Grain refinement by rapid transformation annealing of

Grain refinement by rapid transformation annealing of cold rolled low carbon steels P. Álvarez1,a, C. Lesch2,b, W. Bleck2,c, H. Petitgand3,d, J. Schöttler4,e, J. Gil Sevillano1,f 1

CEIT and TECNUN (Universidad de Navarra), M. de Lardizábal 15, 20018, San Sebastián, Spain

2

RWTH Aachen (IEHK), Institut für Eisenhüttenkunde. Intzestraβe 1, D-52072, Aachen, Germany 3

IRSID, ARCELOR, F-57283, Maizières-les-Metz – Cedex, France 4

Saltzgitter Flachstahl GmbH, D-38223, Saltzgitter, Germany

a

[email protected], [email protected], [email protected], d [email protected], [email protected], [email protected]

Keywords: Grain refinement, microalloyed steels, rapid annealing, α recrystallization, α-γ-α phase transformation.

Abstract. A novel thermal treatment, rapid transformation annealing (RTA), has been applied to six different cold rolled low-carbon (LC) steel sheets with the aim of refining their microstructure. The process involves rapid heating to just above the austenite (γ) to ferrite (α) transformation temperature and subsequent rapid cooling to room temperature. Grain sizes around 2 µm in two different Nb-Ti HSLA steels, 5 µm in a Ti-LC steel and 6 µm in a plain LC (0.037%C) steel have been produced using fast cooling rates (200ºC/s). Non-equiaxed structures are obtained in a Nb-Ti HSIF steel and in a plain LC (0.135%C) (CM) steel due to their higher Mn content. However, very fine equiaxed grains (2 µm) are obtained by rapid intercritical annealing (RIA) in the CM steel. Irrespective of the microalloying concept, the grain growth of recrystallized α grains before their transformation was inhibited in CM and in both HSLA steels. This inhibition is connected with the overlapping of α recrystallization and α-γ transformation processes which is essential in order to achieve extreme grain refinement either by RTA or RIA. Introduction As grain refinement increases both the strength and toughness of steels, special interest has lately arisen in the production of ultra-fine ferrite steel plates and strips, with the aim of overcoming the practical limit grain size of about 5 µm that the optimization of thermomechanical treatments has led to. A complete review of refinement methods and achieved grain sizes was recently given by Howe et al. [1]. Among the rolling-based pilot methods, Deformation Induced Ferrite Transformation (DIFT) and Rapid Annealing are the most promising ones. Very fine structures with grain sizes in the order of 1 to 4 µm have been produced by DIFT in a wide variety of steel chemistries [2,3]. Homogeneous grain sizes of 2 µm have been produced by ultra-rapid annealing in Nb-microalloyed HSLA steels using extremely high heating rates [4]. Leinonen [5] developed a new rapid heat treatment for hot rolled strips by which ferrite grain sizes of 3 to 4 µm were achieved in conventional C-Mn steels. The method consists of a rapid heating to just above the Ac3 temperature, a brief holding time and a rapid cooling to below Ac1. The refinement is enhanced since the transformation range is passed through twice, during heating and during cooling. Rapid transformation annealing (RTA) represents an improvement with respect to such process. In this case, the treatment including α−γ−α transformation is applied to cold rolled strips, so that another metallurgical process promoting grain refinement, recrystallization of ferrite, is introduced. However, previous attempts to develop the RTA have not been very successful despite the use of moderate or ultra-high heating rates, especially with interstitial free steels [6,7]. From those works it could be concluded that in order to obtain very fine α grains in the final

product, α−γ transformation should start before the end of α recrystallization in order to get small enough austenite grain size. Experimental Salzgitter Flachstahl GmbH supplied six different steel grades with cold rolling reductions of 70% and 90% (Table 1): three different Ti-Nb microalloyed steels (HSLA1, HSLA2 and HSIF), one Timicroalloyed steel (ISO2) and two non-microalloyed steels (CM, DC04). Transformation behavior of 70% CR steels was investigated by dilatometry at two heating rates (10 and 100ºC/s). Ferrite (α) to austenite (γ) transformation temperatures were determined at both rates. They are referred as Ac1, Ac3, Ac1(100ºC/s) and Ac3(100ºC/s) in Table 1. Table 1. Chemical compositions (%wt) and α-γ transformation temperatures of the steels. Steel

