THERMALLY INDUCED DEFORMATION IN

IRON - NICKEL - CARBON MARTENSITE

Yu. I. Kogan, G. V. Angelova, and V. P. Makarov

UDC 620.171.3 : 669.111.4 - 621.785.616

Causes underlying the strong temperature dependence of the y~eld point and of the low-temperature brittleness In the case of metals and alloys with a bcc crystal lattice are belng intensively studied. Con- cepts pertaining to this subject relate either to dlslocahons as such (Payerls forces, nuclear structure) or to the presence of interstitial inclusions; diverse test data have not yet produced a defimte answer, al- though there is no doubt that such inclusions have a sigmficant effect on the plastic flow.

As far as the martensite phase in iron alloys is concerned, the problem reduces to two specific as- pects: flrst of all, martensite represents a highly concentrated intersht~al solution and, secondly, it is the incluslon here which produces the high strength [i, 2], so that the question as to whether the carbon in the soluhon constitutes the thermally achvated barrlers to shp becomes one of considerable interest. There are data avmlable which inchcate that nonthermal stresses may be attributed to the carbon in mar- tenstte [4], while the activahon volume and energy of plastic deformation in martensite at low tempera- ture [5] have both been shown not to depend on the carbon concentration. At the same hme, it has also been discovered [5] that the activahon volume decreases with increasing carbon content and, consequently, the latter determlnes the temperature charactemstic of the flow stress.

Our study represents a continuatmn of that research, with the object here to determlne the tempera- ture characterlshc of resistance to plastic deformation and to measure the appropriate parameters of thermal achvahon in alloys with different martenmte morphology and structures. Iron-nickel alloys wlth 0.2% carbon and various amounts of nickel were used for thls study. Almost without affecting the mechan- ical properties [6], the nickel could shift the martensite temperature and change the martenslte type from massive to tvvanned. The same alloys were examlned after two different thermomeehamcal treat- ments. In adchtion to explaining the mechamsm of thermally induced deformahon in quenched steel, these tests also provided data for estimating the temperature characterlstic of the hardening process by one or another method of treatment.

Materlals, Procedure, and Results of Structural

Examination

The alloys were produced by vacuum melting with 0.2% carbon and 0, 16, or 22% nickel, corre- sponchng to a lowerlng of the martensite temperature from 380~ to room level. The specimens were quenched from 970~ in water (with an addition of NaCI, for faster coohng), whereupon they were imme- diately immersed in hquid nitrogen and kept there in storage ready for mechanical testing. In the 22% NJ alloy we recorded 4-5% remdual austemte, whlch was then taken into account by a hnear extrapolation of the flow stress to 100% martensite (as in [2, 3]).

The structure was examined on a model ]~MV-100 electron mlcroscope by passing light through thin f11ms. Typical photographs are shown in Figs. 1 and 2. The plain carbon steel and the 16% Ni alloy have the characteristic structure of massive (acicular) martensite (Fig. i). The martensite needles form stacks 300-600 A thick in plain carbon steel and 500-1000 ~ thick in the 16% Ni alloy. Only in few specimens of the 16% N1 alloy have there been found twin-structure islands containing simultaneously two systems of

Kuznetsov Siberian Inshtute of Physics and Engineering, State Univermty, Tomsk. Translated from Izvestiya Vysshikh Uchebnykh Zavedenit, Fizika, No. 7, pp. 81-86, July, 1974. Original art icle submitted Apml 3, 1973.

�9 76 Plenum Pubhshmg Corporatton, 22 7 West 17th Street, New York, N Y 10011 No part of thts pubhcatton may be reproduced, stored tn a retrieval system, or transmttted, tn any form or by any means, electromc, mechameal, photocopying, mterofilmmg, reeordmg or otherwzse, without written permtsston o f the pubhsher A copy o f this artlele ts avatlable from the publtsher for $15 O0

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Fig. i. Martensite structure in: (a) plain 0.2% car- bon steel, (b) the 16% Ni alloy with 0.2% carbon. Magnification x 65,000.

