On the Relationship between Thermomechanical Processing Parameters and Recrystallization Texture in AA1100 Aluminum Alloy

Abstract

In this study, 48 hot-rolling processing conditions were designed to investigate the influences of thermomechanical processing parameters on the recrystallization behavior and texture development. The hot-rolling experiments were conducted using the thermomechanical simulator Gleeble 3800 at temperatures of 275, 300, and 350 °C with strain rates of 5 and 90 ${\mathrm{s}}^{-1}$ up to 60 and 85% reduction. The microstructure and texture analysis were measured by using the EBSD technique on a large area. Experimental results show that the Cube component maintains a volume fraction between 10% and 20%, below the 40% recrystallization fraction, but the volume fraction of Cube significantly increases between 20% and 50% above the 40% recrystallization fraction. However, the fractions of Rotated Cube (RC) and Goss components remain below 10%.

1. Introduction

The hot-rolling texture of aluminum alloys without dynamic recrystallization is significantly different from their cold-rolling texture. The texture of hot-rolled commercial purity aluminum typically contains the Cube component and the β-fiber rolling orientations of brass (B), copper (C), and S. The dominant texture in recrystallized grains is the Cube component, whereas the β-fiber rolling orientations are dominant in deformed grains [1]. Hollinshead and Sheppard [2] reported that the B component appears at appreciable strengths at low reductions and increases with the reduction for AA 3004 aluminum alloy. Bate and Oscarsson [3] reported for hot-rolled Al–1% Mn–1.2% Mg alloy that C, S, and B have roughly similar densities, with the maximum near copper at a strain of 2, while S and B doubled in intensity, and copper decreased slightly at a strain of 5. Furthermore, the observation that a strong β fiber rolling texture forms during hot deformation is consistent with the experimental observations of Daaland et al. [4], Engler et al. [5], Panchanadeeswaran and Field [6], Vatne et al. [7], Maurice and Driver [8], Duckham et al. [9], and Bate et al. [10].

Bacroix et al. [11] reported for AA 1200 and AA 8979 alloys that the texture strength increases with temperature, and a significant increase in the B component occurs without a complete orthotropic symmetry with the increasing temperature. Samajdar et al. [12] reported that the volume fractions of the C and S components decrease with the increase in temperature or decrease in strain rate. Duckham et al. [9] attributed this strong β fiber rolling texture to the healing of the dislocations within the subgrains and the consumption of subgrains with scattered orientations during the dynamic recovery, which is generally enhanced at high deformation temperatures. Aryshenskii et al. [13] found that increasing the alloying component of magnesium resulted in an increase in Brass and a decrease in S and Copper during hot deformation of aluminum alloys.

The increase in the B component is attributed to the increase in the strain rate sensitivity index by Engler et al. [5], the additional activation of other non-octahedral slip systems at very high deformation temperatures by Le Hazif et al. [14], Le Hazif et al. [15], Maurice and Driver [8], and Bacroix and Jonas [16], and the differential dynamic grain growth by Bate et al. [10]. Bacroix and Jonas [16,17] applied full constraint (FC) and relaxed constraint (RC) models, considering cross-slip systems, to predict the hot-rolling texture. The concurrent activation of both slip and cross-slip systems leads to the presence of three main rolling components, namely, the C, S, and B orientations, in proportions that vary with the CRSS ratios of the different systems. In rolling, the activation of the {100}(110) systems leads to the prediction of a strong S component, whereas the slip on the {110} planes mainly produces a C component, and that on the {112} planes produces a B-type texture. Hirsch and Luecke [18] also used Taylor-type theories under full and relaxed constraints and found that the exact {011}<211> B orientation is found for a Sachs-type RC model in which the grain TD/RD shear is completely relaxed. However, non-octahedral slip has been observed at low strain rates < 10⁻³ s⁻¹ but not yet at high strain rates >10 s⁻¹, which is typical of industrial hot-rolling schedules. Sachs-type grain relaxation generally leads to unreasonably high strain rate incompatibilities between grains.