C

Mn

Ti

Nb

Al

N

Si

HSLA1 HSLA2 CM ISO2 DC04 HSIF

0.065 0.078 0.135 0.050 0.037 0.003

0.75 0.63 1.17 0.19 0.18 0.72

0.018 0.019 0.002 0.026 0.001 0.073

0.028 0.046 0.001 0.001 0.000 0.029

0.051 0.032 0.043 0.048 0.028 0.021

0.0046 0.0056 0.0029 0.0048 0.0034 0.0032

0.026 0.053 0.017 0.012 0.011 0.032

Ac1 [ºC] 730 735 695 725 670 920

Ac3 [ºC] 925 930 885 940 955 1000

Ac1 (100ºC/s) [ºC] 695 700 710 715 715 935

Ac3 (100º/s) [ºC] 920 905 870 940 935 1015

Test strips of 180 x 10 mm were machined from the cold rolled sheets. A resistive heat treatment simulator was employed for the RTA treatments at different heating rates ranging from 10 to 170ºC/s. The peak temperature for each steel grade was fixed just above (10-25ºC) its Ac3(100ºC/s) transformation temperature. For accurate temperature control, two K type Cromel-Alumel thermocouples were spot welded on each specimen. After a programmed holding stage, the samples were cooled to room temperature using a helium gas quench. The cooling rate was calculated as the average cooling rate from peak temperature to 500ºC. Softening behavior during the heating stage of the RTA was studied by heating the samples at different rates and quenching out from different temperature levels. The softening of ferrite was quantified by 1 Kg Vickers microhardness (Hv) measurements. Microstructures of partially and fully softened samples were examined by optical microscopy (2% Nital etching) and by orientation imaging microscopy (OIM) from electron backscattering diffraction (EBSD) measurements. The observation surface was parallel to RD and normal to TD. Grain sizes were determined as the mean linear intercepted length along RD and ND from several micrographs. Results Rapid Transformation Annealing (RTA) Treatments. Fig. 1 shows ferrite grain sizes reached by RTA in 70 and 90% CR HSLA1, HSLA2, ISO2 and DC04 samples. The grain size data correspond to different heating and cooling conditions after a constant holding time of about 1 second at the annealing temperature (Ac3(100ºC/s)+ 10-25ºC). Ferrite grain size decreases with increasing cooling rate. The grain sizes of samples heated at the same rate (10, 100 or 170ºC/s) vs. the logarithm of cooling rate admit a good linear fitting (Equation 1). dα = a · log (Rc) + b. where d is measured by the linear intercept method and Rc is the cooling rate in ºC/s.

(1)

(a)

(b)

