Influence of Microalloying Elements on Recrystallization ... - J-Stage

Materials Transactions, Vol. 51, No. 4 (2010) pp. 625 to 634 Special Issue on Crystallographic Orientation Distribution and Related Properties in Advanced Materials II #2010 The Japan Institute of Metals

Influence of Microalloying Elements on Recrystallization Texture of Warm-Rolled Interstitial Free Steels C. Capdevila1 , V. Amigo´2 , F. G. Caballero1 , C. Garcıá de Andre´s1 and M. D. Salvador2 1 2

Centro Nacional de Investigaciones Metalu´rgicas CENIM–CSIC, Avda. Gregorio del Amo 8, 28040 Madrid, Spain Instituto de Tecnologıá de Materiales, Universidad Politećnica de Valencia (UPV), Camino de Vera s/n, 46022 Valencia, Spain

The addition of microalloying elements such as Ti and Nb to increase the strength of deep drawing quality steels for automotive sheet products might affect the microstructure formed during the annealing after warm rolling in several ways. Firstly, the precipitates can exert a Zener’s pinning on growing recrystallized grain, which leads to a sluggish recrystallization kinetics. Secondly, the amount of microalloying elements control the amount of C, N, S, and P in solid solution, which indirectly affects the recrystallization texture obtained after annealing. In this sense, the work carried out with three different interstitial free (IF) and interstitial free high-strength (IFHS) steel grades allows us to conclude that the increase of microalloying additions delays recrystallization kinetics. Moreover, the abrupt texture change observed between as-rolled and annealed material indicates that the nucleation mechanism for recrystallization is more related to classical nucleation at deformed grain boundaries than subgrain rotation (continuous nucleation). [doi:10.2320/matertrans.MG200909] (Received September 30, 2009; Accepted February 5, 2010; Published March 17, 2010) Keywords: interstitial free (IF) steels, ferrite, recrystallization, texture

1.

Introduction

The excellent deep drawability of interstitial free (IF) steels led to their wide acceptability in the automobile and white goods industries. As the name suggested, IF steels are considered to be free of ‘interstitial’ atoms like C and N. However, it is impossible to get rid of all the C and N from the steel and normally Ti and or Nb are added to completely remove solute C and N content. The IF steels, despite their very good deep drawability, suffer from rather poor strength. For this reason a higher strength version, known as the interstitial free high strength (IFHS) steels have been developed in which Mn and P are added as solid solution strengthening elements.1–3) The addition of Ti, Nb, Mn and P are affecting the microstructure formed during annealing in several ways: firstly, it raises the mechanical strength of steel by precipitation hardening processes. Secondly, the precipitated particles can exerts a pinning effect on recrystallized grains evolving during annealing, thus recrystallization becomes a sluggish process. Additionally, textural evolution in both IF and IFHS grades is strongly dependent on chemistry, morphology, size and distribution of the precipitates, their sequence of occurrence and also their volume fractions. The microalloying additions will also have an indirect action on the microstructure through a change in the recrystallization texture. It is well documented that interstitial elements such as C and N in solid solution has a significant effect on the processes of nucleation of recrystallized grains, triggering grain boundary low energy nucleation processes in detriment of high energy nucleation processes:4,5) {111} or random grains are formed in the early stages of nucleation depending on the amount of C in solid solution. If this amount is low, there is a general and very quick restoration in the deformed grain and it is activated different nucleation sites that leads to a {111} grain texture. Conversely, if the amount of carbon in solid solution is high, the restoration processes are slowed because the C-Mn dipoles immobilize dislocations.6) In this

case, grains with lower strain energy grow at the expense of highly deformed grains.7,8) One goal of this paper is to analyze the effect of various microalloying elements on the complex precipitation behavior and its correlation with textural development in IF and IFHS steels during the different stages of processing. An attempt has also been made to correlate these aspects with recrystallization and texture formation in such high-drawing steels. The steels studied in this work have been produced under warm rolling conditions. ‘‘Ferritic’’ or ‘‘Warm’’ rolling is recognized as a very promising and cost effective way for steel strip processing.9,10) The economical advantages related to the ferritic rolling practice at the hot strip mill are numerous and quite large: energy savings derived from the use of a low reheating temperature; less oxidation in the reheating furnace; possibility to combine ferritic rolling with hot charging, and even with direct rolling, including thin slab casting;11) less work roll wear due to the reduced rolling temperature and less intermediate roll changes; better strip flatness control by rolling and cooling on a pre-transformed and homogeneous microstructure and no more metallurgical limitation for the hot rolling of very thin gauge strip. 2.

