Metallurgical Phenomena during Processing of Cold Rolled TRIP Steel

Materials Technology

[17]C. Donadille, R. Valle, P. Dervin and R. Penelle: Acta Metall., 37(1989), 1547.

[18]Y. Maehara, Y. Ohmori, J. Murayama, N. Fujino, T. Kunitake: Met. Sci,. 17 (1983), 541.

[19]J.A. Jiménez, M. Carsi, F. Peñalba, O.A. Ruano: J. Mater. Sci., 35 (2000), No. 4, 907.

[20]A.F. Padilha, P.R. Rios: ISIJ Int., 42 (2002), 325. [21]J.W. Simmons: Scripta Metall. Mater., 32 (1995), No. 2, 265. [22]Z.Z. Yuan, Q.X. Dai, X.N. Cheng, K.M. Chen: Mater. Charact., 58

(2007), 87. [23]A.L. Schaeffler: Metal Progr., 56 (1949), No. 11, 680. [24]C. J. Long, W. T. DeLong: Weld J., 52 (1973), 281.

[25]F.C. Hull: Welding J. 52 (1973), 193. [26]M. Onozuka, T. Saida, S. Hirai, M. Kusuhashi, I. Sato, T.

Hatakeyama: J. Nucl. Mat., 255 (1998), 128. [27]H. Takahashi, Y. Shindo, H. Kinoshita, T. Shibayama, S. Ishiyama,

K. Fukaya, M. Eto, M. Kusuhashi, T. Hatakeyama, I. Sato: J. Nucl. Mater., 258-263 (1998), 1644.

[28]S. Allain, J.-P. Chateau , O. Bouaziz , S. Migot , N. Guelton: Mater. Sci. Eng.-A, 387–389 (2004), 158.

[29]R. E. Smallman, D. Green: Acta Metall., 12 (1964), 145. [30]S G Chowdhury, S Das, P.K. De: Acta Mater., 53 (2005), 3951. [31]I.J. Dillamore, W.T. Robert: Acta Metall., 12 (1964), 281.

Metallurgical Phenomena during Processing of Cold Rolled TRIP Steel

M. Soliman, B. Weidenfeller and H. Palkowski

Institute of Metallurgy, Clausthal University of Technology, Robert-Koch-Str. 42, 38678 Clausthal-Zellerfeld, Germany E-mail: [email protected]

Metallurgical phenomena taking place during processing of TRIP Steel are investigated and described with the aim of achieving better understanding of the microstructure development throughout the entire integrated processing routes. Different TRIP steel structure sizes were created by controlling the hot rolling process prior to cold rolling. After that the specimens were intercritically annealed under different conditions to obtain prescribed austenite fractions, and subsequently quenched in salt bath at the bainite transformation temperature. The microstructures had been investigated using light optical microscopy (LOM) and the amount of retained austenite was determined bymagnetometry.

Keywords: TRIP-aided steel; non-recrystallization temperature; magnetometry; recrystallisation-precipitation kinetics.

DOI: 10.2374/SRI08SP104; submitted on 5 June 2008, accepted on 11 August 2008

Introduction

In early 1965 Deliry reported that silicon strongly retards the precipitation of cementite during the bainite reaction [1]. The absence of cementite ensures that the carbon will enrich the austenite rather than forming cementite plates. Enriching the austenite with carbon results in lowering its martensite start temperature (MS).Therefore, after the course of bainite transformation, by further cooling to room temperature (RT), a certain amount of austenite can be stabilized. This austenite is known as retained austenite.

More than 20 years later, Matsumura et al. [2] reported the improvement of the mechanical properties of steels due to the transformation of the retained austenite into martensite during deformation (TRIP effect). This was the first report concerning the low alloyed TRIP-aided steels. In these steels, the mechanically induced martensite transformation of the metastable austenite results in geometrical changes at the microscopic level; these geometrical changes have the effect of strain hardening of the surrounding ferrite matrix. Consequently, the failure due to necking is shifted to higher values of stress and

strain. In the past decade there has been a resurgence of interest in these steels, especially for use in the automotive industry [3, 4].

More recently, silicon is completely [5, 6] or partially [7, 8, 9] replaced by elements that play its role. This is because high Si levels that are needed to prevent cementite precipitation do not fit the industrial practice of flat products concerning the surface quality [10, 11]. Al and P are well adapted elements and currently being widely used.

