Multipass hot deformation behaviour of high Al and Al–Nb steels

Materials Science & Engineering A 600 (2014) 37–46

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Multipass hot deformation behaviour of high Al and Al–Nb steels Z. Aretxabaleta, B. Pereda, B. López n CEIT and TECNUN, University of Navarra, P1 de M. Lardizabal, 15, 20018 Donostia-San Sebastián, Basque Country, Spain

art ic l e i nf o

a b s t r a c t

Article history: Received 25 November 2013 Received in revised form 15 January 2014 Accepted 2 February 2014 Available online 8 February 2014

The effect of Nb (up to 0.07 wt%) and high Al content (up to 2 wt%) on the multipass deformation behaviour of steels with 0.2% C and 2% Mn was studied with the aid of hot torsion simulations. From the tests, the critical Non-Recrystallisation (Tnr), Recrystallisation Limit and Stop Temperatures (RLT and RST) and the ferrite phase transformation start temperature (Ar3) were determined. It was observed that an increase in Al content from 1% to 2% or a microalloying addition of 0.03% Nb to 1% Al steel both led to a significant increase in the recrystallisation critical temperatures, which is greater than 100 1C in the case of the Tnr. However, the value of the Tnr was not affected when 0.03% or 0.07% Nb was added to the 2% Al steel. Specimens quenched after several deformation passes were examined by optical and TEM means in order to study the interaction between static recrystallisation, strain-induced precipitation and γ-α phase transformation, and determine the mechanisms leading to strain accumulation in the steels investigated. The results suggest that for the 1% Al steels, the Al and Nb solute drag effect is the main mechanism leading to the increase in the critical recrystallisation temperatures, while for the 2% Al steels the occurrence of γ-α phase transformation at temperatures close to the Tnr is the main mechanism involved in softening retardation, with a limited contribution of Nb. However, γ-α phase transformation taking place at temperatures close to the Tnr resulted in a loss of hot ductility, which can limit the industrial applicability of the 2% Al steels. & 2014 Elsevier B.V. All rights reserved.

Keywords: High Al steels Non-recrystallisation temperature Thermomechanical processing Microalloyed steels

1. Introduction In the past decades, TRIP-assisted steels have attracted increasing interest due to the good strength and ductility balance and excellent formability that they exhibit. These steels present a multiphase room microstructure consisting of ferrite and bainite and a smaller volume fraction of high-C retained austenite (
10%) which transforms into martensite during deformation, leading to high uniform elongation levels. Nowadays, TRIP-assisted steels are produced directly in hot strip rolling mills or as cold-rolled material in continuous annealing lines. In order to obtain the target multiphase microstructure, complex cooling sequences on the runout table and in the coiler or two-step heat treatments after cold rolling must be applied. In addition, these steels usually present relatively high Si addition levels (
1.5%), which helps to stabilise the required austenite fraction, while microalloying with Nb has been suggested as a tool to further increase the steel strength [1,2]. However, high Si levels impair the steel surface

n Correspondence to: Materials Science and Metallurgical Engineering, Department of Materials, CEIT and TECNUN, University of Navarra, P1 de M. Lardizabal, 15, 20018 Donostia-San Sebastián, Basque Country, Spain. Tel.: þ 34 943 21 28 00; fax: þ34 943 21 30 76. E-mail address: [email protected] (B. López).

http://dx.doi.org/10.1016/j.msea.2014.02.001 0921-5093 & 2014 Elsevier B.V. All rights reserved.

quality and galvanizability, and as a consequence partial or complete substitution of Si by other elements such as Al has been proposed [3,4]. Studies on the microstructures and mechanical properties of CMnAl and CMnAlSi steels show that they are comparable to conventional Si alloyed steels [4–6]. However, much less information can be found regarding the hot working behaviour of CMnAl and CMnAlSi steels. Nevertheless, understanding the microstructural evolution of TRIP-aided steels during hot rolling can be a useful tool for improving their mechanical behaviour. Several works have shown that the application of thermomechanical treatments such as Conventional Controlled Rolling, in which the steel is deformed below the Non-Recrystallisation Temperature (Tnr) and results in pancaked austenite microstructures, can lead to microstructural refinement and increased retained austenite volume fraction and stability in TRIP-assisted steels [1,7]. Optimised rolling strategies can also allow the target mechanical property levels to be achieved without the application of additional heat treatments [8]. On the other hand, early Nb(C,N) precipitation during hot rolling can lead to a loss of microalloying efficiency [2]. Most of the works mentioned above were carried out with Si alloyed steels. However, there is evidence that the high Al addition levels (
1.5%) usually employed in TRIP-assisted steels can modify the austenite hot working behaviour. Poliak and Siciliano [9] found

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that high Al addition levels lead to an increase in the stress levels and accelerate the dynamic softening kinetics during hot deformation. Suikkanen et al. [10] studied the static softening behaviour of several types of High-strength TRIP-aided steels and determined that high Al addition levels result in a higher retardation effect than Si on the static softening kinetics. In a previous work we observed that 1% Al addition produced a significant retardation in the softening kinetics due to the solute drag effect, while 2% Al addition promoted γ-α phase transformation at temperatures below 1000 1C, which led to a high retardation of the austenite softening kinetics [11]. However, there is no information regarding how high Al addition levels can affect the strain accumulation potential of TRIP-aided steels during multipass deformation schedules, solely or in combination with Nb. In this work the multipass deformation behaviour of high Al steels with and without Nb microadditions is analysed. The critical Non-Recrystallisation (Tnr), Recrystallisation Limit (RLT) and Recrystallisation Stop (RST) temperatures have been determined, and the microstructural events occurring during multipass deformation have been investigated.

