This is to certify that the thesis entitled THE EFFECT OF PLASTIC DEFORMATIOH ON THE THERMAL STABILlZATlON 0F AUSTENITE-MARTENSITE TRANSFORMATION presented by G l E HONG KANG has been accepted towards fulfillment of the requirements for M. S . degree in Metallurgy 0-7639 MS U is an Affirmative Action/Equal Opportunity Institution lV1531.J RETURNING MATERIALS: Place in book drop to LJBRAfiJES remove this checkout from “—- your record. FINES will “ be charged if book is , returned after the date stamped below. a.‘ “PO 1; u. ‘i I I "z. W. " i, 3" 1 ‘\_ ' \é'pjrf. . t ‘ I ‘ ’ _‘ ‘ I" s .L' ~ I — . -~ | ”“4“" “*" ’V |'>~k‘/ ,‘. r- ' a: w 3 d. ..,‘ « l - THE EFFECT OF PLASTIC DEFORMATION ON THE THERMAL STABILIZATION OF AUSTENITE-MARTENSITE TRANSFORMATION By Gie Hong Kang A THESIS Submitted to Michigan State University in partial fulfillment of the requirements for the degree of MASTER OF SCIENCE Department of Metallurgy, Mechanics and Materials Science 198“ ABSTRACT THE EFFECT OF PLASTIC DEFORMATION ON THE THERMAL STABILIZATION OF AUSTENITE-MARTENSITE TRANSFORMATION BY Gie Hong Kang The influence of plastic deformation on the thermal stabilization of austenite-martensite transformation has been studied in a 16Ni 10 steel. The studies include the effect of i) Extent of deformation prior to martensite formation ii) Rate of straining employed for deforming the specimens iii) Amount of martensite before stabilization and iv) straining subsequent to martensite formation. Deformation prior to martensite formation decreases the MS-temperature significantly. This decrease is very sensitive to the strain rate employed. The temperature interval over which the transformation is inhibited, following stabilizing treatment after partial martensite formation, decreases with increasing prior deformation. However, this trend varies with the percent transformation preceeding the stabilization treatment. In specimens strained after partial transformation, stabilization in- creases with increasing deformation. The transformation both before and after stabilization starts in general with a burst; and the size of a burst increases with deformation. The course of transformation after stabilization is not much different in deformed specimens and the microstructure indicates that martensitic plates formed in the deformed austenite are finer. The observed results appear to be dominated by strainaging effect. ACKNOWLEDGEMENTS The author is much indebted to Visiting Professor Dr. S.C. Das Gupta and Professor Dr. K. Mukherjee for their constant interest in this work, and for many ideas, comment and enlightening discussions. The author also wishes to express his deepest appreciation to his family. TABLE OF CONTENTS Abstract Acknowledgements List of Figures List of Tables 1. Introduction ............................................. 2. Theoretical Aspects ...................................... 2—1 Martensite Transformation 2-2 Principal Characteristics 2-3 The Nucleation of Martensite 3. Previous Work ............................................ Isothermal Formation of Martensite Stabilization of Austenite-Martensite Transformation 3 1 3 2 3—3 Burst Formation of Martensite 3 A Effect of Plastic Deformation of Martensite Formation 1 Materials A 2 Austenitizing A-3 Determination of MS-Point and Measurement of Extent of Transformation A-U Method of Deformation A—S Stabilization Treatment 5 Results .................................................. 5—1 Effect of Strain Rate on MS-Point 5-2 Effect of Prior Deformation on MS-Point 5-3 Effect of Time of Stabilizing Treatment on the Extent of Stabilization 5-4 Effect of Deformation prior to Formation of Martensite on Stabilization 5-5 Effect of Deformation of Partially Transformationed on Stabilization 5-6 Effect of Deformation on Microstructure 6. Discussion ............................................... 7. Conclusions .............................................. Appendix References O-t‘J: 21 21 23 27 31 Figure 10 LIST OF FIGURES a) Schematic Plot of Temp. vs. Time Depicting the StabiliZing Treatment b) Course of Martensite Formation with Decreasing Temp. below MS 1) without Interruption ii) with Interruption Schematic Shape of a Martensite Plate Knapp and Dehlinger Model of the Martensite Embryo Experimental Apparatus a) Schematic Diagram Depicting the Treatment b) Schematic Diagram Depicting the Treatment 0) Schematic Diagram Depicting the Treatment d) Schematic Diagram Depicting the Treatment e) Schematic Diagram Depicting the Treatment Variation of Ms-Point as a Function of Strain Rate: 2% Elongation Effect of Prior Deformation in the Austenitic Condition on MS-Point in a 16Ni -lC steel. a) without b) with Additional Stabilizing Treatment Change in Electrical Resistance as a Function of Temp. in Fe—lei-lC alloy Deformed to Variable Extent prior to Transformation a) Degree of Stabilization as a Function of Stabilizing Time at 39‘C Containing about 30% prior formed Martensite without prior Deformation b) Degree of Stabilization as a Function of Stabilizing Time at 39°C Containing about 30% prior formed Martensite with 2% prior Deformation Variation of Extent of Thermal Stabilization with Varying Amount of Deformation for 1 Hour Aging Treatment at 39‘C. Specimens are First Deformed, then Cooled to Produce (A) 35% and (B) 50% Mantensite and up Quenched to the Aging Temp. of 39 C Page “9 52 53 56 Figure 11 12 13 1“ Variation or the Extent of Stabilization in Specimens Deformed with 30% prior Martensite a) Microstructure of Specimens Containing 30% Martensite b) Microstructure of Specimens Containing 30% Martensite with 0.5% prior Deformation a) Microstructure of Stabilized and Retrans- formed Specimen Showing Effect of 0.5% Deformation following 30% Martensite Formation b) Microstructure of Stabilized and Retrans- formed Specimen following 30% Martensite Formation - Schematic Presentation of g vs. Temp. Page 57 59 60 65 Table LIST OF TABLES Page Effect of Strain Rate on MS Temp. for 2% A5 Elongation Degree of Stabilization as a Function of Time 51 at the Aging Temperature a) undeformed b) deformed 2% prior to Transformation Variation in the Degree of Stabilization as a 55 Function of the Extent of Deformation prior to Formation of Martensite a) 35% Martensite b) 50% Martensite I. INTRODUCTION Thermal stabilization of austenite—martensite reaction is an inhibition of the martensite formation resulting from an aging treatment. The course of formation of martensite in a steel during cooling may be influenced by the cooling procedure adopted. In general, if the cooling is arrested within the martensite transformation range and then allowed to age for a period of time, the transformation does not start on subsequent cooling, until some lower temperature is reached. There is a likelyhood that at subsequent temperatures, the total amount of martensite is somewhat less than that produced by direct cooling to the temperature concerned. This phenomenon is referred to as stabilization. Figure 1 shows a schematic representation of the process. The difference between the temperature of arrest and the restarting of the transformation, following aging treatment as represented in the diagram, is often taken as a measure of degree of stabilization. The stabilization phenomenon will have influence on the amount of retained austenite in the martensitic steels. A cooling rate dependency exists in alloy steels quenched in water and in oil. The thermal stabilizationcfi'austenite in steels is thus of importance because of technological signi— ficance in heat—treatment and scientific implications relative Figure l. LIJ a: :3 F— : ------ AGING TEMP. “J .......... DEFOMATION D. 2 m -..-..- _---ARREST TEMP. *— T I ME UJ i..— U) 2 UJ '— 0: <1: 2 $ Ms: TAMs TEMPERATURE a) Schematic Plot of Temp. vs. the Stabilizing Treatment b) Course of Martensite Formation with Decreasing Temp. below MS 1) without interruption ii) with interruption Time Depicting to the martensite transformation. Presence of retained austenite in a hardened steel is unwanted for various reasons. It may result in a softer marterial and may induce structural as well as dimensional instability. 2. THEORETICAL ASPECTS 2—1 Martensite Transformation Phase transformations in solid state may be divided into two categories essentially: a) Nucleation and growth type (Diffusion controlled) b) Displacive or shear type (Diffusionless) a) Nucleation and growth type transformation In this type of transformations, a new phase grows at the expense of parent phase by relatively slow movement of atoms over multiple of atomic distances. Movements are independent and markedly vary with temperature. Reaction proceeds isothermaly and amount of phase increases with time. b) Displacive or shear type transformation In this type of transformations, cooperative movements of many atoms are utilized instead of independent movement of individual atoms. Rearrangement takes place in an orderly displacive manner. Atomic movements are of minor order and in principle atoms do not change position with neighbors. Martensite may form in a number of alloys of which the transformation in steel is most important and has been studied extensively. 