions would then be machined to produce the cooling ring

configurations. This concept is limited to materials available in sheet form. Method B illustrates a concept

applicable to both multiple material and single material configurations. A reusable mandrel tool would be

lined by a thin, low carbon steel sheet. The material(s) would then be sprayed to a near net shape contour.

The thin low carbon steel sheet would permit removal of the sprayed liner from the reusable mandrel and

could be subsequently leached or machined away. The liner form would then be final machined. Both

manufacturing concepts appear viable and have been projected to be cost competitive in production with

liners machined from forgings.

656
Method A

• Single Material Spray-Up

• Restricted to Ductile

Sheet for Hot Surface

Material

Method B

• Like A but Permits

Selection of Hot Surface

Material

Figure 8 RSPD Combustor Manufacturing Methods

The same basic concept could be used for the fabrication of monolithic combustor liners by RSPD utilizing

high temperature, high strength materials. Many high temperature superalloys were developed for turbine

blade or vane applications and are available in cast form only. They are unsuitable for many structural and

flow path configurations requiring thin cross-sections because of casting process limitations. The RSPD

process permits fabrication of components with thin cross-sections from high strength superalloys. Benefits can be obtained from the use of available alloys, such as Rene'80, and from alloys specially formulated to take further advantage of the rapid solidification rates inherent in RSPD.

The relative tensile and stress rupture strengths of wrought Hastelloy X and RSPD Rene'80 are shown in Figure 9. Of particular interest is the excellent resistance to thermal fatigue demonstrated by high strength

RSPD deposits reported elsewhere*®’

Ultimate Tensile Strength 0.2% Yield Strength Stress Rupture

Temperature (F) Temperature (F) P = T (20 + Log t) x 10 ^

Figure 9 Comparison of RSPD Rene'80 and Wrought Hastelloy x Mechanical Properties

657
RSPD combustor liners offer benefits in terms of increased life and higher operating temperatures. The

increases in life of either a graded material Hastelloy X/A286 (bi-metallic), monolithic MAR M509, or Rene'80 liner relative to a conventional Hastelloy X liner are shown in Figure 10. For this analysis, a relatively severe circumferential hot streak gradient and conventional material properties were assumed. If the

circumferential gradient was eliminated, the life improvement obtained from the bi-metallic design could be

approximately 2.5X that shown because the bi-metallic configuration significantly reduces the hoop stress resulting from axial temperature gradients. Thus, tailoring the material properties to the design problem can

be a powerful design tool. Direct material substitutions such as RSPD MAR M509 or Rene'80 also indicate

significant liner life improvements over wrought Hastelloy X. There improvements can result in significant

savings over the life of the engine. In engine applications where Hastelloy X combustors are not life limited,

improved material substitutions could permit an increase in allowable liner temperature for constant life, as

illustrated in Figure 11. This translates directly into reductions in specific fuel consumption and increases in available power for the gas turbine.

In order to simulate the fabrication and thermal testing of an RSPD combustor configuration, a series of

four rings about ten inches in diameter and three inches long were fabricated from Hastelloy X, Rene'80,

MAR M509, and a bi-metal system, Hastelloy X/A-286. The metal rings were sprayed as cylinders on a steel ring mandrel mounted on a wheel. The sequence of manufacture is summarized in Figure 12 and the

deposition results are given in Table 1.

Detailed measurements were made of the mandrels before and after deposition of the high strength alloys.

These measurements indicated that, while the rings did not lose their concentricity during deposition, the

inside diameter changed as a result of solidification and cooling stresses imposed by the deposits. The

maximum dimensional change was about 0.1 inch reduction in diameter.

After spraying, the mandrels were dissolved using nitric acid and the deposits were heat treated at 2280°F

for four hours. The Rene'80 ring was given the normal aging treatment (16 hrs. @ 1550°F), and then

machined to final cross section shape. All the rings were single-point turned using a carbide bit. The

machining characteristics of the Hastelloy X and the A286 were nearly indistinguishable and were similar to their wrought counterparts. The MAR M509 and the Rene'80 pieces, while requiring more energy, machined

like wrought IN 718 alloy.

During the course of machining the bi-metal Hastelloy X-A286 ring, the A286 layer was inadvertently

removed. This necessitated a respray to replace the A286 deposit. The Hastelloy X ring was machined to a new outside diameter configuration and A286 was deposited in a local area on the remachined ring

corresponding to the position of the A286 section in the final part. This resprayed ring had less distortion as a

result of deposition of the A286 than was experienced with the prior rings. The final machined cross section

in Figure 12D shows the location of the Hastelloy X and A286 deposits in the simulated combustor ring.

Mandrel

Material

Deposited

Weight

Gain

lbs.

Efficiency

(%)

Deposit

Thickness

(Inch)

Steel Rene'80 5.6 43 0.26

Steel MAR R509 6.8 45 0.30

Steel Hastelloy X 8.8 49 0.35

Steel Hast X/A286 16.4 48 0.72*

Machined

RSPD Hast X

A286 5.5 73 0.53

*(—0.36 mils Hast X

-* —.36 mils A286)

Table 1 Results of RSPD Simulated Combustor Ring Spray Runs

658
Life Factor Life Factor

Increasing Firing Temperature

Figure 10 Relative Combustor Liner Life Comparison of Wrought Hast X to Hast X/A286,

MM509, and Rene'80 Assuming a Circumferential

Hot Streak

Figure 11 Increasing Allowable Firing

Temperature by Substitution of Hastelloy X/A286

MM509 or Rene'80 for Hastelloy X in Combustor

Liner

C. RSPD Hastelloy X/

A286 Ring AsDeposited on

Mandrel

A. 10" Diameter Steel

Mandrel

Hastelloy X

A286

E. Microstructure of

Hastelloy X/A286

Bond Line

B. Preheating of

Mandrel Prior to

Deposition

RSPD A286

RSPD Hastelloy X

D. Hastelloy X/A286

Ring Cross Section

Figure 12 Fabrication of Bimetal Simulated Combustor Liner Cooling Ring by RSPD

The microstructures of the fully processed Rene'80 and MAR M509 deposits are shown in Figure 13, and

the densities are given in Table 2. In all. cases, the resultant deposit densities were near theoretical after heat

treatment. Modest (1.0-1. 5%) losses in chromium as a result of evaporation during spraying tend to increase

the apparent density by 0.1 - 0.2%. The accuracy of the density measurement is about ± 0.2%.

659
A. RSPD Rene' 80 B. RSPD Mar M509

Figure 13 Micro Structure of RSPD Rene'80

To explore RSPD processing for modestly sized structures, a simulated combustor inner liner was

fabricated from Hastelloy X. The overall dimensions of the completed ring were 7.5 inches inside diameter at the large end, 6.0 inches at the small end, and 4.2 inches in length. Normally, the ring is machined from a forging preform and is subsequently welded to the combustor dome at one end and to a flanged section at the other. A mandrel conforming to the inside contour of the liner was machined from low alloy steel and

welded to a shaft. The cleaned mandrel was mounted in the low-pressure plasma system and preheated to approximately 1800°F using transferred arc. Reverse transferred arc cleaning was used to enhance bonding

of the Hastelloy X to the substrate. The mandrel was sprayed with Hastelloy X and allowed to cool, then the

surface was machined to form the core of the structure. Close-fitting low alloy steel bands were located at the

cooling ring nuggets and a second layer of Hastelloy X was applied to the ring. After the second deposition,

the part was machined to expose the steel bands and the inner and outer mandrels were leached out using

nitric acid. The liner was then given a 2 hour 2150°F vacuum heat treatment, at which time one of the panel

overhangs separated. Metallographic sections indicated that the separation resulted from porosity generated

by low angle deposition at the bond during the second spray-on cycle. An overall view of the combustor ring

is shown in Figure 14.

Deposit

Material

Rene'80

MAR M509

Hastelloy X

Hastelloy X

with A286

Nugget

Table 2 RSPD Simulated Combustor Rings - Heat Treated Densities

This paper has provided an overview of RSPD including a brief description of the process, some of the

factors to be considered in design and fabrication, and examples of aircraft gas turbine engine components

which could benefit from this approach. Work is continuing on RSPD plasma gun, process and materials

development. Further advances in understanding in these areas will firmly establish the base technology

required for component applications. The coupling of innovative design and manufacturing approaches

utilizing the unique features of RSPD are expected to significantly impact the gas turbine engines of the

1980’s and 1990’s”^'.

After Heat Treatment

Deposit Density % of Conventional

g/cm^ Alloy

8.15 99.8

8.80 99.3

8.25 100.4

8.21 99.9

660
Figure 14 Small RSPD Hastelloy X Combustor Liner

SUMMARY

Recent progress in the development of RSPD technology has demonstrated the feasibility of this process

for the fabrication of aircraft gas turbine components. Small, thin wall parts have been made which exhibit

attractive features such as improved formabiiity and interior surface finish. Larger rings of difficult to process

superalloys and combinations of material have been direct sprayed to near net shape. This flexible

manufacturing process is emerging as a generic fabrication approach to be compared with casting, forging

and conventional PM consolidation. The General Electric Company is pursuing this rapid solidification

technology for potential application to a wide range of engine components.

ACKNOWLEDGEMENTS

Portions of this work were conducted as part of AFWAL Contract F33615-81-C-5156, U.S. Air Force

Systems Command Materials Laboratory, under the direction of Dr. H.L. Gegel, and sponsored by the Defense

Advanced Research Projects Agency. At DARPA the cognizant officer was Lt. Col. L.A. Jacobson. The authors

appreciate the contributions of many other GE employees, particularly R.A. Amato, M.R. Jackson, L. Lievestro

and J.R. Rairden.

REFERENCES

1. Gianinni, G., The Plasma Jet, Sci. Amer., August 1957.

2. The Plasma Spray Handbook, Metco Inc., Westbury, N.Y., 1963.

3. Mash, D.R. and Brown, I. MacP., Materials Engineering Quarterly, P. 18, February 1964.

4. Bittence, J.C., P. 108, Machine Design, April 1973.

5. Muehiberger, E., Proceedings 7th International Thermal Spray Conference, London, P. 245, 1973.

6. Okada, M. and Maruo, H., British Welding Journal 1^8, P., 321, 1968.

7. Rairden, J.R. and Jackson, M.R., Private Communication.

8. Jackson, M.R., Rairden, J.R., Smith, J.S., Smith, R.W., “Production Of Metallurgical Structures By Rapid

Solidification Plasma Deposition”, Jrnl. Of Metals, P. 23, November 1981.

9. Drazkiewicz, T., Proceedings 9th International Thermal Spraying Conference, P. 244, 1981.

10. Smith, R.W., Rigney, D.V., Smith, J.S., and Rairden, J.R., “Plasma Deposition Effects On Structures and

Properties of RSPD Alloys”, this conference proceedings.

11. Jackson, M.R., Smashey, R.W., and Peterson, L.G., “The Mechanical Behavior of RSPD Materials”, this

conference proceedings.

12. Sprague, R.A., "Material and Process Impact On Aircraft Engine Designs of the 1990's”, ASME Publication

82-GT-278, P. 7-8.

661
RECENT APPLICATIONS OF ULTRA RAPID SOLIDIFICATION

TECHNOLOGY (URST) MICROCRYSTALLINE AND AMORPHOUS POWDER MATERIALS

Julian H. Kushnick and Dulari L. Sawhney

Allied Corporation, P. 0. Box 1021R, Morristown, New Jersey 07960

ABSTRACT

Amorphous Ni-Mo-B based alloy powder materials are presently beingz-manufactured

by Ultra Rapid Solidification Technology (URST*) Cast/Quenched at 10°°C/sec. for

a wide variety of commercial applications. Some of these powders are consol i dated /by

hot isostatic pressing into microcrystalline solid blanks for previously described'^"

copper alloy extrusion dies, cutting tools, and aluminum die casting molds. These

applications will be briefly updated. Newer applications are focused on the use of

amorphous powder materials in metal-bonded abrasive wheels, corrosion and wear resistant hardfacing and as an alloy with tungsten for consolidation into experimental high

specific gravity kinetic energy armor penetrators. In each instance, the amorphous

characteristics of the URST powders improve performance.

Amorphous consolidation of METGLAS** alloy magnetic powders and ribbon is being

developed because of the magnetic core properties obtained at high packing factors.

Most recently, ultra rapid solidification is being investigated for the manufacture of

aluminum microcrystalline alloys suitable for consolidation to high specific strength

aircraft components. High temperature resistant alloy additives have been incorporated

in these systems to improve the structure produced after the necessary thermal consolidation processing.

INTRODUCTION

As many of you are aware. Allied Corporation has been a world leader in Rapid Solidification

Technology for more than 12 years. We have researched, developed, and commercialized unique

Metglas amorphous metal alloy products including low loss transformer core materials, magnetic

shielding and ductile brazing foils for gas turbine fabrication. All of these Metglas amorphous

products are made by the rapid solidification of molten metal alloys (Figure 1) into ductile

ribbons or filaments at cooling rates >1,000,000°C. Quenched at these rates, these special

molten metal alloys do not tend to segregate or crystallize and remain amorphous —a frozen liquid

or solid glass.

Technology to produce vast quantities of these amorphous materials has been developed with

the assistance of the Electric Power Research Institute (EPRI) of Palo Alto, Ca. An Allied

Corporation/EPRI pilot plant has been constructed which can manufacture Metglas transformer core materials (2 in., 4 in. or 7 in. wide ribbon) at rates of 3,000 ft. /minute for hours while con- tinuously winding the product. These Metglas core materials have been supplied to major transformer companies who are developing technology for the cost effective production of low core loss

utility distribution and power transformers^. Continuous casting technology for the manufacture

of URST materials is perhaps best exemplified by the facility recently put on line -at Metglas

Products, Allied Technologies Company, in Parsippany, New Jersey. In this facility (Fig. 2) raw materials are melted in an induction melting furnance, the molten metal is transferred to an accumulator/feeder from which it is control fed through a casting head that is capable of depositing a 4 inch wide stream of molten metal in contract with a rotating mold. The molten stream is

solidified within a millisecond to a ribbon approximately 0.0015 in. thick. The dimensions of the

ribbon produced are monitored and adjusted by a computerized feedback system. There is a threadup and winding machine to continuously wind the product at the end of the production machine. This

system is designed to manufacture up to 25,000 tons/year of Metglas alloy amorphous transformer

core material s^. *Trademark of Allied Corporation

**Regi stored Trademark of Allied Corp.

662
FIGURE 1

THE RAW MATERIALS

FIG. 2

663
The Consolidated Metal Products (CMP) group has adapted this versatile Metglas technology for a different application—powder metallurgy via ultra rapid solidification technology.

