Thermally induced deformation in iron-nickel-carbon martensite

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  • THERMALLY INDUCED DEFORMATION IN

    IRON - N ICKEL - CARBON MARTENSITE

    Yu . I. Kogan, G . V . Ange lova , and V . P . Makarov

    UDC 620.171.3 : 669.111.4 - 621.785.616

    Causes under ly ing the strong temperature dependence of the y~eld point and of the low- temperature brittleness In the case of meta ls and alloys with a bcc crystal lattice are belng intensively studied. Con- cepts pertaining to this subject relate either to d ls locahons as such (Payer ls forces, nuclear structure) or to the presence of interstitial inclusions; d iverse test data have not yet p roduced a def imte answer , al- though there is no doubt that such inclusions have a s igmficant effect on the plastic flow.

    As far as the mar tens i te phase in i ron alloys is concerned, the prob lem reduces to two specific as - pects: flrst of all, mar tens i te represents a highly concentrated intersht~al solution and, secondly, it is the incluslon here wh ich produces the high strength [i, 2], so that the question as to whether the carbon in the so luhon constitutes the thermal ly achvated barr lers to shp becomes one of considerable interest. There are data avmlab le wh ich inchcate that nonthermal stresses may be attributed to the carbon in mar - tenstte [4], whi le the act ivahon vo lume and energy of plastic de format ion in martens i te at low tempera - ture [5] have both been shown not to depend on the carbon concentration. At the same hme, it has also been d i scovered [5] that the act ivahon vo lume decreases with increasing carbon content and, consequently, the latter determlnes the temperature charactemst ic of the f low stress.

    Our study represents a cont inuatmn of that research, with the object here to determlne the tempera - ture character l shc of res istance to plastic de format ion and to measure the appropr iate parameters of thermal achvahon in alloys with different mar tenmte morpho logy and structures. I ron -n icke l alloys wlth 0 .2% carbon and var ious amounts of nickel were used for thls study. A lmost without affecting the mechan- ical propert ies [6], the nickel could shift the martens i te temperature and change the martens l te type f rom mass ive to tvvanned. The same alloys were examlned after two different thermomeehamca l treat- ments . In adchtion to explaining the mechamsm of thermal ly induced de formahon in quenched steel, these tests also prov ided data for est imat ing the temperature characterlstic of the harden ing process by one or another method of treatment.

    Mater la l s , P rocedure , and Resu l t s o f S t ruc tura l

    Examinat ion

    The alloys were produced by vacuum mel t ing with 0.2% carbon and 0, 16, or 22% nickel, cor re - sponchng to a lower lng of the mar tens i te temperature f rom 380~ to room level. The spec imens were quenched f rom 970~ in water (with an addition of NaCI , for faster coohng), whereupon they were imme- diately immersed in hqu id nitrogen and kept there in storage ready for mechan ica l testing. In the 22% NJ alloy we recorded 4 -5% remdua l austemte, wh lch was then taken into account by a hnear extrapolation of the f low stress to 100% martens i te (as in [2, 3]).

    The structure was examined on a mode l ]~MV-100 electron mlc roscope by pass ing light through thin f11ms. Typical photographs are shown in Figs. 1 and 2. The plain carbon steel and the 16% Ni alloy have the characterist ic structure of mass ive (acicular) martens i te (Fig. i). The martens i te needles fo rm 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 article 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 of the pubhsher A copy of 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. Martensite structure in the 22% Ni alloy with 0.2% carbon. Magnification 65,000: (a) l ight-field image, (b) dark-f ie ld image, (c) schematic of the e lectronogram de- pzcling the region shown in (a) and (b): the (311) plane of an reverted matr ix -c rys ta l lattice and the (i13) plane of a twin lattice, (112) is the twinning plane.

    daslocations. The martenmte zn the 22% Ni al loy has a typical twin structure (Fig. 2), which has been con- f i rmed by microdif f raet ion analysis. The e lectronograms were decoded by the Meiran-Richman method [7]. An e lectronogram of the region shown zn Fig. 2a has been reproduced in Fig. 2c: it depzcts the indexed (311) plane of an inverted martens i te matr ix -c rys ta l lattice, along wtth supplementary spots re fer r ing to the (113) plane of an inverted twin-crystal lattice; here (112) is the twinning plane. A dark-f ie ld image of that spot in a (110) twin reflex reverses the contrast (Fig. 2b) (the twins are Hluminated), which conf irms our interpretat ion. The twins are 100-200 i thick and located 300-600 A apart.