2-2 Principal Characteristics of Martensitic Transformation There are morphological, thermodynamic and kinetic features which can be considered as typical for martensitic transformations. (1) A plate-like morphology with a small ratio of thickness to other linear dimensions is the typical structural element resulting from martensitic transformations. (2) The habit plane has a definite orientation to the crystalline axes of the initial and final phases. (3) There is an orientational relationship between the lattice of the initial and final phases. (A) The transformation is accompanied by a change in the form of the transformed region which manifests itself in a characteristic relief on the plane surface of the specimen at the site where a new phase plate appears. This macrodeformation is uniform and represents a combination of a simple shear along the plane of the plate and expansion or compression normal to this plane. The macrodeformation parameters are constant for each transformation. (5) The martensitic crystals have a regular internal structure. Often, the martensitic plate is internally twinned, faulted or dislocated. (6) The growth rate of the martensitic crystals is high (lOu-lO5 cm 8.1) and does not display any noticeable temper— ature dependence. (7) According to the operative nucleation kinetics, martensitic transformations are described as athermal or isothermal. In the case of athermal transformation the nucleation rate is high and does not display any temperature dependence. In the case of isothermal transformation the rate of nucleation depends essentially on temperature. (8) Isothermal and athermal transformations can occur in the same material. (9) The amount of transformation is characteristic of temperature, provided other variables, such as grain size, are held constant. Transformation on cooling begins spontane- . ously at a fixed temperature (MS), and as the temperature is changed, more and more material transformations, until the temperature (Mf) is reached, at which the change is complete. (10) Martensitic reactions are reversible in the sense that an initial atomic configuration can be repeatedly obtained. A single crystal of the original phase may transform on cooling into several crystals of the new phase. The reverse change on heating will usually result in a single crystal of the same size, shape, and orientation as the original crystal. The reversibility is associated with a temperature hysteresis and the reverse reaction begins at a temperature above MS (except in some thermoelastic cases). Moreover, in repeated transformations, the plates which form on cooling have the same size and shape, and appear in the same regions of the original crystal. This behavior probably applies in principle to all martensitic reactions. In Fe-C alloys, this, however, is not normally observable since the process is interfered by the tempering phenomenon. By employing extremely rapid rate of heating, tempering phenomenon could be arrested and reversi- bility followed in Fe-C alloys. (11) Plastic deformation is.moreimportant in martensitic transformations than in nucleation and growth changes. Application of plastic stress at any temperature in the trans— formation range usually increases the amount of transformation, and the reaction can often be completed by this means. In some transformations, elastic stresses have a similar effect. When single crystals are used, the direction of the applied stress is important, and some reactions may be inhibited as well as aided by a suitably orientated stress. Deformation above Magmay also result in the formation of the product phase, even though the temperature is too high for spontaneous reaction. The highest temperature at which martensite may be formed under stress is called Md‘ In general, the reverse reaction can be aided in the same way, and a suitable stress will induce transformation below the temperature at which it begins spontaneously. If the original phase is cold-worked in a temperature range where it is stable, e.g., at temperatures sufficiently above Ms’ the resultant deformation may either raise or lower the transformation start temperature depending upon the degree of cold work. When the degree of cold work is relatively small MS may be raised and with relatively larger deformation it may be depressed. (12) The transformation is also inhibited in another way. If the specimen is cooled to a temperature in the transformation range, held there for a period of time, and then cooled again, transformation does not begin again immediately. Supercooling is then necessary to form any further martensite. At all subsequent temperatures the amount of transformation is less than that produced by direct cooling to the temperature concerned. This phenomenon may be referred to as permanent stabilization. The degree of stabilization increases with the time for which the specimen is held at the temperature. Slight variations in amount of transformation with cooling velocity are also presumably to be attributed to stabilization. There is no general agreement whether or not stabilization is produced by halting the cooling above the Ms temperature. Isothermal Martensite Formation--In many ferrous systems, such as Fe-Ni, martensitic transformation can occur both athermally and isothermally. The isothermal transformation kinetics show a typical C-curve behavior indicative of a thermally activated reaction. However, the increase in volume percent of martensite at any isothermal temperature is still due to the formation of new plates rather than the growth of the existing plates. Further, experimental results 1’2 that time of formation of individual isothermal 7 plates is also ~ 10- sec. This process, therefore, is indicate related to thermally activated triggering of embryos, and once triggered the growth no longer requires any thermal activation. During isothermal transformation, it is observed that the initial rate of transformation is greater. This higher initial rate is related to "autocatalysis" where formation of plates stimulates nucleation of other plates. The actual mechanism of this stimulationis not well understood. The subsequent decrease in transformation rate is related to partitioning of the austenite phase, i.e., volume fraction transformed per nucleation event is reduced since these nucleation events occur in progressively smaller volume of austenite. The transformation rate equation, therefore, must take into account both autocatalysis and partitioning effects. If f is the fraction of martensite formed, then the following 3 type of semiempirical rate equation may be used. df 1+—l. -dt- = (ni+pf-Nu) (l—f) vexp (— AWa/RT)mq(l-f) m ...(1) where f = fraction transformed time d ll 10 N1 = number of pre-existing nucleation sites (estimated 107/00) p = number of autocatalytic embryos per unit volume Nu = number of martensitic plates per unit volume V lattice vibration frequency ZkWa = activation energy of nucleation at temperature T m thickness to diameter ratio of plates q average volume per grain of austenite Equation (1) can be integrated numerically by substituting metallographically determined values of q and m and unknown quantitities p andASWa are determined by curve fitting. 3 Experimental results agree well for up toi~10% transformation. Further manipulation of this equation is possible.3 2-3 The Nucleation of Martensite* It has been emphasized earlier that kinetics of martensitic transformation is nucleation dominated. In order to consider nucleation and kinetics of this transformation, let us first consider a simple model of a martensite plate. We shall take an oblate spheroid shape“ as shown in Figure 2. The radius of the oblate spheroid is r and the semi- thickness is c, where r) c. The interfacial free energy for * The review section of nucleation is largely due to: K. Mukherjee, International Summer Course on Martensitic Transformation, Leuven, Belgium, 1982. ll - 4 2 Volume = -3— «r c Surface area i 2nr plane ' Figure 2. Schematic shape of a martensite plate (or an embryo) 12 such a plate is then: VA.gSP+M = 2 nr2x (per plate) ..................... (2) where\/is the volume of the plate andlagSP’M is the surface energy per unit volume and.kis the interfacial energy. Using linear elasticity theory, the strain energy of such a plate is given by: — ll VLig P M = nr2C (—AE—) (per plate) ............. (3) e 3 r where V: nr2C is the volume of the plate and (Ac/r) is the strain energy per unit volume. The factor A has the form: A = {11(2—v)/8(I-v)} ”3+ 9‘s: ..................... (3a) or A = u (r§+ 83,) ............................ (3b) wherev = Poisson's ratio,11= shear modulus and n,and 8, are the shear and dilatational strains associated with shape deformation. In the above equation, elastic constants for the parent and the martensitic phases are assumed to be the same . PaM c as the chemical free energy Using the quantityzig per unit volume of martensite, we have: VASPIM = % nr2cAg59M (per plate) .............. (A) We note that below To,z;gZ’M is negative. Thus the total free energy change associated with the formation of a plate is: l3 AWP’M =.._:_ nrzCagi‘M + 21rr2k + —:—1rr 02A ....... (5) Using the classical homogeneous nucleation theory, we * can determine the free energy of nucleation W corresponding * x to a critical radius r and a critical semithickness C . To do this we set: PeM are r - o and .............................. (6) flzo C We thus obtain: 0* = 2KAagEIM ............................. . ....... (7a) r* = 2Ao*Agg:*M .................................. (7b) V* = —5—1rr*2o* ..... . ......... ......... ........ ...(7c) * «r w = 3217A2x3/3'Qig: M)“ .......................... (7o) 0*2/r* = A/A ...................................... (7e) If the nucleation is completely random then each atom is a potential nucleation site and thus the nucleation rate N is given by: N = (Ng/Vm) vexp (—W*/kT) ......................... (8) where (No/Vm) = number of atoms per unit volume,i:is the lattice vibration frequency (~1012/sec.). The nucleation kinetics predicted by equation (8) should follow an isothermal C-curve behavior, the maximum nucleation rate being at the knee of the C-curve. 1U From the classical homogeneous nucleation model it is possible to calculate V*, c*, r* and W* by using experimental values ofzdgEIM and theoretical values of A and A.“ But classical homogeneous nucleation model can not be applied in martensitic transformations. Recent experimental results as well as theoretical considerations have clearly established that martensitic nucleation is heterogeneous. In classical model, an embryo attains a critical size at a temperature and becomes nucleus. As an alternative to this a "reaction-path" model has been proposed5 in which, through a coordinated movement of atoms within the nucleation volume, the lattice passes through a succession of inter- mediate structures in one region and propagates like a strain wave. Thus, activation is achieved by fluctuations in confi- guration rather than the size. The embryo associated with such a mechanism, therefore, could be a strain center in which ” local arrangement of atoms is part way along the reaction path. Lattice imperfections, such as dislocations, internal surfaces etc. could provide such centers. Above MS, these regions represent centers for high free energy. Using Frank's dislocation model6, Knapp and Dehlinger7 have treated kinetics of athermal martensite formation in steel. This model requires pre—existence of embryos and deals with the propagation rather than nucleation of an embryo. 15 In this model, the martensite embryo is an oblate spheroid surrounded by dislocation loops. These loops are primarily screw dislocations with short edge components joining the positive and negative screws. Growth along [110}, and I225}, can occur by the expansion of the loops and growth along 155A17 requires generation of new loops. This model of the embryos is shown in Figure 3. The formation and propagationcfiTthese loops are equi— valent to the creation dfan interface and the energy required for this process is supplied by the chemical driving force Agg”°l Similarly the strain energy associated with this embryo is also supplied by the chemical free energy change. The embryo in this case is triggered as soon as the overall free energy change per unit volume becomes zero. Writingzaera as -%—13W7Ia' and using our previous convention regarding surface and strain energies, we write: -a' —v ' -o ' -~ ' Aw? ___ _Ag'ci: a +Ag; a +Ag78 a = O ........ (9) Thus for spontaneous triggering raa' -.- 7-oa' 7-.a’ ‘ age (Age +Ags ) .. ............ . ...... (10) \ We can write the last two terms in eq, (10)as Agnon-chem and using equations (2) and (3) for an oblate spheroid: = 3A -Ac Agnon-chem .5-0— + I’ ........................... (ll) \ - \ where r, c,x and A have the same meaning as before. For a radius,¢ognon_chem can be minimized with respect to c by l6 '[ Figure 3. Knapp and Dehlinger Model of the Martensite Embryo. l7 setting (aAg /60) = 0 non—chem 6AA )% _ 3A non-chemfiin( r c ..................... (l2) — Thus Ag where c = (—;A£-)% ................ . .................... (13) 2A ' USlng equations (12) and (13), the minimum lnzxgnon—chem now can be calculated as a function of embryo size. It can be seen from eq. (12), that for a fixed value of c, the minimum value of Ag decreases with increasing radius, non-chem r. That is, larger the embryo, smaller is the non-chemical energy barrier. It also means that as the embryo grows, non-chemical barrier becomes smaller and smaller, thus the propagation of the interface could be very rapid. Below Mg, ) chemical forces acts as a "stress" to move the dislocation idgzfia >Cig for the larger embryos and net non-chem min interface. It is also possible to visualize the interaction of an applied stress with such an embryo both above and below the MS. An estimate of the largest embryo (at MS) can be made by using appropriate experimental and estimated values of various parameters. The above model has been modified by Kaufman.“ In his model, an equivalent dislocation loop, encircling the oblate spheroid embryo, on the habit plane is considered. This circumferential dislocation loop then contains the interfacial energy of the embryo. The radius of this loop is r and its 18 equivalent Burger's vector,§ , is given by cb/d, where c is the semithickness, b is the Burger's vector of the lattice and d is the distance between the loops in Knapp and Dehlinger model (Figure 3). The line tension of this loop is then, P =l‘8/2 , whereu = shear modulus. If the loop is subjected to a net shear stress 7, the corresponding increase in free energy is (starting from zero size) W9 = 2an2—flfafr ........: .................. - ....... (14a) where I‘=u§3/2 ' 9 .................................. (lAb) The shear stresswtis equal to chemical driving force (-zAgC) minus strain energy per unit volume Ac/r. Thus T=_Agc——l-\I—f3—=—Agc—-%-' oooooooooooooooooooo (15) Where Ais the surface energy. We note from eq. (7e) that c2/r = A/A. Thus substituting from eq. (15) and (lAb) in (lAa), we have: Q = «bkr3 ub _£;_% 6 w ___—d [___-Ad +Agc( A) +1] ............. (1) By maximizing W9 as a function of r, we obtain _ 9AA C — Ag2 ooooooooooooooooooooooooooooooooooooooo (l7) 2 3 Q 91rA A and W = -_-—H- ................................... (18) C 23g where rc is the critical radius of the loop at which it can lower its energy by expanding. This growth model is similar to mechanical twinning or kink band formation. Comparing 19 with eq. (7b): rc = -%— r* .... ................................... (19) Since re) r*,.AW* is actually negative at rc It has been proposed (17) that during isothermal trans- formation those embryos with radius r given by r*< rrc are already triggered by the athermal process. In more recent years, some more models have been developed for heterogeneous nucleation of martensite. Olson and Cohen8 have studied a model in which a portion of a wall of dislocations underwent faulting on close-packed planes to produce the close—packed phase in the fcc to hop transformation. A second shearing is required to produce the fcc to bee transformation. No activation barrier occurs in their thermo- dynamic equations, but a barrier is postulated to result, however, from the thermally activated motion of the dislocations associated with the internal slip as first suggested by 9 Magee. The free energy change for nucleation of a martensite embryo attached to and inside a growing dislocation loop has 10 No nucleation been computed by Easterling and Tholen. barrier was found for internally twinned martensite. Their choice of 20 ergs/cm2 for the interfacial free energy is 2O much smaller than the 200 ergs/cm2 estimated by Kaufman and 3 Cohen for a martensite embryo. Use of 200 ergs/cm2 would yield an impossibly large barrier for nucleation in their model. Suezawa and Cook11 suggest that a pile-up of dislocations can initiate heterogeneous nucleation of marten- site and the classical free energy barrier to nucleation can exist for a martensite embryo at a pile up. 21 3. PREVIOUS WORK 3-l Isothermal Formation of Martensite Although martensite formation is more generally a temper- ature-dependent process, its time-dependent nature has also been revealed in some cases. It was observed as early as 12’13 and as much as 1948-A9 by Fletcher, Averbach and Cohen 5% isothermal martensite formation was noted in both plain carbon and low alloy steels. It was, however, considered then as a sort of tailing off effect of the main reaction. 