2. Experimental techniques The chemical composition of the steels investigated is listed in Table 1. A set of steels with 0.2 wt% C and 2 wt% Mn, typical levels of TRIP-assisted steels, were produced in a laboratory unit. In order to investigate the effect of high Al contents, a base CMn steel and two high Al addition levels of 1% and 2% were considered. In addition, some of the Al steels were microalloyed with Nb in the 0.03–0.07% range. In order to investigate the deformation behaviour of the steels, multipass torsion tests were carried out. The geometry of the torsion specimens was a reduced central gauge section, 16.5 mm in length and 7.5 mm in diameter. In all the cases, before the tests, the specimens were soaked in the induction furnace attached to the torsion machine at a temperature of 1250 1C for 15 min. In the tests, after soaking, 20 deformation passes were applied in the 1180–800 1C temperature range at continuous cooling conditions. The temperature drop between passes was 20 1C. In any particular test, the pass strain, strain rate and interpass time were held constant. A strain per pass of ε ¼0.3, a strain rate of ε_ ¼1 s  1 and interpass times of 5, 30 and 100 s were employed in the tests. It must be mentioned that in some of the tests with the 2% Al steels the specimens broke during the test, and as a result the deformation schedule could not be completed. In addition, several specimens were quenched at different stages during deformation in order to investigate the microstructural evolution of the steel during the test. The specimens were examined through Light Optical Microscopy (LOM) and Transmission Electron Microscopy (TEM). A metallographic analysis of the quenched samples was carried out on a section corresponding to 0.9 of the outer radius of the torsion specimen, R, also known as the sub-surface section [12]. For optical microscopy analysis, the specimens were etched either Table 1 Composition of the steels investigated (wt%). Steel

C

Si

Mn

P

S

Al

N

Nb

Ti

C2Mn2 Al1 Al1Nb3 Al2 Al2Nb3 Al2Nb7

0.195 0.210 0.205 0.200 0.205 0.220

0.011 0.010 0.021 0.020 0.010 0.020

1.98 2.04 1.97 1.99 2.03 2.08

0.019 0.018 0.018 0.018 0.018 0.020

0.001 0.001 0.001 0.001 0.001 0.001

0.03 1.06 0.88 2.01 2.02 2.11

0.005 0.005 0.004 0.005 0.005 0.007

0.001 0.001 0.028 0.001 0.030 0.071

0.005 0.001 0.001 0.001 0.001 0.001

with a picric-based etchant in order to reveal the prior austenite grain boundaries or with a 2% Nital solution in order to identify the presence of ferrite in the microstructure. In some of the specimens, the austenite grain sizes were measured with the help of the Leica Qwin v.5.3.1 Image Analysis software in terms of the mean equivalent diameter parameter. From the quenched specimens carbon extraction replicas were prepared and examined in a JEOL 2010 Transmission Electron Microscopy (TEM) operated at 200 KV with a LaB6 filament equipped with an Energy Dispersive Spectrometry (EDS) analysis system. From these replicas, the average diameter of the precipitates was measured with the aid of the DigitalMicrograph software.

3. Results 3.1. Initial microstructure The initial microstructure was analysed from specimens quenched after the soaking treatment (T ¼1250 1C, 15 min). For example, Fig. 1 shows some micrographs of the austenite microstructure obtained for the Al1Nb3, Al2Nb3 and Al2Nb7 steels after reheating. In the three cases a homogeneous austenite microstructure is observed, which appears slightly refined for the Al2Nb steels. On the other hand, as already mentioned in a previous work [11], for the Al2 steel a duplex microstructure with a small amount of ferrite (fv
2%) was obtained after reheating. Similarly, after etching the Al2Nb steels with 2% Nital, some small ferrite grains could be detected within the austenite matrix. These can also be identified after picric acid etching (see black arrows in Fig. 1 (b) and (c)). Nevertheless, as the figure illustrates, in the case of the Al2Nb steels the amount of ferrite present was even smaller, less than 1% according to the metallographic measurements. The presence of this ferrite fraction after reheating in these steels can be attributed to the high Al content, 2 wt%, which has a strong ferrite stabiliser effect [3,4]. This trend is also in good agreement with the results provided by the Thermo-Calc software [13], which predicts that 1% Al leads to an increase in the Ae3 temperature from 780 to 900 1C, while 2% Al results in a further increase up to 1030 1C. However, it should be noted that the soaking temperature employed, 1250 1C, is significantly higher than the Ae3 temperature predicted by the software. This suggests that the time employed in the reheating treatment may be too short to reach equilibrium conditions. From the quenched specimens the austenite grain sizes were measured in terms of Equivalent Diameter (EQD) with the help of a quantitative image analyser. The mean EQD values together with the ferrite fraction detected in each case (if present) are summarised in Table 2. It should be mentioned that the Thermo-Calc software also predicts the presence of the AlN phase in the range of temperatures of interest. This could be of interest, first, because the presence of precipitates can affect the softening kinetics of the steels investigated, and second, because they could lead to a reduced amount of Al in the solid solution. However, according to the software, the amount of AlN which can be formed is limited by the N (50 ppm) content in these steels, and as a result, the amount of Al which can be engaged in the AlN is the same, and it is very small (o0.01 wt%) for all the steels investigated. In addition, the software predicts that for the reheating temperature employed in this work, almost all the AlN that can be formed remains undissolved during soaking. An examination of the microstructure of the high Al alloys by SEM means confirmed the presence of coarse Al and N bearing particles in the high Al steels, larger than 1 μm. It has been reported that relatively coarse particles can act as nucleation sites for new recrystallised grains.