" �9 �9

t]~ jo3 i]~ iel o1~ 7~7 (220)Q o

0:7~ 0~2 07t 00o ~//oeP o~J r 9 o ( ~

Q �9 �9 �9 0 . 0 _ 0 _ ~/ /30 /2/ /TU /03 #4 /25

C

Fig. 2. Mar tens i t e s t ruc tu re in the 22% Ni al loy with 0.2% carbon. Magnificat ion • 65,000: (a) l ight-f ie ld image , (b) da rk - f i e ld image , (c) schemat ic of the e l e c t ronog ram de- pzcling the region shown in (a) and (b): the (311) plane of an rever ted m a t r i x - c r y s t a l la t t ice and the (i13) plane of a twin la t t ice , (112) is the twinning plane.

daslocations. The m a r t e n m t e zn the 22% Ni a l loy has a typical twin s t ruc tu re (Fig. 2), which has been con- f i rmed by mic rod i f f r ae t ion ana lys i s . The e l e c t r o n o g r a m s were decoded by the Me i r an -R ichman method [7]. An e l e c t r o n o g r a m of the region shown zn Fig. 2a has been reproduced in Fig. 2c: i t depzcts the indexed (311) plane of an inver ted m a r t e n s i t e m a t r i x - c r y s t a l lat t ice, along wtth supp lemen ta ry spots r e f e r r i n g to the (113) plane of an inver ted tw in -c rys t a l lat t ice; here (112) is the twinning plane. A da rk - f i e ld image of that spot in a (110) twin ref lex r e v e r s e s the con t ra s t (Fig. 2b) (the twins a r e Hluminated), which conf i rms our in te rpre ta t ion . The twins a r e 100-200 i thick and located 300-600 A apar t .

The quenched 16% Ni a l loy was a lso subjected to mechanica l t r e a t m e n t at high t e m p e r a t u r e (roLling to a 30% reduct ion a t 970~ and wa te r cooling) and to mechanica l t reaianent at low t e m p e r a t u r e (roLling to a 30% reduct ion a f t e r a p recool to 200~ Owing to the low p las t i c i ty of quenched plain carbon s teel , the act ivat ion volume and ene rgy of p las t ic deformat ion could not be m e a s u r e d here . On the o ther hand, the nickel a l loys could not be annealed r e h a b l y and this would have been n e c e s s a r y for compar ing the given proper tzes in spec imens a f t e r quench and a f t e r anneal. Such a compar i son was made w~th annealed plain carbon s teel , t he re fo re , because the l a t t e r in the quenched s ta te dad not d i f fer f rom the a l loys in t e r m s of the t empe ra tu r e -dependen t r e s i s t a n c e to deformat ion (F~g. 4) and contained m a r t e n s i t e of a l m o s t the s ame morpho logy as that in the 16% Ni alloy.

The r e s i s t a n c e to s h e a r T=(r/2 was m e a s u r e d a t a res idual deformat ion of ~= 0.6%, in o r d e r to avoid the "prof i le" effect of in terna l s t r e s s e s zn mar t ens i t e . The t e s t s were p e r f o r m e d on spec imens 20 m m long and 1 m m 2 in c r o s s section, a t an elongation ra t e of ~1=3.3 �9 10-4 /sec -1. The act ivat ion volume v* was m e a s u r e d by suddenly changing the elongation ra te to ~2=1.2 �9 10-2/sec -1 and then calcula~ang v* accord ing to the fo rmula