1A,15 put forth a revolutionary Kurdjumov and Maksimava claim to have suppressed the transformation completely in a 6Mn 0.6C steel by rapid cooling in liquid N and as much as 2 25% martensite formed on reheating and holding isothermally at temperatures. Rate at one temperature was found to have an initial constant rate and followed by a gradual decrease over a greater interval of time. A temperature was found where greatest amount of isothermal martensite formed. Partial suppression of martensite was claimed by them in a 1.6C steel. By quenching first in alkaline iced water (about 20% martensite formed) and then further cooling to liquid N temperature suppressed most of the transformation. 2 0n reheating isothermal transofmration could be studied. Partial suppressioncfi‘martensite formation was also indicated by Das Gupta and Lement16 in a 15 Cr 0.70 steel, although 22 on athermal amount proceeded the isothermal part. A maximum initial rate was noted at an intermediate temperature, (-110°C). At liquid N2 temperature, starting with an athermal amount of about 15% as high as 27-28% amount formed on isothermal holding. Similar type of isothermal formation preceded by 17 athermal amount was also noted by Machlin and Cohen in an Fe-Ni alloy. A few years later, Kulin and SpeichlB, Cech and Hollomon19 reported almost complete isothermal formation in Fe—Cr—Ni alloys and in Fe-Ni—Mn alloys respectively. Subsequently Shih, Averbach and Cohen20 reported complete isothermal formation at -l96‘C in a different but similar composition as Cech and Hollomon. Further evidence of isothermal formation of martensite in ferrous alloy followed and in more recent 21 reported isothermal formation in a year Wasaka and Wayman Fe-Ni-Mn alloy of slightly different composition from that of Shih, Averbach and Cohen. It may be seen from the vairous work that athermal martensite is not essential for isothermal martensite. However, since martensite transformations are strain sensi- tive and autocatalytic, athermal martensite may stimulate the nucleation of isothermal martensite when the cooling is interrupted below MS. Thus nature of isothermal formation is studied either in absence of athermal martensite or more 23 conventionally by introducing a fixed quantity of athermal martensite. Whether the reaction takes place in the absence of athermal martensite or in the presence of a fixed amount, the isothermal kinetics display a C-curve behavior, with the maximum rate coming at some intermediate temperature. In almost all instances the isothermal reaction precedes by the nucleation of new plates, rather than by the growth of existing ones and the plates reach the full size within a small fraction of a second. Kurdjumov has concluded that nucleation rather than growth is the rate limiting step in isothermal formation and metallographic observations confirm this. 3-2 Stabilizationcfi‘Austenite-Martensite Transformation First systematic study was done by Harris and Cohen22 as early as 19A9 in a 1.10 1.5Cr steel. They observed that unless the temperature of holding is below a certain temper- ature below MS (called as) stabilization does not occur and the degree of stabilization occurring an holding below aS increases linearly as the temperature of holding is lowered. Effect of stabilization on isothermal formation of martensite in a Fe 29Ni alloy was reported by Machlin and Cohen23 in 1951. They tried to explain the process as resulting from relaxation of existing martensite embryos. 2“ Working on a 15Cr 0.7C steel Das Gupta and Lement2u observed that when the amount of martensite as formed by interrupting the cooling below MS exceeds a critical amount the stabilization effect induced on subsequent cooling to lower temperature is more marked. However, stabilization effect eventually disappears on continued holding at such temperature by isothermal formationcfl‘martensite. Besides the authors further observed that higher the temperature of intermediate cycling the more permanent is the stabiliza- tion effect. 25 reported for steels containing about 1C Morgan and Ko 0-5Ni that stabilization of austenite occurs during isothermal holding above and below Ms' The rate of stabilization varies with the composition of the austenite, increases with increasing temperature and is greatly increased by the presence of mar- tensite. The stabilization above MS causes a depression of Ms and increases the amount of austenite at a reference temperature near Ms. Edmondson26 examined stabilization in a 10Ni 10 alloy and suggested a theory for stabilization on the diffusion of interstitial solute atoms to preferred sites within the metal to explain the results obtained. The work of Woodilla, Winchell and Cohen27 showed that in a 30.8Ni 0.007C steel the activation energy for the 25 stabilization process is comparable to that for the diffusion of carbon (or nitrogen) in ferrite rather than in austenite. They suggested that the interstitial diffusion controlling the stabilization occurs within the martensite embryos rather than in the matrix of the parent phase. Kinsman and Shyne28 reported that a fixed population of martensite embryos is affected by the segregation of carbon to the embryo-austenite interface. The anchoring effect of the segregated carbon atoms inhibits the activation of the embryos into propagating martensite plates. Thermal stabilization of the athermal martensite trans- formation was evaluated in Fe-Ni-C alloys in the isothermal temperature range of 0-200°C by Kinsman and Shyne.29 They reported a distinct change in the characteristics of thermal stabilization at low and high temperature of aging. At low againg temperature the extent of thermal stabilization varied with time, increasing to a maximum value and the maximum decreased with increasing aging temperature. The phenomenon behaved according to the authors in a manner consistent with a mechanism based on pinning of the interface of the martensite nucleus by a carbon atmosphere formed during aging. At higher temperature of aging beyond 80°C the extent of stabili— zation as a function of time was initially constant and then began to increase at long times. The authors report the 26 observations at high temperature of aging to follow some other mechanism. Jin and Huang30, working on a Fe 29Ni 4.3Ti observed stabilizationcfl7high strength austenite as well as that of annealed austenite and attributed this to the formation of G.P zones or short-range order of less than 10A in size. The initial rate of MS suppression was larger in an ausaged and subsequently transformation-strengthened austenite than in a simply ausaged austenite. Theories on Stabilization of Austenite-Martensite Transformation--Theories of stabilization attribute the phenomenon either to change within the untransformed austenite or to change in the effect which already formed martensite has on the subsequent progress of the transformation. There is a consensus of opinion that thermal stabilization of austenite is caused by relaxation of stress in the austenite around the existing martensite plates. This stress is believed to assist nucleation of fresh plates in an unstabilized austenite. However, how this relaxation takes place has been explained differently by different investigators. Thermal stabilization is a thermally activated process involving mostly interstitial solutes—kinetics being governed by diffusion of interstitials: C or N. 27 (1) Das Gupta and Lement214 are of the view that relaxa- tioncfi‘strain-embryos takes place by diffusion of C-atoms from the martensite present to the strain—embryos in austenite. The driving force for the latter process is provided by the difference in carbon activity which exists. In presence of martensite, activity of carbon in martensite is presumably higher than in austenite. Thus, there will be a tendency for the C-atoms to diffuse from the edges of the martensite plates into strain embryos in the vicinity. 