    The quenched 16% Ni al loy was also subjected to mechanical t reatment at high temperature (roLling to a 30% reduction at 970~ and water cooling) and to mechanical treaianent at low temperature (roLling to a 30% reduction after a precool to 200~ Owing to the low plast ic i ty of quenched plain carbon steel, the activation volume and energy of plastic deformation could not be measured here. On the other hand, the nickel al loys could not be annealed rehab ly and this would have been necessary for comparing the given propertzes in specimens after quench and after anneal. Such a comparison was made w~th annealed plain carbon steel, therefore, because the latter in the quenched state dad not differ f rom the al loys in te rms of the temperature-dependent res is tance to deformation (F~g. 4) and contained martens i te of a lmost the same morphology as that in the 16% Ni alloy.

    The res is tance to shear T=(r/2 was measured at a residual deformation of ~= 0.6%, in order to avoid the "prof i le" effect of internal s t resses zn martensi te. The tests were per formed on specimens 20 mm long and 1 mm 2 in cross section, at an elongation rate of ~1=3.3 9 10-4/sec -1. The activation volume v* was measured by suddenly changing the elongation rate to ~2=1.2 9 10-2/sec -1 and then calcula~ang v* according to the formula

    972

  • OH (c) In -z /

    o-." \7 /~

    The temperature character i s t i cs of %.~ were measured repeated ly , but each t ime on the same spemmen so as to e l iminate uncont ro l lab le d i f fe rences between spec imens . A f te r e= 0.6% had been reached at a g~ven temperature , there fore : f i r s t , the s t ress was measured , then the spec imen was parha l ly unloaded and im- mersed in a hqu id coolant (mixture of l iquid mtrogen and alcohol) , whereupon the spec imen was again re - loaded and the s t resses under a changing temperature were recorded . With the a id of r (T) curves p lotted on th~s bas is , the temperature -dependent component of res i s tance to de format ion r* could thus be ext racted . For th is purpose , then, these curves were d i f fe rent ia ted graph ica l ly and the der ivat ive dr /dT was p lotted

    as a function of the temperature . Since =* =0, d: _ ".~ d:~ (r~ is the nonthermal s t ress component) and dT v. dT d#/dT is c lose to zero when r* = 0, hence an ext rapo la tmn of the d~/dT(T) curve to zero y ie lded the value of To, ~.e., the temperature at which r* = 0; by a fu r ther ext rapo la t ion of the r(T) curve to To, i t was pos - s ib le to separate T~ f rom r*(T) . As ~t tu rned out, T0~ 300~ for a l l t reatments of spemmens m th~s study. The r*(T) re lahons obta ined by th is p rocedure are shown in F ig . 4. These re la t ions and the fo rmula

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

    y ie lded the achvahon energy at the respect ive temperature .

    The temperature character i s t i c of s t resses was p lotted for the -50 to -196~ range, because s t ra in aging had been noted at room temperature . This observat ion was conf i rmed by the anomalous temperature and s t rmn- ra te character i s t i cs of res i s tance to de formahon (F ig. 37, a lso by the appearance of a f low " f inger" when a spemmen is held under load and subsequent ly de formed. The act ivat ion vo lume and energy were measured at severa l chfferent temperatures within the sa id range. An ext rapo la t ion of the v*(r* ) curve to r* = 0 and the H(T) curve to T = T o y ie lded the bru i t ing va lues (v 0 and H0).

    Mechan ism o f Thermal ly I nduced Deformat ion

    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 quenched spemmens with a l l t reatments fit on one curve; the tes t points for the annea led spec imens f it on the same curve. Thus s tee ls with d i f ferent mar tenmte morpho logy and s t ructures have the same T*(T) and v*(r*) character i s t i cs ; th is suggests that the temperature character i s t i c of f low s t ress is in here determined by the res i s tance of the c rys ta l la thce to a shi f t of ind iv idual chs locat ions, i .e . , to over - coming the Payer l s bar rzer .

    The achvahon energy H as a function of the temperature is shown in F ig . 6. The tes t va lues for a l l s tee ls and t reatments in th is study fit, w~th the same d ispers ion , about a s t ra ight l ine through the or ig in of coorchnates. An ext rapo lahon of th is l ine to T = T o y~elds H 0 = 0.5 eV, which zs equal to the energy of a doubly bent d i s locat ion 2U K in i ron [8, 9], 1.e., conf i rms the Payer l s mechan ism of de format ion . Ana lo - gously, the act ivat ion vo lume at temperatures near T o (F ig. 5) is 40-50 V 3, m any case not more than 120 V 3, which agrees wtth the tes t resu l t s in [8] and vnth ca lcu la t ions based on the c r ihca l width of a double bend [10].

