Fine structure of austenite produced by the reverse martensitic transformation

FINE

STRUCTURE REVERSE

OF AUSTENITE

MARTENSITIC G. KRAUSS,

PRODUCED

BY THE

TRANSFORMATION* Jr.t

The fine structure introduced into austenite by the reverse martensitic transformation was investigated by transmission electron microscopy in an iron-33.5 w/o nickel alloy. Direct martensitic transformation induced by cooling to -195’C produces finely twinned martensite surrounded by retained au&en&e containing relatively uncomplicated dislocation arrangements. On rapid heating to 450°C, the reverse transformation replaces these structures with tangled and jogged dislocations in concentrations up to 10” per cm“, and thereby accounts for the marked strength increases observed in reversed austenite. In some previously martensitic areas of the reversed austenite, twins up to 5 microns in width are observed with standard metallographic techniques. Selected-area electron diffraction con6rms the twinned nature of these structures, and it is found that the twins are quite imperfect and contain a high concentration of tangled dislocations. STRUCTURE

D’AUSTENITE FINE PRODUITE PAR TRANSFORMATION MARTENSITIQUE INVERSE La fine structure apparaissant dans l’aust&ite par transformation martensitique inverse dans un alliage fer-33,5% de nickel en poids, a BtB Btudibe par transmission au microscope Blectronique. La transformation martensitique directe provoqube par refroidissementjusqu’ B - 195’C produit une martensite fine et ma&e, entourbe d’austbnite rbsiduelle contenant des arrangements de dislocations relativement simples. Pendant un chauffage rapide jusqu’8 45O”C, la transformation inverse remplace ces struotures par des dislocations enchev%&es et des dislocations crans jusqu’$. des concentrations de 101’ par cme. Ceci explique l’augmentation de la r&stance observee dans l’austbnite inverse. Dans quelques plages de I’aust6nite inverse, qui Btaient antbrieurement martensitiques, des mLcles jusqu’8 5 microns de largeur peuvent Btreobserves par des techniques m&tallogrephiquesnormales. La diffr’actionBlectronique sur des plages &e&ion&es confirme la nature m&&e des ces structures, et on a trouve que les m&cles sont compl&ement imperfaites et qu’elles contiennent un grand nombre de dislocations enchev%r&es. FEINSTRUKTUR

DES

DURCH RUCKUMWANDLUNG AUS MARTENSIT GEWONNENEN AUSTENITS Die Feinstruktur des durch Riickumwandlung aus Martensit entstehenden Austenits wurde an einer Eisenlegierung mit 35 gew. oh Nickel mittels Durchstrahlung im Elektronenmikroskop untersucht. Bei der durch Abktihl ung auf - 195°C in Gang gesetzten direkten Martensitumwandlung entsteht fein verzwillingter Martensit, der von nicht umgewandeltem Austenit umgeben ist und relativ einfache Versetzungsanordnungen enth<. Eine schnelle Aufheizung auf 450°C hat die Riickumwandhmg zur Folge, bei der Netzwerke von Versetzungen mit Sprtigen in Konzentrationen bis 10”/cme entstehen; dies erkliirt die betrlichtliche Festigkeitszunahme, die an diesem Austenit beobachtet wurde. In dem aus den martensitischen Gebieten entstendenen Austenit wurden mit herkbmmlichen metallographischen Methoden Zwillinge bis zu 5~ GriiBebeobachtet. Durch gezielte Elektronenbeugungsversuche wurde der Zwillingscharakter dieser Strukturen best+@ und gefunden, da13die Zwillinge sehr unvollstiindig sind und dichte Versetzungsnetzwerke enthalten.