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OH (c) In -z /

o-." \ 7 / ~

The t e m p e r a t u r e c h a r a c t e r i s t i c s of %.~ w e r e m e a s u r e d r e p e a t e d l y , but e a c h t i m e on the s a m e s p e m m e n so a s to e l i m i n a t e u n c o n t r o l l a b l e d i f f e r e n c e s b e t w e e n s p e c i m e n s . A f t e r e= 0.6% had been r e a c h e d a t a g~ven t e m p e r a t u r e , t h e r e f o r e : f i r s t , t h e s t r e s s was m e a s u r e d , then the s p e c i m e n was p a r h a l l y u n l o a d e d and i m - m e r s e d in a h q u i d coo lan t ( m i x t u r e of l i qu id m t r o g e n and a l coho l ) , w h e r e u p o n the s p e c i m e n was a g a i n r e - l o a d e d and the s t r e s s e s u n d e r a chang ing t e m p e r a t u r e w e r e r e c o r d e d . Wi th the a i d of r (T) c u r v e s p lo t t ed on th~s b a s i s , the t e m p e r a t u r e - d e p e n d e n t c o m p o n e n t of r e s i s t a n c e to d e f o r m a t i o n r* could thus be e x t r a c t e d . F o r th i s p u r p o s e , then, t h e s e c u r v e s w e r e d i f f e r e n t i a t e d g r a p h i c a l l y and the d e r i v a t i v e d r / d T was p lo t t ed

a s a func t ion of the t e m p e r a t u r e . S ince =* = 0 , d: _ ".~ d:~ (r~ i s the n o n t h e r m a l s t r e s s component ) and dT v. dT d # / d T i s c l o s e to z e r o when r* = 0, hence an e x t r a p o l a t m n of the d ~ / d T ( T ) c u r v e to z e r o y i e l d e d the va lue of To, ~.e. , the t e m p e r a t u r e a t wh ich r* = 0; by a f u r t h e r e x t r a p o l a t i o n of the r (T) c u r v e to To, i t was p o s - s i b l e to s e p a r a t e T~ f r o m r * ( T ) . As ~t t u r n e d out, T0~ 300~ fo r a l l t r e a t m e n t s of s p e m m e n s m th~s s tudy . The r*(T) r e l a h o n s o b t a i n e d by t h i s p r o c e d u r e a r e shown in F i g . 4. T h e s e r e l a t i o n s and the f o r m u l a

(O~n~ ~ (0, ~ II= -- gT ~ \ ~ ] r \ - ~ ],

y i e l d e d the a c h v a h o n e n e r g y a t the r e s p e c t i v e t e m p e r a t u r e .

The t e m p e r a t u r e c h a r a c t e r i s t i c of s t r e s s e s was p lo t t e d fo r t h e - 5 0 t o - 1 9 6 ~ r a n g e , b e c a u s e s t r a i n a g i n g had been no ted a t r o o m t e m p e r a t u r e . Th i s o b s e r v a t i o n was c o n f i r m e d by the a n o m a l o u s t e m p e r a t u r e and s t r m n - r a t e c h a r a c t e r i s t i c s of r e s i s t a n c e to d e f o r m a h o n (F ig . 37, a l s o by the a p p e a r a n c e of a f low " f i n g e r " when a s p e m m e n i s he ld u n d e r l o a d and s u b s e q u e n t l y d e f o r m e d . The a c t i v a t i o n v o l u m e and e n e r g y w e r e m e a s u r e d a t s e v e r a l ch f fe ren t t e m p e r a t u r e s wi th in the s a i d r a n g e . An e x t r a p o l a t i o n of the v*(r* ) c u r v e to r* = 0 and the H(T) c u r v e to T = T o y i e l d e d the b r u i t i n g v a l u e s (v 0 and H0).

M e c h a n i s m o f T h e r m a l l y I n d u c e d D e f o r m a t i o n

The thermal component of shear stress r* as a funchon of the temperature is shown in Fig. 4, and the activation volume v* as a function of T* is shown in Fig. 5. It Is quite evident that the test points for the q u e n c h e d s p e m m e n s wi th a l l t r e a t m e n t s f i t on one cu rve ; the t e s t po in t s fo r the a n n e a l e d s p e c i m e n s f i t on the s a m e c u r v e . Thus s t e e l s wi th d i f f e r e n t m a r t e n m t e m o r p h o l o g y and s t r u c t u r e s have the s a m e T*(T) and v*(r*) c h a r a c t e r i s t i c s ; t h i s s u g g e s t s t ha t the t e m p e r a t u r e c h a r a c t e r i s t i c of f low s t r e s s i s in h e r e d e t e r m i n e d by the r e s i s t a n c e of the c r y s t a l l a t h c e to a sh i f t of i n d i v i d u a l c h s loc a t i ons , i . e . , to o v e r - c o m i n g the P a y e r l s b a r r z e r .