After discussing the increasing B component, we now discuss the recrystallized Cube component during hot rolling. Hollinshead and Sheppard [2] reported that the Cube component forms in small quantities during interpass reheat recrystallization for AA 3004 aluminum alloy. Maurice et al. [19] reported that the rotation of the {112}(111) C component is predicted toward the {110}(112) B and {123}(634) S components at high rate sensitivity and TD/RD shear values. Vatne et al. [7] reported that the recrystallized Cube component is enhanced by a high initial Cube fraction in many Cube bands, a low Zener–Hollomon parameter, and a high strain. Liu and Morris [20] reported for AA 5052 aluminum alloy that the initial orientations at elevated temperatures are more easily rotated to the β fiber between the B and S orientations than to the C orientation, whereas those at room temperature are more easily rotated to the S orientation than to the B and C orientations. Alvi et al. [21] reported that Cube recrystallized grains have a strong nucleation advantage at all annealing temperatures, while a growth rate advantage is observed for Cube recrystallized grains at low annealing temperatures. In addition, Cube-recrystallized grains preferentially grow into deformed S grains but grow most slowly into deformed B grains. Wang et al. [22] reported that the cold-rolled texture of Copper and S components is consumed and transformed into the Cube component after annealing.

Vatne et al. [7], Samajdar et al. [12], and Bate et al. [23] studied the stability of the Cube orientation during the hot deformation of aluminum alloys. Vatne et al. [7], Duckham et al. [9], and Bate et al. [23] reported that the stability of the Cube orientation increases with the decreasing Z value. Maurice and Driver [24] attributed the enhanced stability of the Cube orientation during hot deformation to the activation of non-octahedral slip systems. Gatti and Bhattacharjee [25] reported that homogenous recovery eliminates the nucleation advantage of the Cube regions originating from the preferentially recovered structure and significantly weakens the Cube component.

Although numerous experimental studies have been carried out in the past decades on recrystallization kinetics, the experimental materials and testing conditions are different, preventing the comparison of the number of results with the other results. In this work, we designed 12 hot-rolling process conditions to understand the effects of the thermomechanical process history on the recrystallization texture control. The hot-rolling conditions were experimentally simulated by using double-hit compression experiments as the thermal–mechanical simulator.

2. Experiments

2.1. Experimental Materials and Sample Preparation

The dimensions of the annealed AA1100 aluminum alloy provided by China Steel Corporation (Kaohsiung, Taiwan) were reduced to 25 mm × 15 mm × 38 mm for the hot-compression experiment. The corresponding chemical composition is shown in

Table 1. Chemical composition of AA1100 aluminum alloy (wt%).

Alloy	Si	Fe	Cu	Mn	Mg	Cr	Ti	Al
AA1100	0.17	0.38	0.09	0.004	0.002	0.001	0.03	Balance

. The double-hit compression experiment was conducted using the thermal–mechanical simulator Gleeble 3800 (Dynamic Systems Inc., Poestenkill, USA) at temperatures of 275 °C, 300 °C, and 350 °C, with low and high strain rates of 5 ${\mathrm{s}}^{-1}$ and 90 ${\mathrm{s}}^{-1}$, to the thickness of 6.5 mm (60% reduction) and 2.5 mm (85% reduction), respectively. The dimensions of the flat punch’s contact area are 42 mm × 10 mm. Figure 1 illustrates the geometry of the specimen after 60% reduction. After the hot compression, the specimens were isothermally annealed in a salt bath at the corresponding temperatures of 275 °C, 300 °C, and 350 °C for recrystallization. The above-mentioned processes of hot compression and annealing are illustrated in Figure 2. The thermomechanical processing parameters (TMPs), including temperature, strain rate, and reduction ratio, and the annealing parameters, including temperature and time, are listed in

Table 2. Thermomechanical processing parameters (TMPs) of temperature (T), strain rate ($\dot{\epsilon }$), and reduction ratio (ε) tested on a Gleeble thermomechanical simulator and annealing time (t) of AA1100 alloy.