25 proposed that C-atoms can diffuse (2) Morgan and Ko during slow cooling or isothermal holding to form "clusters" or precipitates around the existing dislocation and thereby pinning the dislocations. This would harden the structure and thus make growth of austenite-martensite interface and accompanying deformation more difficult. (3) Based on the model of Knapp and Deblinger, it has been proposed by Kinsman and Shyne28 that thermal stabiliza— tioncfl‘austenite is a result of immobilization of the semi- coherent interface between martensite and austenite matrix because of segregation ofinterstitial solute atoms to the dislocation array comprising the interface. The time and temperature dependence of thermal stabilization must be directly related to the kinetics of solute segregationix>the surface of the nucleus. Thus by developing a time law for interface 28 segregatimitflueauthors develop a time law for thermal stabili- zation. (A) Glover and Smith31 suggest that stabilization may be due to relaxation of elastic stresses by diffusion of carbon and nitrogen inside the martensite, leading to first stage of tempering. (5) According to Maksimova et a132, the relaxation process leaves structural imperfection in its wake and is thus doubly effective in causing stabilization because such imperfections impede the growth of martensite plates. 33 (6) According to Kurdjumov , a readjustment of the stresses and imperfections that ordinarily assist transform- ation is likely to occur in a transformed alloy even at room temperature. The structure that results from this is more stable and is more difficult to transformation. 3-3 Burst Formation In certain cases under suitable condition the progress of martensite formation is accompanied with "burst" formation —-a large volume of transformation taking place in an abrupt manner. The "burst" may occur either at MS or at a temperature below MS. The "burst" may, also occur on long isothermal holding after an incubation period. Recommencement of transformation following a stabilizing treatment may, under certain conditions, 29 be accompanied by a burst. The resultant transformation curves have been shown to be irregular in shape, especially when the size of the burst is relatively small. With progress of transformation the transformation-temperature curve tends to smooth out but may still remain undulating. When the size of the initial burst is large, the course of the subsequent transformation is more regular. The increment in percent martensite immediately following such a burst is small; this process has been named "pseudo-stabilization". The process is autocatalytic in nature and metallographic examination stow these plates forms in a zig-zag pattern in the specimen. The process is somewhat analogus to a chain reaction, the stresses produced by one plate assisting in the activation of the nucleus of another plate. This also means that the plates formed during the burst are such as to reduce the stress in the matrix, the whole group forming a partially self-accommodating system and this is probably why the individual plates are much wider than normal. Machlin and Cohen17 found that about 25% of the whole transformation in an iron-nickel alloy was produced in a burst of very short duration, accompanied by an audible click. Entwisle and Feeney3u suggested that where bursts occur, it appears that the first successful nucleating event triggers 30 off the whole burst, and that Mb truly represents the start of transformation. The burst temperature M is essentially b a function of the thermal treatment to which the austenite has been subjected. The magnitude of the burst is a function of the austenite grain size and the temperature at which the burst has occurred. 35 Interpretation of Burst-—Suzuki and Honma referred the burst transformation akin to deformation twinning. The propagation of such a transformation requires that the dis- location bordering the martensite embryos can multiply along certain favored habit planes. The mechanism of multiplication is such that this process is not favored on another common type of habit planes found in steel viz 225, the transformation then occurs by a process similar to slip. 36 did not agree with this view. They Machlin and Cohen suggest that when martensite forms the austenite is plastically deformed and as a result, new embryos are created or the existing ones made supercritical at a given temperature. Simultaneous activation of a large number of such embryos result in a burst. Cech and Turnbull37 offered a completely different interpretation. They argue that since the burst phenomenon is more prevalent in Fe-Ni-C alloys and as these alloys are 31 relatively free from the heterogenities which nucleate martensite in these alloys, a large degree of supercooling is necessary to nucleate martensite. When transformation starts, it does so form many centers, thus a causing a burst. Philibert and Crussard39 found C and N atoms responsible for stabilization of austenite and the subsequent burst transformation. They proposed a theory according to which the simultaneous activation of many embryos is made possible by the propagation of an elastic wave. 3-A Effect of Plastic Deformation on Martensite Formation Martensite transformations can be affected in a variety of ways by the application of stress and by plastic strain. Patel and Cohen39 demonstrated quantitatively how the initial nucleation of martensite is assisted by applied stress. A mechanical work term consisting of the product of the applied stress field and the transformation strain is added to the free energy change that provides the driving force for the transformation. Low temperature plastic deformation of the parent austenite prior to the martensite transformation influences both the nucleation and the growth of martensite. Nucleation is assisted by small amounts of plastic deformation, while extensive cold working suppresses nucleation. The retarding effect of large plastic deformations on the subsequent martensitic transformation has been reported by 32 Kurdjumov.£40 Breinan and Ansell“1 studied the effect of deformation on the forward transformation in an Airkool-S steel and found that the decrease in the martensite start temperature, MS, with deformation applied above Md (the highest temperature at which deformation can initiate the formation of martensite) correlated with the change in austenite yield strength. They concluded that MS decreased with increasing yield strength as a consequence of reduced dislocation mobility, which may alternately affect either the nucleation of new martensite or the propagation of the martensite-austenite interface. Increasing the yield strength by alloying decreased MS in the same manner. In a similar study, Ankara“2 used transformation cycling of an Fe 30Ni alloy to increase the austenite yield strength and also found a decrease in MS with increasing yield strength, which he attributed to a change in dislocation density. According to Popeu3, the effect of plastic deformation introduced by rolling at room temperature on the austenite start temperature of an Fe 30.3Ni 0.0050 alloy has been determined. The austenite start temperature increases monotonically with deformation. Microhardness measurement show that the austenite start temperature increases with the yield strength of the martensite. 33 Guimaraes and ShyneM4 reported that plastic deformation of austenitic specimens (Fe -28.A5Ni -0.2lC, Fe -3l.15Ni —0.095C) at room temperature caused a rise in the MS temper- ature. The effect was maximum, about 15°C, at an intermediate stage of deformation, 50% reduction by rolling. The so called mechanical stabilization of austenite could only be observed in specimens severely deformed (more than 8% reduction by rolling). Maxwell, Goldberg and Shyneu5 reported that in the Fe-Ni—C alloys two distinctly different types of martensite formed concurrently with plastic deformation. The large difference in morphology, distribution, temperature dependence, and other characteristics indicate that the two martensites form by different transformation mechanisms. The first type, stress—assisted martensites, is simply the same plate martensite that forms spontaneously below MS except that it is somewhat finer and less regularly shaped than that formed by a temperature drop alone. This difference is due to the stress-assisted martensite forming from cold-worked austenite. The second type, strain-induced martensite, formed along the slip bands of the austenite as sheaves of fine parallel laths. Since imperfections play an important role in the nucleation process of martensite and thermal stabilization is 3“ associated with same sort of relaxation process in this austenite around the existing martensite involving rearrange- ment of imperfections and immobilization of the embryo sites present, it is expected that plastic deformation before and after martensite formation would have considerable influence on the stabilization process. Keeping this in view it is proposed to have a new look at the stabilization process in a suitable steel with varying amount of deformation at different stages. 