The a c h v a h o n e n e r g y H a s a func t ion of the t e m p e r a t u r e i s shown in F i g . 6. The t e s t v a l u e s fo r a l l s t e e l s and t r e a t m e n t s in t h i s s tudy fi t , w~th the s a m e d i s p e r s i o n , a b o u t a s t r a i g h t l i ne t h rough the o r ig in of c o o r c h n a t e s . An e x t r a p o l a h o n of t h i s l ine to T = T o y~elds H 0 = 0.5 eV, which zs equal to the e n e r g y of a doub ly ben t d i s l o c a t i o n 2U K in i r o n [8, 9], 1.e. , c o n f i r m s the P a y e r l s m e c h a n i s m of d e f o r m a t i o n . A n a l o - gous ly , the a c t i v a t i o n v o l u m e a t t e m p e r a t u r e s n e a r T o (F ig . 5) i s 40-50 V 3, m any c a s e no t m o r e than 120 V 3, which a g r e e s wt th the t e s t r e s u l t s in [8] and vnth c a l c u l a t i o n s b a s e d on the c r i h c a l width of a double bend [10].

The va l ich ty of t h e s e c o n c l u s m n s can be c he c ke d a g a i n s t the D o r n - R a i n a k r e l a t i o n s be tween the s h e a r

s t r e s s r* , the s t r e s s rp , and the P a y e r l s b a r r i e r 2U k. With a q u a s i p a r a b o l i c b a r r i e r we have H _ 2u~

( ~ ) 2 dH (~_12U~ in ~0/~, a s s u m i n g a s h e a r I - - and, a f t e r d i f f e r e n t i a h o n , v ~ = - - = 2 1 - - . S ince H = k T d~ ~ ~p / "r

modulus independent of the temperature will yield H/2U k = T/T O and, consequently, in order to compare theory and experiment, it is necessary to plot T*/rp versus T/T 0 and V%-p/2U k versus r*/~'p.

The theoretical curves for this particular case are shown in Figs. 7 and 8, also for one case of sl- nusoidal Payerls barriers which has been solved by numerical integration. The test values for all steels and treatments in this study follow the theoretical relahons rather closely. Any deviation in Fig. 8 is not larger than for annealed bcc metals [9].

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Deformatmn % ~

Fig. 3

50 a-t o-4 x-2 v - 5 I - 3 0 - - 6

o 80 /~0 240ToJ2g

Temperature, ~

Fig. 4

Fig. 3. Elongation curves fo r the 16% Ni al loy with 0.2% carbon: 1) tes ted at 20~ with ~=3 .3 .10 -4, 2) tes ted at 20~ with ~= 1.2- 10 -2, 3) tes ted at -60~ with

= 3 .3 .10 -4.

Fig. 4. Thermal component of shea r s t r e s s ~'* as a function of the t empera ture : 1) 16% Ni alloy quenched, 2) 16% Ni al loy with h igh- tempera tu re mechanical t rea tment , 3) 16% Ni alloy with low- tempera tu re m e - chanical t rea tment , 4) plain carbon steel annealed, 5) plain carbon steel quenched, 6) 22% Ni al loy quenched.

o

A - I

IO0 o - $ fl - - -

o - 4 8o ~ o,~

IO 20 $0 ~ ~ tOO 200 $00 Thermal component of ~ Temperature, ~ shear stress, kg/mm ~

Fig. 5 Fig. 6

Fig. 5. Activation volume v* as a function of ~*: 1) 16% Ni alloy quenched, 2) 16% Ni al loy with h igh- t empera tu re mechanical t rea tment , 3)plmn carbon s tee l annealed, 4) 22% Ni al loy quenched.

Fig. 6. Activation energy as a function of the t empera tu re : 1) 16% Ni alloy quenched, 2) 16% Ni al loy with h igh- tempera tu re mechanical t rea tment , 3) plain carbon steel annealed, 4) 22% N, alloy quenched.