We have developed low cost production technology to pulverize the URST materials so that we

can market them as amorphous powders (Fig. 3)^ or consolidate them into new and useful bulk

forms and shapes. URST powders are unique because each particle has identical composition

and is amorphous. In the powder products of most so called rapid solidification processes

(i.e. gas or centrifugal atomization) small particles are sometimes amorphous, medium size

microcrystalline and large particles are usually segregated and dendritic. Each URST particle

is homogeneous regardless of size. The CMP powders and subsequent consolidated products are marketed under the DEVITRIUM trademark^.

COPPER EXTRUSION DIES

In a typical bulk shape application, URST powder produced of a Ni-Fe-Mo-B alloy

(DEVITRIUM*** alloy 7025) is consolidated by hot isostatic pressing (HIPping) at 1100°C and

12,500 psi pressure into microcrystalline solid bar for copper alloy extrusion dies^. During

the HIP cycle, the powder is transformed from the amorphous to the microcrystalline state. Submicron-sized boride particles form at the grain triple points' (as illustrated by the

schematic of grain structure (Fig. 4). Conventional, slower cooled alloys of similar com- position would consolidate as crystals of varying large grain sizes surrounded by a eutectic

boride network, making the material too brittle for most engineering uses.

Devitrium alloy 7025 extrusion dies (Fig. 5) are used in the primary reduction of hot

(850°C) billets of copper and brass alloys to redraw rod and shapes. The principle advantage

of these dies versus dies of steel of high cobalt alloys has been their long service life

(up to 300 pushes have been reported by a major copper producer) and the good surface finish

of the extrudate. This reduces grinding and finishing costs prior to subsequent redraw operations. Dies of DEVITRIUM alloy 7025 have been supplied to more than six extrusion firms and

are utilized when long run campaigns of extruded product are required. Unfortunately,

because of the present recession this industry has not had too many opportunities to use their

DEVITRIUM alloy 7025 dies in recent times.

CUTTING TOOLS

Another application that has been previously reported is cutting tools®. DEVITRIUM

Alloy 3065, a microcrystalline Ni-Mo-B composition in its consolidated form (Figure 6) is

being evaluated for the production of premium wear and chip resistant cutting tools and inserts.

Cutting performance is retained at high rates of speed (up to 500 surface feet per minute)

with tool life in many applications equivalent to uncoated tungsten carbide materials.

DEVITRIUM Alloy 3065 end mills produced by three major tool companies are scheduled for

extensive trials with aluminum and titanium. Hobbs have been fabricated with DEVITRIUM

cutting surfaces for high speed gear production.

INSERT FOR ALUMINUM DIE CASTING MOLDS

Die casting molds for Aluminum are usually machined of alloy H-13, a standard high performance steel that has generally good heat resistance. The life of an H-13 mold in a critical gate area where the hot die casting alloy enters the mold is approximately

75,000-125,000 cycles or "shots". Die casting molds typically fail by erosion and thermal

fatigue. Cyclic thermal stresses cause heat checking cracks and fissures to occur in the

mold surface. The checking cracks gradually worsen until they prevent good release of the

molded part and cause intricate areas of the mold to break. Consolidated DEVITRIUM Alloy

7025 (Figure 6) inserts used in die casting molds have demonstrated service lives of over 250,000 shots and are still in production^. The improved resistance to thermal fatigue or heat checking of Consolidated DEVITRIUM Alloy 7025 is attributed to the following:

1. Consolidated DEVITRIUM Alloy 7025 is structurally homogeneous and isotropic. The

microcrystals are not organized into large crystalline matrices with anisotropic characteristics that can lead to varying stresses. With uniform distribution of thermal stresses, cracks

are less likely to occur during thermal expansion of the insert surface.

***Trademark of Allied Corporation

664
FIG. 3 FIG. 4

FIG. 5

FIGURE 6 Mechanical Properties of Consolidated DEVITRIUM** ALLOYS

DVTM Alloy 7025 DVTM Alloy 3065 Nis3Mo26^®9®2 Ni54Mo^4B02

DENSITY (G/Cm3) 9.06 9.33 HARDNESS ( HRC

47-50 60-62 TENSILE STRENGTH (KSI) 220-220

(0.2%) YIELD STRENGTH (KSI) 175-185 ELONGATION (%) 2-5 1% IMPACT STRENGTH ( FT-LBS .

40-50 2-5 (UNNOTCHED CHARPY)

665
2. DEVITRIUM Alloy 7025 as a consolidated material is a tough, elastic, high strength

composition that is less likely to be torn apart by the thermal strains encountered. These

properties are maintained at elevated temperatures.

3. The thermal strain in the mold material during the die casting operation is a product

of the temperature difference between the bulk of the mold substrate and its surface and the

coefficient of thermal expansion. The thermal expansion of DEVITRIUM Alloy 7025 is approximately 1/3 lower than that of H-13 steel. This minimizes the cyclic strain and fatigue.

METAL-BONDED ABRASIVE WHEELS

DEVITRIUM Alloy 3065 and 7025 are entering use as metal binders for diamond and other

abrasive materials in oil well and mine drilling bits, masonry saws, grinding wheels, etc. Alloy 3065 and 7025 powders are amorphous and contain no crystalline structure. This is

readily accomplished by ultra rapid solidification technology. Typically the DEVITRIUM

powders are mixed with the abrasive material, (i.e. diamonds or silicon carbide), and are then hot pressed in graphite molds. Consolidation at 950°C and 5,000 or 7,500 psi results

in about 97% of theoretical density. This temperature is maintained for relatively short

periods of time to avoid degradation of the industrial diamonds contained in the matrix.

During the hot pressing operation the internal structure of these powders changes from amorphous to microcrystalline at temperatures about 550°C.

Atomic movement and rearrangement occurs during the crystallization phase. There is reason to speculate that this atomic movement may generate clean surfaces and enhance the

adhesion to the abrasive powder mixture. During the hot press consolidation there is rapid

press ram movement as the amorphous materials consolidate and become higher density microcrystalline structures. The finished pressed matrix product is hard and wear resistant with

substantial retention of hardness at elevated temperatures. This helps to increase the cutting

action of the products. The pressed matrix microcrystalline materials have a substantially

higher softening temperature than the original amorphous alloys. This increases the feed

rates and temperatures at which the grinding tool matrix can be utilized.

An offshoot of the metal binder application has been the use of DEVITRIUM Alloys 7025

and 3065 as readily prepared hot pressed molds for the fabrication of other abrasive matrix

materials. Hot pressed consolidated DEVITRIUM molds do not wear rapidly and provide a high

yield of pressed parts. The finish of the pressed parts is significantly better than that

obtained from graphite and other conventional die materials, probably because of the resistance of DEVITRIUM Alloys to wear and galling.

CORROSION RESISTANT PLASMA SPRAY HARDFACING

Amorphous DEVITRIUM Alloys have been developed for a wide range of corrosion resistant

arc plasma hardface coatings. These coatings can be applied by a wide range of plasma spray

equipment (Fig. 7) to achieve uniformly dense, well adhered amorphous coatings of unusual

hardness. The date (Fig. 8) shows that the corrosion, wear resistance and nongalling characteristics of DEVITRIUM coatings are good. Ball and Cup forming tests and other bending tests

indicate that the ductility of the amorphous deposited DEVITRIUM Alloy 7025 plasma coating is substantially better than that of crystalline analog materials of Ni-Mo-B alloy systems.

The DEVITRIUM coatings do not tend to crack or flake from the substrate when they are deformed.

Amorphous DEVITRIUM hardfacing powders have also been used in plasma transferred arc

(PTA) applied coatings and have demonstrated excellent wear resistance. Good coatings without

cracks are routinely achieved. Several extruder "crews and large gate valves PTA coated with

DEVITRIUM Alloy 7025 are undergoing performance trials. The nongalling characteristics of

the DEVITRIUM Alloys is of particular interest in the gate valve application.

KINETIC ENERGY ARMOR PENETRATORS

A program for the development of high density kinetic energy armor penetrator materials

is underway with tungsten-nickel -iron-boron alloys and ultra rapid solidification technology.

CMP has produced URST (Figure 9) materials containing tungsten in the range of 40-65 weight

percent. Alloys containing 40 weight percent of tungsten have been cast completely amorphous.

This demonstrates the beneficial effect of URST in extending the solid solubility limits by

666
FIGURE 7 DEVITRIUM " Alloy 7025

Plasma Spray Powder Coating

Thickness 5-30 mils

Density % of Theoretical 98%

Microhardness VHNioo 990

Bond Strength 1 2,000 psi

Percent Amorphous 80%

Substraight: Mild steel, grit blasted, degreased

Plasma Spray Gun: Baystate PG 100; Metco 3M,

7M; Plasmadyne SG 100

Powder Size: Minus 80 mesh.

FIGURE 8

WEAR AND CORROSION RESISTANCE OF ARC PLASMA SPRAY COATINGS

Coating Alloys

Ncminal Conp. W% DVm 7025 Ni53M03gFegB2

DVTM 3065 Ni54M044B2

DVTM 7060 Ni54”036Ct8®2

Stellite 6B* COBalCr28^4^1.

Corrosion Rate of Arc Plasma

Sprayed Coatings

05O°C -inils/yr.- l%HCi io%h^so2

13

4 20

6 8

Metco 501''' NiBalMo30Cri2Si2. 75®2 .5^.75 Triballoy T-700* NiBal^32^'^16Si3.4 Co-FeO.0;C£.08

Hastelloy C* NiBal^*^17^17^®6'^5 ®

Selfmated

Galling

Threshold Stress 103pSI

Abrasive Wear Vol, loss ASTM G-65 (sand/rubber

wheel) imi^

Abrasive Wear

ASTM D2714-68 (Falex/block & ring) Vol. Loss of Sample nm^ (Blodc)

20 7-10 0.05

10-30 9-10 0.045

7

50-60 9 0.03

8 0.08

50-60 9

* T. M. of Cabot Corp. + T. M. of METCO, INC.

667
FIGURE 9 MECHANICAL PROPERTIES OF CONSOLIDATED URST TUNGSTEN ALLOYS

40 Wt. % 65 Wt. % Tungsten Tungsten

DENSITY (G/CM^) 10.75 13.1

HARDNESS (HRC) 35 40

TENSILE STRENGTH (KSI) 136 162

(0.2%) YIELD STRENGTH (KSI) 105 142

(0.2%) YIELD STRENGTH (KSI) 97.5 150

(IN COMPRESSION)

TRANSVERSE RUPTURE STRENGTH (KSI) 285.5 310

IMPACT STRENGTH (FT-LBS.) 25 10

(UNNOTCHED CHARPY)

OPTICAL MICROSTRUCTURE OF CONSOLIDATED URST TUNGSTEN ALLOYS

40% W

I I 20/AM

65% W

668
comparison with the equilibrium solid solubility of cv) 27 weight percent of tungsten in

nickel. Ultra rapid solidification of alloys with 65 weight percent tungsten produced a composite material, comprised of a uniform dispersion of fine tungsten particles (5-7 micron)

in an essentially amorphous matrix.

Isothermal and isochronal studies of the as cast tungsten alloys, reveal that thermal

treatment precipitates extremely fine tungsten particles in the supersaturated Ni-Fe-W

solid solution matrix. The characteristic feature of both the alloys (i.e. with 40

percent tungsten and 65 percent tungsten) is a very stable microstructure. No substantial

grain growth is observed even with long time exposure at temperatures as high as 1200°C.

Figure 9 illustrates the pertinent microstructural features of URST tungsten alloys and their

physical properties. In addition to the fine and thermally stable microstructure, URST

prevents the formation of any brittle intermetal lie phases, normally associated with either

Ni-W or Fe-W binery systems. Alloys with densities of 11-13.5 has been produced via our

process. Work to produce yet higher densities is in progress.

METGLAS MAGNETIC CORE MATERIALS

Amorphous consolidation of METGLAS magnetic powders is being developed so as to preserve

the unique magnetic properties of the parent Metglas materials in bulk shapes. Amorphous

toroidal cores have been consolidated from various Metglas magnetic materials at bulk

densities ranging from 85% to 95% of theoretical. These cores have been annealed to provide

potential magnetic core materials for such applications as regulator cores in switch-mode

power supplies and stators for small high speed actuator motors.

Technology is also being developed to consolidate layers of thin METGLAS magnetic core ribbon into thick strip bars while retaining the Metglas amorphous structure. METGLAS alloy

POWERCORE* strip bars with good magnetic properties would facilitate the construction of large

transformers and motors. Power transformers typically contain from 12,000 to 300,000 lbs.

of magnetic core material. Fabrication of stacked cores of the size required to provide that

much magnetic flux path sometimes involves stacking the core steel as high as 6 to 8 feet.

This would be very difficult with Metglas ribbon approximately 0.0015 in. thick. CMP has

experimentally produced POWERCORE strip bars 0.070 in. thick with packing factors as high

as 95% that could be used to fabricate large cores. We are presently developing process

technology to produce POWERCORE strip in large quantities with consistent quality.

ALUMINUM MICROCRYSTALLINE ALLOYS

Within the past several years ultra rapid solidification technology has been used for

the manufacture of aluminum based microcrystal 1 ine alloys cast at exceptionally high quench

rates, sometimes in the order of 10°°C per second. The principle advantage of these high

quench/casting rates is that the crystalline structure is so fine that it cannot be discerned

by ordinary light microscopic techniques, (generally regarded as type A structure).^ The

exceptional homogeneity of these microcrystalline materials is retained (with some size growth)

through consolidation processes (typically by vacuum hot pressing and extrusion) into high

specific strength and/or high temperature resistant fabricated aircraft components. As an integral part of these programs alloy additives are incorporated into these aluminum systems

to improve their structure and prevent the formation of precipitate particle free zones during

the thermal processing for consolidation and extrusion.