35 A. EXPERIMENTAL PROCEDURE A-l Material The chemical composition of the steel investigated is as follows: C% — Ni% - Fe% 1 l6 Balance Specimen preparation: Material was obtained in a rec- tangular bar form of an approximate dimension of 2.1" x 0.08" x 0.05". Experimental specimens ranging from 0.040" - 0.060" thickness were obtained by warm rolling in austenite condition in several steps. Austenitizing before rolling was carried out in evacuated and sealed quartz tubes. Decarburization was avoided as much as possible during warm rolling. “-2 Austenitizing Experimental specimens were austenitized for 2 hours at 110000 and then water-quenched to room temperature. The austenitizing temperature was selected on the basis of solution of carbides and reasonable austenite grain size. Prior to austenitizing treatment all specimens were given a stabilizing treatment by cooling to liquid nitrogen temperature. It may also be noted that in order to obtain comparable results a single specimen was repeatedly used while running a definite series of experiments. Ms-temperature did not materially alter on repeated austenitizing of the same specimen. 36 “—3 Determination of MS-Point and Measurement of Extent of Transformation MS—point and the progress of transformation of austenite to martensite was determined by following the change in electrical resistance using a Kelvin double bridge. Specimens used for electrical resitance measurement were approximately of the following dimensions: 0.0“" - 0.06" thickness 0.08" width 2.1" length The specimens were gripped by a suitable holder having knife- edge contact points for the current and potential leads. The specimen dimensions were manipulated so as to maintain electrical resistance ranging between 10-2 and 10-3 ohms. Subzero temperature necessary for following the progress of transformation was obtained by holding the specimens at different levels of the airspace above liquid nitrogen kept in a tall thermal flask as shown in Figure A. The specimens were lowered in steps to different levels manually and the rate of cooling was maintained approximately as 2°C per minute. A range of temperature between +20°C and ~195°C could be followed by the set up. The temperatures were measured by a Leeds-Northrup potentiometer using a copper constant thermo- couple placed in contact with the specimen. 37 Ice Water C. p“ —— _ \ \. I I + - 3 A Kelvin Bridge 5 . , Potentiometer Thermal Flask Holder \\ , Specimen Thermal Flask r’T'Liquid Nitrogen Figure A. Experimental Apparatus 38 A-A Method of Deformation Specimens were deformed by tensile elongation in an Instron Tensile Machine using a friction grip. The rate of loading was maintained normally at 0.05 cm per min. and was varied between 0.2 cm t0 0.02 cm per min. in selected cases. The extent of deformation ranged between 0.5% to A.0% usually. A-5 Stabilization Treatment Specimens containing a definite amount of martensite and deformed in various manners were reheated to +39°C and aged for 1 hour before finally sub—cooling to follow the course of transformation. In selected specimens time of aging was varied from 5 min. to A hours. Schematic drawings of the heat treatments involved are shown in Figure 5. 39 H - heating A - Austenitizing WQ - Water Quenching D - Deforming AA SC - Subzero Cooling S - Stabilizing LLI m D wo *" H < o: D LL] 0.. 2 C Li.) |.... TIME Figure 5—a Schematic Diagram Depicting the Treatment A0 H - Heating A — Austenitizing WQ - Water Quenching D — Deforming AA , SC - Subzero Cooling S - Stabilizing Lu ‘1’- w :> G +— H < o: D S LIJ 0- SC 2 LLI P- TIME Figure 5-b Schematic Diagram Depicting the Treatment A1 H - heating A — Austenitizing WQ - Water Quenching IX D - Deforming SC - Subzero Cooling S — Stabilizing UJ O: 3 H We F- <12 0: S UJ l 0.. S 2 C so LLI .— TIME Figure 5-c Schematic Diagram Depicting the Treatment A2 H - Heating A - Austenitizing WQ - Water Quenching AA D — Deforming SC - Subzero Cooling m o O: S — StabiliZlng 3 W0 F- 4 H o: S UJ a ‘1- SC 2 SC LIJ F- T I M E Figure 5-d Schematic Diagram Depicting the Treatment A3 H - Heating A - Austenitizing SQ - Water Quenching D — Deforming SC — Subzero Cooling AA S - Stabilizing UJ Cl: D we F— < H S a: D “J I 0. 2 SC so In '— TIME Figure 5—e Schematic Diagram Depicting the Treatment AA 5. RESULTS 5-1 Effect of Strain Rate on MS-Point Austenitized and quenched specimens were deformed at room temperature to a constant 2.0% tensile elongation at different rates of straining. Cross-head speed of 0.2 cm/min., 0.05 cm/min. and 0.02 cm/min. were used. The effect of different straining rates on the MS—point has been presented in Table 1. The treatment followed is shown in the schematic diagram (Figure 5-a). An increasing depression of Ms-point with decreasing rate of straining may be noted. Figure 6 shows the variation of MS-point as a function of rate of strain prior to martensite formation. 5-2 Effect of Prior Deformation on MS—Point Austenitized and quenched (above the Ms temperature) specimens were deformed to various extents by tensile elonga- tion prior to martensite formation. The extent of deformation ranged from 0.5% to A.0% (strain). The effect of prior deformation on the MS-point has been studied (1) by cooling subsequent to deformation and (2) by stabilizing at 39°C for 1 hour before transformation. In both cases the MS-point is substantially depressed and with increasing deformation the extent of depression of MS increases consistently. A5 HHI m.o hmpmz ..Hn H pom ooOOHH mHI mo .0 Among» ..HQ H .Hom Do OOHH mm! No.0 hmpmz ..HQ H 90% UoQOHH rim... .2... taxman oofioemeoam em pom whopmpodfioe m2 :0 opmm chnpm mo poomem .H oHnme A6 Ms (0C) 1 J 0.! 0.2 STRAIN RATE (556') Figure 6. Variation of MS—Point as a Function of Strain Rate: 2% Elongation A7 However, the degree of lowering of MS in specimens having stabilizing treatment at 39°C for 1 hour is smaller than that in specimens having no stabilizing treatment as shown in Figure 7. The rate of decrease of MS with deformation appears to be the same in both cases. Resistance change followed as a function of temperature for specimens deformed to various extent has been plotted in Figure 8. Large change in resistance with martensite formation, indicating burst formation, can be noted in both underformed and deformed specimens. The size of burst increases with the extent of prior deformation. It may be further noted that although the undeformed specimen shows two successive relatively smaller burst with beginning of transformation the deformed specimens start transforming with one large burst. The course of transformation as followed by resistance measurement does not change with varying extents of prior deformation. Similar trend of the course of transformation is noted.hispecimens deformed and stabilized at 39°C for 1 hour prior to trans- formation (treatment b). The treatment followed is shown in the schematic diagram (Figure 5—a and 5-b). 5-3 Effect of Time of Stabilizing Treatment on the Extent of Stabilization Specimens with and without prior deformation (2%) were transformed to have 30% martensite and subsequently stabilized Figure I A) B) l j I l l o I 2 3 4 % DEFORMATION EffectcfiTPrior Deformation in the Austenitic Condition on MS-Point in a l6Ni—1C Steel A) Without B) With Additional Stabilizing Treatment 60 56 q 52 m." u 48 2: <1 [—- u: 44 U) u] a: 40 36 32 Figure x6 4 “9 ¢—--l A% deformation ‘/’2 2% 3 1% V§4 0.5% 5 0% l l 1 1 1 -50 -3o -40 IO 30 TEMPERATURE,°C Change in Electrical Resistance as a Function of Temperature in Fe—l6Ni—1C alloy Deformed to Variable Extents prior to Transformation 50 at 39‘C for various lengths of time. The variation in the extent of stabilization have been reported as a function of time in Table 2 and a plot of the data is given in Figure 9. The treatment followed is shown in the schematic diagram (Figure 5—c and 5-d). In the undeformed specimen, the extent of thermal stabilization, 0, varied with time, increasing to a maximum value, and then decreasing at longer aging times. The degree of stabilization for A hours is somewhat less than that obtained for 1 hour stabilization. The deformed specimen behaves similarly, but the degree of stabili- zation is higher in this case. For example, the degree of stabilization for an aging time of 60 minitues is 27‘b for the deformed sample compared with about 20‘C for the un- deformed case. 5-A Effect of Deformation prior to Formation of Martensite on Stabilization Austenitized and quenched specimens were deformed to various extents ranging from 0.5% to A.0% by tensile elongation prior to martensite formation of two different amounts. Results are observed for 35% and 50% martensitic transformation. The treatment followed is shown in the schematic diagram (Figure 5-d). The alteration in the extends of stabilization following stabilizing treatment at 39°C for 1 hour appears to have different trends. This is particularly so when the 51 Fm mm! .EHE Om mml N.mm Rwl : .. q mm moi .gHE om Hal m.