Materials and Experimental Procedure

In order to study the effect of microalloying elements in microstructure evolution of warm-rolled IF steels, three Ti, Ti/Nb and High strength IF steels were studied. The chemical composition of the three steels studied is listed in Table 1. Thermomechanical simulations were performed on a hot rolling pilot. After austenitic reheating at 1373 K, two roughing paths at 1373 and 1323 K, respectively, are first performed with a deformation level of 35%. A last deformation in the austenite domain (1273 K) is performed with a deformation level of 40%. Then, a ferritic rolling is performed with a deformation level of 53% at 923 K. The lubrication efficiency during ferritic hot rolling has a strong influence on the subsequent textures, therefore, we used a beef tallow lubricate during the ferritic pass.12) The steel is

C. Capdevila, V. Amigo´, F. G. Caballero, C. G. d. Andre´s and M. D. Salvador

626

N

Mn

P

Al

Nb

Ti

IF-Ti

0.0039 0.006

0.0043 0.009

0.007

0.028

0.001

0.079

IF-TiNb

0.0018 0.01

0.0021 0.014

0.012

0.027

0.017

0.014

IF-HR

0.0077 0.015

0.0026 0.409

0.031

0.044

0.001

0.118

then directly reheated up to 1023 K. The recrystallization process is monitored through interrupted by quenching cooling cycles at time ranging from 30 to 5400 s. r-values of complete recrystallized steels were obtained following the procedure described in Ref. 13). Specimens, longitudinal and cross-sectional specimens to the rolling direction, were ground and polished using standardized techniques for metallographic examination. A 2% Nital etching solution was used to reveal ferritic microstructure by light optical microscopy (LOM) and scanning electron microscopy (FEG-SEM) in a JEOL JSM6500F field emission gun scanning electron microscope operating at 7 kV. EDX analysis of precipitated were also carried out in this equipment provided with a Oxford INCA detector and operating at 15 kV. Texture measurements were performed by means of the Schulz reflection method using a D-5000 X-ray diffractometer furnished with an opened Eulerian cradle. Details of both the diffractometer and the analysis method have been given elsewhere.14) The pole figures (110), (200) and (211) were measured and a series expansion technique was employed to calculate the orientation distribution function (ODF), along with ghost correction.15) Microtexture analysis of as-deformed and partially recrystallized specimens was performed by the Electron Backscattering Diffraction (EBSD) technique. EBSD patterns were collected at various locations on cross and flat sections carefully polished with colloidal silica in the final phase. Ultrasonic cleansing process in ethanol at 303 K is performed in order to remove all the dirtiness from previous steps. The EBSD patterns were generated at an acceleration voltage of 20 kV and collected using a CRYSTAL detector of Oxford Instruments mounted in a SEM JEOL JSM6300. The indexation of the Kikuchi lines and the determination of the orientations were done with the software INCA developed by Oxford Instruments. The results were represented by means of an inverse pole figure (IPF) maps, which give the orientation of a macroscopic direction with respect to a specific crystal direction. A set of theoretical calculations concerning temperature evolution of different precipitates present, as well as their compositions were performed with the help of a commercial package for thermodynamic calculations in equilibrium named MTDATA.16) 3.

Results and Discussion

3.1 Precipitation state Different precipitates may form during thermo-mechanical treatment of IF and IFHS steels and their stability can be linked to several parameters, in particular the temperature and the amounts of alloying elements in the steels. According

IF-Ti

NbC TiN Ti(C,N) STi

Phase Fraction (%)

S

0.03

MnS

0.02

0.01

0 773

973

1173

1373

1573

Temeprature, T / K

(a) 0.03

IF-TiNb

NbC TiN Ti4C2S2 MnS

Phase Fraction (%)

C

Chemical composition of steel studied (in mass%).