Z. Aretxabaleta et al. / Materials Science & Engineering A 600 (2014) 37–46

39

Fig. 1. Examples of the microstructures obtained after reheating at 1250 1C for the Al1Nb3, Al2Nb3 and Al2Nb7 steels.

Table 2 Initial austenite grain sizes measured from the quenched specimens, and ferrite fraction, if present. Steel

D0 (μm)

α fraction (%)

C2Mn2 Al1 Al1Nb3 Al2 Al2Nb3 Al2Nb7

697 4 1007 3 1027 4 657 2 657 2 567 2

– – –
2% o1% o1%

On the other hand, small particles can retard softening through the Zener pinning effect, this effect being enhanced by increasing the precipitate fraction and decreasing the precipitate size [14]. Due to their scarce presence and large size these particles are not expected to significantly affect the softening behaviour of the steels investigated. Therefore, for these steels only an effect of Al in solid solution can be expected. Moreover all the N present in the steels is expected to be engaged by the AlN precipitates, and therefore, in the case of the Nb microalloyed steels, all the straininduced precipitates that can be formed during deformation must be carbide type, which are expected to precipitate at lower temperatures than nitrides or carbonitrides. Finally, for the Nb microalloyed steels (Al1Nb3, Al2Nb3 and Al2Nb7), replicas extracted from these specimens were examined in the TEM in order to study the precipitation state before deformation. In each case a replica area of approximately 0.25 mm2 was examined. In the case of Al1Nb3, no evidence of precipitates was found. In the case of Al2Nb3 the precipitation was very scarce (only 58 particles were measured, the average size being Dprec ¼93 77 nm). Therefore, for these steels it can be considered that nearly all the Nb is put into the solution during reheating. On the other hand, in the case of Al2Nb7, a significant amount of relatively coarse precipitates with an average size of 12676 nm was detected in the replicas. Some examples of the undissolved precipitates extracted from the Al2Nb specimens are shown in Fig. 2. These findings agree well with the results of the equilibrium calculations performed with different solubility products found in the literature, which are summarised in Table 3. For the two 0.03% Nb steels (Al1Nb3, Al2Nb3), all the solubility products considered, except that given by Palmiere et al. [16], predict equilibrium dissolution temperatures lower than 1250 1C, whereas for the Al2Nb7 steel all the solubility products predict solution temperatures above 1250 1C. Only the Thermo-Calc software predicts a slightly lower temperature of 1236 1C.

3.2. Mean flow stress Fig. 3(a) shows the stress–strain curves obtained in a test carried out with the Al2Nb3 steel with an interpass time of tip ¼100 s. From the curves, the Mean Flow Stress (MFS), defined as the area under the stress–strain curve divided by the pass strain, was calculated for each deformation pass by numerical integration and plotted against the temperature and is shown in Fig. 3(b). Four different regions can be distinguished in the figure: Region I, where complete recrystallisation between passes takes place and the stress increase from pass to pass is only due to the temperature drop; Region II, where recrystallisation between passes is inhibited and as a result the hardening level is increased; Region III, where some degree of softening due to austenite to ferrite transformation occurs, and Region IV, where ferrite hardening starts. Following standard procedure [18], the Tnr was determined as the intersection between the regression lines of the points corresponding to Regions I and II in Fig. 3(b) [18]. The γ-α phase transformation start (Ar3) and finish (Ar1) temperatures were also determined from the MFS data. The MFS vs. temperature plots obtained for the Al1Nb3 steel at different interpass times are shown in Fig. 4(a), while the values obtained at the same deformation conditions (tip ¼ 30 s) for all the steels investigated are compared in Fig. 4(b). Fig. 4(a) shows that for the Al1Nb3 steel, decreasing the interpass time leads in all cases to higher hardening levels and to an increase in the Tnr value. Fig. 4(b) indicates that steel composition also affects the nonrecrystallisation temperature value, although the effect is complex and depends on the alloying level. As the figure shows, 1% Al addition to the C2Mn2 base steel results in higher stress and hardening levels and a slight Tnr increase of
30 1C, from 890 to 923 1C. However, 2% Al addition leads to a significantly higher Tnr increment of
175 1C, up to 1065 1C. The effect of Nb is also complex; while microalloying the 1% Al steel with 0.03% Nb results in a Tnr increment larger than 100 1C, almost no effect is observed when microalloying the 2% Al steels with 0.03 or 0.07% Nb. The values of the Tnr obtained in all the tests investigated have been summarised in Fig. 5 and Table 4. It can be observed that this temperature is a function of deformation conditions and also depends strongly on steel composition, although as shown above, the effect varies with the alloying content. As can be seen from the figure, while 1% Al addition to the C2Mn2 results in a Tnr increment of
30 1C for the longest interpass times (tip ¼30– 100 s), 2% Al addition leads to a significantly higher increase, ranging from
125 1C for tip ¼ 5 s to
200 1C for tip ¼100 s. The addition of 0.03% Nb to the Al1 steel also leads to an important increase in Tnr, in the range of 106–123 1C, even though this is

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Fig. 2. Coarse undissolved Nb precipitates found in the replicas corresponding to (a) Al2Nb3 and (b) Al2Nb7 specimens quenched after soaking.