1. INTRODUCTION

In a recent investigation(l) it was shown that the austenite produced by one or more cycles of the reverse martensitic transformation in iron-nickel alloys is significantly altered in microstructure, annealing behavior and strength characteristics, compared to the virgin austenite. In particular, “reversal twins” were observed in specimens heated approximately 50°C above the At temperature (temperature at which the martensite to austenite reaction is just completed), and the yield strength of such specimens was increased as much as 2.5 times that of annealed austenite. The present investigation was undertaken to determine by transmission electron microscopy the changes in fine structure which occur when the martensite is converted to “reversed” austenite by the reverse martensitic transformation, * This research was conducted by the author as a Research Staff Member in the Department of Metallurgy, MIT. Received August 23, 1962. t Presently at Max Planck Institut fiir Eisenforschung, Diisseldorf, Germany. ACTA METALLURGICA,

VOL. 11, JUNE 1963

499

and to relate these changes to the microstructural and strengthening characteristics of the reversed austenite. 2. EXPERIMENTAL

PROCEDURE

A vacuum-melted iron-33.5 w/o Ni alloy with an M, of -124°C and an Af temperature of about 400°C was selected for this investigation because of the absence of transformation at room temperature. In lower nickel alloys, surface martensite may appear in bulk specimens during metallographic preparation, and martensitic transformation may occur during thinning, t2) No such phenomena were observed in the iron-33.5 w/o Ni alloy. Sheet 0.004 in. thick was prepared by alternate rolling and annealing of a 0.300 in. thick blank which was initially annealed for 64 hr at llOO’%. Discs 0.125 in. in diameter were then punched from the sheet, annealed at 1000°C in evacuated vycor capsules for 1 hr, and water quenched to room temperature without breaking the tubes. The specimens were then subjected to either a martensitic

ACTA

500

transformation reversal

treatment

treatment.

was formed

or a transformation

Martensite

by cooling

reversal

was performed

tensitic

specimens

(about

to -195°C

perchloric

ation. stop”

the mar-

temperature technique

electrolyte

thin the heat-treated

for 30 min, and

room

was

into

using an

developed

discs for transmission

Edges of the discs were lacquered

to

examin-

with “Micro-

to induce uniform thinning of the central parts

of the sample without

selective

Just before perforation,

(1 min at 40 V) the protection

of

the

edges

was

attack

removed.

of the edges.

When

VOL.

and

60 per cent)

by up-quenching

from

molten salt at 450°C for 2 min. A “miniaturized” Bollmannc3) acetic-5%

METALLURGICA,

polishing

was

resumed, there usually developed in the thinned center of the disc a hole which was eventually ap-

11, 1963 3. DISCUSSION

OF

RESULTS

3.1 The$ne structure of marten&e and retained austenite The fine (112), martensitic twinning reported by other investigators(6~6) was also found to be prevalent in the iron-33.5 w/o nickel alloy examined here, the martensite

being

produced

by

cooling

to -195°C.

Numerous examples of well-defined narrow twins were observed; analysis by electron diffraction in these cases frequently showed that either a (IlO), or (113), For

direction

both

these

was normal

orientations,

to the foil surface. the

(112),

planes were normal to the foil surface; in no overlapping

of adjacent

twinning

this resulted

twins

and relatively

proached by openings growing in from the unprotected

few fringes. The twin spacing was approximately 100 A but variations up to 50 A either way were

edges

sometimes

or by

between enough

a second

the holes

internal

proved

for transmission

at this

step

was

hole.

The

examination.

made

regions

to be consistently

possible

process under a microscope.

thin

Good control

by

following

the

This method of thinning

taining

observed.

long

narrow

In addition

to the areas con-

martensitic

twins,

segmented

twins as discussed by Shimizu(7) were also found. Fig. 1 is a transmission micrograph showing very uniform fine twinning over large areas in several

yielded good sample economy and introduced no cold work into the treated discs by any final cutting

martensitic

operation.

of a central twinned core surrounded by an outer heavily dislocated region with few twins. This is

After thinning,

in a Siemens

Elmiskop

the discs were examined

I at 100 kV.

The aperture allowed an area of 0.85 ,u in diameter to be illuminated for selected-area electron diffraction. Some surfaces

were examined shadowed

parlodion

Average ticular

by electron

of chromium-

densities,

were

It

N, of areas of par-

estimated

by

using

the

ex-

(1) where %, and fi, are the number the dislocation

lines with

of intersections

two normal

by

sets of grid

lines and Ll and L, are the respective lengths of the The foil thickness,

but

was

taken

t, was not determined

as 2000 A

experience of Keh and Weissman.(4) the possibility

contrast

that

as shown by the extension of the twins from boundary to boundary

across the martensite

plates in Fig.