0.02

0.01

0 773

973

1173

1373

1573

Temeprature, T / K

(b) 0.06 IF-HR

Ti(C,N) TiN

0.05

(Ti,Nb)(C,N) Ti4C2S2

Phase Fraction (%)

Table 1

MnS

0.04

0.03

0.02

0.01

0 773

973

1173

1373

1573

Temeprature, T / K

(c) Fig. 1 Amount of precipitates formed at different temperatures in (a) IF-Ti, (b) IF-TiNb and (c) IF-HR steels. Phase Fraction stands for the fraction of each precipitate expressed on percentage.

to Hua et al.,17) the required amount of Ti to stabilize IF steels is calculated by Ti = 4C+3.42N+1.5S assuming that C is tied up as TiC, N as TiN and S as TiS. However, if the stabilization occurs by the formation of Ti4 C2 S2 and TiN, the amount of Ti required is given by Ti = 3.42N+3S. It is clear that the amount of Ti in IF-Ti steel (Table 1) is enough by far to stabilize the steel. In fact, is more than enough, and some amount will remain in solid solution. Figure 1 shows the calculated evolution of precipitates by means of MTDATA thermodynamic databank in the three

Influence of Microalloying Elements on Recrystallization Texture of Warm-Rolled Interstitial Free Steels

627

IF-Ti Fe

Fe

Ti N C

Fe

Ti

2

0

4

6

keV (a) IF-Ti1-2 Ti

Ti

Fe Ti

S

Fe Fe

0

2

4

6

keV

(b)

Fig. 2

(a) TiC precipitates in IF-Ti steel, and (b) Ti4 C2 S2 precipitate type in as-warm rolled condition in IF-TiNb steel.

steels studied. Results for IF-Ti steels indicate that C is tied up as (Ti,Nb)(C,N) and NbC. Bearing in mind the thermomechanical processing route followed, these results are consistent with the idea reported by De Ardo and coworkers18) that precipitates containing Nb can only be observed (in Ti and Nb IF steels) in the following conditions: (i) as casted, produced by continuous casting, (ii) reheated below 1273 K and (iii) thermomechanical processing (TMP) below 1273 K. The amount of such precipitates in IF-Ti steel is significantly lower than in IF-TiNb steel (Fig. 1(a) vs. Fig. 1(b)) because of the total Nb. Figure 2(a) shows the SEM micrographs in IF-Ti steel, confirming the presence of very fine Ti(C,N) precipitates. Their small size may affect the progress of recrystallization process during annealing, since these particles can act as pinning points that impede the movement of growing grains. Besides, the precipitation temperature of these particles suggests that recrystallization and precipitation could be coupled processes, i.e. the precipitates are more prone to form in the sub-grain boundaries, and hence will delay or even impede the formation of recrystallize grain by sub-grain rotation or continuous recrystallization mechanism. In the Nb added IF steels stabilization of C is largely independent of the Mn and S contents; while in the Ticontaining steels the stabilization of C depends upon the amounts of Mn and S present.19) In Ti-Nb stabilized IF steel, C is principally removed from solid solution by the formation of Ti4 C2 S2 , shown in Fig. 2(b). The existence of these sulfur compounds indicates that the reheating temperature of 1453 K in the steel has not been high enough to dissolve

before the roughing stage during processing. These precipitates remove carbon from solid solution which promotes the deep drawing properties of the material through the formation of a -fibre texture. However, the formation of these compounds avoid the formation of other precipitates that tied up phosphorus more effectively such as FeTiP, which has a detrimental effect on deep drawing properties by the presence of free P.20) It is important to note here that P is largely responsible for the hardness level of the steel that makes the steel competitive as car body panel material.21–23) On the other hand, precipitation of (Ti,Nb)C is not predicted by calculations in IF-TiNb steel. In addition, Fig. 3 shows the distribution of Nb among the phases present in IF-TiNb steel for a certain temperature. It is clear that most of the Nb is tied up as NbC and only less than 20% of total Nb in the steel remind in ferrite. This result could lead to an improved cold work embritlement (CWE) performance of this steel as compared with similar steels reported in literature, where studies on stabilization of IF steels containing (Ti+Nb) have shown that, depending upon the Ti level, as much as 67% of the total Nb may remain in solid solution.24) This solute Nb has been shown to segregate to grain and sub-grain boundaries thereby improving the CWE otherwise caused by P segregation at the same sites.20) As indicated by Fig. 1(c), precipitation behavior of IF-HR steel is quite different and much more complex than that of IF-Ti and IF-TiNb steels. This can be attributed mainly to the amount of Ti,C. Likewise, as calculations indicate, P will remain in solid solution in IF-HR steel and thus will also result in solid solution strengthening. Figure 4 shows the