(d) (c) Fig. 1. Impact of RTA cycle’s parameters on final grain size in a) HSLA1, b) HSLA2, c) ISO2 and d) DC04 samples. Peak T.: 930ºC (a & b) and 950ºC (c & d). Holding t.: 1s. Cooling rate cannot be increased beyond a limit value that marks the start of the production of low-ductility non-equiaxed (bainitic or acicular) structures. Such limit cooling rate, about 200ºC/s for HSLA1 and 300ºC/s for HSLA2 grades, determines the finest grain size achievable by RTA. For HSLA2 steel, the finest grain size, nearly independent of heating rate, was respectively 2.1 µm and 1.9 µm for 10 and 170ºC/s heating rates (cooling rates close to 300ºC/s). However, faster heating rates led to a significant refinement in HSLA1 samples, where the grain size was refined from 3.4 µm to 2.6 µm by increasing the heating rate (cooling rates close to 200ºC/s). ISO2 and DC04 grades presented equiaxed ferritic structures after cooling rates up to about 200ºC/s. It was not possible to reach faster cooling rates due to sample’s thickness, thus it was not possible to establish the finest grain size achievable by RTA for these two steels. Nevertheless, the grain sizes produced after cooling rates close to 200ºC/s were much coarser than those measured in HSLA steels, 5.1 µm for ISO2 and 6.3 µm for DC04 (heating rate: 100ºC/s). Heating rate did not affect the final grain size after neither fast nor moderate cooling rates. Non-equiaxed and distorted structures (bainite or acicular ferrite) were obtained in CM and HSIF grades from very low cooling rates, of about 10ºC/s. This made impossible the refinement of these steels by trans-critical RTA (peak temperatures above Ac3). Samples cold rolled to 70 or 90% reduction and subjected to similar RTA cycles (cooling rates from 10 to 100ºC/s) led to similar final grain sizes. Rapid Intercritical Annealing (RIA) Treatments. Rapid annealing cycles were carried out without crossing the fully austenization temperature (Ac3). This thermal treatment was very effective in refining grain size of 70 and 90% CR CM samples. Homogeneous grain size distributions with an average grain size of about 1.8 µm were produced (cooling rates: 185ºC/s). Note that this steel could not be refined after annealing in the fully austenization range since distorted structures easily formed, even at low cooling rates.

A constant peak temperature of 840ºC (corresponding to a transformed ferrite fraction of approx. 90%) was selected for studying the effect of the rest of cycle’s parameters. The direct relationship between grain size and cooling rate given by Eq. 1, was also fulfilled in these intercritically annealed samples (Fig. 2a). It is worth mentioning that grain sizes about 3.5 µm were produced in 70 and 90% CR CM samples that were cooled at only 10ºC/s. The final microstructure noticeably refined with increasing heating rate. The impact of cold reduction was negligible.

(a) (b) Fig. 2. Impact of cooling rate, heating rate and cold reduction on grain size in a) intercritically annealed CM and b) subcritically annealed HSIF samples. Peak T.: 840ºC. The reduction of peak temperature to the intercritical domain also prevented the development of distorted structures in HSLA1 steel. This made possible the use of very fast cooling rates without the formation of non-equiaxed structures. Grain sizes as fine as 2.0 µm were reached by RIA at 860ºC (transformed fraction: 60%), using very fast cooling rates (220ºC/s) that would not have given rise to fine ferritic structures after annealing at peak temperatures above Ac3. The increase of holding time at the peak temperature shifted the boundary of bainitic region to lower cooling rates in HSLA1 samples subjected to RTA. Consequently, non-equiaxed structures formed at moderate cooling rates (between 50-100ºC/s) even after holding times as short as 10 seconds. This fact significantly reduces the finest grain size achievable by RTA. By contrast, the intercritical treatment referred as RIA inhibited the formation of non-equiaxed structures to holding times longer than at least 60 seconds when cooling at rates faster than 100ºC/s. The reduction of peak temperature to the intercritical domain did not succeed in refining HSIF steel grade. Distorted structures generated at peak temperatures above Ac1 using moderate cooling rates. It was concluded that grain refinement of this steel by RTA or RIA treatments, comprising phase transformation, is not possible. Different subcritical annealing cycles were then tested keeping the peak temperature at 840ºC (lower than Ac1), i.e., with only recrystallization being active during the entire thermal cycle. Heating rate, cold reduction and cooling rate had not a strong influence on grain size and values between 8.0 and 9.0 µm were measured for a wide range of annealing conditions (Fig. 2b). Investigation of softening kinetics. Softening kinetics, which is related to transformation of cold worked ferrite by either recrystallization or α-γ phase transformation, was investigated performing interrupted annealing cycles at different peak temperatures. Table 2 displays parameters obtained from interpolation of the softening curves of 70% CR samples heated at 10 and 100ºC/s. Softening start (RSTART) and finish (REND) temperatures correspond to temperatures at which the decreases of hardness were, respectively, 10 and 90% of the total softening. In order to keep a similar criterion for comparison, transformation temperatures were defined as the temperatures for 10% (Ac1(100ºC/s)(10%)) and 90% (Ac3(100ºC/s)(90%)) transformed volume fractions of α (measured by dilatometry). They are included in Table 2, as well.