A) are what appear to be stepped dislocation with alternate

a specimen Fig.

2.

segments

paralleling

1

images

t’he boundaries

that

not

at a given orientation

kept in mind

when the

based

on

the

This assumption,

all dislocations

show

of the foil, should be

dislocation

densities

come

under consideration. Tensile tests were performed on standard l-in. gage length, 0.250-m diameter specimens which were given standard heat treatments as discussed above. True stresses were calculated from the recorded load-elongation curves by assuming uniform elongation up to the maximum load.

the

of

at

-195°C

martensite,

A single-surface

by diffraction

smaller

boundary plane.

plates

are present.

made possible to

and crystallographic

transformed

Two

parallel,

and

reported@)

w/o Ni alloy consist

not the general case in the present 33.5 w/o Ni alloy

Other structural

directly

been

of the tine twins.

pressionc4)

grid lines.

has

plates in an iron-30.9

and 2. Also of interest in Fig. 1 (at the points marked

replicas.

dislocation

interest

microscopy

plates.

martensite

plate

corresponded

features

of

are shown

in

approximately trace analysis

from an area adjacent

showed (within

that

the

straightest

lo) to a (2591,. habit

The (112), twinning planes in both martensitic

plates were roughly parallel to (110}, planes. Both of these crystallographic features are consistent with the findings of Kelley and Nuttingc5) for an iron-20% Ni-0.8%

C alloy

rather than b.c.c.

in which A recent

the

martensite

analysis

is b.c.t.

of the marten-

sitic twinning by Cracker@) predicts only a slight difference in the habit planes resulting from the (112)[iil], (lOl)[iOl], shear mode in b.c.c. and b.c.t. structures, although the poles of both habit planes should he nearer the (3 10 15}, than the (259},, planes assigned experimentally. The limited accuracy of a trace analysis on a surface oriented by

KRAUSS:

REVERSE

MARTENSITIC

TRANSFORMATION

FIG. 1. Uniformly distributed fine

twins in adjoining martensitic plates. Note stepped dislocations at Sample transformed at - 195°C. Transmission, 40,000 x .

points marked A.

-

_

FIG. 2. Martensite plates and dislocation arrangements in retained austenite in sample transformed

at - 195°C. Foil is parallel to (112)~. Solid lines indicate directions and dotted lines the traces of planes on (112)~. Transmission, 40,000 X .

501

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502

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FIG. 3. Retained austenite surrounding martensite in a sample transformed at - 195°C. Note dislocations emanating from tip of narrow plate. Transmission, 28,000 x .

Examples

,+

retained

I’

4 ,’

the

dislocation

adjacent

both

1~

photomicrographs

between

log

and

locations

generated

are

or

slightly

which

curved.

the

moderate glide

dis-

by the direct

tend to be either straight

Complex

dislocation

in the reversed

were not

in

plates

densities in

relatively

in the austenite

were found

described,

Dislocation

lOlo per cm2, and

martensitiL transformation

60000

arrangements

to the martensitic

are shown in Figs. 2 and 3.

4E”ERSED 5’ /’ AUSTENITE

of

austenite

observed

tangles,

austenite

to be

in the retained

aus-

tensite. 3.2 Xtrengthening reversed austenite After

reversal

and

dislocation

arrangements

of the martensitic

the strength of the austenite produced higher

than that

of annealed

in

transformation, is appreciably

austenite.

This con-

sequence of the reverse transformation in the iron-33.5 w/o Ni alloy is illustrated in Fig. 4. The yield strength I 010

I

I 0 20

030

I 040

050

I 060

STRAIN

FIG. 4. Stress-strain curves for annealed and reversed au&mite in an iron-33.5 w/o nickel alloy. Reversal temperature = 450°C.

electron diffraction and the somewhat arbitrary selection a straight line corresponding to the slightly curved interface plane in Fig. 2 do not permit a closer

of

evaluation

of the habit-plane

indices than presented.

of the austenite is almost doubled by the reversal treatment, while the capacity to strain harden is greatly reduced. It should be kept in mind that these marked changes occur even though only approximately 60 per cent of the volume of this alloy has undergone the direct and reverse transAlloys in which a larger fraction of the formation. specimen takes part in the transformation are more effectively strengthened W”) by the reversal treatment. In an effort to explain

this strengthening,

thinned

KRAUSS:

foils of reversed

au&mite

mission

microscopy.

electron

figurations tangled

revealed

were

observed

which

were

high

throughout

characteristics

complex,

with intermediate

in

dislocation

the

martensitic.