628 100

100

100

90

90

80

80

70

70

60

60

50

50

Austenite

90 80

Ferrite

Ti Fraction (%)

Nb Fraction (%)

70 60 50 40

40

40 Ti in IF-Ti

30

30 20

20

10

10

Ti in IF-TiNb Ti in IF-HR

30

P Fraction (%)

NbC

20

P in IF-Ti P in IF-TiNb

10

P in IF-HR

0 773

0 773

Fig. 3

973

1173

1373

1573

973

1173

0 1573

1373

Temeprature, T / K

Temeprature, T / K

Distribution of Nb among the phases in IF-TiNb steel.

Fig. 4 Distribution of Ti and P in solid solution in the matrix (ferrite).

IF-HR P

Ti Fe

1

6

11

16

keV

(a) IF-TiNb P

Ti Fe

Nb

0

5

10

15

20

keV (b)

Fig. 5

(a) FeTiP precipitate in IF-HR steel, and (b) Complex Fe(Ti+Nb)P precipitate in IF-TiNb steel.

amount of P and Ti in the matrix (ferrite) for the three steels studied. It is clear that, according with calculations, the whole amount of P and a significant amount of Ti remains in solid solution in the ferrite. However, it has been profusely reported in literature25,26) that in presence of Ti, P forms FeTiP phosphide which leads to the deterioration of both drawability and loss of strength. The FeTiP precipitates present in IF-HR steel (Fig. 5(a)), because of the Nb content, become Fe(Ti,Nb)P in Nb-rich steels (Fig. 5(b)), as has been previously reported in the literature.21,22) The presence of these precipitates depends mainly on the Ti and Nb content of steels. For this reason, the amount of these precipitates in IF-

TiNb steel is very small and almost negligible in the IF-Ti and IF-HR steels. The study of solubility products for the TiC and FeTiP done by Gladman23) shows that both are quite close, and hence those precipitates compete with each other. Thus FeTiP precipitate is more likely to be found in the IFHS steel than in the IF-Ti steel because of the P content. In summary, the precipitation events has been defined in IF-Ti, IF-TINb and in IF-HR steels with the help of MTDATA thermodynamic databank software. By means of ranking the precipitates by decreasing the free energy change during precipitation, the precipitation sequence can be summarized as follows: TiN, Tix S, Ti4 C2 S2 , TiC and FeTiP,


629

ϕ1

φ

3.2

(a)

4

φ

(b)

ϕ1

(a)

4

φ

(c)

(d)

Fig. 7 2 ¼ 45 sections of ODFs for (a) IF-Ti, (b) IF-TiNb, (c) IF-HR steels in as as-warm rolled condition, and (d) major components (after2)).

(b)

20 µm (c)

Fig. 6 As-received optical microstructures of (a) IF-Ti, (b) IF-TiNb and (c) IF-HR steels.

which agrees with thermodynamic studies reported about the precipitation sequence that occurs in these steels during solidification. After the formation of TiN, precipitation of sulfur compounds in the form of TiS, or more complex as the Ti4 C2 S2 is more likely to occur. The stability of the precipitates depends on several parameters, including temperature and composition of steels. 3.2 Crystallographic orientation analysis The deformed microstructures were composed of ferrite grains elongated along the rolling direction, some of which contained in-grain shear bands. Examples of these warm rolled microstructures are presented in Fig. 6 at low magnification. Figure 6 shows the as-quenched microstructure of the IF-Ti and IF-TiNb steels after warm rolling and reheating at 1023 K (samples quenched just after reaching at 1023 K). It can be clearly seen deformed elongated grains in the rolling direction. Deformation bands are also oriented about 45 to the rolling direction. These bands are of great