©m :ml : : p _ m. m.NH m.oml .GHE OH Mal 2.2m m.Hml : = m m MH :ml .GHE m HSI ©.mm m.mml : = m.ea m.ms- .ase s mm- a.mm m.m = = 9 N.ON h.m:l .LQ H m.mml :.Nm m.ml = = (\ m 9: am- .35 cm mom- m.mm Rm ._ = w w... T? ma- .5s m mam- H.mm 3.3 seems a: doe once: .m. p. in m .oEHB. . <9 opHmcmppmz m manumhdeoB <9 2 ‘2 wcHNHHHnmpm .Qsoe pmomp¢ & _z wcHSocmsa wQHNHpHcopms< .COprEpommcmpp on LOHLQ gm cosmomoc ADV weapomocc: Adv .oHSpwpodEoB wcHw< onp pm oEHB mo QOHpoczm m we :OHpmNHHHnmpm mo common .m oHnme 66%) (N 0 IO STABILIZATION, h) 0 Figure 9-a 52 l l l I 2 3 AGING TIME (HOUR) Degree of Stabilization as a Function of Stabilizing Time at 39°C Containing about 30% prior formed Martensite without prior Deformation L) 0V (D «30 2: C) F— <( N20 .J an <1 P-IO U) Figure 9-b 53 I I l l l 1 IO 20 30 4O 50 60 AGING TIME (MIN) Degree of Stabilization as a Function of Stabilizing Time at 39°C Containing about 30% prior formed Martensite with 2% prior Deformation. 5A extent of deformation was rather small such as 0.5%. In case of specimens containing 35% martensite the degree of stabilization increases with 0.5% deformation and with further increase in deformation consistently decreases up A% deformation. 0n the other hand in case of specimens containing 50% martensite the degree of stabilization decreases with increase in deformation up to 2% and then slightly increases when extent of deformation was A%. The trend of alteration in the degree of stabilization having variation in the extent of deformation followed by different amount of martensite formation has been given in Table 3 and Figure 10. 5-5 Effect of Deformation of Partially Transformed Samples on Stabilization Specimens containing 30% martensite, formed by sub- cooling, were deformed at room temperature by tensile elongation of 0.5 and 1% and subsequently stabilized at 39°C for 1 hour. The extent of deformation could not be increased as the specimens failed in brittle manner beyond about 1% strain. When compared as a function of increasing strain, the degree of stabilization increases linearly with deformation (strain) as shown in Figure 11. The treatment followed is shown in the schematic diagram (Figure 5-e). 55 em 9%. moo- mm- 9% s m TH 9%. G. 58. 11.3 m m 5mm 52- em- :17 RE A m men 93.- mm- mi- h? m. m am mm- mm- 3.: fem o more; .2: n8 Deco: m. m.om so- m.mau em- s.mm : m. E me. am. em- mam m m mm as. em- m.mm- e.mm H m mam also- am- mam- 9% m. m. N.mm m.ma- mm- m.m- m.mm O house .aeH hoe eoooaa MW 8 «enmzumw we as as oofinomenoz soaeeMnouoo mofinoeoso mommwwmmwmmmw oofinoooaoz Rom on .opHmcophmz Rmm Adv mpHmnopmz mo QOprEpom ou AOHEQ soamemome mo pcopxm on» mo GOHpossm m mm COHpmuHHHnmum mo oomwmm on» :H COHpmHmm> .m oHnme .. 34 " k) $30. 0) gees { a..- '2‘. 318- ..‘E ml 4 - l l 1 l l o l 2 3 4 ‘ 960EFORMATION (A) 34°“ 0" q)3(3" 532- t— 328— 324— S (020;- l l L l l O I 2 3 4 zgummwnmv (B) Figure 10. variation.ofExtent of Thermal Stabilization with Varying Amount of Deformation for 1 Hour Aging Treatment at 39°C. Specimens are First Deformed, then Cooled to Produce A) 35% and B) 50% Martensite and up Quenched to the Aging Temperature of 39‘E. 04 O m 0 STABILIZATION,9(°C) I 0 Figure 11. 57 O .5 I % DEFO RMATION Variation ofthe Extent of Stabilization in Specimens Deformed with 30% prior Martensite. 58 5-6 Effect of Deformation on Microstructure Optical photomicrographs of specimens transformed to ~30% with 0.5% strain and without prior strain have been presented in Figure 12. Prestrained specimen show an effect of bunching or clustering of martensite plates. The plates are thinner and shorter in the strained situation. It may be referred that earlier Durlu and Christian”6 showed that a decrease of the semi-thickness-to-radius (c/r) ratio of martensite plates resulted in prior plastically strained Fe 26.A Ni 0.2AC alloy. Work of Datta and Raghavanl‘l7 further indicated that martensite plates become progressively thinner with increasing prestrain. When optical micrographs of specimens deformed after formation of martensite (..30%) and then retransformed after stabilization at 39°C is compared with the specimens, similarly treated but undeformed, some noticeable change in this microstructure may also be confirmed. The deformed specimen retransformed with formation of extremely fine plates within the residual austenite space (Figure 13). This effect on the shape and size of the martensite formed after stabilizing treatment is clearly the result of deforma- tion. Martensite formed in strained lattice is much smaller and thinner. 59 Heofiz “a I pronoun Aoozxv oofiposhoeom gonna Rm.o guns moansopnez “om mofiofioosoo mcoEHoQO mo endposnpmonon nImH Oponm 3 Hasz R: I pcmnoum hooaxv oefimcopnez Rom mofiefieosoo mcoaHoodm mo endposnmeAOHz mImH ouocm Heofiz A: - esmroom Aooaxv Aeem I oofimoousez Hapoev COHmepom opHmcophmz Rom wcHBOHHom coEHoon UoEmommcmhpwm Ucm ooufififiomom do onsoosupmonofiz puma ouoed Hmsz a: I pcmcopm Aooaxv gems standouts: Hmuoev COHmemom opHmcwpawz Rom meaaoafiou sofiossnoeoo em.o no oeoeum wcHsonm :oEHoQO Uthommcmhpom use ooufififiomem do ossuosnononofiz eImH ooond 61 DISCUSSION Plastic deformation prior to martensite formation causes substantial depression of MS-point in the alloy (Fe — l6Ni - 1C) under study and the extent of depression consistently increases with increasing deformation within the present experimental range of 0.5% to A% strain. Martensite transformation in deformed specimens is initiated by a large burst. The size of a burst is observed to increase with increasing deformation. Depression of MS-point is also found sensitive to rate of straining. A slower strain rate, for a constant amount of deformation, produces a greater depression of the Ms-point. Such a depression of Ms-point, in plastically deformed specimens, is not usually experienced. Normally small plastic deformation of the austenite is expected to raise the MS temperature. However, we believe that this observed depression of MS is related to high carbon content of the alloy as will be discussed. When austenite is plastically deformed above the Md temper- ature, dislocations are introduced and thereby regions of high internal stress are created. This situation would cause segre- gationcfl‘C-atoms to the region of dislocation arrays comprising the potential embryos which would have otherwise raised the MS if the interstitial content (C, N2 etc.) was zero or very low. Besides segreation of C-atoms to the dislocations, introduced by deformation, in a high carbon alloy, carbon segregation 62 can also occur at preexisting embryonic sites. The latter will imply a depression of MS in the absence of plastic deformation. Both of these mechanisms are operative in a plastically deformed high—carbon alloy. The decoration of dislocations increase the matrix strength and thereby introduces a kinetic barrier to the martensite interface motion. The potency of martensite embryos on the other hand is decreased by carbon diffusion during aging above MS. These factors increase the chemical driving force necessary for transformation and thus MS is depressed. Relatively speaking, if the strain aging due to deformation (i.e. strengthening of matrix) has a greater effect on the Ms, then strain rate will have a noticeable effect. In fact the strain-rate sensitive nature of the depres- sion of MS-point, for a fixed amount of prior deformation (2%), leads to such a conclusion in the present case. Slower rate of straining produces larger depression of MS-point. The burst size during martensite formation is a function of grain size and the chemical driving force at Mb3u. The grain size remaining constant, larger burst with increased depressionof MS-point is quite in order and our experimental result are, therefore, consistent. The degree of stabilization, 0, in partially transformed specimens, with prior plastic strain of varying amounts, seems to have somewhat different trends when the amount of transform— 63 ation is varied. Two different amounts of martensite have been studied, viz (l)~ 50% and (2)'~35%. For stabilization