High specific strength alloys have been developed with greater quantities of low density

alloy elements incorporated than are presently feasible with ingot metallurgy. These alloy

systems can provide a cumulative weight savings of approximately 10-15% in major aircraft

components. In a similar manner, high temperature resistant alloys have been developed that

have maintained high use strength and exceptionally fine microcrystalline structure after

prolonged thermal exposure at temperatures in excess of 450°C. These alloys could also result

in significant weight savings in aircraft fabrication by reduction of Aluminum part size or displacement steel components.

We believe it feasible to manufacture ultra rapid solidification microcrystalline

*Trademark of Allied Corp.

669
aluminum alloys of both the high specific strength and high temperature resistant systems

through integrated production casting systems similar to those capable of going from alloy

raw materials to powder in line described earlier.

SUMMARY

Allied Corporation has developed a proprietary ultra rapid solidification technology

process to manufacture:

1. Materials that are uniform and homogeneous in alloy composition and physical

structure —amorphous or microcrystalline as required.

6 8

2. The process can be operated at cast/quench rates of 10 to 10 °C/second.

3. The URST process combined with pulverization technology provides us with high

yields of products in terms of alloy composition desired powder and particle size.

4. The process is inherently rapid and low in cost to operate.

5. The process is scalable and consistant as has been demonstrated by production of our METGLAS amorphous metal alloys.

670
REFERENCES

[1] S. Pass and J. Dickson, in "A Breakthrough in Powder Metallurgy at Allied Chemical

Corporation", Workshop on Critical Materials, U. S. Dept, of Commerce, NBS,

Gaithersburg, Md., Feb., 1981.

[2] "Amorphous Metals: Cutting Losses in the Core" EPRI Journal, Vol 7, No. 1, Jan/Feb,

1982.

[3] "What if?" Allied Corp., Metglas Products, Bulletin SM, Jan, 1982.

[4] U. S. Patent 4,353,737 Method of Making Metallic Glass Powder From Glassy Alloy.

[5] "Consolidated Metal Products - Introducing a Breakthrough in Powder Metallurgy" - Allied Corp., Nov., 1980.

[6] "New Kind of Alloy Tested in Dies, Cutting Tools" American Metal Market/Metal

Working News, May 18, 1980.

[7] U. S. Patent 4,290,808; Metallic Glass Powders From Glassy Alloys.

[8] "New P/M Material Extends Die Casting Die Life" Precision Metal, Vol. 40, No. 3, March,

1982.

[9] H. Jones, Materials Science & Eng., Vol. 5, 1969, p. 91.

671
SHOCK COMPACTION OF FERROUS ALLOY POWDERS

T. J. Ahrens, Dan Koska, P. Kasiraj and T. Vreeland Jr. California Institute of Technology

Pasadena, CA 91125

A. W. Hare

Northwest Technical Industries, Inc. Sequim, WA 98382

F. D. Lemkey and E. R. Thompson

United Technologies Research Center

East Hartford, CT 06108

ABSTRACT

Experiments on the shock compaction of ferrous alloy powders have been performed using

a 20 mm propellant gun system and explosive charge techniques. Spherical AISI 9310 powders were produced by various rapid solidification rate processes including inert gas, soluble

gas, and centrifugal atomization. The powders were inspected microstructural ly in TEM and

by X-ray diffraction and found to be remarkably similar.

Propellant gun, dynamic compaction of 40% initial porosity, 100pm diameter powdered

9310 alloy produced consolidated specimens with ultimate tensile strengths of 0.7 to 1.3

GPa. Impact with stainless steel projectiles launched at speeds from 1 to 1.6 km/sec

caused initial shock pressures in the compacted powders of 7 to 16 GPa and final shock

pressures of from 19 to 34 GPa. Recovered samples demonstrated diamond pyramid hardnesses

of 370 to 581 kg/mm^ over this pressure range. Metallographic observations showed that the

powder particles in a 2 mm thick specimen are deformed by ~50 to 200% along the direction

of shock compression. Individual grains displayed evidence of inhomogeneous thermal deposition. Thermodynamic analyses indicated that for an initial distension of 1.7 for iron, a portion of the sample will have been melted upon shock recovery from direct shocks of from

7 to 16 GPa.

Flake powders from Markomet (Fe-l.35C-4.25Cr-4.5Mo-4.0V-5.75W-0.65B) and spherical

9310 steel powders were explosively compacted within the same evacuated thin wall cylindrical container. Although complete consolidation was not achieved using an ammonium

nitrate and fuel oil charge (p = 0.443 gm/cm^), improved densification was achieved starting with flake rather than spherical powder geometry.

Introduction

In order to preserve the nonequilibrium nature of rapidly solidified powders and avoid extensive

interactions between powder mixtures, increasing attention is being directed to the examination of

shock consolidation techniques to form bulk structures. Exploratory shock compression and recovery

studies [1-6] have demonstrated irreversible compaction for shock waves propagating through porous

materials with amplitudes in the range of 10-20 GPa and durations of 0.1-1 sec (i.e. energy densities

of -10-100 KJ/cm^). Microscopic examination of the irreversibly compacted, initially fine-grained

samples demonstrate that the irreversible work input to the porous material is nonuniformly deposited.

What appears to be happening is that shock induced heating takes place along particle surfaces

672
presumably via strong surface frictional and deformational processes. After the shock has passed and

the material is released to zero pressure, the thin molten surface layer is rapidly quenched via

conduction from within the particle interiors. Depending on the width of the melted zone, cooling

rates from lO^-lOio K/sec are inferred for melt layer thicknesses of 100 to 0.2um in metals [7]. In this way, the nonequilibrium nature of the rapidly solidified powders may be preserved.

The overall objective of this study is to exploit rapid solidification processing (RSP) technology

for the purpose of fabricating an advanced helicopter gear with improved high temperature (~600°F)

capability. It is believed that RSP technology applied to this situation offers the best prospects

for success for two main reasons. First, rapid solidification of powders is a proven method for

obtaining ultra-refined microstructures, with reduced microsegregation and extensive supersaturation

of compound-forming elements. This is a prerequisite for achieving high strength without impairing

fatigue and fracture toughness properties. Second, dynamic compaction is the only method currently

available, albeit incompletely developed, for making fully consolidated parts from powder without

incurring unacceptable coarsening of the microstructures and chemical reactions between prealloyed

powders of substantially different compositions.

Experimental Program

The dynamic compaction program is divided into two areas, (1) propellant gun and (2) explosive

charge techniques. Initial results on the former have been encouraging and will be published later

this year [8,9]. The ferrous alloy powders compacted in this program were produced by various rapid

solidification techniques. Spherical AISI 9310 (Fe-0.10C-0.65Mn-0.25Si-l.39Cr-0.13Mo-3.2Ni) powders

were atomized by inert gas, soluble gas, and centrifugal processes. The powders were inspected microstructurally by TEM and X-ray diffraction and found to be remarkably similar. All size distributions

contained the assumed high temperature ^-Fe phase as the major constituent, with minor amounts of

f.c.c. Y-Fe and carbide. The cubic lattice parameters for the delta and gamma phases were calculated

as 2.864 and 3.58iA respectively. Actual freezing rates of NiAl-Cr eutectic powders produced by the same RSP techniques were used to calibrate liquid/solid interface velocities as is discussed elsewhere

in these proceedings [10]. These experiments showed the expected dependence of a-Cr interrod spacing

on powder size. Table I, indicating that coupled eutectic growth occurs to -700 mm sec"^. The average

freezing velocities of argon atomized powder was observed to be equivalent to centrifugal ly atomized

powder for equal particle size. Flake powder supplied by Marko Materials of Markomet^M 3.11 (Fe1 .350-4. 25Cr-4.5Mo-4.0V-5.75W-0.65B) was prepared from melt spun ribbon which was subsequently ball

milled to form discrete flake particles.

o Table I. Or Diam. (A) Measurements, x , and v Calculations, from RSP

NiAl-Cr Eutectic Powders

source 70-300/im 44-70/im 37-44/i m <37/tm

RSR

He atomized

d = 263-1052 A range d = 631 A

X =991 A

v~0.9mmsec '

d= 263-789 d=526

X=825

v-1.2

d =105-631 d = 421

X =660

v-1.9

d = 105-421 d=263

X =412

v-4.8

HMI

Soluble gas

atomized

789-1474 d=947

X =1486

v-0.4

421-789 d = 684 X=1074

w-0.9

421-789 d=526

X =826

v-1.2

105-263 d=368

X =577

v-2.7

AMDRY

Argon

atomized

492-1312 d=820

X =1287

w-0.5

263-630 d=526

X =825

v-1.2

420-789 d = 420 X=660

v-1.9

<105-421 d=263

X =412

v-4.8

Avg. values

673
Propellant Gun compaction

A Caltech 20 mm propellant gun was used to impact a 2.5 to 5 mm thick stainless steel projectile

at speeds of 0.96 to 1.8 Km/sec to compact AISI 9310 steel powder. All specimens were initially

pressed to a distention of 1.7 or a volume some 1.7 times that of its crystal volume. Measured densities of the compacted samples were in the range of 7.84 + 0.15 Mg/m^ which is close to the theoretical value of 7.86 Mg/m^. The transmission electron microstructure of a completely compacted

specimen after ion milling is shown in Fig. 1. Dark field (g=110) transmission microscopy indicated

approximately tenfold delta iron grain size reduction to ~l/2ym.

1.8 mm

a) Light micrograph

1/tm

b) TEM, ion milled, refined grain structure

Figure 1. AISI 9310 (-200, +325 mesh) powder shocked to initial

and final shock pressures of 18 and 37 GPa

Some seven tensile test specimens, cut from recovered pellets oriented parallel to the shock

front were pulled to fracture. The ultimate tensile strength (Fig. 2) was observed to increase from

0.8-1. 3 GPa (120-190 ksi) as the projectile velocity increased from 1 to 1.6 Km/s (initial shock

pressure increased from 7 to 16 GPa). Vicker's diamond pyramid hardness (DPH) tests (Fig. 3) demonstrated a hardness increase from 370 + 57 to 557 + 17 at final shock pressures of 20-24 GPa. The

ultimate tensile strength and DPH of dynamically compacted specimens compared favorably to conven- tionally quenched and tempered AISI 9310, i.e. 1.05 GPa and 300-400 DPH respectively.

To estimate the maximum mass fraction of powdered alloy which could be melted at their surfaces,

it was assumed that the interiors of the powder particles become heated only as much as if they were

grains in the interior of the solid alloy. This assumes no work is expended in deforming the solid

particle. The excess thermal energy is assigned to the heat of fusion and the thermal content of the

melt. If L is the mass fraction of the melted film it is given by:

L = (Et - E'T)/(CpATrr, +

where Ej and E'j are the thermal energy deposited in porous and nonporous powders shocked to the same high pressure respectively. Here ATm is the temperature difference between the melting point, and room temperature, Cp is the specific heat, and H^ is the heat of fusion.

674
700

I Ultimate

tensile

I strength

I (ksi)

240

200

160

120

80

40

0.8 1.0 1.2 1.4 1.6 1.8

9310 (42-72/tm)

9310 forged/quenched and

tempered (design typical)

j I I I

I I I I

Vicker diamond

pyramid hardness

(Kg/ mm2)

100

0

72-300/im

12 16 20 24

4

28 32 36 impact veiocity (Km/s) Pressure (GPa)

Figure 2. Ultimate tensile strength versus

impact velocity

Figure 3. Diamond pyramid hardness versus final

shock pressure

Using the values of 1511 K for the difference in the melting point of iron and room temperature,

0.45 MJ/g K for Cp, and 247 MJ/Mg for Hf^, the post-shock temperatures and the amount of melt for iron

for various initial distentions have been calculated (Fig. 4). By way of reference the present ex- periments were all conducted under shock pressures of from 7 to 16 GPa for initial shock propagation

into the sample of distention, m = 1.7. Hence the maximum percentage of melt is expected to vary from

32 to 70% which is overestimated for the present experiment.

Maximum percent melt upon recovery

for different distention of iron samples

Figure 4» Continuum shock temperature increase versus

initial shock pressure

It appears that in addition to the generation of a thin film of molten material at particle

boundaries by the shock, in order to obtain welding of particles coated with molten material it needs

to be held together under pressure for a sufficient time that the excess heat from the localized

melting can be transported by Fourier conduction to within the cooler interiors of the solid particles.

Explosive Charge Compaction

Experiments to explosive compact rapidly solidified powders have been carried out using a cylindrical arrangement shown in Fig. 5. An objective was to determine the design of an assembly which

yields the desired pressure-time history to achieve full densification of the powders. Initial re- sults have not achieved full density using ammonium nitrate plus fuel oil explosives. The successful

explosive consolidation of 9310 powders involves a number of alloy as well as explosive parameters

675
Detonator

Figure 5. Explosive compaction of spherical 9310 and flake Markomet 3.11

which have just begun to be examined. For example flake Markomet 3.11 and spherical 9310 steel powders were compacted within the same evacuated thin wall cylindrical container. Although complete

consolidation was not achieved, as noted in Fig. 5, improved densification was achieved starting with

flake rather than spherical geometry.

Acknowledgements

This work was performed for the Department of Defense under contract DAAG46-82-C-0042 under the

auspices of the Army Materials and Mechanics Research Center, Dr. Saul Isserow, technical monitor.

References

[1] D. Raybould, 'Proceedings of the 15th International Machine Tool Design and Research Conference',

Macmillan Press, London (1975).

[2] D. Raybould, D. G. Morris and G. A. Cooper, J. Mat. Sci

, 14 , 2523-26 (1979).

[3] D. Raybould, Chapter 51 in "Shock Waves and High Strain Rate Phenomena in Metals", edited by M. Meyers and L. Murr, Plenum Press (1981).

[4] C. Cline, J. Mahler, M. Finger, W. Kuhl and R. Hooper in "Rapid Solidification Processing",

edited by R. Mehrabian, B. Kear, and M. Cohen, Claitor's Press (1977).

[5] C. Cline and R. Hooper, Scripta Met. , 1, 1137 (1977).

[6] S. W. Kieffer, 'From Regolith to Rock by Shock', The Moon , 13 , 301-320 (1975).

[7] D. G. Morris, in "Proc. 2nd Int. Conf. on Rapid Solidification Processing; Principles and Technologies", edited by R. Mehrabian, B. Kear and M. Cohen, Claitor's Press (1980).