interest for potential nucleation sites for recrystallized grains. These microstructures are similar to those obtained after cold rolling, i.e. high level of deformation is stored in the microstructure. It can be seen that the percentage of grains containing shear bands, and its intensities, are higher in the IF-Ti steel. On the other hand, the amount of shear bands in IF-HR steel is lower, which is consistent with the inhibition of shear band formation at higher rolling temperatures by the addition of phosphorus.26) Figure 7 shows the ’2 ¼ 45 section of the orientation distribution functions (ODFs) of warm-rolled IF steels, and it is clear that both the and fibres are well developed. The fibre runs from f001gh110i to f111gh110i, with clustering of poles at f001gh110i and f112gh110i. The fibre runs from f111gh110i to f111gh112i and is peaked at f111gh110i. A slightly stronger texture is detected as increasing the microalloying content. Stronger f111gh112i and f111gh110i textures is developed in IF-TiNb and IF-HR steels. The EBSD analysis shows clear differences with the asdeformed microstructure after cold rolling process. Figure 8 shows the presence of sub-grains inside the coarser deformed grains. A more in detail observation of these sub-grains (Fig. 8(b)) reveals an average size of 2 mm in mean diameter. Likewise, Fig. 8(c) shows a misorientation map between neighboring grains, where sub-grain boundaries (those with misorientation lower than 10 ) are clearly seen, meanwhile the grain boundaries of deformed grains have a misorientation higher than 40 . The presence of sub-grains allows us to conclude that the steel has followed a dynamic recovery process during the warm-rolling process. After EBSD analysis the deformed grain types can be classified in four different groups (Table 2). As an example, Fig. 9 shows the type-2 grains present in IF-HR steel. Figure 10 shows the evolution of the recrystallized fraction with holding time during isothermal treatment at 1023 K for the three steels studied. It is clearly seen an increasing delay on the kinetics of recrystallization as


630

(a)

(b)

(c)

(d)

(e) Fig. 8 (a) Presence of sub-grains, (b) sub-grain sizes, (c) sub-grain boundaries, (d) misorientation histogram in the as-quenched microstructure of IF-HR steel. Legend of colour code for 8(a) is presented in (e).

Table 2

Type 1

Deformed grains classification.

Type 2

Type 3

Type 4

ND

ND

ND

ND

RD

RD

RD

RD

ND// RD// ∼

and ND// ∼ Around and RD//

microalloying element content increases. Thus, the IF-Ti steel presents the fastest recrystallization kinetics, followed by IF-TiNb steel and finally, the IF-HS steel. From the precipitation results presented in Fig. 1, we can conclude that

{112}

Around {001}

the differences in recrystallization kinetics are due to the differences in the amount of precipitates. Besides the higher amount of (Ti,Nb)(C,N), the main difference between IF-Ti and IF-HS steels lies mainly on the level of Ti4 C2 S2 and


631

Table 3 Measured average r-values.

(a)

(b) Fig. 9 (a) Type 2 grains present in IF-HR steel. (b) Legend of colour code of 9(a).

100

IF-Ti

IF-TiNb

IF-HR

1.36

1.22

1.08

kinetics. The delay in recrystallization kinetics shown by IFTiNb steel might be due to the combined effect of the precipitates MC-type (such as C0.88Nb) and Fe(Ti+Nb)P precipitates, and to a lesser extent, by solute drag phenomena. Finally, Fig. 11 shows the evolution of microtexture during recrystallization in IF-Ti steel. The deformed state (Fig. 11(a)) is a mixture of Type 2 and 3 grains, and the recrystallized grains has a strong texture with ND == h111i (Fig. 11(c)). On the other hand, Fig. 12 shows the fully recrystallized microstructures of IF-Ti and IF-TiNb steels. It is clearly seen that the coarser final grain size of IF-Ti steel (39 mm in average diameter) as compare with IF-TiNb steel (31 mm average diameter). IF-HR presents the finer recrystallized grain size with a value of 21 mm. Figure 13 shows the recrystallization textures obtained after annealing at 1023 K. The final recrystallization textures developed in the three steels are quite similar in nature. They show a prominent -fiber but maximums is not clearly centered on f111gh112i. Likewise, the texture is decidedly sharper in both IF-Ti and IF-TiNb steels (containing less P) as compared with IF-HR steel, which also possesses a higher value of 4 as compared to 3.2 for IF-HR steel. The higher density of FeTiP precipitates in IF-HR steel appears to produce a less intense {111} texture. Besides, as it was above mentioned, the higher amount of P will decrease the amount of Ti in solid solution in the ferrite which will affect the average r-value. Excess Ti is necessary for attaining a better r-value.28) This is consistent with the r-values measured and listed in Table 3.

