Electron Microscopy Investigations ofFerrous Martensites

GARETH THOMAS

This paper is concerned with electron metallography of bcc or bct ferrous martensite inwhich particular attention is paid to the characterization of substructures. The transforma-tion substructure is complex and new results are reported on multiple {112} twinning. Thefactors controlling the martensitic substructure are evaluated and it is concluded that thestrength and deformation characteristics of martensite are the most important in determiningwhether the transformation deformation occurs by slip or twinning (or both). An appraisal isgiven of in s i tu transformations in the electron microscope particularly with r e g a r d to thenucleation problem and investigations of the interracial structures. The potential role of highvoltage electron micros copy is also discussed and some recommendations are offered forfuture investigations°

THE f i r s t results of transmission electron micros-copy studies of foils made from carbon and alloy mar-tensitic steels were published just over ten yearsago. '-a Since then numerous papers have appearedwhich deal largely with characterizing particularsteels, and relatively few attempts have been made toinvestigate quantitatively the problems associated withmartensite nucleation, the nature of the austenite-martensite interface, and the orientation relationshipsand substructural characteristics of i n d i v i d u a l t r a n s -formation events. Considerable work has been doneon the relationship between submicrostructure andproperties of steels, and recently systematic effortshave been directed a t understanding the role of alloy-ing elements in these relations. Some recent reviewpapers which contain electron microscopy data in-clude those of Wayman, 4 Magee,5 and Nutting.e Re-views of electron microscopy applications to investi-gations of phase transformations in general includethose of Thomas 7-g and Hirsch e t a l . I° These publi-cations describe in detail the important experimentaland theoretical factors involved in the use of electronmicroscopy and diffraction.

The present paper is concerned with electron metal-lography of ferrous martensites, particularly trans-formation substructures. Some discussion is also pre-sented of nucleation and other aspects of the transfor-mation° No discussion will be given of martensiteo t h e r than bcc o r bct ( i . e . , ¢ [hcp] martensite is ex-cluded) nor of s t r a i n induced martensite.

It is hoped that some of the ideas presented hereand the commentaries on both advantages and disad-vantages of electron microscopy may be useful in fu-ture investigations.

1) MARTENSITE NUCLEATION ANDTRANSFORMATIONS IN FOILS

The electron microscope provides information onmorphology, crystallography, and composition. The

GARETH THOMAS is Professor of Metallurgy with the Departmentof Materials Science and Engineering, University of California,Berkeley, Calif.

This paper is based on an invitedtalk presented ata symposium onFormationof Martensite in Iron Alloys sponsored by the IMD FerrousMetallurgy Committee and held at the Spring Meeting of The Metal-lurgical Society of A1ME, May 1970,in Las Vegas, Nev.

METALLURGICAL TRANSACTIONS

analyses are usually c a r r i e d out on foils preparedfrom bulk transformed specimens, but dynamic exper-iments inside the microscope can also be done di-rectly on f o i l s 2 However, even a t 500 to 1000 kv foilsmust be 2 ~ thick or less.1,.2 for electron t r a n s m i s -sion and this is of the same o r d e r of magnitude as thedimensions of single martensite crystals. Thus, infoils, the transformation conditions are essentiallytwo-dimensional in nature and are quite different fromthose involved in bulk transformations. Consequentlythe crystallography and substructure may then also bequite different. These problems were f i r s t demon-strated by Pits chs who studied the martensite t r a n s -formation directly in several ferrous alloys.

Except for the f i r s t direct observations of the fcc-hcp transformation in cobalt ~3 and that of heterogene-ous nucleation of y ' in AI-Ag ~4 and surface precipita-tion in A1-4 pct Cu is in the electron microscope, t h e r eare as yet no other direct observations of nucleationprocesses. In the simple fcc ~ hcp systems, theorypredicts specifically defined imperfections for the nu-cleation site, viz. dislocations and stacking faults, andit is possible to identify and so locate the nucleationsite prior to the transformation process. As a resultit has been possible to follow in detail both nucleationand growth events and the theory of heterogeneous nu-cleation for these transformations has been confirmed.As far as fcc ~ bcc or bct martensites are con-cerned, the nucleation theory is not at present verysatisfactory. There is no clear evidence for nuclea-tion sites except for the case of strain induced mar-tensites where intersecting ~ martensite and slipbands, and twin boundaries have all been observed togive rise to c~ martensite crystals.~s, tv Magees hasemphasized the importance of twin boundary s i t e s ,heterogeneous nucleation in general, and autocataly-sis.

Fig. i indicates that dislocations and stacking faultsmayhave caused the formation of the martensite inthe Fe-8Cr-IC alloy. It is also possible that these de-fects are a consequence of the formation of marten-site. Direct observations of transformation in thissystem are planned and should help to distinguish be-tween these possibilities.

The present thermodynamic theory requires the ex-istence of "embryos:" Pati and CohenIs calculatedthe probability of observing embryos as 2 in 105. War-

VOLUME 2, SEPTEMBER 1971-2373

F i g . 1--(a) to (d) F e - 8 C r - I C bulk t r a n s f o r m e d s h o w i n gs t r u c t u r e of austenite and m a r t e n s i t e a s s o c i a t e d with twinned(225)y plates. The p r i m a r y (112)c~ t w i n s do not extend a l la c r o s s the l a r g e r p l a t e in (d). C o u r t e s y C. M . Wayman(Ref. 20).

limont19 interprets this probability in terms of thenecessity of examiningan area ~ 20 a m 2 in order tofind one nucleus of the most effective size and struc-ture. Althoughthese are almost impossibly smallprobabilities, the fact remains that martensitic reac-tions do occur in thin foils, probably as a result ofstrain induced or autocatalytic events or because con-straints may be removed on thinning.

As far as direct observations of the transformationare concerned, several important factors must he con-sidered. For example:

1) It is essential to have a goniometric specimen(double-tilt or tilt-rotation) holder so that contrast}xperiments and diffraction data can be obtained at alltemperatures. The heating-cooling response of the

2374-VOLUME 2 , SEPTEMBER 1971

s t a g e m u s t be a d e q u a t e to p e r m i t t r a n s f o r m a t i o n t oo c c u r a n d c o n t a m i n a t i o n m u s t be a v o i d e d ( s e e R e f . 21).

2) T h e a r e a b e i n g e x a m i n e d s u f f e r s s o m e b e a mh e a t i n g a n d c o n s e q u e n t l y a s t h e Ms t e m p e r a t u r e i s a p -

METALLURGICAL TRANSACTIONS

proached the transformation will take p l a c e outsidethe area being examined.21 In all experiments t o datenucleation events must have o c c u r r e d outside the fieldof view.