To exploit the full potential of these materials by optimisation of alloy and process design, a clear under-standing of their behaviour along the processing route of reheating, rolling, annealing, cooling and the bainitic transformation is necessary. Many metallurgical phenomena take place during these processing routes, such as recrystallisation, pancaking, carbides dissolution, pre-cipitation, austenite formation and austenite decomposition. Knowledge of the kinetics of these metallurgical phenomena allows control of the microstructure by controlling the processing conditions. This is of prime interest because the properties of steels are, to a great extent, determined by their microstructure which is a result of various phase transformations occurring during manu-

steel research int. 80 (2009) No. 1 57


facturing. This work helps to achieve better understanding of the microstructure development of the TRIP-aided steel throughout the entire integrated processing routes aimed at further development of these steels in terms of strength and ductility.

Experimental Procedure

Investigated material. The Si-Al-Mo-Nb steel studied in this work was melted in a high frequency furnace in the laboratory. The chemical composition is given in Table 1.The selection of the composition was based upon current researches and understanding in the field. Nb is the most efficient micro-alloying element in TRIP steel. Nb increases tensile strength due to grain refinement [12, 13] and fine precipitates [13,14,15]. Furthermore, it improves the strength-ductility balance due to the increased volume fraction of retained austenite [16, 17]. The complex addition of Nb and Mo results in further strengthening of the steel [18, 19].

Rolling Process. The material was subjected to several hot rolling conditions designed to generate different structure sizes as will be described later. The surface oxide scale was then removed from the slabs using shot-blasting, and finally the 4 mm thick hot rolled plates were cold rolled to a thickness of 2.5 mm.

Heat treatments. The heat treatment cycles shown in Figure 1 were conducted in two salt baths, one for annealing followed by one for austempering. The intercritical annealing temperatures were chosen on the basis of the dilatometric measurements as will be described later. The annealing had been performed in Durferrit GS 540/R2® and austempering in Durferrit AS 140® salt baths. The inertor R2 had been added to prevent any oxidation or decarburisation during the austenitisation process.

Metallographic study. For the metallographic study, heat treated specimens were sectioned in rolling direction. Micrographs had been taken perpendicular to transverse-rolling direction. The specimens were prepared by mechanical grinding followed by polishing up to a 0.25 µm-grade alumina solution. In order to reveal phases and grain boundaries, LePera’s etchant was used, which is a 1:1 mixture of 1% sodiummetabisulfite in distilled water and 4% picric acid in ethyl alcohol [20] with an etching time between 20 and 30s. LePera’s etchant stains ferrite brown and/or blue, bainite dark-brown to grey, while martensite and austenite remain white.

Magnetic measurements. The magnetic measurement technique, using a hysteresis recorder (Figure 2), was used to measure the amount of the retained austenite (V ). As compared to x-ray diffraction (XRD), the advantages of the magnetic technique are that magnetic measurement is performed on the whole volume, the specimens require no special preparations, and the measurements are fast, reliable and more sensitive to retained austenite [21, 22].The measurements had been done by exposing the

Figure 1. Schematic of the thermal cycle used in the currentinvestigation.

Hysteresis recorder

Figure 2. Schematic of the hysteresis recorder.

Table 1. Chemical composition of the steel studied, in wt.%.

C Si M S n Al Mo Nb P N

0.25 0 1 2 .86 .37 0.64 0.38 0.04 0.03 0.0 0.008

Table 2. Measured and calculated saturation polarisation Jsalloy.

Measured Theoretical [23] Theoretical [24]

J alloy 1.948 1. 881s in T 984 1.



specimens to an alternating magnetic field of magnetising frequency fmag=1 Hz. The generated magnetic flux density includes not only magnetic properties of the sample but also magnetic properties of the air. To eliminate this part of the induction an identical coil system is connected in series (compensation coils). The induction coils are connected anti-parallel to get a resulting signal which only delivers the change of the magnetisation in the sample and therefore, the polarisation J of sample is directly measured. A maximum magnetic field of approximately Hmax = 45 kA/m was applied to the sample, the saturation polarisation, Js, was determined using the equation

)1( 2Hb

Ha

JJ s (1)

under neglecting higher order powers. The amount of retained austenite is calculated by

comparing the saturation polarisation of the specimen with retained austenite (Js

A) with that of an austenite-free specimen containing the same amount of alloying elements (Js

ref ), through the equation

refs

As

refs

JJJ

V (2)

The austenite-free specimen had been prepared by heating the specimens to a temperature of 950°C, holding there for 20 min and followed by very slow cooling, by switching off the furnace, down to RT. The cooling rate did not exceed 1.5 K/min. The measured Js

ref was 1.948 T. The Js

ref can also be theoretically calculated using a linear function of the weight percent alloying additions, rather than measuring saturation on a reference sample using the equation

Jsref = Js

Fe – xn Jselement (3)

JsFe is the saturation polarisation of pure iron Js

Fe = 2.158 T [24]. xn is the amount of alloying element and Jn

element is the decrease of polarisation of 1% of a certain alloying element. The values of Jn

element are given elsewhere [23, 24]. Table 2 compares the measured and calculated values.