Irvine [15]

C2Mn2Al1Nb3 1233 C2Mn2Al2Nb3 1241 C2Mn2Al2Nb7 1376

Koyama [17]

Thermo-Calc (TCFE6) [13]

1285 1295 1443

1129 1136 1281

1127 1132 1236

Region IV Region III Region II

tip=100 s

50

0

1

2

3

4

5

Ar3 =920ºC

200

100

Tnr =1010ºC

50

Tnr=1029ºC

Al1Nb3

0 1190

1090

990

6

Region IV

C2Mn2 Al1 Al1Nb3

200

890

790

Al2 Al2Nb3 Al2Nb7

Tnr Al2-Al2Nb≈1065ºC 150 100

Tnr=1022ºC

Tnr C2Mn2=890ºC

50

Tnr Al1Nb3=1029ºC Tnr Al1=923ºC

0

1190

1090

990

890

790

T (ºC)

150 100

Fig. 4. (a) MFS plots obtained for the Al1Nb3 steel at different interpass times . (b) Effect of steel composition on the MFS for tests carried out with tip ¼ 30 s.

Ar1=840ºC

50 0

Ar3 =800ºC

Tnr=1055ºC

150

250 100

Strain

Mean Flow Stress (MPa)

200

Region I

150

0

Ar3 =820ºC

tip=5s tip=30s tip=100s

T (ºC)

Mean Flow Stress (MPa)

Stress (MPa)

200

Palmiere [16]

Region I 1150

1050

Region III Region II 950

850

C2Mn2 Al1 Al1Nb3

1150 750

T (ºC) Fig. 3. (a) Stress–strain curves obtained in a multipass torsion (Al2Nb3 steel, tip ¼ 100 s, ε¼ 0.3) and (b) corresponding mean flow stress values plotted against the deformation temperature.

lower than that due to 2% Al addition. On the other hand, microalloying the Al2 steel with 0.03 or 0.07% Nb does not influence the Tnr values. Only for the longest interpass time (tip ¼100 s) is a Tnr decrease of
40 1C observed. The figure also indicates that for the C2Mn2, Al1 and Al1Nb3 steels the Tnr decreases with increasing interpass time, whereas in the case of

Tnr (ºC)

Steel

250

Mean Flow Stress (MPa)

Table 3 Calculated dissolution temperatures (1C) of Nb precipitates for the Al1Nb3, Al2Nb3 and Al2Nb7 steels.

Al2 Al2Nb3 Al2Nb7

1050 950 850

0

50 100 Interpass time (s)

Fig. 5. Non-recrystallisation temperatures obtained for the steels investigated at the different deformation conditions.

Z. Aretxabaleta et al. / Materials Science & Engineering A 600 (2014) 37–46

Table 4 Tnr, RLT, RST, Ar1 and Ar3 temperatures determined for the steels investigated. The specimens which broke during the tests are also indicated in the table. tip (s)

Tnr (1C)

RLT (1C)

RST (1C)

Ar3 (1C)

Ar1 (1C)

Specimen broken during test

o 800

– – –

5 30 100

930 890 862

958 901 873

870 – –

o 800

5 30 100

949 923 887

1060 939 917

– – –

o 800

Al2

5 30 100

1056 1067 1064

– 1067 1064

950 – –

920 940 920

820 – 820

– X –

Al1Nb3

5 30 100

1055 1029 1010

– 1060 1016

976 951 951

820 820 800

o 800

– – –

Al2Nb3

5 30 100

1061 1066 1022

– 1107 1030

997 970 –

920 940 920

– – 840

X X –

Al2Nb7

5 30 100

1066 1065 1028

– 1101 1041

1019 980 –

920 940 920

– – 840

X X

C2Mn2

Al1

0.3

o 800

RLT=939ºC RLT=901ºC Tnr=890ºC Tnr=923ºC

100 80

Tnr Al2~RLT=1067ºC 60 RLT=1060ºC

FS (%)

ε

– – –

Tnr=1029ºC

40 20 0 1200

C2Mn2 Al1 Al2 Al1Nb3

1100

RST=951ºC

1000 T (ºC)

900

Tnr≈RLT=1067ºC

100

800

RLT=901ºC

80 FS (%)

Steel

41

RLT≈1100ºC

60 40 20

the Al2 steel this temperature remains approximately constant in the range of interpass times studied. For the Al2Nb steels the Tnr only decreases at the longest interpass time (tip ¼100 s).

Tnr=890ºC

Tnr≈1065ºC

0 1200

C2Mn2 Al2 Al2Nb3 Al2Nb7

1100

RST=980ºC-970ºC

1000 T (ºC)

900

800

3.3. Anisothermal fractional softening RLT=1101ºC

FS ¼

sim  siyþ 1 ðsi0 =si0þ 1 Þ sim  si0

ð1Þ

In the equation, si0 and si0þ 1 are the yield stresses of a fully recrystallised material for the i-th (at temperature Ti) and i þ1-st (at temperature Ti þ 1) passes, respectively, sim is the maximum level of stress achieved before unloading at the i-th pass, and finally siyþ 1 is the yield stress at the iþ 1-st pass. The sim as well as the siyþ 1 values are determined from pass to pass flow curves, while si0 and si0þ 1 values are calculated by the relationship obtained with the yield stress data measured in the stress–strain curves corresponding to the range of complete recrystallisation (initial passes). The yield stresses are determined through the 2% offset method. Fig. 6(a) and (b) shows the values obtained for all the steels at the same deformation conditions (tip ¼30 s), while in Fig. 6(c) the values determined for the Al2Nb7 steel at different interpass times are plotted. The figure shows that the data corresponding to each test can be approximated by three different linear slope segments. The temperatures at which they intersect are denoted as Recrystallisation Limit Temperature (RLT) and Recrystallisation Stop Temperature (RST). RLT is theoretically defined as the lowest temperature above which recrystallisation between passes is complete (criteria of 85% or 95% recrystallisation are used) and RST represents the highest temperature at which recrystallisation is completely absent (usually the 5% recrystallised fraction is taken; the higher softening level observed in the graphs should