3) Because the transformation occurs very rapidly,it is impossible t o study both the morphology ( i . e . , t h eimage) and the crystallography (diffraction pattern)during the transformation. It is helpful, therefore, touse s ing l e crystals of known orientation, for which thebulk crystallography is already documented. In mostsituations it is necessary t o repeat experiments,studying one particular feature at a t i m e . It will alsobe necessary t o use high speed, high resolution movietechniques to r e c o r d the transformation.

4) It is difficult, if not impossible, to mount a foil inthe specimen holder without introducing some strains.The strained a reas may then become preferential nu-cleation sites.

5) Some idea of when a transformation is occurringcan be obtained by careful observance of the t e m p e r a -ture rise associated with the r e l e a s e of the drivingforce. Our observations show a temperature rise of+30°C for Fe-24Ni-4Mo-0.28C and +60°C for Fe-25Ni-0.3V-0.3C during cold s tage experiments on themartensitic transformation in foils of these steels.

6) Because of elastic relaxation, the macroscopicshape change in foils may not r e q u i r e the amount oflattice invariant s h e a r (slip or twinning r e q u i r e d bythe phenomenological theories) necessary in bulk sam-ples. Direct observations of martensite formation andgrowth will enable analysis t o be made of the o r ig in ofthe martensite substructure, i . e . , contributions fromtransformation deformation and subsequent accommo-dation deformation may be distinguished by comparingfoils of bulk transformed s t e e l and transformed foils(see Sec. 2).

7) It is not yet c l ea r that the transformation in foilsis always different from that in bulk, i . e . , differing incrystallography and morphology. T h e r e is simply notenough data from bulk and foil specimens in the samealloys.

8) An important factor in the transformation offoils is the apparent change in M s temperature withspecimen thickness and cooling r a t e , e . g . , due to pre-cipitation or carbon segregation in austenite. The up-per temperature l imi t of M s has been observed to de-c r e a s e when the specimen thickness is less than about1 ~.~9 Warlimont19 has recently studied the functionalM s dependence of thickness in F e - N i alloys. He foundtwo opposing effects which produced a scattering ofM s to lower as well as to h ighe r temperatures. Heexplained t h e s e results in t e r m s of the distribution ofnuclei per unit volume and t o the effects of imageforces on martensite nuclei (assuming heterogeneousnucleation). On the b a s i s of the embryo theory, sup-posing t h e r e are 107 embryos per cu cm at the f i r s tnucleation stage, ~a then for a foil of volume 10-7 cucm (1000 ~ thick, by 1 mm by 1 mm) probably only oneembryo would be present. Due t o the g r e a t e r penetra-tion at high voltages a foil 2 ~ thick should be t r a n s -parent and in this case the number of embryos in-c r e a s e s to 20.

The work of Easterling and Swannee showed thats m a l l austenite particles in copper did not transformto martensite except a f t e r elastic deformation even

though the bulk M s is 600 °C. Kinsman2s showed, how-e v e r , that if constraints a r e removed, e . g . , by polish-ing the surrounding copper away, the austenite thentransforms spontaneously.

Finally it is worth emphasizing that more attentionshould be paid t o electron diffraction. Although e l ec -tron diffraction patterns a r e widely used for orienta-tion determinations, no attempts seem t o have beenmade to investigate diffraction patterns before andduring the y - a transformation. Such work would befacilitated by u s i n g s ing l e crystals of appropriate ori-entation. Examination of the diffraction patterns, par-ticularly the diffuse reflections, could provide i n f o r -mation such as 1) what happens t o the austenite at thetransformation star t , 2) the structure of the f i r s tformed martensite, 3) the shape of the f i r s t formedcrystal. Double diffraction effects must be accountedfor 9, 10

2) SUBSTRUCTURE OF MARTENSITE

1) General Morphology

T h e r e are two genera l morphologies of f e r r o u s m a r -tensites, viz. dislocated laths and twinned plates. Thedislocated laths are composed of elongated bundles ofslightly misoriented subgrains, within which the d i s l o -cation density may vary with composition and coolingrate (specimen thickness) and is associated with the( l l l } r habit.2 An example is shown in Fig. 2. Thelaths can also be twin related. In low carbon s t e e l sadjacent laths appear to adopt different variants of theK - S orientation relationship.61 The dislocation s t r u c -ture is very complicated and s i m i l a r t o that in ~ - i r o nand metals of high stacking fault energy 24 a f t e r highspeed deformation. P l a t e martensites contain finetransformation twins. Those within which the twindensity va r i e s and in which the twins do not extendcompletely a c r o s s the plates are associated with{ 225} ~, { 259} y habits, w h e r e a s fully twinned platescorrespond t o the {3, 10, 15} ~ habit. A s u m m a r y in-cluding recent, as yet unpublished data, is given inTable I, and examples of t h e s e twinned structures areshown in F i g s . 1, 3 to 6, 9, and 11. In all ca se s thesubstructure within a g iven martensite g ra in is com-plex and in many fully martensitic steels the m i c r o -structures are often complex mixtures of dislocatedand twinned martensites, Fig. 6, so that detailed mi-crostructural analyses a r e very difficult t o c a r r y out.

The martensitic substructures may a r i s e from twofactors: 1) the lattice invariant s h e a r of the transfor-mation s t r a i n itself and this may be slip, twinning ora combination of both, and 2) plastic deformation ofthe martensite a f t e r i ts formation, due to accommo-dations resulting f r o m the constraints of the s u r r o u n d -ings. It follows that surface martensite or martensitef o r m e d in the thin foils should be less d e f o r m e d thanbulk martensite b e c a u s e of the a b s e n c e of bulk con-straints at the free surface. U n l e s s the d e v e l o p m e n tof the substructure can be o b s e r v e d directly duringthe transformation, see Figs. 7, 8, it is not possibleto distinguish b e t w e e n transformation and deformationsubstructures. This distinction is important in r e g a r dto relating o b s e r v e d structures to the s h e a r s p r o p o s e din the crystallographic theories.