Results

Hot rolling. It order to generate different structure sizes before the heat-treatment process, different hot rolling schedules were applied. These schedules were selected according to the non-recrystallization temperature TnRX. By deforming the material in the recrystallization region the austenite grains are being refined. This is important because the grain size of the austenite strongly affects both, the kinetic of subsequent transformation and the ferrite grain size, namely smaller austenite grains consequently lead to the refinement of ferrite grains. When deformations are applied at temperatures below the temperature of non-recrystallization TnRX the austenite grains become elongated and deformation bands are

pancaking. As the amount of deformation in this region increases, the number of nucleation sites at the austenite grain boundaries and within austenite grains increase. Because of that, transformation from deformed austenite yields in m h finer ferrite grains than that from recrystallized, strain-free austenite. Therefore, TnRX is a very important parameter and its determination represents a crucial step in designing rolling schedules. Accordingly, the present hot rolling schedules had been selected according to this temperature.

TnRX can be determined by

introduced within the grains. The process is called

the method proposed by Jo

uc

nas and co-workers [25, 26] which is based on multi-stage torsion test. In the present study this method is used, but the deformation was conducted by means of multistage compression test. Since the maximal strain of Bähr 805D deformation dilatometer is 1.2, six strain steps were selected each one of them = 0.2 with a strain rate

. =

0.6 s 1. The cycle included a prior heating to 1250°C and holding for 300 s. The stress-strain curves obtained are shown in Figure 3a and indicate that the level of stress depends on deformation temperature. Actually, wider temperature range can be covered by applying multistage compression tests on sequences of temperature ranges. Applying the same strain rate

. at each deformation-step

is vital because any deviation rom this value affects the stress-strain curves and results in a misleading result. Figure 3a shows that the stress increases as temperature

f

0

50

100

150

200

250

300

350

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 1.2

True strain

True

stre

ss (M

Pa)

1050°C 1000°C 960°C 915°C 870°C 825°C

(a)

120

140

160

180

200

220

240

260

280

300

0.75 0.77 0.79 0.81 0.83 0.85 0.87 0.89 0.91 0.93

1000/T (K-1)

MFS

(MP

a)

T nRX

(b)

Figure 3. (a) The true strain-true stress curves obtained from thesix-stage compression test in a temperature sequence denoted inthe figure. (b) Dependence of the MFS on the inverse absolutetemperature.



decreases, but this increment is higher at the last two deformation steps. On the other hand, there is higher work hardening rate observed on these two curves.

The procedure of determination of TnRX consists of ca

of TnRX can be determined by fi

stimated temperature, the deformation part fo

Intercritical annealing. In order to define the inter-cr

ng temperature was chosen on th

icrostructure characteristics. The different hot-ro

tionco

lculating the mean flow stress (MFS) that corresponds to each deformation step. MFS is the area under the given stress-strain curve for selected interval divided by pass strain. The mean flow stresses for all the deformation steps had been calculated by numerical integration and the results are plotted as a function of inverse absolute temperature on Figure 3b.

From this graph the valuending the intersection between the regression lines of the

points that corresponds to each part of the curve with two different slopes. The estimated value of TnRX from Figure 3b is 885°C.

With this er the schedules was determined in such a way that all the

three possibilities were covered, namely all deformations conducted above TnRX, deformations below TnRX and deformations mixtures of below and above TnRX. The cast ingot was reheated to 1250°C to dissolve the NbCN before hot rolled in four passes from a thickness of 19 mm to 4 mm, and with a true strain value = 0.38 at each pass before cooling in air. The different schedules resulted in microstructures shown in Figure 4. The microstructure formed from the recrystallized-austenite is denoted by “L”, whereas the microstructure formed from pancaked austenite is denoted by “S”. The “M” microstructure results from the recrystallized and then pancaked austenite. The structure is of quite duplex nature consisting of both polygonal ferrite and acicular ferrite. Pronounced variation in structure size with varying the hot-rolling schedule can be seen. The polygonal ferrite grain sizes have average values of 8.5, 6.9 and 4.5 m for the L, M and S microstructures, respectively.

itical region, dilatometric measurements had been applied by heating specimens up to 1100°C. For better tracing the equilibrium points, a heating rate of 0.05 K/s had been used. The variation of the relative change in length as a function of temperature was measured from

which the transformed austenite fraction (f ) was calculated employing the lever rule [27, 28]. The result is shown in Figure 5. In this figure a clear demarcation of two slopes can be seen. The point where the slop is changed is marked by TC.