100

Al2Nb7 RLT=1041ºC Tnr=1028ºC

80 FS (%)

In order to further investigate the strain accumulation behaviour of the different steels, the interpass softening fraction under anisothermal conditions was also calculated from the strain–stress data, using the following equation proposed by Liu and Akben [19]:

Tnr=1065ºC

60 40 20

Tnr=1066º RST=1019ºC RST=980ºC

0 1200

1100

1000 T (ºC)

tip=5s tip=30s tip=100s

900

800

Fig. 6. (a and b) Anisothermal fractional softening obtained for all the steels investigated for tip ¼30 s and (c) effect of tip on the softening for the Al2Nb7 steel.

be related to the recovery contribution). The Tnr values have also been indicated in the figures. It should be noted that in some of the tests the RST temperature could not be determined since, as mentioned above, the specimen broke before the test was finished. The RLT and RST temperature values determined for the different steels and deformation conditions are listed in Table 4. Fig. 6 (a) and (b) shows that the evolution of the interpass softening fraction depends on steel composition. For the C2Mn2 and Al1 steels, high interpass softening levels are obtained during the first
12–14 deformation passes. At temperatures of
900 1C for the C2Mn2 steel and of
940 1C for the Al1 steel, the softening values begin to fall more rapidly, indicating that strain accumulation processes are starting. For the Al2 and Al1Nb3 steels, the temperature at which strain accumulation initiates is significantly higher, at
1060 1C. As shown in Fig. 6(b), for the two Al2Nb steels, very similar softening behaviour is observed, with RLT

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values of
1100 1C, which are slightly higher than for the Al2 steel. Fig. 6(c) indicates that interpass time also affects the critical temperatures; longer interpass times lead to lower RLT temperatures and higher softening levels. Moreover, for the shortest interpass time (tip ¼5 s) the softening levels are below 80% during the overall deformation schedule. This denotes that for this steel, at tip ¼ 5 s strain accumulation is taking place from the beginning of deformation, thus starting at temperatures significantly higher than the Tnr. It can also be observed that in all the cases the Tnr is between the RLT and RST temperatures and corresponds to softening levels which range between
50% and 90%. 3.4. Ar1 and Ar3 temperatures The Ar1 and Ar3 temperatures determined from the tests are summarised in Table 4. The tests in which the specimen broke are indicated in the table. From the data it can be observed that for the C2Mn2 and Al1 steels, the Ar1 and Ar3 are below the test finish temperature, 800 1C, for all the conditions investigated. For the Al1Nb3 steel, Ar3 values ranged from 800 to 820 1C. Although Nb is known to retard the equilibrium austenite to ferrite phase transformation, this slight increase in the Ar3 compared to the Al1 and C2Mn2 steels can be related to a higher degree of strain accumulation taking place in the tests on Al1Nb3, which is known to accelerate phase transformation. Finally, for all the 2% Al steels significantly higher Ar3 temperatures ranging from 920 to 940 1C were determined for all the conditions investigated. For Ar1 temperature values ranging from 820 to 840 1C were calculated in all the cases where the specimen did not break during the test. Finally, it is interesting to note from the table that while in the case of the C2Mn2 and 1% Al steels all the tests could be completed, in more than half of the tests performed with the 2% Al steels the specimen broke during the test. This denotes a significant loss of hot ductility for these high Al steels. For the three steels all the specimens tested at tip ¼30 s failed during the test. So did the specimens tested at tip ¼5 s for Al2Nb3 and Al2Nb7 steels, however, the specimens tested at tip ¼100 s did not break in any case. The results seem to indicate that hot ductility behaviour improves for longer interpass times.

4. Discussion During austenite multipass hot deformation, strain accumulation start is related to the inhibition of the softening processes which take place during the interpass time between deformation passes. In the case of Nb microalloyed steels, both the solute drag effect due to Nb in solid solution and the Zener pinning drag force due to Nb(C,N) strain-induced precipitation can retard or even

Al2, Tquench =1020ºC t ip=5 s, Tnr =1056ºC

stop austenite recrystallisation. As a result, both mechanisms can lead to an increase in the recrystallisation critical temperatures, depending on steel composition and deformation conditions [18,20]. The data shown in Figs. 5, 6a and b indicate that the addition of 1% and 2% Al to the C2Mn2 steel also leads to an increase in the recrystallisation critical temperatures. For the 1% Al alloyed steel, the increase in the recrystallisation temperatures compared to the base C2Mn2 material can be attributed to the retarding effect exerted by Al in the solid solution in austenite recrystallisation. This agrees well with the results of other works [10,11] where it was observed that Al in solid solution resulted in a retardation of the isothermal austenite softening kinetics. It also correlates well with the decreasing trend of the Tnr with the interpass time observed in Fig. 5 for both C2Mn2 and Al1 steels. This behaviour is usually associated with those conditions where the Tnr value is mainly determined by a solute drag effect since increasing the interpass time in the absence of strain-induced precipitation allows higher interpass softening levels to be reached. For the 1% Al steel, strain accumulation is further enhanced by 0.03% Nb addition, resulting in a significant increase in the Tnr. On the other hand, for the 2% Al steels (Al2, Al2Nb3, Al2Nb7) the increase in the recrystallisation critical temperatures is significantly higher (Tnr increment
125 1C–200 1C for the Al2 steel vs.
30 1C for the Al1 steel) and almost independent of the Nb content. In order to investigate this behaviour, several tests were interrupted at different stages during multipass deformation and the specimens were quenched for microstructure analysis. Examples of the microstructure observed in specimens quenched after two passes below the Tnr in the 2% Al steels (2% Nital etching) are shown in Fig. 7. In the micrographs the dark regions correspond to martensite, product of the quenched austenite, and the white regions to ferrite. It can be observed that although the ferrite fraction tends to be clearly higher for the tests carried out at the longest interpass times (Fig. 7b and c), in the three conditions the amount of ferrite is higher than the 2% volume fraction measured in the initial microstructure. This indicates that for the 2% Al steels γ-α phase transformation is initiated during the multipass torsion tests at temperatures close to the Tnr. The presence of ferrite at these high temperatures would explain the low hot ductility shown by these steels, leading some specimens to break during deformation at relatively high temperatures. The reason why these steels present such high recrystallisation critical temperatures seems to be related to this early appearance of ferrite. In a previous work it was observed for the Al2 steel that after deformation and isothermal holding γ-α phase transformation took place at temperatures below 1000 1C, resulting in a high delay of the softening kinetics [11]. Since both ferrite and recrystallised grains share the prior austenite grain boundaries as preferential nucleation sites,