Name	TMPs			Annealing Time t (s)
	T (°C)	$\dot{\mathit{\epsilon }}$ (s−1)	ε (%)
T275-1 (1)	275	5	60	0	1800	4800	10,800	28,800
T275-2 (2)		5	85
T275-3 (3)		90	60
T275-4 (4)		90	85
T300-1 (5)	300	5	60	0	300	900	3000	9000
T300-2 (6)		5	85
T300-3 (7)		90	60
T300-4 (8)		90	85
T350-1 (9)	350	5	60	0	30	150	300	900
T350-2 (10)		5	85
T350-3 (11)		90	60
T350-4 (12)		90	85

. A total of 48 specimens were prepared to systematically investigate the influence of thermomechanical processing parameters on the recrystallization behavior.

After thermomechanical processing, the middle area of the ND-RD section with a 10 mm width was first cut from all H-shaped test specimens, as shown in Figure 1, using the low-speed precision cutter Plusover S150 (Yu-Chang Instruments Co., Kaohsiung, Taiwan). Then, the cut specimens were rotated 90 degrees, and the cut middle area of the ND-TD section with a 10 mm width cut to obtain the 10.0 mm × 10.0 mm sample. In the end, the specimens were ground to the middle surfaces at $s=0$. Later, the specimens were ground using #1500, #2500, and #4000 and cleaned with ethanol. Subsequently, the specimens were polished using silica suspensions with particle sizes of 3, 1, and 0.25 μm. Finally, electropolishing was performed using a perchloric acid solution and ethanol at 30 V and −15 °C.

2.2. EBSD Measurements

After the sample preparation, the EBSD measurements taken using Hitachi SU5000 (Hitachi High-Tech Corporation, Tokyo, Japan) with EDAX VELOCITY PLUS (AMETEK EDAX, LLC, Mahwah, USA) were used to examine the microstructure and texture of the AA1100 aluminum alloy on the surface parallel to the normal direction (ND), as shown in Figure 1, at an accelerating voltage of 20 kV under the tilt angle of 70°, where the measured area was 2.0 mm × 1.6 mm, with a step size of 1.0 µm. Here, the microstructure and the texture were analyzed using the OIM Analysis 8.5 software (version 8.5, AMETEK EDAX, LLC, Mahwah, USA), and the volume fraction of the components was calculated using the LaboTex 3.0 software. The volume fraction of the major components shown in

Table 3. Miller index of texture components in hot-rolled aluminum alloy.

Texture Component	Miller Index
C	{112}<111>
S	{123}<634>
B	{110}<112>
Cube	{001}<100>
RC20°RD	{013}<001>
Goss	{110}<001>

was calculated by integrating the orientation intensity of f(g) within a Euler space with $∆{\phi }_1=∆\mathsf{\Phi }=∆{\phi }_2=10°$.

3. Results and Discussion

3.1. Effects of Thermomechanical Processing Conditions on Recrystallization

We measured 3,200,000 points over a large area with 2.0 mm × 1.6 mm with a scanning step of 1 μm for each specimen in order to collect a large dataset to calculate the orientation distribution function, which shows the same distribution by calculation from X-ray pole figures. The magnified inverse pole figure maps of all the hot-rolled microstructures are shown in Figure 3 and Figure 4 after the thermomechanical processing conditions of the hot rolling, as shown in Table 2, without further annealing, which is called the hot-rolled state. However, there are some black points in Figure 3 and Figure 4, which belong to non-indexed points due to heavily deformed structures or dislocation structures. After recrystallization, recrystallized microstructures are revealed in Figure 5 and Figure 6, while some black points are found in Figure 5a,e and Figure 6b,d due to partial recrystallization.

The kernel average misorientation (KAM) calculated within approximately the seventh nearest neighbor was utilized to distinguish between recrystallization and deformation in terms of the KAM value. The simultaneous Gaussian fitting using two peaks was applied on the KAM distribution curve to separate two peaks between low and high KAM values, as shown in Figure 7d. The intersection of the two curves was defined as the threshold of 1.58° in Figure 7d between recrystallization and deformation. The value of KAM can be considered the index of dislocation density. Thus, the KAM value of deformation shows a high value in Figure 7b, and that of recrystallization indicates a low value due to a low dislocation density in Figure 7c. In addition, the recrystallization fraction was estimated by the integration of the area fraction within the range of KAM from zero to the threshold of 1.58°. Here, we took the condition of 85% hot-rolled AA1100 alloy at 350 °C, for instance, to demonstrate the determination procedure of the threshold value of KAM. The threshold value of KAM lies in the range between 1.23 and 1.5, and the mean value is 1.33 for the 12 conditions. The results of the different values reveal the reproducibility question, and thus, it is necessary to keep improving this technique.