Table 2. Parameters obtained from softening and transformation curves of 70% CR samples. Heating rate: 10ºC/s

Heating rate: 100ºC/s

Steel

RSTART [ºC]

REND [ºC]

Range [ºC]

RSTART [ºC]

REND [ºC]

Range [ºC]

∆RSTART [ºC]

∆REND [ºC]

Ac1(100ºC/s) (10%) [ºC]

Ac3 (100ºC/s) (90%) [ºC]

HSLA1 HSLA2 CM ISO2 DC04 HSIF

690 705 645 675 610 745

785 790 670 735 670 785

95 85 25 60 60 40

755 760 680 705 640 760

845 840 760 760 685 800

90 80 80 55 45 40

65 55 35 30 30 15

60 50 90 25 15 40

775 775 750 815 795 940

880 885 840 910 905 955

With increasing heating rate, anisothermal softening shifted to higher temperatures. Both, RSTART and REND raised. The increment of RSTART was more pronounced in both HSLA steels (approx. 60ºC) than in the rest of steel grades (15-35ºC). It was also remarkable the slowdown of softening kinetics of CM at the fastest heating rate. Softening temperature intervals of 70% CR ISO2, DC04 and HSIF steels heated at 100ºC/s were notably shorter (40-55ºC) than those of 70% CR HSLA1, HSLA2 and CM steel grades (80-90ºC). RSTART can be regarded as the recrystallization start temperature because it was lower than Ac1(100ºC/s)(10%) in all the studied steels. However, for CM and both HSLA steels annealed at 100ºC/s, REND was notably higher than Ac1(100ºC/s)(10%). This means that by the time the softening finished, a certain amount of ferrite had already transformed to austenite and, therefore, there was recrystallization-transformation interaction during RTA of the CM and the two HSLA steels. The double effect of increasing heating rate, that raises the temperature at which recrystallized ferrite nucleates and enables the recrystallization-transformation interaction, has its reflect on a greater microstructural refinement at temperatures around REND. The former effect increases recrystallization nucleation rate encouraging the formation of more nuclei [6,8,9] and the latter one inhibits the growth of these nuclei before their transformation to austenite [6,10]. For instance, Fig. 3 shows microstructures of 70% CR CM steel that were annealed at temperatures just above REND using different heating rates. The refinement of ferrite grain size, from 6.1 µm to 2.4 µm, is evident. This fact was not observed in 70% CR DC04, ISO2 or HSIF steels where the heating rate influence on RSTART was much less important and where there was no recrystallization-transformation interaction. For these grades, the ferrite grain sizes at REND were between 7.0-8.0 µm independently of heating rate.

(a) (b) Fig. 3. Microstructures of 70% CR CM samples annealed at temperatures just above REND; a) Heating rate: 10ºC/s, Peak T.: 696ºC; b) Heating rate: 100ºC/s, Peak T.: 789ºC/s. ↑ RD. With increasing cold reduction (heating rate: 100ºC/s), softening was shifted to lower temperatures and was completed in a shorter temperature interval. Consequently, REND greatly reduced (50-75ºC) limiting the recrystallization-transformation interaction. However, ferrite grains

about 2.0 µm were still produced in 90% CR CM and both HSLA steels when annealing at peak temperatures just above REND (740 and 790ºC, respectively). The increment of recrystallization nucleation rate due to the higher stored internal energy is thought to be the reason explaining those results as the transformed ferrite fractions are not higher than 15% at these temperatures. As Fig. 4 shows, ferrite grain sizes achieved after completion of softening do not practically change in the entire range of peak temperatures from REND to just above Ac3 where α-γ-α transformation is ongoing.