At the present time t h e r e is a g r e e m e n t b e t w e e n the

METALLURGICAL TRANSACTIONS VOLUME 2, SEPTEMBER 1971-2375

Fig. 2--Lath martensite in Fe-II.3Ni-7.5Co-0.24C-0.44Cr-0.3Mn (Ms~ 307°C). Laths occur as bundles of complex dis-located regions slightly misoriented from region to region.The diffraction pattern in (c) shows superposed [100]ce and[111]ce patterns which are a common feature. The regionscorresponding to their orientations are revealed in the darkfield images of spots A, B (Ref. 31).

phenomenological theory and substructure only for thecase o~ the {3, 10, 15}~ habtt~ a n d { U 2 } < t l ~ ) ~ tw~n-ningo~ The {225}r partially twinned martensites are

very complex as will be shown in more detail in Sec-tion 2.2. Physically, of course, s i n c e the twinning dis-location must r e m a i n in the habit plane, and if the twindislocation is of the a//6 <111> type, then for c u b i c al-loys the h a b i t plane must be actually (5, 10, 15}. In{ 259 } martensites (bc tetragonal) t h e r e is one reportof {t10} as well as 4112} twinned martensite plates 4' 49( F e - l . 8 pct C steel).

It is likely that martensites in specimens preparedfrom bulk transformed alloys have undergone bothtransformation and accommodation deformation, sothat it is not c l ea r in case s w h e r e one finds both dis-located and twinned substructures, e . g . , Fig. 5, whicheffect has dominated. Twins are found in lath marten-

2376-VOLUME 2, SEPTEMBER 1971 METALLURGICAL TRANSACTIONS

Table I. Structure of BodyCentered Ferrous Martensites Formed on Cooling Through Ms-M~

Alloy System Ms Substruct ure/Composition'~ Reference

Fe-C >350°C C < 0.3 mainly dislocated laths 1, 2 , 28

<250°C C > 0.6 mainly twinned plates 40

>-50°C Ni < 25 dislocated laths 1,2, 39, 41,42 ,43

44 ,45, 48 ,53

Fe-Ni

- 30° to - 150°C andbelow

Fe-Ni-C as above

Fe-25Ni-O.3C-4.5Mo <-77°C+4.7 Cr - 16°C+1.85V -3°C

Fe-Ni-Co-C (~Mo-½Cr) - 260°C

150° to 260°C

Fe-Ni-Co as above

Fe-Cr-C 350° to 400°C

Fe-5Cr

Fe-SCr-IC -30°C

Fe-5Ni-0.25C 315°C

Fe-3Mn-0.25C 315°C

Fe-5Ni-7Mn-0.25C 65°C

Fe-7Mn-0.25C 190°C

Fe-Ni-Ti

Fe-Cr-Ni

Fe-Mn-Cr-Ni

m e t a s t a b l e a u s t e n i t e s

Nii> 30 midrib twinned only

Extent of twinning increases with pct Ni, and lower Ms

I n c r e a s e in carbon for same nickel content andvice versaenhances twinning

A l l mainly twinned. However ausforming causes precipitationof carbides, so that the resulting martensite is less twinned.

C < 0.3 dislocated laths

C > 0.4 twinned cobalt does not decrease twinning

Partially twinned

Cobalt raised Ms-does not decrease percent twinning

Dislocated l a t h s

Twin density decreases with increasing plate size

Mainly dislocated

Mainly dislocated

Mainly twinned

Mainly twinned

For a given carbon level, manganese promotes twinning

Dislocations and twins depends on Ni-Ti andheat treatment

Low stacking fault energy dislocated ~ from e

1,2,40,43,28,48,39,61

29,30

31

69

35

43

20,25

36

36,47

36

36,47

33

55,56,57,21

*All data here refer to nominal cooling rates. Thems temperature and structure can be variedby varying quench rate (see text).1"Composition refers to wt pct.

s i t e s in s e v e r a l a l l o y s26 ,54( F e - N i - C , F e - N i - C o - F ,F e - N i - M n - C , F e - N i - C r - C ) a n d d e b a t e i s o c c u r r i n g a st o w h e t h e r o r n o t t h e s e t w i n s a r e due t o p l a s t i c d e f o r -m a t i o n o f t h e t r a n s f o r m e d m a r t e n s i t e , ze

2.2) D o u b l y T w i n n e d M a r t e n s i t e s

T h e t w i n n e d s t r u c t u r e s o f m a r t e n s i t e s e x h i b i t i n g{ 225 } y a n d { 259 } y h a b i t s a r e p a r t i c u l a r l y p u z z l i n gt o t h e c r y s t a U o g r a p h e r s , a n d t h e r e i s n o s a t i s f a c t o r ye x p l a n a t i o n f o r t h e s e h a b i t s in t e r m s o f t h e o b s e r v e ds t r u c t u r e s a n d l a t t i c e i n v a r i a n t s h e a r s . H o w e v e r , u n -t i l n o w i t h a s not b e e n s h o w n t h a t t h e t w i n s t r u c t u r e i si t s e l f c o m p l e x . T h e f o l l o w i n g r e s u l t s s h o w t h a t thet w i n s t h e m s e l v e s m a y be t w i n n e d .

T h e p r o o f o f t h i s p h e n o m e n o n r e q u i r e s p a i n s t a k i n ge x p e r i m e n t s by e l e c t r o n m i c r o s c o p y a n d i n v o l v e s d e -t a i l e d d a r k f i e l d a n a l y s e s o f a l l t h e s p o t s in t h e p a t -t e r n . I t i s a l so h e l p f u l in t h e a n a l y s i s t o s t u d y o r i e n -t a t i o n s w h e r e m a r t e n s i t e t w i n s a r e p a r a l l e l to t h ee l e c t r o n b e a m , e . g . , ( 1 1 3 > a f o i l s , a n d p l a t e s in w h i c hthe t w i n s a r e i n c l i n e d , e . g . , ( 1 2 2 > a . T h e s e h a v e b e e ns t u d i e d in F e - 3 3 N i a n d F e - 2 5 N i - 0 . 3 V - 0 . 3 C a l l o y s . T h er e s u l t s a r e s h o w n in F i g s . 8 to 11 .