The recrystallization behavior is illustrated in Figure 8, where the recrystallization fraction determined by the above-mentioned KAM method is plotted with respect to the annealing time at annealing temperatures of 275 °C, 300 °C, and 350 °C after the double-hit compression in the AA1100 alloy. Except at 300 °C, the typical recrystallization curve is a sigmoidal curve with an S shape at 275 °C and 350 °C. Furthermore, it is observed that the fitting curves do not fit the data well in Figure 8a, especially for the condition of 5 and 60% at 275, 300, and 350 °C, which means a low strain rate and a low reduction ratio.

The Johnson–Mehl–Avrami–Kolmogorov (JMAK) equation is usually used to describe the recrystallization fraction as a function of time as follows [26]:

$$X_v=1-exp{(-0.693({\displaystyle \frac{t}{t_{0.5}}})}^{n_R}),$$

(1)

where $X_v$ is the recrystallized fraction, $t_{0.5}$ is the time required for the recrystallization fraction of 50%, $t$ is the time, and $n_R$ is the JMAK exponent. The fitting parameters are shown in

Table 4. Fitting parameters of n_R-value, t_0.5, and the volume fraction of recrystallization at the beginning X₀ for different thermo-mechanical processing conditions of reduction ratio (ε), strain rate ($\dot{\epsilon }$), and temperature (T) in AA1100 alloy.

$\mathit{T}$ (°C)	$\dot{\mathit{\epsilon }}$ (s−1)	$\mathit{\epsilon }$ (%)	Z (s−1)	nR	t0.5 (s)	X0 (%)
275	5	60	5.75 × 1015	0.69	5894	0
		85		0.39	2806	0.03
	90	60	1.04 × 1017	0.36	1685	0.05
		85		0.48	6238	0.01
300	5	60	1.27 × 1015	0.32	2683	0.05
		85		0.15	584	0.22
	90	60	2.28 × 1016	0.23	2336	0.10
		85		0.41	1395	0.04
350	5	60	8.85 × 1013	0.68	500	0
		85		0.80	108	0
	90	60	1.59 × 1015	0.98	136	0
		85		1.03	169	0

. The n_R value of the recrystallization exponent lies in the range between 0.15 and 1.03 in Table 4. It is noted that the recrystallization fraction X₀ at the beginning is larger than 0 at the seven conditions in Table 4, where the recrystallization fraction X₀ is predicted by using the fitting curve at t = 0. This finding indicates that different thermomechanical processing conditions possibly lead to the partial recrystallization during the thermomechanical processing. Therefore, this partial recrystallization during the thermomechanical processing reduces the driving forces for the following recrystallization, and the n_R value of the recrystallization exponent decreases.

Meanwhile, the t_0.5 value is usually expressed in terms of Z and annealing temperature $T$ on t_0.5 as follows [27]:

$$t_{0.5}=A{\epsilon }^{-a}{Z}^{-b}exp\left({\displaystyle \frac{Q}{RT}}\right),$$

(2)

where a is 1.5, b is 0.75, Q is 220 kJ mol⁻¹, and k is 1.5 × 10⁻⁴. Thus, according to Equation (2), t_0.5 decreases as Z increases at a constant annealing temperature, as shown in Table 4. This shows a decrease in t_0.5 as T increases at constant Z and strain. According to Sellars and McTegart [28], the hot working conditions of the temperature and the strain rate are empirically combined in the Zener–Holloman factor as

$$Z=\dot{\epsilon }{e}^{Q/RT},$$

(3)

where $\dot{\epsilon }$ is the strain rate, Q is the deformation activation energy, R refers to the ideal gas constant, and T is the deformation temperature. The deformation activation energy of Q is 156 kJ/mol, as previously reported [29]. Therefore, we can plot the n_R value of recrystallization obtained from the experiments as functions of log Z and reduction rate $\epsilon$ in the 3D diagram, where the Z values are listed in Table 4 in descending order with respect to temperature (T) and strain rate ($\dot{\epsilon }$).