The intercritical annealie basis of the dilatometric measurement. A phase content

of 50 % and 30% polygonal ferrite (PF) was required at the end of intercritical annealing. Figure 5 shows that this phase distribution is obtained at intercritical annealing temperatures (TA) of 810°C and 850°C, respectively.

Mlling schedules resulted in pronounced differences in the

hot-rolled structure-size (Figure 4). Due to the latter variation different final TRIP-aided steel structure size is observed as shown in Figure 6. Making use of the beneficial effect of the hot rolling below TnRX, a pronounced finer cold-rolled TRIP-aided steel structure was produced. On the other hand, the observed effect of the prior hot-rolling schedule on V was very limited.

For the “M” structure size, microstructural evolurresponding to PF content of 30% was studied along

isothermal holding at 400°C. Depending on the bainitic holding time, various mixtures of bainite, martensite and retained austenite can be found in the microstructure after quenching. Figure 7 illustrates the different micro-structures with micrographs corresponding to bainite holdings of 60, 120, 300 and 480s. The LePera etchant revealed the ferrite blue or light brown bainite brown while the martensite and the retained austenite (MA) remained white. During the bainite holding the intercritical austenite progressively transforms to bainite, bringing about a larger stabilization of austenite. According to the stability of the austenite, a certain part retains at RT, whereas the remaining part transforms to martensite. As both, retained austenite and martensite appears white colour, the white areas of Figure 7 represent the non-transformed austenite during the bainite holding stage. Accordingly, the gradual decrease in the white area of the figures corresponds to the progress of the bainite-reaction. In addition to the metallographic investigations, the retained austenite volume percentage (V ) was measured

L

SRT: 1150°C, FRT: 910°C

M

SRT: 1000°C, FRT: 780°C

S

S RT: 850°C, FRT: 710°C

Figure 4. Different structure sizes (large “L”, medium “M” and small “S”) resulted from applying different rolling schedules (etchant: nital).Micrographs taken at the middle section of the hot-rolled material. SRT: start rolling temperature; FRT: finish rolling temperature.



0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1 1.40

-2

-1.5

-1

-0.5

0

0.5

1

1.5

2

-50 -30 -10 10 30 50

Magnetic field H (A/m)

Pol

aris

atio

n J

(T)

V = 0%

V = 18.4%

700 730 760 790 820 850 880 910 940 970 1000Temperature (°C)

Aus

teni

te fr

actio

n (f g

)

0.70

0.80

0.90

1.00

1.10

1.20

1.30

Rel

ativ

e ch

ange

in le

ngth

(%)f

Relative change in length (%)

TC

Figure 5. Dilatation versus temperature curve observed during continuous heating together with the calculated formed austenite fraction (f ) employing the lever role.

Figure 8. The change in hysteresis curve due to the variation ofthe austenite content from 0% to 18.4%

(a) (b) (c)

(d) (e) (f)

Figure 6. Different TRIP steel microstructures resulted from different prior structure sizes of Figure 4: (a) and (b) from the L-size , (c) and (d) fromthe M-size and (e) and (f) from the S-size. Material was intercritically annealed at 810°C for 480s and isothermally transformed at 400°C for 480s.

(a) (c)

(b) (d)

Figure 7. Microstructure observed after annealing at 850°C for 480 s and isothermal transformation at 400°C for (a) 60 s, (b) 120s, (c) 300s, (d) 480s.



using the magnetic method by comparing the hysteresis curves of the samples after heat-treatment with that of the austenite free sample. Figure 8 shows a representative example comparing the hysteresis curve of austenite free sample with one having 18.4% retained austenite. Figure 9shows V as a function of the austempering time for the two steels. This figure shows that V decreases by increasing the austempering time. It is also noted that the dependence of V on austempering time decreases by increasing the transformation-time.