Al2Nb3, T quench =980ºC t ip=100 s, Tnr =1022ºC

Al2Nb7, Tquench =1020ºC tip=30 s, Tnr =1065ºC

Fig. 7. Microstructures obtained for the Al2 steels quenched two passes below the Tnr.

Z. Aretxabaleta et al. / Materials Science & Engineering A 600 (2014) 37–46

this retardation in the mechanical softening was attributed to recrystallisation prevention when phase transformation was previously initiated. It should be noted that the temperature at which γ-α phase transformation was first observed after deformation at isothermal conditions (E1000 1C) is in the range of the Tnr values determined here for the three 2% Al steels in the different conditions investigated (between 1022 and 1067 1C (Table 4)) and close to the Ae3 ¼ 1030 1C predicted by Thermo-Calc [13]. Therefore, it could be thought that a similar effect occurs during multipass deformation under cooling, that is, when the temperature decreases sufficiently phase transformation starts and prevents austenite recrystallisation from taking place, and, as a result, the non-transformed austenite accumulates deformation. The hardening produced by this retained strain may counteract the expected softening due to the presence of ferrite in the microstructure, at least until a certain ferrite fraction is developed. It is interesting to note that although for the 2% Al steels the initiation of the γ-α phase transformation has been identified at temperatures that are close to the Tnr, the Ar3 temperatures determined mechanically (920–940 1C) are significantly lower. Examples of the microstructures obtained for the Al2 steel in specimens quenched at the Ar3 at different test conditions are shown in Fig. 8. When comparing Figs. 7a and 8a, which correspond to the microstructures obtained at tip ¼ 5 s after quenching at the Tnr and Ar3, respectively, an increase in the ferrite fraction is clearly evident, although the transformation is far from being completed. At the Ar3 temperature a relatively low ferrite fraction is still present in the microstructure. At longer interpass times of tip ¼ 30 and tip ¼100 s (Fig. 8b and c), although the transformed fraction is higher, mainly at the longest interpass time, the microstructure observed at the Ar3 continues to be partially transformed. Regarding Fig. 8(c), which corresponds to tip ¼ 100 s, the amount of ferrite is significantly larger compared to the amount present at shorter interpass times (Fig. 8(a) and (b)), although the experimental Ar3 ¼ 920 1C is similar to that observed at the shortest time (Fig. 8(a)), where the amount of ferrite is significantly smaller. Jonas et al. have suggested that a minimum ferrite fraction must be developed in order to observe a decrease in stress during multipass torsion tests [21]. This is in good agreement with the results obtained in the present work, even if in this case the ferrite fraction changes with the test conditions. It must be remembered that the Ar3 is detected as the temperature at which the Mean Flow Stress starts to deviate from the hardening trend observed below the Tnr (see Figs. 3 and 4). However, the rate at which the MFS increases depends on the test interpass time. As shown in Fig. 4(a), higher hardening rates are observed for the tests carried out with shorter interpass times. This implies that in these tests a smaller ferrite fraction might be needed to lead to the softening level necessary for the Ar3 detection. On the other hand, the observed behaviour could also be explained by the occurrence

tip=5 s, Tquench =A r3 =920ºC

43

of ferrite hardening phenomena that could counteract the softening associated with this phase. As ferrite fraction increases, the applied strain can be distributed more homogeneously between both austenite and ferrite phases, thus giving rise to the ferrite phase to strengthen as well. From the results it can be said that in the Al2Nb steels, γ-α phase transformation plays a very important role in strain accumulation (softening retardation), which seems to be independent of the Nb content (no effect of Nb on the critical recrystallisation temperatures is observed for the 2% Al steels). The softening delay associated with Nb microalloying arises from the solute drag effect due to dissolved Nb and the pinning effect produced by fine Nb(C, N) precipitates formed during deformation. In order to investigate the contribution of precipitation in the present steels, replicas extracted from Al2Nb3 (tip ¼100 s) and Al2Nb7 (tip ¼30 s) specimens quenched two passes below the Tnr were examined. In both cases, some Nb precipitates were identified both in ferrite and martensite (previously austenite) regions. In the case of the Al2Nb3, the precipitates were scarce and difficult to find in both regions. Some examples of the particles observed in this specimen are shown in Fig. 9. As the figure shows, the size of the precipitates found in ferrite, with a mean value of dαprec ¼74 713 nm, was significantly coarser than those present in martensite, dγprec ¼47 77 nm. For Al2Nb7, the amount of precipitates in the replicas was higher, the mean precipitate size being similar in ferrite, dαprec ¼ 997 17 nm, and martensite, dγprec ¼98 78 nm. It should be noted that some of the Nb precipitates found in the latter steel could in fact correspond to particles not dissolved during the reheating treatment. These particles cannot be distinguished from the new precipitates formed during deformation and may contribute to the above measurements to some extent. Nevertheless, the size of the precipitates found in both steels is too large to have a significant effect in retarding austenite recrystallisation [22–24]. On the other hand, for the 1% Al steel, a significant increase in the recrystallisation critical temperatures is observed when 0.03% Nb is added. As mentioned above, Nb can increase the austenite recrystallisation critical temperatures due to solute drag effect, or to strain-induced precipitation. In order to investigate this, replicas extracted from a specimen quenched two deformation passes below the Tnr (tip ¼ 30 s) were also examined via TEM. A very limited number of particles, only 32, was detected, with a relatively large average size of dγprec ¼ 76 718 nm. Unfortunately, little information is available in the literature about Nb straininduced precipitation kinetics in high Al steels, although some evidence that Al can retard the kinetics of strain-induced precipitation was reported. Wang observed that for a 0.04% Nb microalloyed steel the addition of 0.08% Al resulted in a delay in the strain-induced precipitation start time by a factor from 1.25 to 2 at temperatures ranging from 900 to 950 1C [25]. Moreover, it