Tangled

by trans-

dislocation

networks

that

the

are shown

in dis-

areas of lower dislocation

resulting

formation

interspersed

of

observed

in many other f.c.c.

and have been extensively Wilsdorfol)

et al.

the present

that

of ways

different

from

dislocations with

the

dislocation factor

of

Clearly, significant

in

the

direct

retained

martensitic

densities 10 greater the

straight

reverse

dislocation

also in

or

of complex

the

tending

reversed

transformation multiplication

be

a

specimens. has

caused

and interaction,

of jogged

dislocations, austenite

to

the

close-packed

the marked

f.c.c.

increase

of high concentrations

of tangled

dislocations.

These

configurations

distributed

generally

throughout

by the reverse martensitic

of the

transformation,

to inhibit not only the unpinning

but also their subsequent

movement.

3.3 General aspects of reversal twins Another prominent feature of the austenite produced

curved

to

jogged

of dislocations

associated The transformation.

differ,

of the introduction

would be expectedol)

austenite

FIG. 5. Networks 2

of jogged

smoothly

lattice

On the basis of this study,

and

in the reversed austenite are quite the

martensite

of plastic flow can be viewed as a direct consequence

tangles.

The dislocations

the vacancies

here

metals

stress the impor-

creation

was In

in the resistance of reversed austenite to the initiation

loops which can contribute

to the

case, it is possible

the et al.

with

by Kuhlmann-

in the formation

and dislocation

in a variety

discussed

Their arguments

tance of excess vacancies

dislocation

tangles

loops similar to the type described

are frequently

dislocations

dislocation

In

arrangements. Kuhlmann-Wilsdorf

are generated as a consequence of the sudden volume decrease which occurs during the transition from the austenite.

and

by

of vacancies or irradiation.

b.c.c.

regions

dislocation

discussed

the excess

in producing

the necessary supersaturation introduced either by quenching

dislocation nature

which seem to be important

examples

density are shown in Fig. 5, while Figs. 6 and 7 show loops and other debris scattered throughout The of tangled and jogged dislocations.

503

and perhaps in such a way as to provide vacancies

regions of the

of jogged

TRANSFORMATION

con-

densities

extensive Some

of these arrangements

Figs. 5, 6 and 7. locations

The quite

and

formerly

MARTENSITIC

were examined

were

dislocations

REVERSE

by

the

reverse

martensitic

transformation

formation

of the islands of structure

twins”.(r)

The identification

was

recently

vations, the

but

twinning

deduced now

of these regions as twins

from

electron

relationship

is the

called “reversal

metallographic diffraction

is cited

dislocations in austenite formed by reversal ot 450°C. Transmission, 80,000 x .

obser-

evidence below.

Fig.

for 8

504

ACTA

FIG. 6. Dislocation

FIG. 7. Tangled

METALLURGICA,

VOL.

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1963

arrangements in austenite formed by reversal at 45O’C. Note dislocation at points marked A. Transmission, 120,000 x .

dislocations in au&mite formed by reversal at 45O’C. debris scattered throughout the tangles. Transmission,

Note unresolved 80,000 x .

loops

dislocation

KRAUSS:

REVERSE

MARTENSITIC

TRANSFORMATION

FIG. 8. Reversal twins and surrounding diffuse-etching regions in a sample reversed at 450°C. HNO,-HCl-H,O etchant, optical photomicrograph, 200 x .

FIG. 9. Reversal twins in sample reversed at 450°C. Grooves are from an orientation dependent etching effect. Note how they change direction at twin boundaries. Chromium-shadowed replica, 16,000 x .

505

506

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METALLURGICA,

VOL.