F i g . 9 (c) t h r o u g h ( f ) s h o w s a s e t o f d a r k f i e l d i m -a g e s c o r r e s p o n d i n g t o t h e s p o t s in F i g . 9 ( b ) i n d e x e d inF i g . 1 0 ( c ) . I f t h e [ 2 2 1 ] m a r t e n s i t e w e r e s i n g l y t w i n -

METALLURGICAL TRANSACTIONS

n e d o n t h e ( i 1 2 ) ~ t h e [ 0 0 1 ] ~ t w i n o r i e n t a t i o n w o u l d bep r o d u c e d a n d t h e p a t t e r n w o u l d be s u p e r p o s e d [ 0 0 1 ] ~a n d [ 2 2 1 ] ~ . I f t h i s w e r e t r u e , t h e n d a r k f i e l d i m a g e so f both 1 1 0 a n d 110 d i f f r a c t i o n s p o t s w o u l d r e v e a l t w i nc o n t r a s t . H o w e v e r , a s t h e d a r k f i e l d i m a g e s o f F i g s .9 ( c ) a n d (d) s h o w , t h e t w i n r e v e r s a l only o c c u r s f o rone a n d not both o f t h e s e r e f l e c t i o n s . T h e e x p l a n a t i o ni s that t h e t w i n s o n ( 1 1 2 ) ~ a r e t h e m s e l v e s t w i n n e d o n(112)c ~ so t h a t t h e [001Ice_twin of [ 2 2 1 ] ~ n o w t r a n s -f o r m s l o c a l l y i n t o t h e [ 2 2 t ] ~ o r i e n t a t i o n . H o w e v e r ,the e x t e n t o f t h e d o u b l e t w i n n i n g i s l i m i t e d . T h e c o n -t r a s t a t S in F i g . 8 i s c o n s i s t e n t w i t h s e c o n d a r y t w i n son (119.)(~ in t h e p r i m a r y t w i n s . T h e d a r k f i e l d i m a g eo f t h e s p o t m a r k e d " T " , F i g . 8 ( c ) , r e v e r s e s t h e c o n -t r a s t o f b o t h p r i m a r y a n d s e c o n d a r y t w i n s . T h i s s p o ti s a c t u a l l y t h e s u p e r p o s i t i o n o f b o t h ( 1 1 0 ) ~ p r i m a r yt w i n a n d (110)~ s e c o n d a r y t w i n r e f l e c t i o n s .

T h e s e c o n c l u s i o n s r e g a r d i n g t w i n n e d t w i n s a r e f u r -t h e r c o n f i r m e d by a n a l y s i s o f t h e e d g e - o n t w i n s in t h efu l l y t w i n n e d [ 1 1 3 ] q m a r t e n s i t e p l a t e o f F i g . 11 . H e r et h e p r i m a r y t w i n p l a n e i s ( 2 1 1 ) ~ w h i c h p r o d u c e s t h e[ 1 1 3 ] ~ --- [ f f 3 ] ~ t w i n n i n g t r a n s f o r m a t i o n a s v e r i f i e dby t h e d i f f r a c t i o n p a t t e r n a n d d a r k f i e l d a n a l y s i s . H o w -e v e r , t h i s e x a m p l e i s e v e n m o r e c o m p l e x . F a i n ts t r e a k s a r e r e s o l v a b l e o n t h e n e g a t i v e o f F i g . 11 (b)p a r a l l e l t o t h e <110 > d i r e c t i o n s i n d i c a t e d o n t h i s p a l -t e r n . T h e s e s t r e a k s c a n n o t be e x p l a i n e d by d o u b l e

VOLUME 2 , SEPTEMBER 1971-2377

{112} twinning as in Figs. 8 and 9, but they must becaused by planar defects on {110}, e . g . , twins al-though this is not expected in cubic martensite. How-ever, since the diffraction streaks are so weak, it is

Fig. 3-- Fe-25Ni-0.3V-0.3C bulk-transformed in liquid nitro-gen. The lens-shaped plate with (112)ce twinned midrib coex-ists with dislocated 3' and untransformed 7• Notice the densedislocation tangles, i.e., work-hardened appearance in theaustenite.

not possible to resolve their morphological origin inthe image. It is important to emphasize here that al-though Fe-Ni martensites are well characterizedrnacroscopically, it is clear that there are complica-tions when each individual martensite crystal is ana-lyzed. Differences can exist locally in compositionand Ms temperature, and this could result in marten-site plates of different habits and transformation sub-structures. The complexity of the situation is empha-sized by the fact that Figs. 9 and 11 were obtained onthe same specimen and within I0 p of each other.

It is interesting to note that the interface in Fig. 11is parallel to (739)a. Assuming a Bain correspond-ence this plane wouldbe parallel to the (529) of aus-tenite. Although this is a speculative correlation atthis stage, it is possible that {II0}~ secondary twin-ning may be a characteristic of low temperature (259)7habits. As more detailed analyses of martensite sub-structure are obtained, it is hoped that further clari-fication of the habits in terms of transformation shearswill be possible.

It is noted that the twins are often displaced or trun-cated along <III >, i . e . , parallel to {011}~ traces con-sistent with slip (or possibly twin) deformation. Dislo-cations can be resolved between the twins in Fig. 5, butin general the contrast requirements are different forrevealing twins and dislocations separately.

There is surface structure resolvable on the broadfaces of the twins as canbe seen in Fig. 1(d), andFigs. 7 and 9. Some of the complex fringe patterns ob-served in the dark field images Figs. 9 (c) to (f) maybe related to the secondary twinning or to variations intwin thickness, e . g . , at ledges. These fringes mayalso be associated with regular arrays of twin/twin ortwin/matrix interface dislocations, but it has been dif-ficult to obtain definitive images even after using sev-eral different reflections, Figs. 9 (c) to (f). The factthat these fringes do not change orientation (direction)

Fig. 4--Fe-33Ni-0.01C. Bulk t r a n s -f o r m e d by q u e n c h i n g r a p i d l y to l iqu id h e -l ium v i a i c e d b r i n e a n d l iqu id nitrogen.Mostly t w i n n e d p l a t e s and r e t a i n e d aus-tenite. The la rge p l a t e ( top left) i s a l soh e a v i l y d i s l o c a t e d . 650 kv .

2378-VOLUME 2 , SEPTEMBER 1971 METALLURGICAL TRANSACTIONS

when different reflections a r e used, Fig. 9, ru l e s outthe possibility of Moirg contrast. Other separate ex-periments show that complex dislocation arrangementsa r e present in addition to twinning. Fig. 5 is an exam-ple. The dislocations shown here cannot be a / 6 (111}twinning dislocations s i n c e all are invisible in the

1g = [110] reflection (g" b = T or 0).It may be a rgued that the double twinning may r e s u l t

from plastic deformation a f t e r transformation. How-e v e r , the fact that double twins are observed in m a r -

Fig. 5--Fe-33Ni-0.01C. As Fig. 4, same specimen, dark fieldimage of matrix 110 spot showing twin fringe contrast as wellas dislocations in two adjacent plates. 650 kv.

tensites f o r m e d directly in thin foils, F i g s . 7 and 8,lends good support for the belief that double twinningis characteristic of the transformation itself. Doubletwinning has also been detected in lath martensite,see Ref. 26.