t ip =30 s, Tquench =Ar3 =940ºC

t ip =100 s, Tquench =Ar3 =920ºC

Fig. 8. Microstructures obtained for the Al2 steel after quenching at the Ar3 in tests carried out with different interpass times.

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Z. Aretxabaleta et al. / Materials Science & Engineering A 600 (2014) 37–46

Fig. 9. Example of the precipitates found in (a) martensite and (b) ferrite in a sample of the Al2Nb3 steel quenched at 980 1C after 10 passes þ100 s (Tnr ¼1022 1C).

1150

C2Mn2 Al2 Al1 Al2Nb3 Al1Nb3 Al2Nb7 Steel 1 Steel 2 Steel 3 Steel 4 Steel 5 Steel 6 Steel 7 Steel 8 Steel 9 Steel 10

1050 Tnr (ºC)

must be remembered that in the present steels nearly all N is engaged in the form of coarse undissolved AlN particles, and therefore, only niobium carbides can precipitate during deformation with the subsequent retardation of precipitation kinetics. The limited amount of relatively coarse Nb particles observed indicates that the Nb solute drag effect should be the main mechanism responsible for the Tnr increase in this steel at these conditions. This behaviour is also confirmed by the shape of the interpass softening evolution curve shown in Fig. 6a for the Al1Nb3 steel. Within the RLT–RST temperature interval, a gradual decrease of softening from pass to pass is observed for this steel. This gradual softening drop is usually related to those cases where solute drag effect appears as the dominant mechanism in retarding softening, in contrast to a sharper softening drop typically found when strain-induced precipitation occurs and becomes the main retarding mechanism [26]. Summarising, the experimental results indicate that in the case of 1% Al steels, the solute drag effect due to Nb and Al additions leads to an increased potential for strain accumulation, while for the 2% Al steels, early γ-α phase transformation highly contributes to this effect. This could be of interest for the application of conventional controlled rolling treatments. In these rolling strategies, the last deformation passes are applied at temperatures lower than the Tnr, which results in high strain accumulation levels in the austenite and refined final product microstructures after phase transformation. In the case of HSLA steels, usually microalloying with Nb is necessary in order to obtain the high recrystallisation critical temperatures needed for applying this thermomechanical treatment. In Fig. 10, several Tnr values reported in the literature for Nb microalloyed HSLA steels [18,20,27–30] are compared with those obtained in the present work for the TRIP-type steels under study. The composition of the steels and the deformation conditions applied in each case are summarised in Table 5. From the figure it can be observed that the non-recrystallisation temperatures obtained for the HSLA steels show a relatively large scatter. This is expected since the value of this temperature depends on deformation conditions and steel composition. For Nb microalloyed steels, the Tnr is affected by the amount of Nb in the solid solution and by Nb(C,N) strain-induced precipitation, and as a result, increasing the Nb, C or N contents can result in higher Tnr values. For example, when comparing the data corresponding to Steels 2 and 4, which were similarly deformed, it can be observed that for similar Nb content (
0.03%) increasing the C content from 0.05% to 0.1% leads to a slight increase in the Tnr. Steel 1, with a lower Nb (0.02%) but a higher C content (0.14%), also shows a Tnr value similar to Steels 2 and 4. On the other hand, for Steel 5 (0.06% Nb), a significantly larger Tnr was obtained (1074 1C vs. 1001 1C for Steel 4), even though it has a low carbon level of 0.06%. This suggests that under the

950

850

0

50 100 Interpass time (s)

150

Fig. 10. Values of Tnr determined for the steels studied in this work and for several HSLA steel types with composition listed in Table 5.

Table 5 Composition of the HSLA steels with data shown in Fig. 10. Steel Steel Steel Steel Steel Steel Steel Steel Steel Steel Steel

1 2 3 4 5 6 7 8 9 10

%C

% Mn

% Nb

%N

ε per pass

ε_ (s  1)

Ref.

