PHASE TRANSFORMATIONS IN STEEL DURING RAPID HEAT TREATMENT

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1 PHASE TRANSFORMATONS N STEEL DURNG RAPD HEAT TREATMENT 1. PROHASZKA nstitute of Mehanial Tehnology and Materials Siene, Tehnial University, H-1521 Budapest Reeived April 19, 1987 Abstrat Phase transformations in low-arbon steels during rapid heat treatment proeed very differently than under equilibrium onditions. The differenes an suessfully be utilized for improving the properties of the steels. ntrodution During the past deades many efforts were made to inrease produtivity of tehnologial operations, i.e. amount of produt per unit time, simultaneously maintaining or improving the quality of the produt. This eonomitehnologial demand led to important results in the manufature of metals and alloys, inluding aeleration of heat treatment proesses: instead of traditional indiret heat transfer in furnaes, the metal is diretly heated by the Joule heat of the eletri urrent passed through it. This tehnique allows to raise the heating rate from traditional K/min by several orders of magnitude, to K/s [1, 2J, thus inreasing produtivity of the heat treatment operation in a similar extent. The substantial inrease of the heating rate largely hanges the harater of the strutural transformations taking plae in metals and alloys during heat treatment. This paper deals with the strutural hanges taking plae in unalloyed arbon steels during rapid heating. First of all, it should be onsidered that at suh high heating rates, diffusion proesses, i.e. plae hanges of atoms owing to strutural hanges will onern pratially only atoms interstitially dissolved in the phases partiipating in the transformations. By way of example, let us mention that in dual-phase steels - even though their heat treatment time is muh longer, in the order of minutes, than that of rapidly heated steels - phase hange onditions differ so greatly from equilibrium onditions that the paraequilibrium onept has been introdued, indiating that when dual-phase steels are being heated in the y region to the so-alled interritial temperature in order to attain the desired a/y proportion, instead of aiming at thermo- 1*

2 88 J. PROHASZKA dynamial equilibrium, the aim is to reah the state where the appropriate a/"; phase proportion is realized and the arbon ontent in the phase reahes the value orresponding to the given onditions. During this time substitutionally dissolved atoms pratially do not hange their plaes, sine their diffusion rate is lower by several orders of magnitude than that of interstitially dissolved atoms. n Table 1 the diffusion oeffiients of some interstitially and Table 1 Coeffiient D=Do(exp(-Q/RT)) Phase Element Q Do D723'C D0000C D: D " D: D " [kal/mol] [m2js] [m2/s] [m2/s] 723 'C 1000 C o:fe C o:fe Fe ' o:fe Co ' ' 10 6 o:fe Cr 5:' ' o:fe Mn ' 10 5!XFe Ni ' o:fe V ' ' 10 5 o:fe W ' ' 10 6 i'fe C Fe Fe ' ' ' ' 10 5 i'fe Co ' ' ' 10" 2.7' 10 6 {Fe Cr ' , Fe Mn ' ' i'fe Ni ' ' Fe V ' ' 10'13 2.6' 10" 3.5' 10 6 i'fe W ' 10- '" 4.45' ' substitutionally dissolved atoms at 723 QC and 1000 QC in the ferriti phase (ex) and in the austeniti phase (y) are listed [3]. These values an also be interpreted as the time required for the atom in question to over unit path length in the steel. This is one of the reasons why the harater of strutural hange will be different in rapid heating. Some remarks should be made regarding Table 1. First, it should be noted that the diffusion oeffiient of arbon is higher by 5-7 orders in both ferrite and austenite than that of substitutional atoms in the same phase. Seond, it should be remarked that the diffusion oeffiients of both interstitial and substitutional atoms are higher by several orders of magnitude in ferrite than the values or alulated values in austenite. This phenomenon is due to the muh looser struture of ferrite, allowing easier plae hanges for both interstitial and substitutional atoms as well as for same-speies atoms (f, self-diffusion oeffiients). The very low values for substitutional atoms also indiate that even homogenization of several hours will not be suffiient to ahieve uniform distribution of these atoms in the alloy.