The a b s e n c e of reports of double twinning in e a r l i e relectron microscopic studies of martensites is s o m e -what surprising, but s e r v e s t o emphasize the necessityof doing complete dark field and diffraction analysest o characterize microstructure.

2.3) Factors Affecting Substructure

Recently Nutting~ reviewed some of the f a c t o r sthought t o be important in controlling the martensitesubstructure, and in view of more r e c e n t data, TableI,it is useful to s u m m a r i z e t h e s e factors so as to at-tempt to gain a comprehensive understanding. T h e s efactors are:

1) Composition: Increasing the total solute contenttends t o change the morphology from dislocated lathst o twinned plate martensite, T a b l e I. C a r b o n appearst o have the strongest effect in promoting twinning.

2) M s temperature: Concurrently with compositiont h e r e is usually a d e c r e a s e in M s temperature withincreasing solute content: the exception is cobalt--however, cobalt does not d e c r e a s e twinning, sl, 3~ In agiven alloy s e r i e s the M s can be kept constant andthen the composition determines whether laths orplates form21

3) The strength of the martensite: S ince the latticeinvariant s h e a r involves plastic deformation, the sub-structure must depend on the yield strength of m a r -tensite. This property also depends on compositionand temperature. The strength of austenite and d e f o r -mation substructure in austenite ( e . g . , Fig. 3) mayalso influence the morphology, but this possibility hasnot yet been studied in any detail.

4) Austenite stacking fault energy: T h e r e is no sim-ple relationship. High stacking fault energy (SFE) is

Fig. 6--Fe-28Ni-0.3C ( M s ~- 8°C) Bulktransformed in liquid nitrogen showingmixed complex structure of twinned anddislocated martensite. (Courtesy O. Jo-hari.)

METALLURGICAL TRANSACTIONS VOLUME 2, SEPTEMBER 1971-2379

s u p p o s e d t o p r o m o t e l a t h s i f M s i s r e l a t i v e l y h i g h , buti f Ms i s low, t w i n n e d p l a t e s a r e f o r m e d .6 H o w e v e r ,i n c r e a s i n g t h e concentration o f a l l s o l u t e s s e e m s t op r o m o t e t w i n n i n g , e . g . , m a n g a n e s e a n d n i c k e l , y e tm a n g a n e s e l o w e r s S F E a n d n i c k e l r a i s e s S F E o f a u s -t e n i t e . I t i s d i f f i c u l t to s e e w h a t r e l a t i o n s h i p e x i s t sb e t w e e n s t a c k i n g f a u l t e n e r g y o f a u s t e n i t e a n d t r a n s -f o r m a t i o n s u b s t r u c t u r e o f m a r t e n s i t e u n l e s s t h e l a t t e ri s i n f l u e n c e d by t h e m o d e o f p l a s t i c d e f o r m a t i o n o fa u s t e n i t e d u r i n g t r a n s f o r m a t i o n . S F E o f a u s t e n i t e i si m p o r t a n t , h o w e v e r , in d e t e r m i n i n g w h e t h e r ~ m a r -t e n s i t e i s p r e f e r r e d to c~ 16 , 17, 56, 57

5) C o o l i n g R a t e : C o o l i n g r a t e a f f e c t s Ms, 2s,54 a n din m a r t e n s i t e s t h e d i s l o c a t i o n d e n s i t y v a r i e s ; ~ in a l -l o y s t e e l s p r e c i p i t a t i o n o n s l o w c o o l i n g r a i s e s Ms sot h a t t h e m a r t e n s i t e i s d i l u t e d a n d h e n c e i s l a t h t y p e .F a s t c o o l i n g r e t a i n s s o l u t e s in s o l u t i o n a n d t w i n n e dp l a t e s r e s u l t .

6) T h e r m a l - m e c h a n i c a l : P l a s t i c d e f o r m a t i o n c a nc a u s e p r e c i p i t a t i o n in a u s t e n i t e ( a u s f o r m i n g ) , ~ , so so

Fig. 7--Fe-25Ni-0.3V-0.3C. Thin foil t r a n s f o r m e d in l iqu idnitrogen s h o w i n g p a r t i a l twinning (compare to the s a m e alloy,Fig. 3). The p r i m a r y t w i n s and f r i n g e c o n t r a s t a r e r e v e r s e din this dark f i e l d i m a g e of the ll0c~ twin spot . The t w i n s a r etruncated p a r a l l e l to (110)~ matrix. The b r o a d f a c e s of thet w i n s contain s t r u c t u r a l d e t a i l s , cf. Fig. l ( d ) , w h i c h may bel e d g e s , s e c o n d a r y t w i n s , or interface dislocations, s e e a l soF i g . 9 . Orientation n e a r [3 , 1 , 11]c~. 650 kv .

Fig. 8--Fe-25Ni-0.3V-0.3C. Thin foil t r a n s f o r m e d in the colds t a g e at - 80~C showing the development of the doub ly twinneds u b s t r u c t u r e . The twin f r i n g e c o n t r a s t i s s t r o n g n e a r the ex-tinction contour. The d a r k f i e l d i m a g e in (b) c o r r e s p o n d s toposition C in Fig. 10(c). Notice fine twin f r a g m e n t s , a r r o w e d ,and s e c o n d a r y contrast at S (consistent with s e c o n d a r y twint r a c e ) . This r e s u l t p r o v e s that the t w i n s a r e twinned.

2380-VOLUME 2 , SEPTEMBER 1971 METALLURGICAL TRANSACTIONS

F i g . 9--Fe-33Ni-0.01C As Fig. 4, showing dark f i e l d experiments to r e v e a l twinned {112} twins . (a)_Bright f i e l d image.(b) Selected a r e a diffraction of c e n t r a l part of A, see Fig. 10. (c) Dark f i e l d i m a g e of spot C (l l0)T 1 + (l l0)T 2. The twins showwedge f r i n g e contrast at W. (d) Dark f i e l d i m a g e of spot D ( l l 0 ) m a t r i x + (l l0)T1. (e) Dark f i e l d i m a g e of s p o t E w h i c h i s 200T 1, ordoubly d i f f r a c t e d m a t r i x and T 2 spo t s . (f) Dark f i e l d i m a g e of (240)T1 + ~024) matr ix . In C the s t r o n g f r i n g e c o n t r a s t i s due tow e d g e thickness c h a n g e s . T h e o t h e r fringes are parallel to the t r a c e of (112) in the (001) Ti, a n d c o u l d r e p r e s e n t d o u b l e twin in-terface. 650 kv.