treatment conducted after ~50% martensite formation, the degree of stabilization, 0, has been found to decrease monotonically with the amount of prior strain upto the uppermost strain employed in the present study (~'A%). However, if the stabilization treatment is given after forming ~35% martensite formation, 0 is found to increase first (upto ~0.5% strain) then decrease monotonically (beyond 0.5% strain upto A% strain. It has been suggested that internal stress due to plastic deformation or partial transformation can aid martensite formation by a process of stimulation of the existing embryos.l‘l8’u9 On the other hand, plastic deformation results in strain hardening of austenite. Dislocation accumulation become insurmountable obstacles for the propagation of the martensite plates. The strain-hardening of austenite is expected to depress the Ms—temperature. A decrease in MS-temperature is expected to take place with increasing yield strength of austenite at Ms-temperature. A comparison of the energy necessary to deform surrounding austenite (in order to accommodate transformational volume and shear strains) with chemical driving force at MS in a series of Fe-NiHCr-C alloy demonstrated that the resistance to plastic deformation represents a substantial portion of 6A 50 non-chemical free energy for this alloy series. Earlier work on Fe 26.ANi 0.2AC and Fe 2ANi 0.A5C has shown that Mb temper- ature first rises with deformation and then falls with higher deformation. If interstitials such as carbon etc. are present then an additional effect of hardening due to strain aging must be considered as has been pointed out earlier. Strain-aging effect appears to be dominating in this alloy even when the amount of strain is rather low (0.5%). This is evident in the nature of plot showing the effect of depression of Ms-point as a function of increasing strain (Fig. 7). It is now necessary to inquire why the stabilization increases when the amount of prior deformation is small (-0.5%) and the stabilization is conducted after a small volume % transformation (e.g.'~35%). The chemical driving force, Age, for martensitic transformation can be defined as A gc. = (T)50%M } and (AEC)35%M< (AEC)50%M where (T)35%M and (T)50%M are temperatures at which 35% and 50% martensite respectively form. Thus for small deformation and small % of martensite the driving force is relatively small due to two factors: 1) Suppression of MS is small. and 2) (To — Mg) is small. Therefore we may conclude that the small increment of driving force Mose) = a (.Ms' - Ms) is not sufficient to overcome the strain-aging induced kinetic barriers as discussed earlier. If, however, the % transformation, after which stabilization is conducted, is large, then the sample must be cooled down to a lower temperature (e.g. T50%M) and hence the driving force is larger. If this driving force is large enough to overcome kinetic barriers due to strain aging, then stabilization will decrease. As the % strain increases, (TO - Mg) increases and thus the stabilization decreases again. This tentative model seems to explain the observed variation of stabilization with % transformation as 67 well as % plastic strain. Our last result concerns the effect of deformation after partial transformation. When a partially transformed sample ('~30% martensite) is deformed, the degree of stabili- zation increases linearly with increasing strain up to 1% (Fig. 11). Unfortunately, such an experiement could not be conducted at larger strains because of the brittle manner of samples after giving 1% strain. Since martensite was formed prior to deformation, the effect of deformation is highly localized in the untransformed austenite volume. Under this circumstance the result obtained can be easily explained by assuming, strain aging and strengthening of the remaining austenite during stabilization aging at 39°C. This can be viewed as if the MS temperature of the remaining austenite has been depressed by strain aging and further transformation can only occur at a lower temperature than the original arrest temperature. It has been shown in section 5—6 that the martensite formed, in the sample deformed after 30% martensite, is smaller and deformed after 30% martensite, is smaller and thinner than the previous martensite plates. This result can now be explained in terms of kinetic barrier for interphase motion as discussed above. One further evidence that C diffusion is responsible for the various stabilization effects, is the dependence of 68 stabilization on the aging time. In deformed and partially transformed specimens, the extent of stabilization varies with time of aging at 39‘b increasing to a maximum value and then decreasing at longer aging times. This result confirms earlier 28:29 on plain C—steels and Fe-Ni alloys without any work deformation. The degree of stabilization at relatively longer time of aging has been considered analygous to overaging in solid state precipitation. 69 7. CONCLUSIONS (1) Plastic deformation prior to transformation of austenite causes depression of MS-point in a Fe l6Ni 10 steel. (2) Depression of MS-point caused by prior plastic deformation increases with the extent of deformation. The increase in the depression of MS-point, however, is compara— tively small with higher deformation in the range of deform- ation studied. (3) Depression.ofMS-point is found to be sensitive to rate of straining. Slower strain rate for a constant deformation produces more depression of MS. (A) Martensite formation is found to start always with burst of a reasonable size. The size of burst following prior deformation is larger with higher extent of deformation and consequently lower temperature of martensite formation. (5) Number of bursts in undeformed specimens is observed to be more than those for deformed specimens. For the latter a single large burst is the most common mode of transformation. (6) In deformed specimens, after forming 50% prior martensite, the degree of stabilization,0 , decreases consis- tently with the amount of prestrain up to about 2% and then slightly increases with further strain up to the upper limit of strain employed in the present study ('~A%)- 70 (7) In deformed specimens, after forming about 35% martensite, increases for small prestrain (0.5%) but decreases monotonically (beyond 0.5% strain) up to A% strain. (8) When a partially transformed sample ( ~30% martensite) is deformed the degree of stabilization increases linearly with increasing strain up to 1%. Unfortunately such an experiment could not be conducted at larger strains because of the brittle failure of samples beyond 1% strain. (9) In deformed and partially transformed specimens, the extent of stabilization varies with time of isohermal aging at 39°C, increasing to a maximum value and then decreasing at longer aging time. 71 APPENDIX Procedure for obtaining a specific amount of martensite during resistance measurement. 1) Initially resistance data from RT to liquid N2 temperature is obtained starting from austenitic condition. After maximum transformation (cooled to liquid N2) change of resistance on heating is also to be noted. On the basis of above the following are calculated: Average resistance drop per l‘b in austenite = R7 Average resistance drop per l‘b in martensite = Ra Resistance drop for 1% martensite formation = RM Maximum amount of martensitic is found (based on X-ray diffraction) as 80%. Then, R = EMS - R195 - Ra (Ms - 195) M 80 where EMS is the resistance at MS and R195 at -l95‘b. For avoiding effect of size difference in specimens all the values are divided by RMS and the ratio Ry/RMS, Ra/RMS and RM/RMS obtained. It is expected that these ratios would not be altered much from specimen to specimen. 2) For stabilization experiments the steps followed are: (A) Resistance is measured up to the transformation point and average value of R7 is determined. 72 (B) Taking the resistances before and after transformation as RB and RAl and corresponding temperature as TB and TAl Resistance at Ms, RMS may be calculated as RMS = RB - R7(TB - MS) U3)% martensiterat TAl is then calculated as follows: 1 — RAl/RMS - Ry/RMS (MS - T11) RM/RMS % martensite = (D) Amount of transformation on following lower temperature (TA2) is then calculated as % martensite at TA2 = X2 X1 X1 RAl/RMS - RAP/HMS - [Ry/RMS ( 100 ) + Ra/RMS (l - —ffififl(TA2 - TAl) RM/RMS +Xl (E) Amount of transformation on successive lowering of temper- ature is calculated accordingly. 10. 11. 12. 13. 1A. 15. 16. 17. l8. 19. 73 REFERENCES .F. Bunshah and R.F. Mehl, Trans. AIME, 1251 (1953). Mukherjee, Trans, AIME, 232, 1A95 (1968). Kaufman and M. 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