(b) (a) Fig. 4. Microstructures of 90% CR HSLA2 samples rapidly annealed at a) 804ºC and b) 934ºC. Cooling rate > 200ºC/s; Holding t.: 1s. ↑ RD. Discussion α-Recrystallization and α-γ Transformation. RSTART (recrystallization nucleation temperature) is related to the concentration of Nb and Ti microalloying elements as the following sequence has been established for 70% CR samples at 10 and 100ºC/s heating rates (Table 2): Rstart → HSIF (Nb+Ti) > HSLA1, HSLA2 (Nb+Ti) > ISO (Ti) > CM, DC04 (C-Mn) The extremely strong effect that solute Nb has in retarding recrystallization of cold worked ferrite has already been studied [11]. Characterization of the precipitation state in HSLA2 samples showed that at least 50% of Nb is in solid solution at RSTART. On the other hand, 90% of Ti had precipitated before annealing cycle to form titanium nitrides and did not evolve during annealing. It is also well known that a fine dispersion of second phase particles (size~50Å) delays recrystallization [12]. Equation 2 has often been employed to determine the recrystallized fractions in sub-critical annealing treatments of cold rolled steel strips [10,13,14]. It has been employed here to determine the softened (recrystallized or α-γ-α transformed) ferrite fractions at each temperature. % X REX =

H 0 − HT ·100. H0 − H F

(2)

where H0 is the hardness before recrystallization starts, HF is the hardness after recrystallization finishes, and HT is the hardness at an intermediate stage. The fractions of softened (non-deformed) ferrite at different temperatures have been calculated using Eq. 2 and the hardness measurements after interrupted treatments. There is a good agreement between deformation-free ferrite fractions measured by EBSD and those calculated with Eq. 2. Curves corresponding to deformation-free ferrite fraction (dashed curves) together with transformation curves that give the double-transformed ferrite fraction at each temperature (dotted curves), are drawn in Fig. 5 (heating rate: 100ºC/s). It is again evident that recrystallizationtransformation interaction occurs in 70% CR CM and both HSLA steels. On the assumption that

transformation of deformed ferrite (with higher stored energy) takes place prior to transformation of recrystallized ferrite when both structures coexist, the fraction of recrystallized ferrite remaining in the microstructure at a given temperature can be calculated as: % Recrystallized ferrite = % Deformation-free ferrite - % Double-transformed ferrite.

(3)

Solid curves in Fig. 5 represent the fraction of recrystallized ferrite at each temperature calculated according to Eq. 3. In the cases of 70% CR HSLA1 and HSLA2 steels, this fraction does not exceed 50-55%. So we can say that recrystallization is progressively substituted by α-γ transformation, the later playing an important role on the replacement of the cold worked matrix of these steels.

(a)

(b)

(c)

(d)

(e) (f) Fig. 5. Evolution of deformation-free ferrite (dashed lines), recrystallized ferrite (solid lines) and double-transformed (dotted lines) with peak temperature in 70% CR a) HSLA1, b) HSLA2, c) CM d) ISO2, e) DC04 and f) HSIF steel grades. Heating rate: 100ºC/s. Common to the steels that are greatly grain refined by RTA or RIA (HSLA1, HSLA2 and CM) was the observation that α-γ transformation concomitant with recrystallization inhibited the growth of recrystallized grains before their transformation to γ, as confirmed by measurements of ferrite grain size during the different stages of rapid annealing in 70 and 90% CR samples. On the contrary, recrystallized α grains rapidly grew before their transformation in ISO2, DC04 and HSIF samples.