METALLURGICAL TRANSACTIONS VOLUME 2, SEPTEMBER 1971 2381

A A A a A A h

A h A a b h A

l iO h

A A A

A 5 A a A A A

A A ~ h A A h(a)

dl 0 0 i 0 a 8

A A A A A A A

A a Ea a a s'

am 0 co • 0 I~ 8

& & a he a a A

h A & h A & A

ii 0 ~ i a D g( c )

D I~ D a O O w

A A A A A A A

A A A _,~o,,h A A

a O co.o,' ' , O O B~iOr2 IIOT~

A A A . A A A A

5 5 h A h A A

D n O O • Ci n(b) lI4Tz

• Matrix Reflection M

A 1st Twin Reflections

El 2nd Twin Reflections

X In absence of 1"1, these

are double diffraction

spots from Matrix

and T2

Fig. 10--Explanation of Fig. 9. (a) The pat-tern if the 211~ matrix is singly twinned on(112); superposed [221}a matrix + (001)atwin. (b) The pattern if the matrix is com-pletel_.y converted to the double twin [001]T1+ [22117-. (c) The pattern if the [221]~ ma-trix twins to [0Ol] a and the [001]a partiallytwins to [221]a. This pattern explains Fig.9(b) upon analysis of the dark field imagesand comparison to the possibilities shownhere, proving that the structure of Fig. 9cannot be a single twin.

that on transformation lath martensite is formed.Likewise p r i o r ausaging can affect the subsequentmartensitic structure. 3a

7) Transformation morphology is determined bypressure. In Fe-C alloys twinned plates can be pro-duced at 40 kh whereas at atmospheric pressure, themartensite is lath type.~9, 6o

All of these results lead t o one consistent pattern,viz. the h ighe r the solute content the g r e a t e r is theprobability of f o r m i n gtwinned plates. Conversely, ifthe austenite is diluted as a r e s u l t of precipitation,e . g . , by slow cooling, or by plastic deformation ( a u s -forming) a normally twinned martensite can be changedt o dislocated laths. ~, 3o Results obtained recently onalloys with s i m i l a r M s temperatures and on the ef-fects of cobalt clearly show that M s temperature a l o n eis not the controlling parameter determining the mar-tensitic substructure, s~,3~ T h e s e results bare led t othe suggestion32 that the strength of martensite (and ofthe austenite from which it forms) is probably the sin-gle most important factor which determines whethermartensite is dislocated or twinned or both.

This idea can be developed qualitatively by consid-e r i n g Fig. 12 which suggests schematically a variationin critical resolved s h e a r s t r e s s with temper a t u r e forslip and twinning. This variation will depend on com-position, and detailed knowledge of t h e s e relationshipsis needed. Increasing the s o l i d solution compositionand strength may r a i s e the CRSS for slip above thatfor twinning so that twinning is the p r e f e r r e d s h e a rmode. S imi la r relationships might be expected fors t r a i n rate dependence. At low temperatures or when

2382 VOLUME2, SEPTEMBER 1971

solutes l o w e r CRSS for twinning, twinning is expectedto be the p r e f e r r e d mode of plastic deformation. Also,s i n c e martensite f o r m s over a range of temperatureM s t o M f [ ~100 ° t o 200°C] and if this range includesthe slip ~ twin transformation, then mixed substruc-t u r e s a r e expected. Likewise the loca l temperaturerise at M s and M b could allow slip t o o c c u r whenotherwise twinning dominates. It is also possible thatdifferent twinning or slip modes may operate at verylow temperatures.

Carbon (and nitrogen) seems t o be the most potentof the alloying elements in promoting twinning. Car-bon is also the most potent strengthener of s t e e l .Thus, the s t r o n g e r the s t e e l the more difficult it is forslip t o o c c u r and then twinning is preferred. This isalso confirmed by the experiments of R i c h m a n66 andBevis e t a l . 67 who observed mechanical twins in de-formed c a r b o n martensites.

It may also be mentioned that the martensite sub-structure closely r e s e m b l e s that in shock-loaded ma-t e r i a l s .~ In fcc m e t a l s as the shock p r e s s u r e is in-c r e a s e d twins are f o r m e d when the dislocation densityhas attained some critical value, sT,ss By analogy it ispossible that the twins observed in martensite l a t h s31may have been induced because the limiting d i s l o c a -tion density had a l r e a d y been attained. This wouldmean slip s h e a r s always p r e c e d e twin s h e a r s .

Another factor which may be important is the d r i v -ing f o r c e associated with the transformation. Bell andOwen 34 suggested that when a critical d r i v i n gf o r c e isexceeded ( ~ 315 col per m o l e ) the change from dislo-cated to twinned martensite is expected. P a s c o v e r and

METALLURGICAL TRANSACTIONS

Fig. 11--Fe-33Ni-0.01C. Bulk trans-formed [ 113]a orientation. (a) Brightfield. (b) Selected area diffraction patternshowing strong streaks due to (121)a pri-mary twinning; double diffraction spots~white arrows) and very faint streaks along110 matrix and 110 twin. [113]c~ primarytwin spot pattern is superposed on [113]~matrix. (c) Dark field image of 110 T1spot (which also includes the faint 110streaks). (d) Dark field image of 110 ma-trix. 650 kv.

Radcliffe 43 extended this concept to F e - N i and F e - C rw h e r e the driving f o r c e was in the range 300 to 370 calper mole for the Fe-29Ni alloy which is twinned. How-e v e r , twins were not observed in t h e i r F e - 5 C r alloyalthough the driving f o r c e was 300 to 350 cal per m o l e .This can be explained on the b a s i s that the CRSS forslip vs. that for twinning are the important parameters.