Fraction Recrystallized (%)

ReX_IF-HR ReX_IF-Ti

80

ReX_IF-TiNb

60

40

20

0 1

10

100 Time, t / s

1000

10000

Fig. 10 Hardness and recrystallized volume fraction evolution with time for isothermal holding at 1023 K.

FeTiP compounds. Therefore, the latter two should be responsible for the slower kinetics in IF-HS steel, not the Ti(C,N) precipitates more abundant in IF-Ti as compared with IF-TiNb steels. It is also well established that the presence of microalloying elements in solid solution may retard the recrystallization by solute drag phenomena.27) In this sense, because of the largest concentrations, IF-HS steel is more prone to present the most sluggish recrystallization

3.3 Recrystallization mechanism In the following paragraphs the mechanism of recrystallization will be described. Previous works reported that deformation bands are noticeable features in the microstructure of the present warm worked IF steel. Their directions to the rolling direction of 35 are independent of strain, temperature and initial grain orientations and are very similar to the banded structures reported in Al alloys.29) Likewise, it was observed that a decrease in recrystallization kinetics as compared with cold-rolled IF steels due to thermally activated processes such as dislocation cross slip and climb occur at a greater rate during warm rolling, and thus lower the overall dislocation density.30) Meanwhile the recrystallization of the warm-rolled low carbon steel occurred mainly from deformed f100gh110i oriented grains (low energy nucleation), the main nucleation sites in the warm-rolled IF steels were the deformed {111} grains (high energy nucleation).31) The low stored energy nucleation occurs typically by the migration of a ‘bulge’ in an existing grain boundary from low energy grain into a high energy grain. It is easy to envisage site saturation occurring in a low stored energy material, when nucleation involves the bulging of existing grain boundaries, than when a subgrain growth or coalescence mechanism is operating.

632


(a)

(b)

(c)

(d) Fig. 11 Microstructural evolution in IF-Ti steel after annealing at 1023 K for (a) 0 s, (b) 60 s and (c) 180 s holding time. ND color maps have been used. Legend is presented in (d).

On the other hand, recrystallization nuclei in IF steels form preferentially within grains containing shear bands, probably because of the higher stored energies of these grains.32) The presence of shear bands of moderate intensity in warm-rolled IF steels, which tended to be insensitive to rolling temperature, where reported in literature.33) The development of shear bands was insensitive to the presence of microalloying precipitation, but shear band formation is inhibited at higher rolling temperatures by the addition of phosphorus.32) This is consistent with microstructures shown in Fig. 6, where shear bands are more abundant in IF-Ti and IF-TiNb steels than in IF-HR steel. As to why low stored nucleation occurs, which is not the dominant mechanism in warm-rolled IF steels, it is due to the strain induced precipitation that occurs after some conditions of warm working of IF steel, which also is responsible of the retardation of recrystallization in warm-rolled material as compared with cold-rolled material.29) Senuma et al.34) have published results suggesting that the presence of Ti-precipitates prevents bulging. Therefore, the continuous recrystal-

lization could be the controlling recrystallization mechanism in IF warm-rolled steels. Continuous recrystallization is referred as a mechanism in which a physical nucleus is not formed but rotations of two neighboring sub-grains coalesce in a new grain. If the recrystallization mechanism follows a sub-grain rotation or continuous recrystallization process, the texture of the new grain should not be very different than the previous deformed grain. However, if the predominant recrystallization mechanism is by grain-boundary nucleation, the texture of the new grain could be very different than the one of previous deformed grain. These are consistent with the experimental results obtained for IF-Ti, IF-TiNb and IF-HR steels (Fig. 14), and it allows us to conclude that grain-boundary nucleation could be the responsible recrystallization mechanism in IF-HR steel, meanwhile continuous recrystallization of grains nucleated in the shear bands is the responsible mechanism for IF-Ti and IF-TiNb steels. The micrographs in Fig. 11 show the evolution of the microstructure of the IF-Ti steel during the