The influence of the intercritical annealing temperature on the microstructure is given in Figure 10. Increasing the

oncurrent increase in the ount of MA phase was observed. On the other hand, the

intercritical annealing temperature had no significant effect on the grain size of ferrite and retained austenite. A similar observation was reported by Shi et al. [29]. The grain sizes of ferrite and retained austenite varied between 3-7 m and 0.7 - 3 m, respectively. This study detected no significant influence of austempering temperature on the location and morphology of the retained austenite.

Additionally, Figure 10 shows that the structure is fully recrystallized. Consistent with this result, Petrov et al. [30] observed no interaction between recrystallization and

transformation phenomena because the static recrystal-lization was already completely finished before the start of the phase transformation.

Discussion

Effect of aluminium in TRIP steel. In order to under-stand the aluminium effects in TRIP steel, it is instructive to consider the Fe-Al phase diagram. With the com-position shown in Table 1, the equilibrium phase diagram shown in Figure 11 was calculated using Thermo-Calc©.

he thermodynamic calculations were performed using the

In this diagram the intercritical region lies between curve “1” (Ae1) and curve 3 (Ae3). The cementite is com-pletely dissolved at curve “2”.

During continuous heating of steel, the transformation takes place initially by the growth of austenite into the carbide-rich areas [31, 32]. This stage is representative in the equilibrium diagram between curve “1” and curve “2”. When the cementite is completely exhausted (at curve “2”), another mechanism of transformation takes place. In this mechanism th austenite grows into ferrite and simul-taneously the carbon is redistribute tween the former

and the latter phases [31, 32]. This stage of transformation corresponds to the region between curve “2” and the lower part of curve “3” in the equilibrium diagram.

A clear demarcation of these two transformation mechanisms is also observed in the dilatometric result (Figure 5). The change in slope of the line represents how f depends on the annealing temper e (T ). It corres-

t is marked in the measured kinetics (Figure 5) by TC. This corresponds in the equilibrium diagram to the point of intersection between the line represent-ing the studied alloy (dashed line) with curve “2”. Table 3 compares the pre-dicted critical temperatures with those measured by dilatometry for the invest-igated alloy.

It is noted from Figure 11 that increasing the Al-content widens the intercritical region, which renders the inter-critical annealing more control-able. However, increasing the Al-content beyond 1.2 wt% restricts the formation of -iron and causing the -area of the diagram to contract to a small area referred to as the gamma loop. This means that the Al motivates the formation of ferrite and results

intercritical annealing temperature resulted in decreasing the PF content to reach approximately the expected percentage (30% and 50%). A cam

Tdatabase TCFE4 and considering the ferrite, austenite, AlN and cementite.

ed be

atur Aponds to the change in the austenite formation mechanism [27]. This poin

finally in a continuous ferrite and

0

5

10

nite

vo

15

20

25

30

0 60 120 180 240 300 360 420 480 540Austempering time (s)

Aust

elu

me

perc

enta

ge (V

)

T = 850°CT = 810°C

A

A

Figure 9. Volume percentage of retained austenite as a function of the austemperingtime at 400°C.

Figure 10. Microstructure observed after annealing at 810°C (left) and 850°C (right).Isothermal transformation at 400°C for 480 s.



austenite phase field. The alloy is therefore not amenable to control in the intercritical region. Furthermore, ferrite will remain in the matrix when solution heat treated. This could cause fusion welding problems by forming a ‘‘ferrite necklace’’ along the fusion line [33].

Curves “4” and “5” represent the locations of temperatures where NbCN and AlN, respectively, are completely dissolved in austenite.

In order to achieve the desired metallurgical advantages of Nb, the solubility of the chunky NbCN particles during casting is required. Depending on these calculations, a reheating temperature of 1250°C before hot rolling was selected.

On the other hand, it is noted from Figure 11 that increasing the aluminium content reduces the solubility of AlN in austenite, resulting in higher solution temperature of AlN particles, as is evident from the relationship between solubility product of AlN and temperature [34]. Beyond 0.06 wt% Al, AlN particles are stabilised such that the required reheating te

(i) Positive effect by pinning the austenite grain boundary with AlN particles during reheating and consequently the control of prior austenite grain size. Furthermore, if AlN particles quickly dissolved during reheating, then the Al would play a harmful role as an austenite grain size controller due to the fact that the quickly dissolved AlN particles cause a notable drop in the pinning forces, leading to a considerable expansion of the grain bound-aries where these particles are found [35].

(ii) Negative effect by restricting the precipitation of finer AlN particles during the subsequent processes.

The mutual effects of Nb and Al on controlling the TRIP steel structure need to be experimentally investigated.