The f requency fac tor g0 can be de te rmined e i ther f rom the graph in Fig. 6 or f rom the formula

0= ~ (O In ~ / �9 F o r the l~ayerls mechanism we have t 0 = ~w where L is the d is - In.=r ~ K ~ - ~ - ~ / ~

location segment set in m o r o n by any impact fluctuation, w is the cr i t ica l width of a double bend, and p is the density of slip dislocations. The cr i t ica l width can be calculated according to the formula in [9, 11] for a quasiparabol ic b a r r i e r

w 2 ~b--~-p / "

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0,8 ~.

0.4 ,

O.2

A-/ n-2

x-,~

I

~m

~2

' o-.~

Q-41 I /

%,

o ~,e 4~ ~ o,8rTge

Fig. 7 Fig. 8

Fig. 7. Theoretical relatlon "r*/Tp versus T/T0: solid hne refers to a sinusoidal barrler; dashed hne refers to a quasiparabolie barrier. Test points: i) 16% Ni alloy quenched, 2) 16% Ni al- loy with low-temperature mechanical treatment, 3) 16% Ni alloy with high-temperature mechani- cal treatment, 4) 22% Ni alloy quenched.

Fig. 8. Theorehcal relahon v*~-p/2U k versus T*/Tp: sohd line refers to a smusoidal barrier; dashed line refers to a quasiparabolic barrier. Test I)oints: i) 16% Ni alloy quenched, 2) 16% Ni alloy with high-temperature mechanical treatment, 3) plmn carbon steel annealed, 4) 22% Ni alloy quenched.

Calculations based on test data have yielded w = 18 V (almost the same as m [8, 9]) and ~0 = 10G/sec" One can now estimate p: for the 16% NI alloy p = 10l~ -2, whlchis close to values obtained by x-radiography for low-carbon martenslte.

Eleetron-ml croscope examinahon indicates that the dislocation density in martensite with a twin structure is low. This is, probably, why the value of the frequency factor for such a steel is somewhat lower. Assuming, as was done in [2], that in this case L is equal to the intertwin distance (400 i), we ob- tain according to the same formula p = 107/cm -2. The contribution of the given disloea~on density to the nonthermal component T ~ ape V~ ylelds approximately 6 kgf/mm 2, and the corresponding eontribuhon of the carbon can be estimated according to the formula in [12]: AT= 7.2" I0-2L-I/2 (wt. % C)I/3 = 30 kgf/ ram2; in total, this is almost equal to the test value T~t = 35 kgf/mm 2, which confirms the estimate.

These results confirm those in [3] and inchcate that the Payerls mechanism of deformation governs baslcally the temperature characterishc of flow stress.

1.

2.

3o

4.

5.

6.

7.

8.

LITERATURE CITED

G. V. Kurdyumov, in: Problems in Metallography and Metal Physics [in Russian], VoI. 4, (1955). P. G. Wmchell and M. Cohen, The Structure and Hardness of Martensite. Electron Microscopy and Strength of Crystals, New York-London (1963). M. J. Roberts and W. S. Owen, J. Iron and Steel Institute, 2.~.06, 4 (1968). M. A. Shtremel', Dissertation [in Russlan], Moscow (1972). P. Guiu, Phil. Mag., 21, 170, 365-374 (1970). R. Fleischer and D. Hibbard, Structure and Mechanical PropertLes of Metals [Russian translation], ~etallurgiya (1967). E. S. Melran and H. IV[. R1chman, Trans. AIIVIE, 227, 1044 (1963). H. Conrad, High-Strength Materials, New York-London-Sydney (1964).

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9. P. Guiu and D. Dorn, Impor t an t P r o b l e m s in the Theory of Dis locat ions [Russian t ransla t ion] , Ml r (1968).

10. A. Seeger, Phil. Mag., i , 651 (1956). 11. J .E. Dorn and S. Rainak, Trans. AIME, 230, 1052 (1964). 12. M. Cohen, J. Iron and Steel Inst., 201, No. 10 (1963).

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