[8] P. Kasiraj, T. Vreeland, Jr., R. B. Schwarz and T. J. Ahrens, 'Mechanical Properties of Rapidly

Solidified Iron Based Powders Compacted by Shock Waves', to be presented at APS Meeting, March

1983.

676
[9] R. B. Schwarz, P. Kasiraj, T. Vreeland, Jr. and T. J. Ahrens, 'Shock Wave Velocity and Temperature Measurements During Consolidation of Rapidly Solidified Iron-Based Powders', to be presented

at the APS Meeting, March 1983.

[10] W. J. Boettinger et al . , 'Rapid Solidification Processing', edited by R. Mehrabian et al .

1983.

[11] D. Raybould, J. Powder Metallurgy and Powder Technology , 16 , 9-19, 1980.

677
ULTRASONIC GAS ATOMIZED POWDERS OF MAGNETIC ALLOYS

R. C. O'Handley, P. Domalavage, B. W. Corb, Y. Hara, and N. J. Grant

Department of Materials Science and Engineering,

MIT, Cambridge, MA 02139.

ABSRTACT

Glassy and microcrystalline powders of FeNiBSi and CoNbB alloys have been fabricated

by atomization with ultrasonic He gas jets. Differential and cumulative weight

distributions vs. mean particle diameter (by sieve analysis) for the ^®g 5 Nii qB. ^S^q alloy show the particles below 63 pm to be fully amorphous (as judged by Mo h X-ray

diffraction) . X-ray diffraction analysis shows the crystalline particulate fraction of

the FeNiBSi alloy to be composed of a-Fe-like and Fe^B-like phases. Thermomagnetization

curves indicate the additional presence of an Fe^g-like phase. The same three phases oc- cur - in different ratios (less a phase) - upon crystallizing the fine particulates which were glassy as atomized.

Introduction

Several reports of powder fabrication in alloy systems with favorable glass-forming tendency

have appeared in the recent literature (see Table I) [1]. Miller and Murphy have gas-atomized amorphous CuZr and FeBSi alloys [2] and Yamaguchi and Narita [3] have atomized an FeBSi alloy, in both

cases into the path of a water jet (effectively steam quenching). A high voltage,

electrohydrodynamic (EHD) technique [4] , and a spark erosion technique employing an organic medium

[5,6], both making extremely fine powders, have also been reported in the literature. The EHD technique has been used to fabricate powders of several glass-forming alloys with the result that amor- phous structures appeared for diameters less than 0 .1 pm. Spark eroded amorphous c powders were produced up to 30 pm in diameter. Patterson et al have used the Pratt and Whitney RSR

technique to atomize FeBSi and NiBSi alloys. Their powders were amorphous up to approximately 15 pm

[7]. Nanao has employed sub-sonic, dry gas atomization using approximately 200 psi tank pressure of Ar to produce powders of FeBSi which were amorphous below 25 pm [8]. We have recently demonstrated

the suitability for glass formation of the ultrasonic gas atomization (USGA) technique [9] by the

production of glassy CuZr powders up to 63 pm in diameter [10]. We report here the extension of

this technique to magnetic alloys with favorable glass-forming tendency.

The ultrasonic gas atomization was effected by delivering the melt stream to an array of

ulrasonically pulsed (100 kHz), high-speed (mach 2 or greater) jets of dry He gas released at a gauge pressure of approximately 1200 psi. The powder was collected at a pressure slightly above one atmosphere in a steel vessel fitted with a cyclone collector. The technique is described in more

detail in Ref. [9] where it is applied to aluminum alloys.

Results

Structural Qiaracterization . The powder yield showed a unimodal size distribution (Fig. 1) similar to our CuZr atomization

run [10] as well as to numerous microcrystalline copper and aluminum alloy runs done with this apparatus [11] . The differential weight distribution (by sieve analysis) peaks near 70 pm in the present case, higher than the peak of the distribution in our CuZr run which was at 44 pm. This difference may be due to the differences in density and viscosity between the two alloy melts as well as

to differences in pour speed and gas pressure at the USGA die during atomization.

Mo K X-ray diffraction analysis of the sieved powders (Fig. 2) showed the particles to be

fully amorphous below a diameter d = 63 pm, just as in the case of our CuZr run [10]. Above d = 125

pm no broad diffraction peak could be discerned but only sharp crystalline peaks indicating a predominantly a-Fe structure with some peaks present. A quantitative analysis of the

intensities suggests a 70/30 weight ratio for a to Fe

2 B-like phases. In the intermediate size range

the X-ray pattern is mixed. Crystallization of the originally amorphous powder fraction d < 44 pm

gave a diffraction pattern revealing a greater fraction of phase than exists in the

678
Table I

Powder fabrication in readily-glass-fonning systems. Glass Cutoff refers to the particle diameter

below which predominantly glassy phase was observed.

ALLOY TECHNIQUE T (°C) m

GLASS CUTOFF REF.

subsonic, gas/water 890 40 pm [2]

*^®81.5®14.5^^4 subsonic, gas/water 1160 10 pm [2]

^®74®11^^15 subsonic, gas/water 1120 15-20 pm [3]

FeSiB EHD (<10“^ Torr) 0.1 pm [4]

FeNiPB EHD (<10“^ Torr) 0.1 pm [4]

*^®75®10^^15 spark erosion 1130 30 pm [5]

*^®75®10^^15 spark erosion 1130 >20pm [6]

Fe79BigSi5 RSR 1150 <5-10 pm [7]

Ni75Bi7Sig RSR 1100 15pm (a) [7]

*^®75®10^^15 subsonic. Argon 1130 25 pm [8]

USGA, hel ium 890 63pm (b) [10]

*^®65^h0®15^^10 USGA, helium 1150 63 pm (c) present

CoNbB USGA, helium 1250 present

(a) For Argon, 50 wt% of product is finer than SOpm whereas for He, 50 wt% of product is finer than

100pm. However, at a given d. He gives greater amorphous fraction than argon.

(b) Distribution peaks near 44 pm.

(c) Distribution peaks near 70 pm.

PARTICLE DIAMETER (yitm)

EA

<

z

LU

CC

LU

Fig. 1. Cumulative and differential particle size distributions by weight determined by sieve

analysis.

679
ze

Fig. 2. X-ray diffractometer scans for four powder fractions. Crystalline peaks for d > 125 pm

index to reflections for a-Fe and tetragonal Fe

2 B.

crystalline-as-atomized fraction. Quantitative analysis here sets the a to Fe^B-like weight ratio

at 40/60. This difference is brought out in Fig. 3 where the major a-Fe and Fe_B peaks are identified. The several unidentified peaks of low intensity in the crystallized d < 44 pm sample do

not match those of either orthorhombic or tetragonal Fe-B which are often stabilized by impurities

[ 12 ] . ^

Scanning electron micrographs (Fig. 4) of particulates show approximately spherical particles.

Some of the larger particles are lightly decorated with satellites, but to a lesser extent than was

the case with the CuZr powders [10] which were atomized in an earlier model apparatus. Scanning

electron micrographs of subsonic atomized powders [2,7,8] show no satellites, probably because of

the lower level of turbulence in that process. While our particulates greater than 125 pm in diameter show clear signs of crystallinity on their surface, some surface segregation is evident even in

the size fraction less than 44 pm. Electron microprobe analysis showed no compositional differences

between the different size fractions (chemical analysis is being done) and marginal differences between the regions of different contrast on the surface of the crystalline particulates.

Magnetic Characteristics .

Fig. 5 shows the magnetization at 6 kOe as a function of temperature for three size fractions

d <44 pm, 53 < d < 63 pm, and d > 125 pm. The two smaller size fractions, predominantly amorphous

as atomized, show identical magnetization curves. The finer fraction was heated above crystallization and, upon cooling, revealed a two—magnetic phase behavior.

680
The magnetic data can be used to supplement the X-ray phase information given above. Turning

first to the d > 125 |im sample, glassy as atomized, we assume all of its phases to be of approximately the same composition as the parent alloy because of the rapid solidification rate. Both of

the phase types identified by X-ray analysis have high Curie temperatures (T^(a-Fe) = 1044 K and

Tp(Fe2

B) = 1010 K) which should be moderately suppressed with small Ni or Si substitutions

(T^a-Fe gyNi = 970 K and T^( (Fe g^Ni 2 o) 2 (BSi) = 920 K) . Thus, the high-Curie-temperatur e phase

is probably* a-FeNi(B, Si) plus (FeND^tBSi) and the low-Curie-temperatur e phase, not seen in the

X-ray diffraction pattern, remains to be identified. Its Curie temperature (750 K) , while below

that of (T„ = 810K) is reasonable for that of (Fe g^Ni j^^s^FSi) which can be estimated to be

in the range of 730 E. Consideration of the magnetization' of ’these phases (using the 70/30 ratio

for the two high temperature phases determined by X-ray intensity analysis) suggests a 2/ 1 volume

ratio of high temperature phases to low temperature Fe^B-like phase. See Table II.

Table II

Relative fraction of various phases present in two powder samples.

Because the densities of the iron borides are close to that of a-Fe,

the difference between volume and weight fractions is less than the

uncertainty in the assigned fractions.

Sample a FesB

Crystalline, d > 125 pm 0.47 0.20 0.33

Crystallized, d < 44 pm 0.22 0.33 0 .45

29

Fig. 3. Comparison of X-ray diffraction patterns for crystallized-af ter-atomization d < 44 jjun fraction and crystalline—as—atomized d y 125 pm fraction. The major peaks are identified

as a Fe like and Fe^B— like phases. The lengths of the arrows indicate expected relative

intensities.

681
MAGNETIZATION

(emu/g)

125 |jm sample is not typical (it is the only one seen in thousands scanned) but is an example of an interesting but undesirable feature which can be found in some larger

particluates

Fig. 5. Specific magnetization as a function of temperature at an applied field of 6 kOe. For d < 44 (irn (solid curve) data is shown for heating the glassy state through crystallization and

cooling in the crystallized state. The behavior of the 53 < d < 63 pm glassy fraction

(dashed line) is essentialy identical to that of the d < 44 |jm sample. The

crystalline-as-atomized fraction d> 125 pm (solid line) shows a high Curie temperature due

to the dominance of a-FeNi (probably with appreciable boron in solution)

682
Similar analysis for the crystallized d < 44 (un sample (using the 40/60 weight ratio for a to

Fe^B-like phases determined by X-ray intensity analysis) points to an approximately 55/45 volume ra- tio of high temperature to low temperature phases. We can only speculate why the inferred Fe-B-like

phase would not appear in X-ray diffraction: it may exist in finely dispersed, microcrystallite

form.

Magnetization-field loops on the packed powders revealed only minor differences between the

amorphous and crystalline samples. Figure 6 shows an M-H loop for the d < 44 pm fraction corrected

for demagnetization of the contiguous spheroidal particles. Results derived from these loops for

all samples are listed in Table III. The coercivities of the compacted glassy powders are three

orders of magnitude greater than those of glassy ribbons of comparable compositions. Magnetic

anisotopies of the different samples were estimated from the area above the magnetization curve. Although the anisotropies are two orders of magnitude greater than values obtained in the same way

from amorphous ribbons [13], their principal value is as a comparison between the various size

fractions. The comparison shows that in the compacted powder form the properties appear to be

dominated by surface effects that mask the intrinsic magnetic softness of the amorphous magnetic

state and blur the differences between the crystalline and amorphous samples.

Table III

Iilagnetic characteristics of compacted powder samples from three size fractions. Specific saturation

magnetization a was determined by extrapolation to infinite applied field. Values in parentheses

are for crystallized particulates.

5 3 Size fraction a (emu/gm) 4nM (kG) Coercivity (Oe) Anisotropy (10 erg/cm )

d>125 pm 141 13.3 70 1.35

53<d<63 pm 132 12.4 45 (45) 1.4 (1.65)

d<44 pm 135 12.7 50 (85) 1.29 (1.92)

a (a) A mass density of 7.5 g/cm^ was assumed for the alloy.

Hi„,(kOe)

Fig. 6. Magnetiztion vs. applied field for compacted particulates of d < 44 pm fraction. The

correction for the demagnetization of the contiguous spheres is illustrated as is the

method for estimating the magnetic anisotropy. Differences among the various samples are listed in Table III.

683
Discussion

Implications fil magnetic measurements

The composition is hypereutectic according to the ternary FeNiB diagram [14] . The

mode of crystallization of our alloy with ®j^ 5 Sij^Q into a-FeNi(B.Si) + (FeNi),(BSi) with finely

dispersed Fe^B-like microcrystals possibly also present, suggests eutectic crystallization in the

hypereutectic region [15] . The difference between the crystalline-as-atomized powders (d > 125 pm)

and the crystallized-af ter-atomization powders (d < 44 pm) lies not in the phases nucleated, but in

the times available for growth of existing nuclei. The crystal line- as- atomized sample shows a smaller ratio of iron boride phases to a phase than does the sample crystallized after atomization.

Crystallization after atomization at a heating rate of a few degrees per minute allows more time for

diffusion of the metalloids from the a phase to form more of the iron boride phases than exists in

the d > 125 pm fraction.

The increased concentration of Fe B (n = 2,3) phases in the crystallized d < 44 pm sample probably contributes to its greater magnetic hardness relative to the d > 125 pm sample which contains more a phase. It is known that a-Fe crystallites do not have as deleterious an effect upon magnetic

softness as do the lower symmetry phases Fe^B and Fc2 B [16].

Comparison of atomization te chnioues

Convective undercooling AT^ of a droplet, per unit time At, can be expressed [17,18]:

AT /At = 6h(T-TQ)/(dpC )

( 1 ) where h is the heat transfer coefficient, T^/\ isIS the LHC temperature LcmperaLuxc ofui the lc cool ing medium, d is the

droplet diameter, p is its mass density and C is its specific heat as it cools. The heat transfer

coefficient has been observed to vary roughly linearly with the droplet velocity [18] and can be

approximated empirically as [17,18]:

h = 0.6k ^^^(vp/d)^^^C ^^^(1/n )^^^

o o o (2)

where k , p , C , and n are respectively the thermal conductivity, mass density, specific heat and

viscosity of ^he cooking medium and v is the relative velocity of the droplet in the medium. The

factors have been arranged in order of their importance.