Kinetics of precipitation and recrystallisation. he contrast between the precipitation and recrystallisation

ed steels. It can be determined experimental- by torsion testing [25, 26] or by hot compression testing

as suggested here. TnRx can also be predicted on the basis of the relative rates of precipitation and recrystallisation [36, 37]. The equations proposed by Dutta and Sellars are commonly used for this [36]. These equations describe the recrystallisation-precipitation-time-temperature (RPTT) kinetics after a single step deformation. The precipitation kinetics is described by a precipitation-time-temperature (PTT) diagram. The PTT diagram is characterized by a C-shape associated with the mechanisms that are controlled by nucleation at higher temperature nd by diffusion and

mperature to dissolve NbCN is expected to be lower than that required to dissolve AlN. This has two effects:

Tkinetics is the reason for the occurrence of TnRx possessed

y microalloybly

agrowth at lower temperatures [38]. The predicted time required for 5% of precipitation can be calculated by

5.0.

11

RT400000expNb05.0 13t 60

(4)

2s

3

10

)k(lnT105.2

expT000

gas constant, T is the temperature in K f Nb super-saturation, given by:

R270

exp

R is the universal and ks is the rate o

)T77014

N12

(5) 6

26.2(s

C Nbk

10

Figure 11

Table 3.

. F d alloy.

Pre

e-Al equilibrium diagram of the studie

dicted versus measured critical temperatures.

Ae1 (°C) TC (°C) Ae3 (°C)

Predicted 725 745 921

Measured 734 760 932

650

750

850

950

1050

1

Tem

pera

ture

(°C

)

T nRx

10 100Time (s)

tp0.05

tx0.05

tx0.95

Figure 12. Predicted RPTT diagram calculated based on the rolling simulationparameters (Figure 3).



The predicted PTT diagram for the studiedgiven in Figure 12. This time-temperature cunose at about 925°C in the case of NbCN precipitatio

By contrast, recrystallisation-time-temperatucurves do not have any noses; instead, the decreases continuously as the tempera

steels is rve has a

n. re (RTT) start time

ture increases. This fe anism is

egard to to Dutta

[36] the time necessary for 5% recrystallization ca

ature can be attributed to that this mechcontrolled by diffusion mechanism, both with rnucleation as well as to growth [38]. Accordingand Sellars

n be calculated by

05.0xt = 6.75×10-20 420d ×

exp Nb1851075.2exp000 300 5

TRT(6)

en before se p cess. hi th t, we

unction % C, th

nt. Befo quenched

his [41], this 12.7 m.

recrystal-igure 3). f time as

ation was

(7)

perature as shown

id than e plies

fortakes place, the re

li higher

(ii) The high Si-content of the steel, which according to Boratto et al. [39], tends to decrease TnRx.

The temperatures at which the material is deformed with respect to TnRx are decisive for the austenite grain morphology from which the ferrite is formed. Deformation below TnRx produces grains from work hardened (pancaked) austenite through austenite to ferrite trans-formation. The sizes of these ferrite grains are smaller than those produced from the recovered and recrystallized austenite grains (deformed above TnRx). Thus, the ferrite grain size developed after the transformation strongly depends upon the austenite grain structure that appears just before the start of transformation.

Microstructure formation. The current result (as well as pervious report [29]) that the annealing temperature has no significant effect on ferrite and retained austenite grain sizes (Figure 10) infers that, in the current study, the

t only on the hot-rolling conditions.

as twofold effects on its mechanical properties. In addition to the well-known effect of efinement proving ical

operties, t reater stability of the smaller retained stenite grains has a further i ving effect. Brandt et al.

[40] pointed out that smaller retained austenite particles contain less potential nucleation sites for the trans-formation to martensite and consequently require a larger total driving force for nucleation of martensite. Thus, controlling the deformation temperature and the degree of deformation below TnRx during the hot-rolling process has a decisive effect on structure refinement of the final product of the cooled-rolled TRIP-aided steel. The deformation below the TnRX results in distinct refinement of the final TRIP-aided steel microstructure.