3 PHASE TRANSFORJfATOSS N STEEL 89 Figure 1 represents the part important for pratie of the isothermal setion at 740 QC of the state diagram of Fe-Mn-C steels [4, 5]. n the figure, the broken lines indiate the true thermodynami equilibrium for a steel ontaining 0.06% C and 1.5~~ Mn. The figure demonstrates that the equilibrium phase ontains more arbon and manganum than the paraequilibrium phase formed in rapid heating. The same statement is represented in Fig. 2, showing Mn distribution in ferrite and austenite after 1 hr heating to 740 C in the above omposition steel. Aording to Fig. 1, in the equilibrium state, at 0.06% C ontent, austenite will dissolve 3% Mn. This is onfirmed in u 6 Q; t. al <Jl C o g>2 o ::E Equilibrium -- True --- Para O~----'---~-----'~---r~ o Carbor., perent Fig. 1 Fig. 2. Ferrite, however, in the equilibrium state, dissolves around 1 ~~ Mn, whereas Fig. 2 indiates that ferrite, after 1 hr heat treatment at 740"C, still ontains about 1.3~~ Mn at the ferrite-austenite phase boundary. The figures onfirm that heat treatment for 1 hr is insuffiient for rearrangement of substitutional elements. The duration of heat treatment in modern tehnol- C ~ u ex ~ 2 ~ ' <Jl 0 g>1 0 :;;E: lh 740 QC 0, Distane, A Fig. 2

4 90 1. P ROH.4SZKA ogies is in the order of seonds, that is, onentration distribution of substitutional elements will be retained pratially unhanged. However, the fat that the distribution of substitutional elements will remain pratially unhanged will - to a ertain extent - affet the distribution of interstitial elements, too. This is learly shown by Fig. 1, demonstrating that in the given alloy, austenite ontains about 0.5~~ C in the para-equilibrium state, while in the equilibrium state C ontent is 0.43%. Phase transformations are of greater importane than onentration distribution of the omponents. Before disussing these phenomena, let us have a look at the rerystallization proesses in plasti steel deformation. Diffusion will play no important role here, sine the transformation whih as to its harateristis belongs to rerystallization proesses without onentration hanges in the alloy will proeed differently under the effet of rapid heating than in the ase of slow heating. Rerystalliza tion To desribe rerystallization exatly, the frequeny of nuleus formation, i.e. the number of nulei apable of growth formed in unit time and unit volume N N = No exp ( - QN/ R T) (1) (where No and QN are assumed to be onstant values for the given material) and linear growth rate of the nulei G G= Go exp (- QG/RT) (2) (where Go and QG are materialspeifi onstants) should be ombined with the Avrami-Mehl equations. However, the required data No, Go, QN' QG are not at disposal for the majority of metals and alloys used in pratie. Therefore, the relationship for rerystallization rate whih has proved a very good approah for pratie will be made use of: dv v = Tt = A exp ( - Q/ R T) (3) where A is a onstant harateristi for the material and its deformation, Q is well approahed by the self-diffusion oeffiient of the material and dv is the volume transformed during the interval dt. t follows from Eq. (3) that if a given material rerystallizes at the temperature Tl during the time t l, a time interval t will be required for rerystallization at the temperature T, i.e. (4)

5 PHASE TRANSFORMATONS N STEEL 91 Making use of Eq. (4), we listed, in Table 2, the time intervals required for rerystallization of three metals: A, Cu and Fe, assuming that at the homologous temperature 0.5, rerystallization time is 0.5 hr. The data indiate that time requirement for Al rerystallization is redued to 1.5 s by raising rerystallization temperature by 110 QC. To attain similar redution in rerystallization time for opper, a temperature inrease of 160 QC is neessary. The temperature inrease neessary for steel does not exeed 220 QC. However, suh an inrease will raise the temperature of steel above the temperature A: it is more than 100 QC higher than the temperature of perlite-austenite transformation. None the less, one may disregard this transformation, sine - as will be shown in the following - the transformation annot proeed at the given temperature in suh a short time. Drastially redued rerystallization time results not alone in important time saving of the heat treatment proess, but also in average rystallite size far below the value attained with traditional heat treatment methods. For instane, a value of 8 to 9 by the ASTM standard is onsidered a good result for steels rerystallized in furnaes. With rapid heating, during frations of a seond, an ASTM value of an readily be obtained. Table 2 T Al Cu Fe ym.p. T[K] T' [C] t [s] T[K] T' [C] t [s] T[K] T' [OC] t [s] S The above data are equivalent to average rystallite sizes of mm for traditional heat treatment and mm for rapid heat treatment. Aording to the Hall-Peth equation, strength values of the steel will inrease importantly, for the above numerial values by at least 50%, as a result of smaller rystallites. The reason why suh fine rystallites are formed in rapid heat treatment is that despite very high rystallization rate, the frequeny of nuleus formation also greatly inreases, and the short meeting intervals will hinder growth. Owing to the short time there is no opportunity either for the rystallites to grow at the expense of one another.