The difference between our USGA results (1200 psi) and those of Nanao (200 psi) (see Table I)

points to the importance of droplet velocity and atomization gas in determining heat transfer and

amorphous alloy formation as indicated by Eqs. (1) and (2). Clearly the greater pressure used in

USGA leads to greater average relative droplet/gas velocity. The heat transfer of He relative to argon is affected by the eightfold greater thermal conductivity of He, the tenfold greater specific

heat of He and the tenfold greater mass density of Ar. Based on relation (2) these differences lead

to a heat transfer for He approximately 2.7 times that of argon.

Patterson found Ar gave a finer particulate distribution than He even though He gave a greater

yield of amorphous material at a given droplet diameter [7]. This more efficient atomization by ar- gon relative to helium can be understood in terms of the greater density of the argon relative to hel ium.

Spark erosion and the EHD process seem well suited to the production of super-fine

distributions but appear to have no advantage over gas atomization techniques, especially USGA with

He, in terms of heat transfer and glass forming capability. However, some of the finer particulates

produced in spark erosion are believed to be quenched from the vapor state and thus experience

enhanced quench rates [6]

Other differences in the upper limit to amorphous particulates (Glass Cutoff in Table I) can be

understood in terms of ease of glass formation as gauged roughly by differences in alloy melting

temperatTires

It is interesting to note that our cutoff diameter for glassy atomized powders is greater than

the critical thickness above which ribbons of any FeB Si alloys can no longer be made glassy by melt

spinning: d . . < 42 urn [19]. crit

Aclrnowl ed pement

We are pleased to acknowledge the support of this research by the U. S. Army Research Office

under contract No. DAAG-29-80-0088 . Helpful discussion with C. Ashdown and E. Ting are appreciated.

684
References

[ 1 ]

[ 2 ]

[3]

[4]

[5]

[ 6 ]

[7]

[ 8 ]

[9]

[ 10 ]

[ 11 ]

[ 12 ]

[13]

[14]

[15]

[16]

[17]

[18]

[19]

H. H. Liebermann and J. L. Walter in Ranidlv Solidified Amorphous and Crystalline A1 1 ov

. op. cit. p. 111.

S. A. Miller and R.J. Murphy, Scripta Met. 1^, 613 (1979) and in Rapid Solidification

Proce ssine : Principles and Technologies II (RSP II), ed. R. Nehrabian, B. Rear and M. Cohen (Claitor's Pub. Div., Baton Rouge, LA, 1980), p. 385.

X. Yamaguchi and E. Narita, Appl . Phys. Lett. 468 (1978).

J. Perel, J.F. Mahoney, P. Duwez, and B.E. Ealansher in RSP II op. cit. p. 287, J. Perel, J. F. Mahoney, S. Taylor, Z. Shanfield, and C. Levy in Rapidly Solidified Amorphous

and Crystalline Alloys , Ed. B. H. Eear, B. C. Giessen, and M. Cohen, (North Holland, New

York, 1982) p. 131. also p. 258 in RSP I.

T. Yamaguchi and E. Narita, IEEE Trans. MAG-13 , 1621 (1977) .

A. Berkowitz and J.L. Walter in RSP II op. cit. p. 294 and Mat. Sc. and Eng. 5^, 275

(1982). Further results are described by A. Berkowitz, J.L. Walter and E.F. Wall, Phys.

Rev. Lett. 46, 1484, 1981.

R. J. Patterson II, D. L. Ledwith and J. C. Dwyer, in Processing of Metal and Ceramic

Powders , ed. E. W. Lay and R. M. German (The Metallurgical Society, Warrendale, PA, 1982)

in press.

Prof. S. Nanao, Dniv. of Tokyo, unpublished work.

V. Anand, A.J. Eaufman and N.J. Grant in RSP II, op. cit.,p. 273.

P. Domalavage, C. Ashdown, R. C. O' Handley and N. J. Grant, J. Mat. Sc. and Eng. 57 . 143 (1982)

P. Domalavage and N.J. Grant, internal report, and N. J. Grant in Advances in Powder

Technology . (ASM, 1982) in press.

R. Hasegawa, R. C. O'Handley and L. I. Mendelsohn in AIP Conf. Proc. (AIP, New York,

1976) p. 298.

R. C. O'Handley, AIP Conf. Proc. 29 . (AIP, New York, 1976) p. 206.

H. H. Stadelmaier and C. B. Pollack, Z. Metall. 60 , 960 (1969) .

U. Herold and D. Eoster, in Ranidlv Quenched Metal

III, ed. B. Cantor (The Metals Society,

London, 1978) p. 281 and R. Hasegawa and R. Ray, J. Appl. Phys. £9, 4174 (1978).

A. Datta, N. J. DeCristofaro and L. A. Davis, in Ranidlv Quenched MetalIV, ed. T. Masumoto and E. Suzuki (Japan Inst, of Metals, Sendai, 1982) p. 1007.

R. Mehrabian, in Ranid Solidification Processing ; Principles and Technologies I ( RSP I) , ed. R. Mehrabian, B. Eear and M. Cohen (Claitor's Pub. Div., Baton Rouge, LA, 1978), p. 9. R. J. Patterson, in Ranidlv Solidified Amornhous and Crystalline Alloys , op. cit. p. 123.

N. J. Grant, in RSP I, op. cit. p. 230.

F. E. Luborsky, H. A. Davies, H.- H. Liebermann, and J. Reeve, J. Appl. Phys., Nov. 1982.

685
PRODUCTION APPLICATIONS OF RAPIDLY SOLIDIFIED TOOL STEELS,

SUPERALLOYS, TITANIUM ALLOYS, AND CORROSION-RESISTANT ALLOYS

J. H. Moll

Crucible Research Center

Pittsburgh, Pa. 15230

F. J. Rizzo

Crucible Compaction Metals Operation

Oakdale, Pa. 15071

ABSTRACT

While extensive work on rapidly solidified materials is being conducted in the laboratory, a number of these materials are now in production or have near-term production status.

These materials include tool steels, superalloys, titanium alloys and corrosion-resistant

alloys. Rapid solidification of tool steels has resulted in refined microstructures,

improved performance, and advanced alloy development. Several tool steels have been

developed which can be made only by rapid solidification techniques. Superalloy and

titanium alloy products made from rapidly solidified powders are currently in production

or near production for a number of aerospace applications. The materials offer improved

quality, property capabilities and uniformity. Rapid solidification is now being evaluated

for corrosion-resistant alloys. Some unique physical metallurgy aspects and applications

for these materials have been established.

Introduction

In the past 20 years, processes have been developed to utilize rapid solidification technology

(RST) to produce upgraded or entirely new high performance alloys. The principal advantages of these

processes are (1) the ability to improve property uniformity and alloy development flexibility due to the elimination of macrosegregation, and (2) the ability to make near-net shapes which results in

reduced input material and less machining to produce a finished part. To avoid segregation, solidification rates faster than those used in conventional ingot cooling are required. Fast cooling rates

(1Q3 to 10^ C/sec) are attainable by making small particles with a large surface area to volume ratio

This criterion is met with powder particles ranging in size from several hundred microns to less than

37 microns. Spherical powders of this type are currently in production and are being used to produce

shapes ranging from bars to complex near-net shape parts using hot-isostatic-pressing. The following

are some typical production applications of rapidly solidified tool steels, titanium alloys, superalloys and corrosion-resistant materials.

RST Tool Steels

The Crucible Particle Metallurgy (CPM) process for high speed tool steels has been commercially

applied since 1970. The process enables production of superior quality grades of conyentional high

speed steels and also lends itself to the production of new higher alloyed grades which are not

practical with conventional steel making. These new grades provide unique combinations of higher red

hardness, wear resistance and toughness and are proving cost effective in many commercial applications. The process involves gas atomization of prealloyed induction melted molten metal. The resulting powder is dried, screened, and loaded into steel containers which are subsequently evacuated and

sealed. The powder is consolidated by hot-isostatic-pressing (HIP). The resulting compact, weighing

up to 2000 lb, can be rolled or forged on conventional mill equipment or used in the as-HIP condition

after the can is stripped off.

One of the most detrimental characteristics of conventionally produced high alloy tool steels is

their tendency to form carbide segregation in slowly cooled Ingots. Rapidly solidified powder, on the

other hand, insures a very fine uniform distribution of carbide. Figure 1 compares the microstructural

686
appearance of conventional T15 steel with CPM T15 steel. Virtually all the carbides in the CPM

product are less than 3 microns in size. In the conventional product, more than half of the carbides

are larger than 6 microns and sizes up to 34 microns have been observed. The uniform distribution of

fine carbides in CPM steels results in superior grindability in the hardened condition (Figure 2)

without sacrifice in tool wear. For example, in form turning screw machine parts from Armco iron,

CPM T15 had a tool life of 180 minutes compared to 80, 50 and 30 minutes for titanium carbide,

tungsten carbide, and conventional T15, respectively. In a punching operation on T430, CPM T15 outperformed D2. Tool life (pieces/grind) for D2 was 70,000; tool life for CPM T15 was 500,000.

Rapidly solidified powder also opens development for new tool materials. A cold work steel

recently developed is CPM lOV which contains 2.4% C and 10% V. The alloy could not be practically

produced by conventional ingot technology. The alloy gives an unusual combination of wear resistance

(Figure 3) and toughness (Figure 4). CPM lOV is outperforming conventional tool materials in such

applications as cores used in molding ceramics, dies for cold compaction of steel powder, dies for

cold drawing Ml high speed steel, flattening dies used in the production of drills, and punches for

slots in iron oxide coated tags.

Figure 1. Comparison of Carbide Structure of Conventional T15 (left) and CPM T15 (right).

GRINDING RATIO GRINDING RATIO

/ VOLUME OF METAL REMOVED \ V VOLUME OF WHEEL WORN )

- A

- 2

Figure 2.

Figure 4.

Comparison of Grindability

of Conventional Tool Steels

with CPM Tool Steels.

Comparison of Toughness of Conventional

Tool Steels and CPM

Tool Steels.

Grade HRC02

62 M264 CPM M4 CPM 10V CPM 10V 64 60 63

Figure 3. Wear Properties of Conventional Tool Steels and CPM

Tool Steels.

35

30

25 Charpy C-Notch 20

Impact taat

(tt.-ib) 15

10

CFM M4 CPM 10V CPM 10V 64 60 63

687
RST Titanium Alloys

Titanium alloys are used in the aerospace industry where high strength-to-weight ratios at moderate temperatures are required. This involves both airframe and engine applications. Conventional forgings offer attractive properties but high buy-to-fly ratios result in high cost. Casting

technology has been developed to produce near-net shapes but properties of castings are often inadequate. A process, utilizing rapidly solidified powder, has been developed and is now in pilot production for producing fully dense near-net shaped parts with properties equivalent to forgings. The

P/M process involves the use of a ceramic mold made by the lost wax process, powder made by PREP

(Plasma Rotating Electrode Process), and hot-isostatic-pressing.

Figure 5a shows an engine compressor rotor made by using the ceramic mold process and Ti-6A1-4V

powder. Using conventional processing, the part is machined from a pancake forging which weighs 30

lbs. The as-HIP near-net shape weighs only 6.2 lbs. The process is also capable of making large

complex airframe parts such as the arrestor hook support fitting shown in Figure 5b. Normally made

from a forging weighing 175 lbs, this as-HlP part weighed only 60 lbs. In some instances, parts made

by the ceramic mold process need not be machined on all surfaces. For the fuselage braces shown in

Figure 5c, only the top and bottom surfaces will be machined. The webs will require only light chemical milling to remove a thin alpha case before Installation into an aircraft.

The combination of rapidly solidified powder and hot-isostatic-pressing results in an isotropic

microstructure and property capability (Table I). With the advent of clean titanium alloy powder,

the fatigue properties of P/M titanium have improved markedly in recent years. Figure 6 compares the

fatigue properties of P/M Ti-6A1-4V with conventionally produced material. The outstanding capability

of current P/M material is evident.

Coll IndusIriM

(a) (b)

Figure 5. Hardware Produced Using Titanium Alloy Powder and Crucible's Ceramic Mold Process.

(a) Compressor Rotor, (b) Arrestor Hook Support Fitting, and (c) Fuselage Braces.

Table I RT Tensile and Toughness Properties

of HIP T1-6A1-4V (Ref. 1)^

Test

Orien- YS TS El RA Kq°

tat ion (ksi) (ksi) (%) (%) (ksiy'^in)

LT 132 139 13 38 73

TL 130 138 12 35 74

ST 132 139 12 36 72 AMS

Spec.

4928H

120 130 10 15

^1300°F/8 hr Vacuum Anneal.

^Predicted from precracked Charpy tests. Figure 6. Room Temperature Fatigue Properties

of Ti-6A1-4V (Ref. 2, 3, and 4)

688
RST Superalloy

Rapid solidified powder has been successfully utilized over the past five years in the production

of as-HIP superalloy turbine engine hardware. A prime example of the application of powder to as-HIP

superalloys is the General Electric T700 engine which powers the Army Blackhawk helicopter. Six

critical rotating turbine components per engine have been produced from as-HIP Rene 95 since the onset

of its production.

Rapid solidification technology has addressed the problems of critical material shortages by

using less input weight. There has been a 40-lb weight savings of purchased hardware per engine over conventionally produced components. Other advantages are: (1) Shorter lead time since many processing steps are eliminated, (2) no need for heating and forging steps that could introduce detrimental

conditions, (3) less finish machining because of the reduced input weight, (4) improved machinability

because of greater product uniformity, and (5) lower testing costs because of mechanical property

reproducibility; no parts have been rejected for failure to meet mechanical property requirements and

testing has proceeded from a 1 in 25 cut-up plan to a 1 in 100 cut-up plan.

Rapidly solidified material is produced at Crucible using the inert gas atomizer depicted in the

schematic in Figure 7. This material is then processed and qualified by a series of characterization

tests. Once it is found acceptable, it is loaded into stamped shaped metal containers. Figure 8 shows a cut-away of the T700 disk container which was designed to use the minimum amount of input

material. After the loading process, the container is outgassed and hot-isostatically-pressed to 100 percent density. This is followed by heat treatment, non-destructive testing and rough machining.

A sampling of parts is cut-up for mechanical property evaluation and quality assurance. Figure 9 shows a picture of an as-HIP disk, a rough machined disk and the finish machined components made for

the T700 engine.