Influence of Microalloying Elements on Recrystallization Texture of Warm-Rolled Interstitial Free Steels ϕ1

633

ϕ1

φ

3.2

(a)

φ

4

(b)

ϕ1 (a)

3.2

φ

(c)

(d)

Fig. 13 2 ¼ 45 sections of ODFs for (a) IF-Ti, (b) IF-TiNb, (c) IF-HR steels in fully recrystallised condition after 1023 K annealing, and (d) major components (after2)). (b)

(Ti, Nb) (C, N) and Fe(Ti+Nb)P in IF-TiNb steel, and FeTiP in IF-HR steel have been observed. It is clearly seen an increasing delay on the kinetics of recrystallization as microalloying content increases, probably as a result of solute drag phenomena. Thus, the steel with faster recrystallization kinetics is the IF-Ti steel, followed by IF-TiNb and finally, the IF-HR steel. The presence of sub-grains in the microstructure as compared with cold-rolled steels indicates the existence of dynamic recovery processes during warm rolling of deepdrawing IF and IFHS steels. The sub-grains found in IF-HR 40 µm steel are 2 mm in size. (c) Grain-boundary nucleation could be the responsible recrystallization mechanism in IF-HR steel, meanwhile Fig. 12 Full recrystallized microstructures in (a) IF-Ti, (b) IF-TiNb and continuous recrystallization of grains nucleated in the shear (c) IF-HR steels. bands is the responsible mechanism for IF-Ti and IF-TiNb steels. The observed deformed grains in Ti-IF, IF-TiNb and IFisothermal annealing at 1023 K. The recrystallized grains nucleate first into certain deformed grains, while others do HR steels can be classified into four distinct groups: Type 1 not exhibit any signs of recrystallization. The growth of these (ND == h111i and RD == h112i), Type 2 (ND == h111i and firstly recrystallized grains seems to be restricted to the same RD == h110i), Type 3 (Around (112)h110i) and Type 4 deformed grains where they nucleated. At later stages, the (Around (001)h110i). The most important alloying element in IFHS steels, recrystallization affects all the initial deformed microstructure. This picture is similar to those observed in IF-TiNb which distinguishes this grade from the IF grade, is P, which steel. In those steels, it was observed that the {111} grains do is added intentionally for solid solution strengthening. A not have significant number advantage over ‘‘other’’ orienta- major problem with these steels is the formation of FeTiP tions but are systematically more numerous than those which degrades the {111} texture in three different ways: belonging to {110} or {100} components. For this same These precipitates take out a substantial amount of stabilizing sample, the number of the {111} grains remains almost element Ti from the steel, leaving behind much less Ti in the constant, while the {100} and {110} grains decrease matrix to combine with the interstitial elements. Secondly, removal of P atoms from the matrix will result in a material continuously in number during recrystallization. of poorer strength. And finally, FeTiP particles, often present along the grain boundaries, can produce Zener drag and 4. Conclusions retard the growth of favorable recrystallized grains, giving Precipitates such as TiC and Ti4 C2 S2 in IF-Ti steel, rise to a poor texture.


634

{–8 7 8}

{–9 16 –1}

{1 2 –2}

(a)

(b)

(c) Fig. 14 (a) Recrystallized grain (see arrow) after 180 s holding time in IF-HR steel, (b) its orientation to the matrix, and (c) legend of colour code.

Acknowledgements The authors acknowledge the economical support of European Commission through the RFCS Programme (RFS-PR-03136) and the Spanish Ministerio de Ciencia e Innovacioń through ENE2006-15170-C02-01 project. C. Capdevila Acknowledge the Spanish Ministerio de Educacioń, Cultura y Deportes for economic support through the ‘‘Salvador de Madariaga’’ Program (PR2006-0045). REFERENCES 1) 2) 3) 4) 5) 6) 7) 8) 9) 10) 11) 12)

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