γ-α Transformation. The grain refinement achievable by RTA depends mainly on cooling rate through the control of the γ-α transformation. Increasing cooling rate shifts transformation to lower temperatures affecting nucleation (nucleation site density increases) and growth of ferrite grains (controlled mainly by the C diffusion) and consequently, decreasing the final grain size [15,16]. The maximum cooling rate that enabled the production of fine equiaxed ferrite grains was peak temperature or holding time dependent and seems to be mainly influenced by Mn content. The following sequence was established after 1 second holding time RTA cycles: Max. cooling rate → HSLA2, DC04, ISO1 (>200ºC/s) > HSLA1 (200º/s) > CM, HSIF (10ºC/s) Mn is known to retard the onset and kinetics of γ-α transformation, enhancing the formation of non-equilibrium phases [7,16]. This is attributed to segregation of Mn atoms along the α⁄γ phase boundary causing a strong solute drag effect. Conclusions Very fine ferrite grain size (2 µm) can be produced by rapid trans-critical annealing of Ti-Nb microalloyed and non microalloyed low-C steels using cooling rates of about 200ºC/s. However, such cooling conditions promote undesirable non-equiaxed structures when the Mn amount is larger than 0.75%, especially after long holding times at the peak temperature. Such behavior can be prevented by reducing the peak temperature to the intercritical domain. Neither trans-critical nor intercritical rapid annealings were successful with ultra low-C HSIF steels. The finest grains are produced when α recrystallization and α-γ transformation are concurrent (HSLA1, HSLA2 and CM). Such interaction inhibits grain growth of recrystallized α grains before their transformation to γ and promotes nucleation of austenite in the deformed ferrite matrix, both effects leading to finer γ structure. The strongest interaction takes place if recrystallization is effectively hindered (by using rapid heating rates, Nb in solid solution and fine Ti precipitates) and if the steel has low transformation temperature (Ac1). Despite its refining of the α recrystallized grain size, an increment of cold reduction has little influence on the final α grain size because it accelerates recrystallization. References [1] A.A. Howe et al.: “Ultra-fine Steel Flat Products” ECSC final report, 7210.PR/167 (2000). [2] M.R. Hickson, P.D. Hodgson et al.: Metall. Mater. Transactions A 33A (2002) 1019-1026. [3] R. Priestner and A.K. Ibraheem: Mater. Sci. Tech. 16 (2000) 1267-1272 [4] R.C. Hudd et al: “The Ultra-Rapid Heat Treatment of Low Carbon Strip” ECSC final report, 7210.MB/818,203,819 (1997). [5] J. Leinonen: “New Process Improves Properties”. Adv. Mater. Processes 8 (1999) H29-30. [6] M A.C.C. Reis, L. Kestens et al: ISIJ Int. 43 (2003) 1260-1267. [7] N. Yoshinaga, K. Ushioda, A. Itami and O. Akisue: ISIJ Int. 34 (1994) 33-42. [8] J.J. Lebrun, G. Maeder and P. Parniere: Proc. ICOTOM VI (1981) 787. [9] D. Muljono, M. Ferry and D.P. Dunne: Mater. Sci. Eng. A 303 (2001) 90-99. [10] J. Huang, M. Militzer et al.: Metall. Mater. Transactions A 35A (2004) 3363-3375. [11] R.E. Hook and H. Nyo: Metall. Transactions A 6A (1975) 1443-1451. [12] T. Gladman, I.D. McIvor and F.B. Pickering: J. Iron and Steel Inst. May (1971) 380-390. [13] J. Stockemer and P.V. Brande: Metall. Mater. Transactions A 34A (2003) 1341-1348. [14] J.L. Bocos, I. Gutierrez et al.: Metall. Mater. Transactions A 34A (2003) 827-839. [15] M. Militzer, R. Pandi and E.B. Hawbolt: Metall. Mater. Transactions A 27A (1996) 1547-1556. [16] M. Thompson, M. Ferry and P.A. Manohar: ISIJ Int. 41 (2001) 891-899.