Over 6000 parts have been produced using this technology and over 800 engines have been

delivered to the field for service. There are over 200,000 operating hours accumulated on as-HIP

Rene 95 parts. This engine experience is ongoing proof that rapidly solidified powder has been and

will be beneficial to superalloys. Probably the most important aspect of this application of the

technology is that it has resulted in a $3000 per engine savings to the ultimate customer.

Figure 7. Schematic of Crucible's

Production Inert Gas

Atomizer

Crucible Compoction Metols

Colt Industries 0 r

2"

Figure 8. Sectioned Powder Container for a T700 Turbine Disk Sonic Shape.

Figure 9. P/M Rene 95 Turbine Disks.

(Top left-as-HIP shape. Top right-machined sonic

shape. Bottom left-final machined 1st stage disk.

Bottom right-final machined 2nd stage disk.)

689
RST Corrosion-Resistant Alloys

Powder is being considered for corrosion-resistant alloy applications. With increased exploration in sour oil and gas wells, the need for corrosion-resistant materials for use as tubing, connectors and valving has increased. The alloy steels commonly used by the oil industry do not hold up in

these highly corrosive environments. Conventional production of high alloy tubing stock is expensive

and time consuming. Since most sour wells are 15,000 feet and deeper, this makes obtaining the product from these wells a very costly proposition. There is currently no wellhead equipment made from

corrosion-resistant alloys. So, in addition to the cost problem, there is still a materials problem.

To address the sour well materials problem, rapid solidification technology is being applied

using two approaches: (1) solid tubing and near-net shape wellhead equipment, and (2) clad tubing

and clad wellhead equipment. The advantages to the solid approach are tied to material usage. The

application of powder for solid tubing results in higher yields from starting material than conventional processes. For wellhead equipment, powder allows the manufacturing of near-net shape parts

which have never been attempted by conventional forging processes. A photograph of solid alloy C-276

tubing made by the powder process is shown in Figure 10. A solid alloy 625 valve body made using

rapidly solidified powder and the Crucible ceramic mold shape process is pictured in Figure 11.

The clad approach offers the same advantage as the solid plus the economic advantage accompanying

the use of less high alloy corrosion material to do the job. Processing with the clad approach may be more costly, but it offers a base material of alloy steel with a high experience level with a coating

of corrosion-resistant material on surfaces which contact the hostile environment. Figure 12 shows a picture of an alloy steel valve body clad on the inside with alloy 625 powder. Clad tubing is also a feasible product using powder. Figure 13 shows a cross-section of alloy 625 clad on the outside of a lower alloy material using hot-isostatic-pressing.

Without the use of powder, it would be virtually impossible to clad a shape such as a valve body.

In addition, because corrosion-resistant alloys generally develop their strength through cold work

(such as in solid tubing), even forging of a valve body would not be effective in producing a component with minimum specified mechanical property levels. RST produced solid and clad valve bodies

achieve property levels consistent with those specified for valve bodies made of alloy steel. Some

mechanical property levels possible with rapidly solidified materials are shown in Table II.

Figure 10. Alloy C-276 Tubing Produced

from Hot-IsostaticallyPressed Powder.

Crucible Inc Reeeerch Center

Coll Indutiries @

Figure 11.

Alloy 625 Valve

Made from HotIsos tat icallyPressed Powder

(Top-as-HIP Part,

Bottom-Machined

Part)

690
Figure 12. Alloy Steel Valve Body Clad on the Figure 13. Cross-Section of Alloy 625 Powder

Inside with Alloy 625 Powder using Clad on the Outside of a Lower

Hot-Isostatic-Pressing. Alloy Material using HotIsostatic-Pressing .

Table II

Mechanical Properties of RST

Corrosion-Resistant Alloys

Requirement

Wellhead Equipment

API

Type 2 API

Type 3 Typical RST

Alloy 625

TS (ksi) 90 100 144

YS (ksi) 60 75 83

El (%) 18 17 38

RA (%) 35 35 38

Hardness (Rc) 35 Max 35 Max 22

Typical RST

Tubing Requirement Alloy C-276

TS (ksi) - 172

YS (ksi) 150- 170 161

El (%) - 24

RA (%) - 69

Hardness (Rc) - 36

References

(1) V. C. Petersen and V. K. Chandhok, "Manufacturing Process for the Hot-Isostatlc-Pressing of Large

Titanium PM Shapes," 1st Interim Report, Contract F33615-77-C-5005, June 1979.

(2) F. H. Froes et al, "Developments in Titanium Powder Metallurgy," Journal of Metals, Vol. 32,

February 1980.

(3) F. H. Froes et al, "Fatigue Properties of Hot-Isostatically-Pressed T1-6A1-4V Powders," Proceedings of 1982 National Powder Metallurgy Conference, Montreal, Canada, May 1982.

(4) F. C. Teifke et al, "Effect of Processing on Fatigue Life of T1-6A1-4V Castings," Advanced

Processing Methods, Proceedings of 1981 TMS-AIME Fall Meeting, Louisville, Ky., 1981.

"The material contained in this paper is intended for general information only and should not be used

in relation to any specific application without independent study and determination of its applicability and suitability for the intended application. Anyone making use of this material or relying

thereon assumes all risk and liability arising therefrom."

691
COMMERCIAL RAPID SOLIDIFICATION (CRS) PROCESSING

FOR ADVANCED TURBINE DISK APPLICATIONS

R. Beck, Teledyne CAE

Toledo, Ohio 43612

W. Kent, Universal Cyclops

Bridgeville, Pennsylvania 15017

C. S. Kortovich, TRW Inc.

Cleveland, Ohio 44117

ABSTRACT

Rapid solidification rate powder metallurgy processing was applied to the development of enhanced

superaUoys for jet engine turbine disk applications. Argon atomization employing commercially available

facilities was used to produce the rapidly solidified powders. Consolidation involved the preparation of forging

billets by the CAP (Consolidation at Atmospheric Pressure) process which were then isothermaUy forged and

heat treated to develop the desired microstructure and mechanical properties in candidate alloy compositions.

The series of superaUoy compositions was processsed into CRS powder and was consolidated by CAP and

isothermal forging techniques. The alloys included AF-115, IIB-llE, 312Hf, AF2-1DA-2, AF2-1DA-5, AF2-1DA6, Rene' 95 and IIB-7. The CAP material exhibited excellent forgeability in channel die tooling which reduced

the cylindrical shaped billets to rectangles featuring a reduction in thickness of approximatley 50%. Subsequent

to forging, heat treatments were developed to produce an ASTM 5-6 grain size. Mechanical property evaluations

including elevated temperature tensile, stress rupture and low cycle fatigue tests were conducted on these

materials. The results of these evaluations will be presented. This work was performed under Contract F33657-

80-C-0124.

Introduction

Ultra-pure prealloyed superaUoy powders solidified at very high rates of cooling have exhibited a strong

potential to limit the aUoy segregation in cast/wrought aUoys to a microscopic scale. Elimination of this segregation has resulted in a stronger material exhibiting a more uniform range of mechanical properties. In order to more fuUy evaluate this class of materials for advanced turbine disk applications, a materials properties

screening study was conducted on powder aUoys produced by the commercial rapid solidification (CRS) process. The scope of this effort included the evaluation of CRS powder consolidated at atmospheric pressure (CAP),

foUowed by isothermal forging, to provide base material for screening mechanical property tests. The program

mechanical property goals included a 170 ksi 0.2% offset yield strength at 1100°F, 100 hour rupture lives at 1100°F/170 ksi and 1220OF/140 ksi, and 2400 cycles low cycle fatigue life when tested at 1100°F, with a strain

range of 0.64% and a strain ratio of 1 and were selected to meet Teledyne CAE requirements for their advanced

turbine disk designs.

Experimental Procedures

The nominal chemistries of the aUoys chosen for evaluation are listed in Table 1.

692
Table 1

Nominal Chemistry (Weight Percent) of CRS AUoys(l)

Alloy C Cr W Mo Co A1 Ti Cb B Ta Zr Hf V

AF2-1DA-2 .04 12.0 6.0 2.75 10.0 4.6 2.8 _ .015 1.5 .10 _ _ AF2-1DA-5 .04 12.0 6.0 2.75 9.0 4.6 2.8 - .015 2.2 .10 - - AF2-1DA-6 .04 12.0 6.0 2.75 10.0 4.6 2.8 - .015 1.5 .10 - - EB-llE .06 8.5 7.1 1.90 9.0 4.5 0.7 - .013 6.4 .04 0.7 0.5 XN-312-Hf .08 11.5 3.0 3.30 20.0 2.9 4.2 2.0 .020 2.0 .05 1.3 - AF-115 .15 10.7 5.9 2.80 15.0 3.8 3.9 1.7 .020 - .10 2.0 - BB-7 .13 9.0 7.5 2.00 9.0 3.5 .8 - .020 10.0 .10 1.0 0.5 Rene' 95 .06 13.0 3.5 3.50 13.0 3.5 2.5 3.5 .10 - .05 - -

(1) Balance Nickel

The first three are modifications of AF2-1DA, a disk alloy developed by Universal Cyclops under Air Force

sponsorship. The remaining alloys chosen for study were existing aUoys currently in advanced development

stages. The HB-IIE alloy was developed by Universal Cyclops under NASA sponsorship, the XN-312-Hf was

developed by Climax-Molybdenum and the Rene' 95 and AF-115 were developed by General Electric under Air

Force sponsorship.

AU of the powder meteil stock was produced by Universal Cyclops from high purity vacuum induction

melts. The alloys were argon gas atomized using the CRS process. This process was developed to yield a greater

percentage of fine powder particles. The atomized powder was screened to a -150 mesh size and blended for

uniform distribution. Sieve analyses indicated generally greater than 50% of particles finer than -325 mesh

(approximately 40 microns).

Powder consolidation was accomplished at Universal Cyclops using the CAP process. This is a high

temperature sintering operation which is conducted on powders sealed under vacuum in glass containers. The

powders are specially treated prior to encapsulation to enhance their sinterability. The powders were

consolidated into 4 inch long, 1 inch diameter cylinders for subsequent isothermal forging trials.

Isothermeil forging was accomplished by TRW in a channel die tooling arrangement which forged the

initial cylindrical CAP preforms into rectangles exhibiting approximately 50% reduction in thickness in relation

to the diameter of the CAP preforms. The dies were made from Mo-TZM and forging was conducted in a single

blow under a protective argon atmosphere at temperatures ranging from 2000-2200°F.

Mechanical property testing was conducted by TRW and Teledyne CAE and consisted of IIOQOF tensile

tests, stress rupture tests conducted at a number of conditions and IIOQOF low cycle fatigue tests.

Results and Discussion

MetaUographic analysis of the CRS powders indicated that they were generally spherical in shape, with

associated satellite particles and were found to be free of foreign material contamination. An example of

typical CRS gas atomized powder as observed by SEM is shown in Figure la. The powders were characterized by

dendritic and microcrystalline types of structures, with the smaller powder particles (100 microns or smaller)

generally exhibiting a finer dendritic structure or a microcrystalline structure. CAP consolidation resulted in structures exhibiting approximately 2% residual porosity. An example of a typical as-CAP structure is shown in Figure lb. The grain sizes of the CAP materials varied somewhat,

depending upon the temperature required for consolidation. The grain sizes ranged from ASTM 4-8.

The isothermal forging results over the temperature range 2000-2200°F indicated that aU of the alloys

exhibited excellent forgeability. Complete densification of the CAP preforms was accomplished with no forge

cracking. Differences were observed, however, in terms of resultant grain size as well as time required to accomplish the forging. In general, forging at 2200°F resulted in a large (greater than ASTM 2) grain size in the

material. This type of grain size would be expected to exhibit poor strength properties at 1100°F. Forging at the lower temperatures resulted in either no recrystaUization or a partially recrystallized structure. This

situation offered the opportunity to manipulate the structure through subsequent heat treatment operations. In terms of forging time, deformation weis accomplished within 5-10 minutes at 2200°F, often taking approximately one hour at 2000°F. Long deformation times would result in poor production rates during actual parts

manufacture. Because of the resultant structures and relatively short deformation times, 2100°F was selected

693
as the forging temperature for the production of specimens for mechaniceil property testing. An example of the

typicsil as-forged structure for material forged at 21000P is shown in Figure Ic. In order to develop optimum

mechanical properties in the 1100-1220OF temperature range, solution heat treatments were selected which

resulted in grain sizes in the range ASTM 5-7. A typical heat treated structure is shown in Figure Id.

Mechanical property screening tests were conducted on forged material to characterize the effects of

the various heat treatments on tensile, stress rupture and low cycle fatigue properties. Analysis of the tensile

data indicated that only the AF2-1DA-2 and IIB-iiE alloys did not meet the yield strength goal of 170 ksi at 1100°F. In terms of reduction of area only the 312 Hf aUoy exhibited values (3-4%) below 12%. The stress rupture results indicated that aU of the alloys screened in this program exhibited good stress rupture life characteristics in the intermediate temperature range. A Larson MiUer summary plot of the stress rupture data

shown in Figure 2a indicates that aiU of the materials tested met or exceeded the Teledyne CAE design goals.

The low cycle fatigue results were equally promising. The results for aUoys IIB-7, AF2-1DA-6 and AF-115 are shown in the plot of Figure 2b compared to the program goal. These three alloys exhibited the highest llOflOF

yield strength values. As shown in Figure 2b, the data indicated that AF2-1DA-6 and AF-1115 have very similar

fatigue life characteristics. The heat treatments for both of these alloys included the following:

AF2-1DA-6 AF-115

2100°F/2 Hours Salt Quench to 1100°F 2100°F/2 Hours Rapid Air Cool

1400OF/16 Hours Air Cool 1400°F/6 Hours Air Cool

These heat treatments were selected to maintain a fine recrystallized (ASTM 5-7) grain size that was considered

desirable for optimum tensile properties. A fine gamma-prime matrix and grain boundary precipitate was also developed throughout the microstructure.

Conclusions

As a result of the material property screening conducted as part of this effort the following conclusions are offered:

1. Processed by the CAP/isothermal forge method, both CRS AF2-1DA-6 and AF-115 wiU meet the

Teledyne CAE design requirements for advanced turbine disks.