3) DIFFRACTION CONTRAST

Whilst t h e r e are no difficulties associated with dis-tinguishing between the slip structure and twin s t r u c -ture within martensite, so far t h e r e have been veryfew quantitative analyses (see for example Pattersonand Wayman41 and Warlimont 44). It is now c lea rthat the twin structure itself is not simple. Althoughthe twin planes can be characterized without too manydifficulties, the twinning s h e a r must be found by con-t r a s t analysis of the twinning dislocations. If the twinsfully t r a v e r s e the plate the dislocations terminate atthe interface ( y - M s or M s - M s ) and the contrastanalysis is then quite difficult. For {112} twins theB u r g e r s vectors of the twinning dislocations are p r o b -ably of the a/6 (111> type. For these t o be observed,in general7 g .b > ~-, w h e r e g is the operating r e f l e c -tion. Thus a/6 <111> dislocations will not be visiblewith g = <IL0) nor <200>. Fr inge contrast is observedonly when g .R is non-integral (R is the twin shear),

METALLURGICAL TRANSACTIONS

ioe., when the phase shift c~ = 2~gR is nonzero (orn 27r). For {112}< 111> twinning in bcc metals, if R= ~<111> and s i n c e h + k + l is always even, then

is =~2 ~/3 or zero and the contrast is identical t othat for {111} <112> stacking faults in fcc crystals,as can be seen in F i g s . 5 and 8 (see also Ref. 41).Hence the fringes are symmetrical in intensity inb r igh t field and asymmetrical in dark field. Thus, thesuggestions of Krauss and Pitsch5s r e g a r d i n g stackingfaults in F e - 3 3 pct Ni can be explained in t e r m s ofthin twins although the a rgument could be more con-vincingly settled with respect t o the evidence from thediffraction patterns (not available). Contrast analysisof C l i 0 } twinning has not yet been published. How-e v e r , if the twinning displacement is also a/6 <111>the f r inge contrast from {112}~ and {110}~ m a r t e n -sitic twins would be identical s i n c e the f r inge c h a r a c -teristics depend upon ~ = 2 $ g .R . Thus, more de-tailed studies of the f r inge character may be useful.Distinction between these different twin planes can inprinciple be obtained by t r a c e analysis and fromchecks of the twinning reflections with the appropriatetwinning ru l e s ( i . e . , the twin reflection can be pre-dicted from twinning analysis l°, 51, 52). The distinctionis important with rega rd t o the phenomenological the-ory of the transformation.

The above r e m a r k s on contrast apply, of course, totwo-beam diffracting conditions and only for orienta-

VOLUME 2, SEPTEMBER 1971-2383

fffl

I,LI

I--

~ ~ MARTENSITESUB STRUCTUREDISLOCATED

WINNEDI

I ~ " \ ~ I - ----._ ¢

I = ~NL

IM,SLIP

IMs

TEMPERATURE ---,>Fig. 12--Schematic suggested variation of the stress for slipand twinning as a function of temperature. Composition vari-ations are expected to affect the twin-slip cross-over andpossible changes are indicated by arrows. The slip or twinmodes may also vary with temperature and composition.

tions when the m a t r i x and twin diffraction vectors co-incide. If the l a t t e r condition is not met, the fringecontrast will be modified due to the orientation differ-ence a c r o s s the matrix-twin interface (so-called exc i -tation e r r o r , or AS fringes, e . g . , Ref. 68). From apractical point of view and when examining c o m m e r -cial steels, it is virtually impossible t o obtain two-beam orientations because of the s m a l l crystal size.This difficulty increases with increasing voltage, be-c a u s e of many beam excitations.

Except for Fe-Ni alloys the dislocated martensitesare very complex and have not been analyzed, i . e . , nodetermination of B u r g e r s vectors have been made. Ino r d e r for this t o be done effectively, experimentsshould involve preoriented single crystals and w h e r etransformation is incomplete so that constraints duet o impinging martensite crystals can be avoided. InF e - N i alloys the dislocation structure 4b 44 is oftensimple with long a r r a y s of s c r e w s lying in <111>.However, in fully transformed steels the structure isfar more complex and difficult to analyze, e . g . ,Fig. 6.

Transformation dislocations must r e m a i n with theinterface and will not necessarily have simple Bur-gers vectors, i . e . , a dislocation in austenite cannot ofitself produce a martensitic structure unless it decom-poses into partials which can accomplish the transfor-mation. The simplest example of this is in fcc ~-- hcptransformations which can be accomplished by spl i t -ting of the ½<ll0>fec dislocations into 61-- <112> or1-< 111> partials, thus producing stacking faults and3the hcp structure. Furthermore, if the habit plane is(111), successive s h e a r s by partials ~ , ,~[112] produce net zero shape change. Although thequestion of s i m i l a r dissociations have been consideredwith respect t o bcc ~ { f c c , hcp} transformations, ~°specific attempts t o identify the transformation dis-locations, including a detailed study of the N/Ms in-

2 3 8 4 - V O L U M E 2 , S E P T E M B E R 1 9 7 1

t e r r a c e contrast, do not appear t o have been publishedso far.

In studying the substructure of partially transformedalloys the substructures and interfaces of both phaseswill not generally be in contrast simultaneously. Ino r d e r t o r e s o l v e the relationship between the i n t e r -face structure and the martensite substructure, it willbe necessary to orient the foil for a s t r o n g martensiticreflection. Examples of the two contrast conditionsa r e shown in Fig. 1 (Ref. 20). F i g u r e 1 (a) is orientedfor martensite contrast and it is c l ea r the martensiteis made up of closely parallel bcc (twinned) blocks. In(b) which is oriented for austenite contrast the d i s l o -cation and faulted austenitic structure can be s e e n .Comparison of (a) and (b) shows that the d i s l o c a -tions a r e associated with the austenite phase. In (c)both are in contrast simultaneously.

It would be advantageous in such circumstances( i .e . , overlapping phases) t o do stereomicroscopy. Infact, g r e a t e r use of this technique would be valuable.

4) SUMMARY AND CONCLUSIONS

It is c l ea r that of the various physical features ofmartensitic transformation, two main a reas a r e inneed of elucidation, viz. nucleation, and detailed c h a r -acterization of the transformation substructure and itsrelation t o the transformation mechanism and crystal-lographic theory. Both of these problems a r e poten-tially solvable by d i r e c t studies of the transformationin the electron microscope. However, the fact r e m a i n sthat t h e r e are no conclusive r e s u l t s yet showing thenucleation site itself. Also, s i n c e the nucleation theoryis unsatisfactory in its present fo rm, the design ofcritical experiments t o focus on these problems hasbeen limited.

It is usually a s s u m e d that each martensite platecomes from one embryo. However, it is possible t ohave simultaneous nucleation of many embryos, whichgrow together t o form plates. In other words t h e r emay not be a s i m p l e relationship between the numberof nuclei and the number of martensite crystals. Di-rect observations a r e obviously necessary, and it willbe useful t o devote more time t o experiments u s i n ghigh voltage electron microscopy in o r d e r to take ad-vantage of the increased probability of finding embryos,if they e x i s t . Due t o the i n c r e a s e in penetration theprobability is an o r d e r of magnitude h ighe r at 650 kvthan at 100 kv. So far no attempt has been made inthis direction except for copper alloys, 11 and t h e s eexperiments were not r e l a t e d t o the nucleation p r o b -lem.