6 92 J. PROHASZKA Phase transformations in low-arbon steels during rapid heat treatment During rapid heat treatment phase transformations will differ very muh from equilibrium, sine there is not enough time at disposal to proeed in the manner as austomed in traditional heat treatment. For simpliity's sake let us follow the hanges of two ferrite rystallites and a ementite rystallite in between (Fig. 3), sine phase transformation proeeds similarly in the whole bulk. Above the x-x line, perpendiularly to it we plotted the C onentration of the phases. The orresponding part of the Fe-FeC state diagram is shown in Fig. 4. ex Fe 3 C ex x -~ = f-- x.e 2 u g U Fig. 3 0 (;;; " Q. E '/ )[ T2 G '-:-:--:~ f._:_::_----- T, E - P S J,Q Fe Fig. 4 --

7 PHASE TRAVSFORMATlOSS N STEEL 93 During rapid heat treatment heating rate is suh that the speimen reahes the temperature of 1000 GC in 1 s. Let the temperature-time funtion of the heating proess be T= K,/t. The speimen will then reah the eutetoid temperature in 0.88 s. During this time the amount of arbon defined by the P-Q line should go into solution from the tertiary ementite. Owing to the short time at disposal, this annot take plae, and if the steel was in equilibrium at the start of heating, only the arbon ontent of the ferrite layers in diret viinity to the ementite rystallites will grow aording to the P-Q line. but nothing will hange in the interior of the ferrite partile. Equilibrium onditions hange when the temperature A 1 is reahed. When heating rate is traditionally low, austenite will appear. At rapid heating, however, the short time at disposal does not allow it. Both ferrite and ementite are stable, ferrite, depending on its C ontent, up to 910 C, and ementite up to even higher temperatures. n the momentary state only the onditions of austenite newly formed must be onsidered. n this situation phase hange is aused by the fat that ferrite and ementite annot o-exist at this temperature under equilibrium onditions. Separately, however, both of them would be in equilibrium. Free enthalpy is high only on the ferrite-ementite phase boundaries; suh phase boundaries should disappear and be replaed by ementite-austenite and austenite-ferrite phase boundaries. However, the preondition for suh boundaries is that beside ementite an austenite nuleus to be formed and grow. For suh a nuleus to be formed it is neessary that in the ementite phase and in the ferrite phase, in a region orresponding to the ritial nuleus size, arbon onentration of 0.8~;'; should be established at 723 C. This needs time, sine the C onentration of the existing two phases is rather far from it, one being 6.67/~, the other 0/ 0, f the total time of heating up to 1000 C is only 1 s, it is pratially impossible for austenite to be formed during this interval. t beomes possible only when temperature will exeed the intersetion of the GOS line and the vertial line haraterizing the C ontent of ferrite in the steel. n this moment the free enthalpy state of the system will hange to a state that will fore this transformation. Up to reahing the GOS line, the two phases beside one another, ferrite and ementite, are to be onsidered stable, owing to what has been said. Only the phase boundary layer is labile, having a free enthalpy value muh higher than the equilibrium value. The differene in free enthalpy will further inrease with temperature and will reah a fairly high value at 900 C. The differene will derease only above 900 C, where austenite an be formed at arbon onentrations as low as 0.1 %. t must be underlined, however, that austenite with suh low C onentrations an be in equilibrium with ferrite only; the phase boundary ferrite-austenite orresponds to equilibrium onditions. n ontrast, the phase boundary ementite-austenite an, at 900 GC, be in equilibrium only with austenite ontaining lose to 1.6%