11, 1963

illustrates

the

metallographic

appearance

regions and the diffuse-etching them.

The boundaries

be slightly the islands

of these

areas which surround

are well-defined

but tend to

curved rather than straight. Some of are relatively large, up to 0.11 mm in

length and 5 ,u in width, although distribution of sizes below this higher magnifications

there is a wide maximum. At

made possible

of surface

replicas

orientation

difference

by examination

in the electron between

microscope,

can be seen (Fig. 9) by the change in direction grooves

resulting

etching

not visible

from

an

of the

orientation-dependent

at lower magnifications.

higher magnification,

At this

more of the boundaries

to be straight-sided

the

the twins and matrix

but curved

boundary

appear

segments

are still present. The

metallographic

twinning Fig. in

relationship

10 where photomicrographs the

of

to the prior martensitic martensitic

condition

condition

are presented.

the reversed

structure

on the martensitic

the

reversal

areas is shown in of the same area

and

In Fig.

in the reversed

10(c) a tracing

of

in Fig. 10(b) is superimposed

structure

of Fig. 10(a) inside the

large annealing twin common to both microstructures. The reversal twins never appear in areas that were not

martensitic

in

the

previously

transformed

structure, and the twins derived from a given martensitic

plate

groups,

are

one martensitic Figs.

arranged

the orientation

8 and

in more

plate

to the other.

10(b) shows

clusters are roughly twin boundaries

or less parallel

of the groups changing Inspection

that the twins

from of

of certain

parallel to the { 11 l}y annealing

and trace analysis

of the boundary

(marked A) of the reversal twin in Fig. 11 places it within 5” of a {ill),, 3.4 Electron

microscopic

Transmission the reverse 0.004 in.

plane. evidence

examination

transformation

thick

specimens

of reversal

twins

of thinned foils in which had been carried out in confirmed

the

previous

metallographic evidence for the presence of reversal twins.(l) Fig. 11 contains a cluster of islands separated from their surroundings

by well defined

boundaries

and Fig. 12 shows one of the tips at a higher magnification. diffraction

Care

was taken

patterns

to

secure

either entirely

selected

area

within the island

of Fig. 12 or solely in the adjacent region. Analysis of the resulting patterns showed that there was indeed a f.c.c. twinning relationship between the Fm. 10. (a) Martensitic area after transformation at -195°C. Annealing twin boundaries are emphasized by dashed lines; (b) Same area after reversal at 450°C; (c) Superimposed tracings of structures within the common annealing twin of (a) and (b). HNO,-HCl-Hz0 etchant, optical photomicrograph, 200 x .

two adjoining structures. The beam was parallel to a [3&l], direction in the twin and a [i05], direction in

the

surrounding

superimposed

indexed

matrix.

Fig.

diffraction

13

shows

patterns

for

the the

KRAUSS:

REVERSE

MARTENSITIC

507

TRANSFORMATION

FIG. 11. Cluster of reversal twins in a sample reversed at 450°C. Dashed line is (ill)y is within 5” of boundary marked A. Transmission, 20,000 x .

trace which

FIQ. 12. Tip of a reversal twin shown in Fig. 11. Foil is parallel to (3Tl)y within twin and to (iO5)?, in surrounding matrix. Note dislocation tangles within twin. Transmission, 80,000 x .

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642

533

315

524

424

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FIU. 13. Superimposed electron diffraction patterns of the twin end matrix regions of Fig. 12. Matrix spots of the [iO5]y zone axis sre indicated by triangles and twin spots of the [341]y zone axis by circles. Solid symbols represent spots sctudly observed in the patterns.