Questions can be r a i s e d r e g a r d i n g the v a l u e of thinfoil observations for the interpretation of bulk t r a n s -formations, and it is essential t o compare foils withbulk materials. Nevertheless, useful information canbe obtained and we have encouraging results that showat l e a s t substructural similarities, viz . , double twin-ing in martensites f o r m e d in f o i l s and in bulk s p e c i -mens of the same alloy. It should be noted that whilsthigh voltage microscopy p e r m i t s examination of thick-er specimens, the upper useful l imi t of thickness forsteels is still only of o r d e r 2 # at 500 t o 1000 kv, andthis thickness is hardly representative of the bulk.

High voltage microscopy offers other advantages be-

M E T A L L U R G I C A L T R A N S A C T I O N S

s ides that of greater penetration. T h e s e advantagesinclude the improvements in resolution due to de -creased chromatic aberration, the great reduction inthe minimum s e l e c t e d a r e a of diffraction, the abi l i tyto examine the same a r e a for extensive t imes withoutcontamination being a limitation (the contaminationrate i s lowered by increas ing e lectron energy , and thevacuum in high voltage microscopes i s ~ 100 t imesbet ter than at 100 kv), and the advantages of workingwith fo i l s having large transparent areas availablefor analys i s , e . g . , Figs . 7, 9, and 11. It shou ld benoted, however, that radiation damage i s a compl i -cat ing factor at high voltages . Iron will su f f er primaryknock-on damage in e lec tron microscopes operatedabove 500 kv.

It has been sugges ted e a r l i e r and substantiated herethat the strength and deformation characterist ics ofmartensite (and poss ib ly austenite) control the marten-site substructure . In all cases the substructure i scomplex and probably involves several shear modes.There i s a need for detailed data on the crit ical r e -so lved shear s tress dependence for s l i p and twinningas a function of composition, s tra in rate and tempera-ture , before t h e s e suggestions can be properly eva l -uated.

Finally i t can be sa id that apart from the consider-able attention which has been paid to crystal lographicrelationships and the structure-properties approach,electron microscopy studies of martens i te have , ex-cept in rare cases , neglected detailed contrast anddiffraction analyses . Admittedly many martens i tes ,particularly those in commercial s t e e l s , are far toocomplex for this to be facilitated. However, exper i -mental a l loys can be made so that part ia l ly trans-formed s t e e l s are available whereby crystal lography,morphology and substructure can all be analyzed rig-orously . If the same attention to detail, as far aselectron metallography i s concerned, i s paid to fer-rous martens i tes as has been paid to non=ferrousphase transformations we can look forward to a morecomplete understanding of this fascinating and impor-tant subject .

ACKNOWLEDGMENTS

It i s with great pleasure that I acknowledge thestimuli and help provided by past and present graduatestudents, part icularly I. Cheng, S . Das, P. R. Okamoto,G. Stone, and M. Raghavan. I wish to thank theU.S.A.E.C. for continued financial support and Mr. D.Jurica for continued attention to the high performanceof our 650 kv e lectron microscope.

R E F E R E N C E S1. P. M. Kelly and J. Nutting: Proc. Roy. Soc., 1960, vol.259A, p. 45.2. P. M. Kelly and J. Nutting: J. Iron Steellnst., 1959,vol. 192,p. 246;

Z 1ton Steellnst., 1961,vol. 197,p. 199.3. W.Pitsch: Z lnst. Metals, 1958,vol. 87,p. 444;Phil Mag., 1959, vol. 4,

p. 577.4. C. M. Wayman: Recent Developments in Diffraction Methods inMater. Sci.

Proc. NATO InL Summer School, Antwerp, Belgium, p. 187, North-HollandPubl., The Netherlands, 1970.

5. C. L. Magee: Phase Transformations, p. 115, ASM, Metals Park, Ohio, 1970.6. J. Nutting: Z Iron Steellnst., 1969,vol. 6, p. 872.7. G. Thomas: Transmission Electron Microscopy ofMetals, Wiley, New York,

1962;also ThinFilms, ASM, 1964, p. 227.8. G. Thomas: Trans. TMS-AIME, 1965, vol.233, p. 1608.

METALLURGICAL TRANSACTIONS

9. G. Thomas:Recent Developments in Diffraction Methods in Mater.Sci. Proc.NA TO Int. Summer School, Antwerp, Belgium, p. 131, North-Holland Publ.,The Netherlands, 1970.

10.P. B. Hirsch, A. Howie, R. B. Nicholson,D. W. Pashley, and M. J. Whelan:Electron Microscopy ofThin Crystals, Butterworths, London, 1965.

11.H. Fujita, Y. Kawasaki,E. Furubayashi, S. Kajiwara, and T. Taska: Japan J.AppL Phys., 1967, vol. 6, p. 214.

12.G. Thomas:Phil. Mag., 1968,vol. 17, p. 1097.13. E. Votava: ActaMet., 1960, vol.8, p. 901 ;J. Inst. Met., 1961, vol. 90, p. 129.14. J. A. Hren and G. Thomas: Trans. TMS-AIME, 1963, vol. 227, p. 308.15.G. Thomas and M. J. Whelan: Phil. Mag., 1961, vol. 6, p. 1103.16.P. L. Mangonon,Jr. and G. Thomas:Met. Trans., 1970,vol. 1, p. 1577.17. J. A. Venables: Phil. Mag., 1962, vol. 7, p. 35.18. S. R. Pati and M. Cohen: Acta Met., 1969, vol. 17, p. 189.19. H. Warlimont:Met. Trans., 1971, vol. 2, pp. 1847-50.20. K. Shimizu, M. Oka and C. M. Wayman: ActaMet., 1970,vol. 18, p. 1025.21. G. Thomas:High Temperature, High Resolution Metallography (AIME),

P. 217, Gordon and Breach, New York, 1967.22. K. E. Easterling and P. R. Swann: Inst. MetalsMon. 33, 1969, p. 152,23. K. Kinsman: 1970 EMSA Conference, Houston, p. 520, ClaitorsPublishers,

New Orleans, La., 1970.24. W. C. Leslie, D. W. Stevens, and M. Cohen: HighStrengthMaterials, p. 382,

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