0.14 0.05 0.05 0.1 0.06 0.07 0.08 0.09 0.08 0.12

1.2 0.3 1.6 1.4 1.0 0.6 0.9 1.35 1.5 0.4

0.019 0.029 0.029 0.035 0.06 0.027 0.035 0.040 0.041 0.05

0.004 0.004 0.005 0.005 0.010 – 0.005 0.007 0.005 0.005

0.3 0.3 0.3 0.3 0.3 0.4 0.35 0.2 0.3 0.3

1 1 1 1 1 2 2 3.6 2 2

[20]

[27] [28] [29] [30] [18]

conditions investigated, the Nb content seems to have a larger effect on the Tnr than C. For Steels 6, 7, 9 and 10, the Tnr values tend to be lower in general. Nevertheless, it must be pointed out that the test conditions applied to determine the Tnr in these steels involve higher strains and/or strain rates, both of which are reported to lead to a Tnr reduction when being increased [18,20]. It can also be noted that among these steels, Steel 10, with the higher C (0.12%) and Nb (0.05%) content, shows the highest Tnr levels. Finally, Steel 8 also shows a high Tnr value, which may be attributed to the high C (0.09%), Nb (0.05%) and N (0.007%) content of this steel. The figure illustrates that in almost all cases, the nonrecrystallisation temperatures reported for the HSLA steels are higher than the values determined in this work for the Al1 steel but lower than those obtained for Al1Nb3. For example, the Tnr corresponding to Steels 2 to 4, which were determined at deformation conditions similar to those employed in the present

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45

Fig. 11. RLT, Tnr, RST and Ar3 maps showing the different processing regimes for the steels investigated.

work, exceed by more than 50 1C the values obtained for the Al1 steel. However, the Al1Nb3 steel shows Tnr values 30 to 55 1C higher than those found in Steels 2, 3 and 4, although all of them show similar Nb levels (
0.03%). This could be attributed to the Al solute drag effect. However, it must also be remembered that the Al1Nb3 steel shows a C level (0.2%) that is significantly higher than that of HSLA steels, and this could lead to an enhanced contribution of strain-induced precipitation and therefore an increase in Tnr. It can also be observed that the 2% Al steels shows Tnr values that are in the range of those obtained for Steel 5, which has the highest Nb content (0.06%). Finally, although the Tnr gives an approximate indication of the strain accumulation potential of the steels investigated, it must be remembered that the softening results shown in Fig. 6 denote that strain accumulation may start at temperatures significantly higher than the Tnr. In order to compare the strain accumulation potential of the different steels, the critical temperatures (RLT, RST, Tnr) have been plotted in Fig. 11 as a function of the interpass time. From these graphs, the microstructural evolution regimes can be identified as a function of processing temperature: at temperatures above the RLT the steels recrystallise completely, between the RLT and the RST partial recrystallisation occurs and strain accumulation processes start, and below the RST recrystallisation is prevented and all the applied strain is retained. Although it was noted previously that for the 2% Al steels γ-α phase transformation may start at temperatures that are significantly higher than those determined mechanically, the Ar3 determined from the multipass torsion tests has also been included in the plots. A comparison of the data corresponding to the C2Mn2 and Al1 steels shows that the solute drag effect due to 1% Al addition has a small but noticeable effect in increasing the potential for strain accumulation, an effect that is enhanced for the shortest interpass times. The potential for strain accumulation is clearly increased when 0.03% Nb is added (Al1Nb3 steel). For this steel it can be observed that below 950 1C recrystallisation is completely prevented for all the interpass times studied, leading to a wide temperature processing window for strain accumulation. In this case the risk of having partially recrystallised austenite microstructures, which could lead to heterogeneities in the final product, would significantly decrease. On the other hand, as was shown above, precipitation of niobium during deformation is scarce, denoting that in addition to providing strain accumulation,

Nb may be left in solution after hot rolling, making it available for further precipitation after deformation. The figures also indicate that the temperatures for strain accumulation onset are also very high for the 2% Al steels. In this case, it seems that γ-α phase transformation starting at temperatures close to the Tnr significantly contributes to this behaviour. However, it must be remembered that more than half of the 2% Al specimens broke during the multipass tests, which indicates that the occurrence of γ-α phase transformation during deformation leads to a significant loss of ductility. This clearly limits the industrial applicability of the 2% Al steels. The formation of a thin film of ferrite at the austenite grain boundaries is well known to be one of the embrittlement mechanisms leading to the ductility drop when the deformation temperature decreases below Ae3 [31]. The strain concentration in this softer ferrite film promotes the formation of cracks at the interface between prior austenite grain boundaries and ferrite thin films. As the ferrite film thickens, the ductility recovers. This might explain why in the present multipass tests the specimens deformed with the longest interpass time (100 s) have not broken in any case. The higher ferrite fractions observed at these conditions compared to those found for shorter interpass times (5 and 30 s) could be the reason.

5. Conclusions - The addition of 1% or 2% Al to a plain CMn steel results in an increase of the Tnr, RLT and RST recrystallisation critical temperatures. However, the mechanisms contributing to this increment are strongly dependent on the Al addition level. For 1% Al addition, a slight increase in the recrystallisation critical temperatures is observed, which is related to the softening retardation produced by the Al solute drag effect. 2% Al addition results in a significantly greater increase in these critical temperatures. In this case, in addition to the solute drag effect of Al, the occurrence of γ-α phase transformation starting at temperatures close to the Tnr produces a significant delay of recrystallisation kinetics, and this significantly contributes to strain accumulation. - For the 1% Al steel, the Tnr was further increased by
100– 120 1C with a 0.03% Nb addition. This was related to the solute drag effect of Al and Nb, with no contribution of strain-induced

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precipitation. However, the Tnr was nearly unaffected by a 0.03% or 0.07% Nb addition to the 2% Al steel. - Temperature processing maps showing the different microstructural evolution regimes of the steels investigated were built. The maps indicate that 1% Al addition leads to an increased potential for strain accumulation during multipass deformation schedules, although this effect is clearly enhanced when microalloying the 1% Al steel with 0.03% Nb or increasing the Al content up to 2%. However, γ-α phase transformation taking place at temperatures close to the Tnr resulted in a loss of hot ductility, which limits the industrial applicability of the 2% Al steels.

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