(341), twin and (f05), matrix orientations, and the (001) stereographic projection in Fig. 14 illustrates the twinning relationships about a (ill), plane which are consistent with the indexed diffraction patterns. The projection indicates that twinning symmetry brings the [34lJ, pole into coincidence with the [TO&J,pole, and also reveals the origin of some of the major poles which appear on the diffraction pattern of the twin. The reversal twins shown in Figs. 8, 9, 10 and 11 differ significantly from the commonly observed annealing twins formed in f.c.c. metals during recrystallization and grain growth. One striking difference is their highly imperfect internal structure, the predominant feature of which is a very high ~onoentration of tangled dislocations. The reversal twin considered in Fig. 12 for example, contained within its boundaries an estimated dislocation density of 10” per cm2, one of the highest observed in this investigation. The mechanism by which the reversal twins form is not clear at this time. The twins are not present to any great extent in specimens reversed just above the A,(l) the structure after such a treatment consisting predominantly of diffuse-etching regions in the prior martensitic areas. Well defined reversal twins are first observed in specimens heated approximately 50°C above the Af. It is therefore likely that the formation of twins in the sizes and shapes described comprises the initiation of recovery after reversal. The strength of the reversed austenite is unaffected by the appearance of the twins, a result

FIG. 14. Standard [OOl] S~reogr&phic Projection of the orystellographic relationships between the reverse1twin and sustenitic matrix of Fig. 12. Twin symmetry about the (ill)r plane is shown. Plus signs indicate poles of the twinned region. Other indices are for poles of the surrounding matrix.

no doubt of the high dislocation densities within the twins and their bounda.ries. Once formed the reversal twins appear to be quite stable and many persist along with the high strength levels of the reversed austenite until recrystallization reestablishes a more perfect austenite. 4.

CONCLUSIONS

(1) Fine martensitic twins of the (112) type are frequently observed in specimens of an iron-33.5 w/o nickel alloy transformed at -195’C. The dislocation arrangements in the retained ausenite surrounding the martensitic plates are relatively uncomplicated and the dislocations tend to be either straight or gently curved. (2) The austenite which is formed by reversal of the martensitic transformation is quite imperfect, containing high concentrations of tangled and jogged dislocations with interspersed loops. The dislocation densities are about a factor of 10 greater than those observed in the retained austenite of transformed specimens. (3) The introduction of the complex dislocation configurations by the reversed martensitic transformation are considered to cause the marked strengthening of austenite subjected to cyclic ma~ensitio transformation. (4) The relatively large islands of structure observed metallographically after reversal 50°C above the A, temperature have been verified to be twins by

KRAUSS:

REVERSE

MARTENSITIC

selected-area electron diffraction. These reversal twins, although formed in the very early stages of recovery of the reversed austenite, are also quite imperfect and contain high concentrations of tangled dislocations. ACKNOWLEDGMENTS

This research was made possible by the Office of Naval Research under contract NONR-1841-35. The author is extremely grateful for the continued encouragement and advice of Professor Morris Cohen throughout the course of this investigation. Thanks are also extended to Walter Fitzgerald for his help in the initial stages of sample preparation, to Miriam Yoffa for her help in preparation of the manuscript and to the Ford Scientific Laboratory for supplying the vacuum-melted iron-nickel alloy.

TRANSFORMATION

509

REFERENCES

1. G. KRAUSS,

JR. and M. COHEN, Trans. Amer. Inst. Min. (Metall.) Engrs. to be published. 2. H. WARLIMONT, Trans. Amer. Inst. Min. (Metall.) Engrs.

221, 1270 (1961). n W. BOLLMANN. Phvs. Rev. 102. 1588 (1956). d. 4. A. S. KEK ani S. %EISSXANX, Confe;ence’on

the Impact of Transmission Electron Microscopy on Theories of the Strength of Crystals, Berkeley, California, July 1961. 5. P. M. KELLY and J. NUTTING, J. Iron St. Inst. 197, 199

(1961). 6. Z. NISHIYAMA, K. SHIMIZU and K. SUXNO, Mem. Inst. Sci. In&&r. Res. 18,71 (1961). 7. K. SHIMIZU, J. Phys. Sot. Japan 17,508 (1962). H. WARLIMONT, private communication.

:: A. G. CROCKER, Acta Met. 10,113 (1962). 10. K. A. MALYSHEV, N. A. BORODINA and V. A. MIR~ELSTEIN, Trud. Inst. Fisiki Metallov U~alskii FGial29, 399 (1958). 11. D. KUHLMANN-WILSDORF, R. MADDIN and H. G. F. WILSDORF, Strengthening Mechanisms in Solids p. 137. Amer. Sot. Metals, Ohio (1962).