    The val ichty of these conc lusmns can be checked aga ins t the Dorn-Ra inak re la t ions between the shear

    s t ress r*, the s t ress rp, and the Payer l s bar r ie r 2U k. With a quas iparabo l i c bar r ie r we have H _ 2u~

    (~)2 dH (~_12U~ in ~0/~, assuming a shear I -- and, a f te r d i f fe rent iahon , v ~ = - - = 2 1 - - . Since H=kT d~ ~ ~p / "r

    modulus independent of the temperature will yield H /2U k = T /T O and, consequently, in o rder to compare theory and exper iment , 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 Payer l s barr iers wh ich has been solved by numer ica l integration. The test values for all steels and t reatments in this study fol low the theoretical re lahons rather closely. Any deviation in Fig. 8 is not larger than for annea led bcc meta ls [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 for the 16% Ni alloy with 0.2% carbon: 1) tested at 20~ with ~=3.3.10 -4, 2) tested at 20~ with ~= 1.2- 10 -2, 3) tested at -60~ with

    = 3.3.10 -4.

    Fig. 4. Thermal component of shear stress ~'* as a function of the temperature: 1) 16% Ni alloy quenched, 2) 16% Ni alloy with high-temperature mechanical treatment, 3) 16% Ni alloy with low-temperature me- chanical treatment, 4) plain carbon steel annealed, 5) plain carbon steel quenched, 6) 22% Ni alloy 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 alloy with high-temperature mechanical treatment, 3)plmn carbon steel annealed, 4) 22% Ni alloy quenched.

    Fig. 6. Activation energy as a function of the temperature: 1) 16% Ni alloy quenched, 2) 16% Ni alloy with high-temperature mechanical treatment, 3) plain carbon steel annealed, 4) 22% N, alloy quenched.

    The frequency factor g0 can be determined either from the graph in Fig. 6 or from the formula

    0= ~ (O In ~ / 9 For the l~ayerls mechanism we have t 0 = ~w where L is the dis- In.=r ~K~-~-~/~ location segment set in moron by any impact fluctuation, w is the crit ical width of a double bend, and p is the density of slip dislocations. The crit ical width can be calculated according to the formula in [9, 11] for a quasiparabolic bar r ie 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. Theoret ical relatlon "r*/Tp versus T /T0 : solid hne refers to a sinusoidal barrler; dashed hne refers to a quasiparabol ie barrier. Test points: i) 16% Ni alloy quenched, 2) 16% Ni al- loy with low- temperature mechan ica l treatment, 3) 16% Ni alloy with h igh- temperature mechan i - cal treatment, 4) 22% Ni alloy quenched.

    Fig. 8. Theorehca l re lahon v*~-p/2U k versus T*/Tp: sohd line refers to a smuso ida l barrier; dashed line refers to a quasiparabol ic barrier. Test I)oints: i) 16% Ni alloy quenched, 2) 16% Ni alloy with h igh- temperature mechan ica l treatment, 3) p lmn carbon steel annealed, 4) 22% Ni alloy quenched.

    Calculations based on test data have yielded w = 18 V (a lmost the same as m [8, 9]) and ~0 = 10G/sec" One can now est imate p: for the 16% NI alloy p = 10 l~ -2, wh lch is close to values obtained by x - rad iography for low-carbon martensl te.

    E leet ron-ml c roscope examinahon indicates that the dislocation density in martens i te with a twin structure is low. This is, probably, why the value of the f requency 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 accord ing to the same fo rmula p = 107/cm -2. The contribution of the given d is loea~on density to the nonthermal component T ~ ape V~ ylelds approx imate ly 6 kg f /mm 2, and the cor respond ing eontr ibuhon of the carbon can be est imated accord ing to the fo rmula in [12]: AT= 7.2" I0-2L-I/2 (wt. % C)I/3 = 30 kgf/ ram2; in total, this is a lmost equal to the test value T~t = 35 kg f /mm 2, wh ich conf i rms the estimate.

    These results conf i rm those in [3] and inchcate that the Payer l s mechan ism of deformat ion governs baslcally the temperature character i shc of f low stress.

    1.

    2.

    3o

    4. 5. 6.

    7.

    8.

    L ITERATURE C ITED

    G. V. Kurdyumov, in: P rob lems in Meta l lography and Meta l Phys ics [in Russian], VoI. 4, (1955). P. G. Wmche l l and M. Cohen, The Structure and Hardness of Martensite. E lectron Mic roscopy and Strength of Crystals, New York -London (1963). M . J. Rober ts and W. S. Owen, J. Iron and Steel Institute, 2.~.06, 4 (1968). M . A. Shtremel' , Dissertat ion [in Russlan], Moscow (1972). P. Guiu, Phil. Mag. , 21, 170, 365-374 (1970). R. F le ischer and D. Hibbard, Structure and Mechan ica l PropertLes of Meta ls [Russian translation], ~eta l lurg iya (1967). E. S. Me l ran and H. IV[. R1chman, Trans . AIIVIE, 227, 1044 (1963). H. Conrad, H igh-St rength Mater ia ls , New York -London-Sydney (1964).

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  • 9. P. Guiu and D. Dorn, Important Prob lems in the Theory of Dislocations [Russian translation], Mlr (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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