The difference between the MA% in the photo-micrographs (white areas in Figure 7) and the V values (Figure 9) represents the martensite phase. Because the austenite is less stabilized at the beginning of the bainite-reaction, more martensite phase is expected which increases the difference between the white areas in the figures and the V . However, proceeding in the bainite-

The latter is gradually stabilised and simultaneously decreases in quantity during the course of bainite-reaction. The former effect results in decreasing the martensite phase and hence decreasing the difference be een the MA% and the V

Conclusions

This work addresses various metallurgical phenomena

ena is vital to reach an optimum structure for specific applica-tions. Practicing controlled rolling process treatment of plates and strips requiphenomena. It is agreed tis insignificantly affected by the intercritical annealing te peratures, but this study found that it is decisively

where d0 is the austenite grain size of the specimthe deformation process. During the deformationd0 decreases; this motivates the recrystallisationRevealing the austenite grain boundaries by etcnot successful most of the time, so instead ofmeasured martensite packet size being a linear fthe prior-austenite grain size [41]. Below 0.43 wtfunction is slightly affected by the carbon contethe deformation step at 870°C the material wasto RT to form martensite. The

quence structure fineness (Figure 6) is dependanronag was This grain refinement in TRIP steel h

of is prre au

measured packet size of tmartensite was 4.7 m. According to Maki et al.martensite packet size corresponds to d0 of about This d0 affects the subsequent kinetics where nolisation occurs (deformation at 870°C, see FAccordingly, the tx0.05 is plotted as a function oshown in Figure 12.

The time necessary for full (95%) recrystallizcalculated by

tx0.95 = 7.66 tx0.05

The outcome of this divergence in temdependence of recrystallisation and precipitationin Figure 12 is that recrystallisation is more rapprecipitation at high temperature while the reversat relatively low temperatures. Thus, after dewhen static recrystallisation

apmation, crystal-

reaction is accompanied by carbon enrichment of the austenite.

sation process precedes precipitation attemperatures, while precipitation occurs first at lower temperatures. Once the latter takes place, it prevents the nucleation and growth of new grains, i.e. prevents the recrystallisation process. According to this approach, TnRx

is defined as the intersection of the curves corresponding to the time necessary for the occurrence of 95% of austenite recrystallisation and for the occurrence 5% of carbonitride precipitation as shown in Figure 12. The recrystallisation is prevented by copious precipitation below TnRx. According to Dutta and Sellars model, the TnRx

occurs at about 927°C. Notice that the predicted TnRx ishigher than the measured one. A possible explanation for this observation is:

(i) The decrease of Nb in solution due to the precipitation that occurs at the former deformation processes. Lower contents of solved Nb decelerate the precipitation process and thus result in lower TnRx.

grain r on im mechanhe g

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like au enite decomposition, recrystallisation, precipita-tion, etc. which occur during processing of cold rolled TRIP-aided steel. Comprehension of these phenom

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res the management of these hat the TRIP steel structure size

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Concepts, Technical Transfer Dispatch #6, ISIJ, Brussels, May 2001.

[4] W. o RIP-Aided High Strength Ferrous Alloy p.13.

[5] . Metens, P. Jacques, Y. Houbaert, B. Verlick: Scr. Mater., 44 (2001), 885.

[6] P.J. Jacques, E. Girault, A. Mertens, B. Verlinden, J. van Humbeek, and F. Delannay: ISIJ Int., 41 (2001), 1068.

Bleck, and K. Hulka: Proc. Int. Conf. on TRIP-Aided High Strength Ferrous Alloys, 2002, Gent, pp. 199-206.

[1

ent Technologies Inc. IEEE

[2

ng, R. Fu, L. Wang, and P. Wollants: Mater. Sci.

[4

fected by the hot rolling parameters. A finer cold rolled TRIP steel structure resulted from controlling the deformation temperature and the degree of deformati

elow non-recrystallisation during hot-rolling, leading to improved strength of the material, too.

References

[1] J. Deliry: Mem. Sci. Rev. Metall., 7-8 (1965), 524 [2] O. Matsumura, Y. Sakuma, H. Takechi: Trans. ISIJ, 27 (1987), No. 7,

570. [3] ULSAB Advance Vehicle

Bleck: Proc. Int. Conf. n Ts, 2002, Gent,

E. Girault, A nden, and J. Van Humbee

[7] M. De Meyer, D. Vanderschueren, and B.C. De Cooman: ISIJ Int., 39 (1999), 813.

[8] W. Bleck, E. Kechagias, J. Ohlert, J. Christen, A. Moulin, H. Hofmann, and T.W. Schaumann: Optimisation of Microstructure in Multiphase Steels Containing Retained Austenite, Final Report, Technical Steel Research Series, Luxembourg, 2004.

[9] S. Traint, A. Pichler, K. Hauzenberger, P. Stiaszny and E. Werner: Proc. Int. Conf. on TRIP-Aided High Strength Ferrous Alloys, 2002, Gent, p.121.