8 94 J. PROHAsZKA arbon. (This is why the temperature A t an be disregarded in rerystallization at rapid heat treatment.) f austenite is formed in the rapidly heated steel in the temperature interval At -A3' in the equilibrium state its arbon onentration on the phase boundary with ferrite will be given by the setion of the GOS line, while its arbon onentration on its boundary with ementite will be given by the setion of the SE line valid for the given temperature. These onditions and the ritial nuleus size will define the formation of austenite in that interval. Fig. 5 shows arbon distribution at the temperatures Tt and T 2, resp., along the x-x line for the ase where austenite rystallites are formed between ferrite and ementite..--~, %.'2 e '" u o u T, Fig. 5 As soon as the temperature reahes the intersetion of the GOS line and the vertial line orresponding to the arbon ontent of ferrite, a ritial thermodynami situation fully novel in its harater and related to the total ferrite amount will arise. Austenite an be formed at any point of the ferrite phase, if the ritial nuleus size is attained. t is not limited by any onentration requirement, sine the C ontent of the ferrite and the austenite being formed is idential. From this moment on, transformation proeeds rapidly; it is further promoted by the fat that with inreasing temperature ritial nuleus size dereases, and ferrite must disappear very swiftly from the whole struture. n this stage of rapid heating we have arrived to a point where the struture onsists of austenite and of the remaining ementite rystallites of the perlite. n this moment, the arbon ontent of the austenite is still largely inhomogenous. Maximum arbon ontent will be at the ementite boundaries orresponding to the C~r value in Fig. 5. Minimum C ontent, i.e. pratially zero, will be where the ferrite-austenite transformation has only just taken

PHASE TRANSFORMATONS N STEEL 95 plae. The mirophotograph of an un alloyed arbon steel with 0.08 C ontent tempered in water from the transition stage is represented in Fig. 6.

9 PHASE TRANSFORMATONS N STEEL 95 plae. The mirophotograph of an un alloyed arbon steel with 0.08 C ontent tempered in water from the transition stage is represented in Fig. 6. The struture is very similar to that of dual-phase steels. n austenite the diffusion of C atoms proeeds very rapidly and the arbon onentration graduailz equalizes. n the meanwhile, however, the boundaries of austenite in ontat with ementite will progressively be extended in the diretion of ementite, onsuming the ementite amount. This proess will ontinue till the size of Fig. 6 the ementite rystallites does not finally reah the ritial nuleus size for the given temperature. This is again a ritial point of the transformation, sine when this size or a smaller size is attained, the ementite rystallites must immediately disappear. Their disappearane an only proeed by their transformation into austenite or eventually into some other phase with a smaller rx nuleus size. This latter possibility is, however, very improbable. Austenite formed in this manner would, at the plae of ritial nuleus size, have a arbon onentration of 6.67%, whih - evidently - annot be kept up. For the same reasons whih led to the arbon onentration orresponding to ementite in austenite during ooling within the ritial nuleus size under the effet of thermal flutuation, now this small region must disappear very rapidly. Under suh onditions, the arbon ontent will equalize rapidly. n Fig. 7, a steel with the above-mentioned C ontent, heated for 5 s only and tempered in water is shown. These phase transformations differing from

10 ..-- -_.-._ _.._-_.._----.._ _ J. PROH.4SZKA Fig. 7 the traditional an readily be utilized for hanging mehanial properties [2]. The mirostruture shown in Fig. 6 (also obtained after 5 s heating and tempering) orresponds to 700 MPa tensile strength and A 5 = 15~~ elongation, whereas the mirostruture represented in Fig. 7 orresponds to 1000 MPa tensile strength and an elongation of 7%. Referenes 1. PROHASZKA.1.: H. nternational Congress on Heat Treatment of Materials offht. Florene. taly, 20-24, Sept p PROHASZKA, J.: Neue Hiitte, 30 (1985) WEAST, R. C. and M. 1. ASTLE: Handbook of Chemistry and Physis. C.R.C. Press, Florida ( ). 4. RV~, K. 1.: Strong Tough Strutural Steels. Spe. Rep. The ron and Steel nstitute, London (1967). 5. HLLERT, M. and W ALDENSTROM, M.: C alphad, 1. (1977) p n: Fundamentals of Dual Phase Steels. Ed. by R. A. Kot and B. L. Bramfit. Conf. Pro. The Met. So. of AME Prof. Or. Janos PROHASZKA H-1521 Budapest

ESTIMATION OF FRACTURE TOUGHNESS OF THE OXIDE SCALES. Jozef HRABOVSKÝ, Petr LOŠÁK, Jaroslav HORSKÝ

ESTIMATION OF FRACTURE TOUGHNESS OF THE OXIDE SCALES Jozef HRABOVSKÝ, Petr LOŠÁK, Jaroslav HORSKÝ Heat Transfer and Fluid Flow Laboratory, Faulty of Mehanial Engineering, VUT Brno, Tehniká 2, 616 69 Brno,