It is observed in Table 4 that the value of n_R is high with Z and deformation reduction, which corresponds to low deformation temperature, high strain rate, and high reduction. By contrast, the values of n_R are small at high Z and low deformation reduction or low Z and high deformation reduction, which correspond to low deformation temperature, high strain rate, and low reduction and high deformation temperature, low strain rate, and high reduction, respectively.

3.2. The Kinetic Behavior of the Texture Components

To understand the kinetic behavior of the texture components, we also adapted the concept of the JMAK equation, and the fraction of texture component $X_t$ at time $t$ is expressed in terms of time as follows [30]:

$${\displaystyle \frac{X_t-X_i}{X_f-X_i}}=1-exp{(-0.693({\displaystyle \frac{t}{t_{0.5}}})}^{n_T}),$$

(4)

where $X_i$ and $X_f$ are the volume fractions of the component at the beginning and after recrystallization, respectively. In addition, we assumed that the constant of 0.693 is orientation-independent in this study. The developments of the recrystallization textures are plotted as a function of annealing time for the 12 hot-rolling processing conditions in Table 2, as shown in Figure 9 and Figure 10. The fitting parameters of the n_T value and t_0.5 for the C, S, B, Cube, RC, and G components are shown in

Table 5. Fitting parameters of n_T-value and t_0.5 for C, S, B, Cube (Cb), RC_20°RD (RC), and Goss (G) components at 12 thermo-mechanical processing conditions in AA1100 alloy.

	Name	1	2	3	4	5	6	7	8	9	10	11	12
C	n	-	2.20	0.73	3.15	1.49	3.00	1.71	-	7.01	0.80	0.85	1.00
	t0.5 (s)	-	1919	1513	7070	458	693	541	-	221	39	28	34
S	n	0.95	2.02	0.77	2.07	2.77	3.00	1.04	3.00	4.29	1.06	1.20	1.46
	t0.5 (s)	3910	2831	1200	4524	745	610	513	900	200	39	68	56
B	n	0.61	1.95	0.34	2.20	1.20	1.52	1.43	1.60	1.65	0.87	3.50	2.00
	t0.5 (s)	2355	3470	107	6187	418	316	659	936	131	37	108	62
Cb	n	1.09	0.87	1.82	1.28	1.16	2.50	-	2.50	2.60	1.18	2.06	1.88
	t0.5 (s)	6000	4348	2613	3920	890	554	-	742	218	57	72	86
RC	n	-	-	-	2.25	2.00	2.37	1.21	-	4.00	1.74	2.37	2.40
	t0.5 (s)	-	-	-	4335	398	900	780	-	190	50	76	65
G	n	-	-	-	1.59	-	-	0.43	2.14	-	-	3.75	2.68
	t0.5 (s)	-	-	-	12,300	-	-	101	3922	-	-	550	255

for the 12 thermo-mechanical processing conditions.

In this study, the deformation texture is composed of three main components, C, S, and B, and the recrystallization texture consists of Cube, RC, and G components. It was observed that the three deformation texture components, C, S, and B, decrease, but three recrystallization texture components, Cube, RC, and G, increase with increasing annealing time at 275, 300, and 350 °C, as shown in Figure 9 and Figure 10. The recrystallization component of the Cube with a volume fraction larger than 30% increases significantly in comparison with the RC and G components. Furthermore, the initial volume fraction of the Cube component is larger than 10% at the beginning of recrystallization. This observation is in agreement with the occurrence of partial recrystallization during the thermomechanical processing, and the Cube is retained in the thermomechanical processing.

3.3. Relationship between Recrystallization and Texture

As discussed above, the 12 hot-rolling processing conditions reveal the different fractions of recrystallization after the hot-rolling process. During the hot-rolling process, we take the retained recrystallization fraction into account, which is called the initial fraction for the following annealing process. In addition, we separate the deformation and recrystallization textures in order to understand the competition between both textures during the annealing process. The texture components of C, S, and B belong to the deformation texture, and the Cube, RC, and G components are categorized to the recrystallization texture, as shown in Figure 11.