[10] J. Maki, J. Mahieu, and B.C. De Cooman, S. Claessens: Mater. Sci. Technol., 19 (2003), 125.

[11] L. Tosal-Martiez, D. Vanderschueren, S. Jacobs, and S. Vandeputte: Steel Res., 72 (2001), No. 10, 412.

[12] K. Hulka, W. Bleck, K. Papamantellos: Relationships between heat treatment conditions, microstructure and properties of Niobium microalloyed TRIP steel, 41st MWSP Conf. Proc., 1999, ISS, Vol. 37, p.67.

[13] E.J. Drewes, E.F. Walker: Proc. Int. Symp. Niobium 2001, TMS, Warrendale, PA, 2001, p.873.

[14] X.D. Wang, B.X. Hunag, L. Wang, and Y.H. Rong: Metall. Trans. A, 39 A (2008), 1.

[15] A.J. DeArdo: Int. Mater. Rev., 48 (2003), 371. [16] J. Ohlert, W.

7] A.Z. Hanzaki, P.D. Hodgson, S. Yue: ISIJ Int., 5 (1995), 324. [18] S. Hashimoto, S. Ikeda, K. Sugimoto, and S. Miyake: ISIJ Int.,

2004, 44, 1590. [19] M. Soliman: Phase Transformation and Mechanical Properties of

New Austenite-Stabilised Bainite Steels, Piepersche Druckerei & Verlag GmbH, Dissertation, Clausthal-Zellerfeld, 2008.

[20] E Girault, P Jacques, Ph. Harlet, K Mols, J Van Humbeeck, E Aernoudt, and F.Delannay: Mater. Char., 40 (1998), 111.

[21] E. Wirthl, A. Pichler, R. Angerer, P. Stiaszny, K. Hauzenberger, Y.F. Titovets, and M. Hackl: Proc. Int. Conf. on TRIP-Aided High Strength Ferrous Alloys, 2002, Gent, p.61.

[22] J.Z. Zhao, C. Mesplont, B.C. DeCooman: ISIJ Int., 41(2001), 492. [23] S. Traint: Phase transformations and mechanical properties of low

alloyed dual-phase and TRIP-steels, Dissertation, VDI Verlag, Düsseldorf, 2002.

[24] R.M. Bozorth: Ferromagnetism, LucPress, New York, 1978.

5] L.P. Karjalainen, T.M. Maccagno, J.J. Jonas: ISIJ Int., 35 (1995), 1523.

[26] D. Q. Bai, S. Yue, W. P. Sun, J.J. Jonas: Metall. Trans., 24A (1993), 2151.

[27] M. Soliman and H. Palkowski: ISIJ int., 47 (2007), 1703. [28] M. De Meyer, J. Mahieu, B.C. De Cooman: Mater. Sci. Technol., 10

(2002), 1121. [29] W. Shi, L. Li, C. Ya

Eng. A, 429 (2006), 247. [30] R. Petrov, L. Kestens, and Y. Houbaert: ISIJ Int., 41 (2001), 883. [31] W.C. Jeong and C.H. Kim: J. Mater. Sci., 20 (1985), 4392. [32] M. De Meyer, J. Mahieu, and B.C. De Cooman: Mater. Sci.

Technol., 18 (2002), 1121. [33] C. Conrady and N. Kapustka: Sheet Metal Welding Conf. XII,

Livonia, MI, May 9–12, 2006. [34] T. Gladman and D. Dulieu: Grain Size Control in Steels, The Metals

Society, London, 1974, p. 167. [35] M. Chapa, S. F. Medina, V. López and B. Fernandez: ISIJ Int., 42

(2002), No. 11, 1288. [36] B. Dutta, C. M. Seller: Mater. Sci. Technol., 3 (1987), 197. [37] L. P. Karjalainen: Mater. Sci. Technol., 11 (1995), 557. [38] J.J. Jonas: Proc. of the TMP Conference, Liège, 2004, pp 35-42. [39] F. Boratto, R. Barbosa, S. Yue and J. J. Jonas: Proc. Int. Conf. on

Phys. Metall. of Thermomechanical Processing of Steels and Other Metals (Thermec-88), Tokyo, Japan, June 1988, pp. 383-390.

[40] M.L. Brandt and G.B. Olson: Ironmaking and Steelmaking, 20 (1993), No. 5, 55.

1] T. Maki, K. Tsuzaki, I. Tamura: Trans. ISIJ, 20 (1980), No. 4, 207.

Metallurgical Phenomena during Processing of Cold Rolled TRIP Steel

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