2. The CAP/isothermal forging process is effective and has the potential to reduce powder metal forging

processing costs from current industry levels.

3. In the alloys screened, a 50% reduction in preform thickness at 2100Op was adequate to eliminate

powder porosity and develop the desired mechanical properties.

4. The solution heat treatments applied to these alloys exhibit sufficient flexibility to produce grain sizes

that wiU optimize mechanical properties in the intermediate temperature range.

694
Af, %

200,

20

10

a, KSI

ico _nrv f L - \ A, A

\ A

V 140 KSI/1220'’F/100»HR

A

i

f 1 - 1 ?r\ l^'ClMi/V^oC/inO UD Q A9

O-IIB-7 0-IIB-11E A-AF2-1DA-2

-AF2-1DA-5

4--312 Hf o -RENE’ 95

-GOAL

35 36 37 38 39 40 41 42 43 44 45 46 47 P = (T 460HC + LOG t)

(a) Stress rupture results

(b) Low cycle fatigue results

Figure 2. Summary of stress rupture and low cycle fatigue results of

CAP/IsothermeQ forged CRS alloys.

695
Figure 1. Photomicrographs of CRS P/M alloy AF2-1DA-6.

696
PANEL/WORKSHOP ON

CRITICAL QUESTIONS IN RAPID SOLIDIFICATION PROCESSING

Three overall topics were discussed in this session as follows:

1. Roadblocks to RSP, organized by M. Cohen and J. Moore

Reporters: M. Greenfield and J. Manning

2. Alternatives to RSP, organized by J. Williams and R. Sprague

Reporters: A. Adair and L. Jacobson

3. New Horizons for RSP, organized by B. Kear and L. Davis

Reporters: D. Polk and B. Rath

Various speakers were selected to address specific questions under each topic and to stimulate

discussions on questions having broad implications for the field of rapid solidification. A summary

of the short presentations and the resulting discussion is presented here. In the interest of

completeness, conciseness, and uniformity of style, the Reporters have taken the liberty occasionally

to rephrase statements made by the speakers and discussers.

TOPIC 1. Roadblocks to Rapid Solidification Processing

A. Availability of RSP Materials for R&D - T. Tietz

Results gathered during a recent NMAB study on the availability of rapidly solidified material

for research and development purposes were summarized. This NMAB report will soon be available for

distribution as NMAB-411. The study was undertaken to determine the current status of organizations

supplying RSP materials and what problems might exist. Thus, a survey of existing facilities for

production and consolidation of RS particulate is included in the report. In addition, a survey was made of R&D organizations to determine (1) past use and projected requirements for RSP materials,

(2) problem areas experienced in obtaining RS particulate and/or consolidation services and

(3) sources able and willing to supply RS particulate and/or consolidation services to others.

Problem areas reported by researchers included (1) most facilities produce 100-200 lb powder lots

with very few facilities being available for producing small lots or custom orders for other organizations, (2) excessive delay times in obtaining test materials, (3) contamination during processing

both from refractory furnace components and from cross contamination, (4) inability to obtain

specific alloy compositions, (5) lack of compositional control, (6) lack of process control, (7) cost

of particulate, (8) cost of consolidation, (9) difficulty or inability in obtaining larger lots of RS material (50 to 200 lb) and (10) inability to locate suitable consolidation facilities. Atomization

processes appear well-developed and sufficient, although melt facilities and melting practices often

are sources of problems.

With respect to the roadblock provided by an inadequate supply of rapidly solidified material,

interest and approval was expressed by workshop participants in the creation of new rapid solidification facilities at the National Bureau of Standards as a source of small lots for cooperative

research activities. Large commercial firms with RSP facilities are likely to see little economic

benefit in supplying rapidly solidified material to other investigators. However, there is some industrial capacity for production of RSP materials by small companies. Provision of samples to universities by industrial concerns might be expected if patent and licensing questions could be

clarified. Some university sources of rapidly solidified material are available, for example, melt

spun alloys, but that facility provides only ribbon and not powder.

In further discussion of the production of RSP steel products, control of composition, particularly carbon and nitrogen, was emphasized as a significant problem. Since powder processing can influence the composition, redesign of starting compositions may be necessary in some instances.

It was emphasized that considerable time is required to develop systems to go from powder to high

quality consolidated products. Cleanliness during the powder production and consolidation stages

continues to be very important.

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B. Problems in Consolidation of RSP Particulates - M. DeCrescente

The current methods for the consolidation of RSP particulate were reviewed. The alternatives

available to todays metal processor provide the opportunity to select a consolidation process

tailored to individual need. There are still some problems preventing wide spread application,

however. Concern for contamination during compaction is an issue. The increased awareness of its

detrimental effect has not only spurred research toward solutions in that area, but has also created

some hesitancy among current and potential users of this technology. Similar comments can be made

in the areas of toxicity and safety.

In most cases, existing consolidation equipment and technology appear suitable for RSP powder

having a refined microstructure. In fact, some property improvement is evident at even slightly

higher, more effective pressing temperatures. Further, some of the as HIP property combinations

appear unique enough to warrant greater study, post process heat treatment may not be necessary.

The situation for melt spun product is currently not as positive. Additional work will be needed to efficiently process ribbon into final billet or component shape. There does not appear to be a great

deal of activity addressing this problem.

Dynamic compaction, a novel low heat input technique, offers promise. In this area, questions

concerning size limitations and frictional heating effects must be addressed before wide scale usage

is realized.

Other techniques such as containerless processing and fluid die hot pressing are characteristic

of the innovation which exists today in industry. Overall, the future is expected to bring low cost,

high volume production processes yielding a superior product.

C. Surface Treatments Involving Rapid Solidification - S. Copley

Four main uses of directed energy beam (laser or electron beam) systems were described:

(1)

melting to produce modified surface microstructures after rapid resolidification, (2) alloying

by adding solute elements in the surface melted layer, (3) cladding and (4) dispersing foreign

particles in the melted layer, such as hard particles to provide improved wear resistance. Overlapping trails can be used to produce complete surface coverage. When combined with material

deposition, these surface melting and resolidification processes also can produce bulk material.

To date, surface melting techniques have been demonstrated to produce microstructure refinement

and modified epitaxy. Undercooling in the melt before resolidification can allow extended solubility

and the production of amorphous phases in the resulting surface layer. Alloying techniques have been

demonstrated to produce homogeneous alloyed layers when there is sufficient convection occurring.

The feasibility of cladding and dispersion techniques have been demonstrated.

Roadblocks to further application of these surface treatment techniques include:

(1) Residual strains that can affect the stability and integrity of the modified layer, as by

causing surface cracking.

(2) Surface relief can be produced such that an originally smooth surface loses its smoothness.

(3) Methods for synthesis of materials need to be improved and the range of alloy systems that

can be treated by the various methods still need to be defined.

(4) Fundamental understanding of the surface processing needs to be improved, including

establishment of kinetic theories, development of metastable phase diagrams and control of

convection effects.

(5) Surface absorption processes need to be understood and controlled.

D. Impediments to the Application of RSP Materials - T. Kearns

This topic began with an overview of current factors which can limit rapid implementation of RSP

technology and hence decrease the opportunity for early benefit to industry. Impediments exist in all four areas of the general cycle toward advanced material utilization. Roadblocks exist in new alloy and process development, material characterization, scale up and production, and system

application. The problems facing RSP, however, are not particularly unique.

University involvement is generally limited to suitable thesis topics and suffers from small new equipment budgets. Industry on the other hand is almost entirely driven by perceived economic

benefit and can not justify long term basic research. The government, through its in-house

698
laboratories and contract programs, does offer some relief to this problem, but unfortunately must

spread its limited resources and manpower among many potential advanced technologies.

One promising solution is the current increased interaction between the industrial and university

community to develop the needed technologies. Although problems can exist with proprietary rights

and the freedom to publish, solutions can be achieved to the benefit of all.

There has been progress in the past few years in increased availability of advanced RSP material.

However, having materials to test and initially characterize is only a first step. Increased

utilization of a new material requires an expensive engineering data base. In order to generate one efficiently, user needs and wants (dreams) must be made clear. This is not a simple task when a new technology may offer radically new opportunities. Innovative concepts must be developed in a parallel effort.

Scale up and production are always a major industrial risk. What is the best production process?

Today, there are a number of highly touted alternatives which make a decision on large capital

equipment expenditures extremely complex, not to mention the current economic state of the industry.

Will there be a lasting demand? Or will a new technology (like carbon-carbon composites and fiber

reinforced superalloys for high temperature gas turbine applications) supplant RSP? Certainly a heady question.

Final material application, whatever the material, rests with the designer. Industrial confidence in RSP advantages will not be achieved overnight. Generally ten to fifteen years are required

to advance a technology from initial feasibility to having components in production. Opportunities

for earlier implementation sometimes can occur, however, but only as a "fix" for poor initial design.

In summary, there is some successful production and application of RSP alloys in place today.

Furthermore, new experimental alloys certainly appear tantalizing. What is required is more alloy

development, greater awareness of the potential advantages of RSP, increased designer demand and

industrial commitment. The recognition of RSP as a major government thrust area is certainly

beneficial to the eventual application of this new technology.

TOPIC 2. Alternatives to Rapid Solidification Processing

A. Single Crystal Turbine Blades - C. Kortovich

Similarities between current cast single crystal turbine blades and candidate RSP blade alloys

were discussed. For the cast single crystal blades (SC), compositional variations across dendrites

are minimized due to the absence of grain boundary strengthening elements. However, there is more segregation in SC blades than is observed in RSP alloys; also, there are some low angle boundaries

present in the SC blades. Some of the inhomogeneities found in SC blades may be beneficial to fatigue behavior in service.

The experimental RSP blade alloys with directional grain structures having high modulus

orientations are superior in stress rupture to cast directionally solidified blades (DS) and cast

SC blades having the <100> orientation. Based on equiaxed cast blades, and testing in stress rupture

at 1800°F and 30 KSI stress level, the temperature capability increases as follows: DS + 30°F;

SC + 40°F, RSP + 110°F. An appropriate reminder was made regarding the experimental status of todays

RSP blade alloys (concerns of cost, inspectability, reproducibility, etc.); the same was true for

cast SC blades ten years ago and they are now an established production item.

Additional competition for RSP alloy blades comes from oxide dispersion strengthened alloys,

directionally solidified eutectic alloys, and from fiber reinforced superalloys.

B. Advanced Composites - C. Blankenship

Organic matrix composites (OMC) are a very strong competitor for RSP aluminum alloys. The most

significant disadvantage of OMC is a low strain capability prior to fracture. Todays OMC have a strain limit of 0.004; a capability of 0.008 to 0.01 is needed. By comparing 1980 aircraft construction (about 78% al uminum, 3% OMC) with two different scenarios for the 1990's, one (which offers about

a 30 percent Weight reduction using OMC) has about equal use of OMC and aluminum (40% of each, with

about 10% titanium), the other is more optomistic and favors OMC at 60% with aluminum and titanium

at about 20% each.

Both OMC and RSP Aluminum suffer from lack of innovation in manufacturing components. For OMC,

we need to use methods of carrying loads around corners, and stop using methods we use to build

"metal" airplanes. For RSP Aluminum, we need to have new concepts for higher structural efficiency.

699
such as the superplastic forming-diffusion bonding methods available for titanium alloy construction.

Regarding automation in manufacturing, rapid progress is being made on OMC; they could easily

surpass the efficiency of manufacturing using RSP Aluminum and eventually be cheaper! Thus, the

pressure is on for new fabrication concepts and automated manufacturing practices for RSP Aluminum

to be competitive with OMC, While organic matrix composites are a real challenge to RSP Aluminum,

the jury is still out.

The major concerns for aircraft structures intended for long term service are environmental

effects and structural integrity. Current experience for OMC is quite satisfactory, however, newer matrix resins are probably suspect. The clear message is that OMC components are flying and are becoming accepted.

Joining technology state-of-the-art is not equal to fabrication technology for both OMC and RSP

Aluminum. Laminated metal structures would offer some structural advantages; however, the data base

that would allow designer acceptance does not exist for these bonded structures.

A metal matrix composite (MMC) of short, discontinuous fibers in an aluminum matrix probably

would be a competitor for OMC and RSP Aluminum. Some MMC are currently using short fiber silicon

carbide in an RSP Aluminum alloy matrix.

C. Magnetic Materials - L. Johnson

Amorphous magnetic materials (AMM) for use in power distribution transformers have considerably

lower core losses than anything else available today. The material in these transformers represents

about 35 percent of the cost; there is a potential replacement cost for transformers in the 20 to 30

billion dollar range. The likelihood of this replacement taking place depends on both performance

and cost. While the performance of the AMM is quite superior, improvements have been made in the

grain oriented steels, so the performance of the competition is also improving. The cost outlook

depends on whether you are "buying or selling,"

The industry is very conservative; they have been making transformers the same way for 80 years.

Twenty units of the 20 KVA power level have been made from AMM, at least three of them have been in

service on investor owned utility poles for over one year, and they are performing very well. The

historical trend is for the cost of AMM to be coming down, and they are expected to impact the

market when their cost decreases to a level three times that for M4 steel; the lower the cost the

larger the impact. The beginning of the earliest market penetration is expected to be 1985,

The economics for using AMM is smaller transformers, motors, etc. is not favorable. Transformers

in the 25 to 50 KVA range are considered "small" and they are the market target. Motors have been

built and their performance demonstrated, however, the 30 cents/lb material cost for the conventional

material and the simple stamping process for laminated sheet production make it very difficult for

the AMM to compete.

The cost of manufacturing transformers using AMM is estimated to be higher than for todays

practice. Some investment in capital equipment will probably be required, but the future manufacturing cost is not considered a significant problem. The current down turn in housing construction and

general business level makes it difficult to justify capital expenditures at this time.

D. The VADER Process - W. Boesch

The VADER process is a "ceramic-free" melting process with a current capacity of 150 lb ingots

(billets) formed in a static casting mode. The process is also capable of a dynamic casting mode

where the mold is slowly withdrawn giving a constant height of fall for molten droplets. The mold

can also be rotated to form hollows. The process is now a fully programmable melt process using

servos and a microprocessor. A 20 Ib/min melt rate is in use; this is considerably higher than for

other vacuum arc remelting processes (VAR).

Since no water cooled crucible is used, the VADER process requires only about 40% of the energy

used for normal VAR processes. Up to 7 inch diameter ingots of IN 100 alloy have been directly

ultrasonically inspected. The ingot microstructure is equiaxed and fine grained, typically ASTM 3. There is a very low residual stress present, the ingots are free from macrosegregation, the ingots are isotropic, and they have shown significantly improved hot workability. There are some ingot

surface quality problems where it comes into contact with the mold. Ingots of several different

alloys have been produced so far, including steels, nickel-base alloys, and aluminum alloys.

The maximum size ingot produced so far by the VADER process is 10 inches in diameter. While the

process is "ceramic-less," care will be needed to minimize ceramic content in the electrodes that

700
are melted, since ceramic inclusions in the electrodes are likely to fall into the ingot.

The VADER process billets can be hot die forged without difficulty. Ingot sections have also

been forged in a single blow from a height of 2 inches to 1/8 inch with no cracking. The same melting action, i.e., the formation of droplets and build-up of the cast ingot is available in

electroslag remelting (ESR) also.

TOPIC 3. New Horizons for Rapid Solidification Processing

A. Surface Melting

1. What steps can be taken to keep residual stresses under control in surface melting treatments? - R. Schaefer

Both concerns and suggested solutions to the residual stress issue prevalent in surface melted

structures were discussed. Surface melting frequently results in substrate warping, plastic

deformation, hot tearing and crack initiation in the substrate. The nature of these problems have

been demonstrated in 2024 aluminum alloys using an electron beam for surface melting. The extent of

the problem is strongly dependent on alloy selection, substrate temperature, details of beam

manipulation such as oscillations rastering patterns, and melt-zone shape. For example, it was shown that lowering the traverse rate of the electron beam or oscillating the beam during melting

reduced crack density and size. Overlapped melt passes also helped in minimizing crack densities.

The importance of shifting the thermal stress to the free surface by control of processing variables

was also identified during the discussion.

2. Can we develop an integrated surface treatment processing method for shaped parts ?-P. Strutt

The methodology was presented for the rapid processing of materials using a computer controlled

electron beam system. The system incorporates features such as computer control, real-time

inspection using sensors such as video, infra-red and x-ray detectors. This procedure has been

used or can be used for rapid solidification, surface alloying, rapid heat treatment and joining

of novel structures. For example, this procedure has been used for 52150 bearing steels to produce

an ultra refined grain and martensitic microstructure in less than 0.1 sec, which would take more

than an hour using bulk processing methods. The integrated materials processing allows controls

over the nature and scale of solidification structure, size and distribution of high temperature

carbide phases, scale of martensitic structures and degree of tempering.

B. Melt Spinning

1. Can we devise methods of strip casting in the 0.005 inches to 0.050 inches thickness range?

- R. Maringer

Strip casting methods have extreme flexibility and provide a number of innovative ways of

fabricating structures. With proper selection of nozzle geometry, and the distance separating the

molten stream outlet and the rotating wheel, it is quite possible to build up microcrystalline

strips and subsequently hot roll them for full densi fication if necessary. It may also be possible

to directly cast multilayered structures by this method.

2. What are the opportunities for materials synthesis by rapid melt mixing/quenching

techniques? - B. Giessen

An approach was presented for fabricating microdispersions of high melting particles in lower

melting matrices by mixing melts in suitable proportions. The microalloying approach permits rapid

mixing of reaction components as solutes in the molten matrix followed by quenching. To obtain an end product of finely dispersed compound, such as AB in a matrix of C, the method proposes the

mixing of two liquids, one consisting of C + A and other C + B such that both are complete solutions.

Rapid quenching of the mixture of these two liquids will produce a fine dispersion of AB in a matrix of C. Examples of the liquid pairs and the reaction products in A1 , Cu, Ni , Co, Fe, Ag and

Pb - base alloys are:

A1 Al(Si) +Al(Mg) —> A1 + Mg

2 Si

A1 + Cu(0) —> Al(Cu) + AI

2 O

3

Cu Cu(0) + Cu(Al) —> Cu + AI

2 O3 (Cu, Nb) + Cu (Ge, A1 ) —> Cu(Nb) + Nb

3 (Ge, A1

701
Ni Ni(Th) + Ni(0) —> Ni + ThO^

Ni(Y) + Ni(0) —^ Ni + Ni(B) + Ni(Mo) —> Ni + Mo^^B Ni(Al) + Ni(Nb) —> Ni + Nb3(A1,Nb)

Co Co(C) + Co(W) —> Co + WC

Fe Fe(C) + Fe(Mo, W) —> Fe + (Mo, W)C

Ag Ag(Cd, Bi ) + Cu(0) — Ag(Cu) + CdO, Bi^O^

Pb Pb(Sb) + Pb(In) Pb + InSb

These fabricated products may be used as dispersion strengthened alloys for elevated temperature

application (A1-, Cu-, Ni-base alloy systems), wear resistant alloys for tool bits (borides,

carbides, etc.), corrosion resistant alloys (Cu-Ni base alloys), electric contact materials (Ag-base

alloys), superconductors [Cu + Nb 3 (Ge, A1 ) systems], magnetic alloys, bearing materials (Cu-Sn and

Zn-base alloys) and metal -ceramic composites.

Some concerns expressed for these concepts were (a) need for complete mixing, (b) thickness of

melt spun ribbons, (c) different melt temperature requirement for each molten alloy, (d) need to maintain proper ratio of mixture of liquids, and (e) existance of other methods for specific systems

for which it may be simpler and easier to control the process parameters.

C. Powder Metallurgy

1. What techniques are available, and what might be tried for consolidation of metastable

powders? - H. Fraser

New consolidation methods are needed in order to preserve the refined microstructure of

rapidly solidified particulate. These methods should limit thermal excursions while providing a good metallurgical bond between particles. Consolidation methods such as dynamic powder compaction

to include shock consolidation and vibratory compaction, and incremental consolidation such as layer glazing, in- situ consolidation of splats or osprey, VADER, and plasma spraying should be

considered. Some of these processes are not without problems. These compacts have remanent voids

and high dislocation densities which have to be considered in defining the fabrication parameters.

Most of the discussion on this subject dealt with process variables such as the effect of vacuum level on powder compaction, and the useful Iness of an electron beam to identify the quality

of the compacted product. The direct rolling of powder compacts to consolidated forms was presented

as a new commercial method. Sheets and plates of 7091 aluminum alloy and titanium alloys have been

fabricated directly from RSP powders.

2. Can we develop processes for optimization of undercooling before nucleation? - M. Flemings

In reference to J. Perepezko's prior oral presentation, it was suggested that RSP methods can be

developed to enhance undercooling. Processes such as emulsification in molten salts or glasses can

be commercially developed to improve undercooling. Undercooling up to 500°C has been achieved by

such processes. Molten metals can be injected either into colder liquids or processes developed in

a manner similar to electroslag remelting. Electroslag remelting has the advantage of protecting

the melt from environmental contamination.

Selection of proper fluids are important for this process. Water atomization, although used

extensively for stainless steel powder fabrication, is not satisfactory because of steam which acts

as a barrier to rapid cooling.

702
Author Index

Abbaschian, G. J. , 129

Adair, A. M. , 609

Adam, C. M. , 70, 629

Adler, R. P. I. , 448

Ahrens , T. J. , 672

Aigeltinger, E. H. , 186

Amato, R. A., 547

Anderson, I. E. , 178

Anderson, R. E. , 488

Andrus, J. S. , 590

Apelian, D. , 105

Ast, 0. G. , 384

Ayers, J. 0., 360

Aziz, M. J. , 113

Bass , M. , 303

Beck, D. , 303

Beck, R. , 692

Beckerman, L. P. , 579

Biancaniel 1o, F. S. , 390, 476

Billman, F. R. , 532

Blankenship, C. , 699

Boesch, W. J., 259, 700

Boettinger, W. J. , 45, 390, 476

Bourdeau, R. G. , 640

Broderick, T. F. , 585

Brosius, K. S. , 229

Cavin, 0. B. , 264

Chang, K. M. , 337

Chang, W. H. , 500

Chatterjee, 0. K. , 240

Cheeks, T. L. , 118

Chin, H. A. ,488, 590

Clauer, A. H. , 291

Clough, R. B. , 403

Cohen, M. , 1, 124, 171

Cook, M. M. , 597

Cooper, K. P. , 90

Copley, S. M. , 303, 698

Corb, B. W. , 678

Coriel 1 , S. R. , 45

Cote, P. J. , 429

Dabkowski, F. , 435

Das, S. K. , 559

Davis, L. A., 559

DeCrescente, M. , 698

Dickson, J. , 609

Dimiduk, D. M. , 186, 276

Domalavage, P. , 678

Domingue, J. A., 259

Doraivelu, S. M. , 519

Ehlers, S. K. , 240

Eisel stein, L. E. , 246

Ekvall , J. C. , 615

Fatemi , M. , 603

Flemings, M. C. , 28, 702

Fraser, H. L. , 56, 702

Froes, F. H. , 585

Gegel , H. L. , 519

George, D. B. , 640

Giamei, A. F. , 186

Giessen, B. C. , 435, 439, 701

Glasgow, T. K. , 377

Glicksman, M. E. , 83, 118

Gourdin, W. H. , 565

Grant, N. J. , 209, 215, 678

Green, H. M. , 500

Griffith, W. M. , 597

Gunasekera, J. S. , 519

Hahn, S. , 443

Hall , E. L. , 118, 337, 377

Hara, Y. , 678

Hare, A. W. , 672

Hi 1 deman, G. J. , 532

Hirth, J. P. , 135

Horowitz, E. , 416

Horton, J. A., 264

Hsu, S. C. , 448

Huang, S. C. , 337, 377

Hughes , J. R. , 365

Hughes, S. E. , 488, 590

Isserow, S. , 323

Jackson, A. G. , 585

Jackson, M. R. , 118, 198, 270, 547

Jacobson, L. A. , 609

Johnson, A. M. , 650

Johnson, L. , 700

Kang, S. , 209

Kapoor, D. , 559

Kasiraj, P. , 672

Kattamis, T. Z. , 476

Kear, B. H. , 483

Kearns, T. , 698

Kelly, T. F. , 124

Kelm, J. S. , 650

Kent, W. , 692

Khan, A. S. , 348

Kim, Y.-W. , 597

Kissinger, R. D. , 259, 510

Koch, C. C. , 264

Koczak, M. J. , 223

Kortovich, C. S. , 692, 699

Koska, D. , 672

Kuli, J. C. , Jr. , 532

Kumar, S. , 223

Kurup, M. , 312

Kushnick, J. H. , 662

Langenbeck, S. , 629

Larson, J. L. , 590

Lawley, A., 223, 229

LeBeau, S. E. , 90

Lederich, R. J. , 281

Lemkey, F. D. , 672

Levi, C. , 458

Lewis, B. G. , 483

Lewis, R. E. , 615

Liebermann, H. H. , 365

Lipsitt, H. A. , 234, 240

Liu, C. T. , 264

Lu, Y. Z. , 286

703
Mahoney, J. F. , 458

Malas, J. C. , 519

Maringer, R. , 701

Martin, P. L. , 252

Masumura, R. A. , 178

Matson, L. E. , 519

Maxwell, R. E. , 365

McDevitt, D. D. , 129

Mehrabian, R. , 1, 96, 397, 403, 416

Meier, G. H. , 348

Meisel , L. V. , 429

Mendiratta, M. G. , 240

Meschter, P. J. , 281

Miller, J. A. , 640

Miracle, D. B. , 234

Mohan, S. , 229

Moll , J. H. , 686

Morgan, J. t. , 519

Moteff, J., 585

O' Handley, R. C. , 678

Olson, G. B.,. 171

O'Neal, J. E. , 281, 452, 579

Paik, J. S • > 28

Paliwal, M • > 105

Palmer, I. G. , 615

Panchanathan

,

, V. , 323, 443

Pande, C. S.,, 360

Patterson, R. J. , II, 70

Peng, T. C. , 452, 579

Perel , J. , 458

Perepezko, J. H. , 28, 90

Peterson, L. G. , 198

Peteves, S. D. , 129

Petit, J. I. , 532

Pettit, F. S. , 348

Pickens, J. R. , 150

Quinn, D. M. , 429

Qui st, W. E. , 615

Rairden, J. R. , 468

Rath, B. B., 178, 603

Ray, R. , 323

Rice, S. L. , 483

Richmond, J. J. , 90

Ridder, S. D. , 96

Rigney, D. V., 468, 650

Ritter, A. M. , 270

Rizzo, F. J. , 686

Rosen, M. , 397, 416

Rowe, R. G. , 365

Ruano, 0. A. , 246

Shiohara, Y. , 28

Siemers, P. A. , 547

Simon, J. W. , 629

Smashey, R. W. , 198, 650

Smith, C. H. , 215

Smith, J. J. , 397

Smith, J. S. , 468

Smith, R. W. , 105, 468

Smugeresky, J. E. , 565, 572

Sprague, R. A. , 500

Starke, E. A. , Jr. , 150

Strutt, P. R. , 312, 464, 483, 701

Taquir, A., 312

Taylor, S. , 458

Thompson, E. R. , 672

Tien, J. K. , 259, 510

Tietz, T. , 697

Tucker, T. R. , 291

Tuli , M. , 464

Vander Sande, J. B. , 56, 124

van Reuth, E. C. , 609

VanStone, R. H. , 365

Void, C. L. , 178

Voorhees, P. W. , 83

Vreeland, T. , Jr. , 672

Wadley, H. N. G. , 403, 416

Wadsworth, J. , 246

Wakeman, T. G. , 650

Walker, J. A., 532

Wang, J. R. Y. , 559

Wayne, S. F. , 483

Whang, S. H. , 286, 435, 439

Williams, K. A. , 234

Williams, J. C. , 135, 252

Woodyatt, L. R. , 229

Yurek, G. J. , 171, 424

Zielinski, P. G. , 384

Sakata, I. F. , 615

Sastri , S. , 443

Sastry, S. M. L. , 452, 579

Sawhney, D. L. , 662

Scarbrough, J. 0. , 264

Schaefer, R. J. , 390, 397, 701

Sekerka, R. F. , 45

Shanfield, Z. , 458

Shechtman, D. , 96, 397, 476

Sherby, 0. D. , 246

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