At the beginning of the annealing process in Figure 11a, the initial fraction of the deformation texture ranges from 3% to 12%, while the fraction of the recrystallization texture lies below 13% in Figure 11b. It is noted that the deformation texture decreases, but the recrystallization texture increases with increasing recrystallization fraction, that is, the annealing time at a fixed temperature. In the case of the recrystallization texture, the volume fractions of RC and G are below 10%, and these initial fractions are close to zero. However, the initial fraction of the Cube is about 12%, which suggests the Cube is retained during the hot-rolling process.

In order to investigate the effects of the 12 hot-rolling processing, we plot the volume fraction in terms of the recrystallization fraction for the six components of C, S, B, Cube, RC, and G, as shown in Figure 12. Firstly, considering the deformation texture of C, S, and B components, it is observed that the initial fraction of B, S, and C is 13%, 9%, and 3%, respectively, as in Figure 12a–c. By increasing the recrystallization fraction, the components of B and S decrease significantly more than the C component.

For the recrystallization texture of Cube, RC, and G, the volume fractions of Cube and G are enhanced with the increasing recrystallization time in Figure 12d,f, while the RC component does not show a significant change in the volume fraction in Figure 12e. Furthermore, the Cube component plays an important role in the annealing process because the 50% fraction of Cube is larger than the 12% fraction of Goss. Thus, we take the Cube component into account for the following discussions between the 12 hot-rolling processing conditions and the texture components.

The 12 hot-rolling processing conditions in Table 4 present the various Z value defined by Equation (3). The smallest Z value was found for the strain rate of 5/s and the temperature of 350 °C, and the largest Z value was found for the strain rate of 90/s and the temperature of 275 °C. By increasing the strain rate and decreasing the temperature, the Z value is enhanced. Figure 13 shows the volume fraction of the Cube component with respect to the recrystallization fraction for the 12 hot-rolling processing conditions. In comparison with the strain rate of 90/s in Figure 13b, the fraction of Cube is larger than that at the strain rate of 5/s in Figure 13a, and it has a larger fraction at 350 °C than at 275 °C. The conditions of 5/s and 350 °C correspond to the smallest Z value. It was also reported that the stability of the Cube orientation increases with the decreasing Z value [7,9,23].

After the hot-rolled state before annealing, the 13% volume fraction of the Brass component is larger than the S and Copper components in Figure 12a–c. As increasing the recrystallized fraction, the β-fiber components decrease. The hot-rolled state before annealing already reveals a 10% recrystallization, in which the recrystallization texture component is dominated by the Cube component with a 10% volume fraction in Figure 12d. The Cube fraction increases markable in the range between 20% and 55% depending on the thermomechanical process conditions above 40% recrystallization. The presence of the {001}<100> orientation in the hot bands or shear bands of the deformation matrix resulted in the development of the Cube-dominating recrystallization texture [31,32,33,34]. The severe thickness reductions involved (above 98%) elevated temperatures and large inter-pass times between two consecutive rolling passes may induce static recrystallization in the breakdown rolling and lead to the evolution of Cube orientation [35]. Sidor et al. [36] claimed that the hot-rolling texture is composed of β-fibre and recrystallization texture of very weak Cube and Goss components at low finishing temperature (T < 240 °C). At low roll finishing temperatures (T ~250 °C), the formation of recrystallization texture components could be attributed to the enhanced interpass time by triggering the recrystallization process, which leads to β-fibers mixed with recrystallization components. Purely RX texture components were observed at the increasing roll finishing temperatures (T > 300 °C) [37]. The boundary relation between the recrystallization nuclei and a deformed matrix in Al has the rotation axis of the <772> axis with a rotation angle of 52–70° [38].

Then, after annealing, the maximum volume fraction of the Cube component is 55% under the reduction of 85% with the strain rate of 5/s or 90/s at a high temperature of 350 °C for complete recrystallization. However, it is about 50% under the reduction of 60% with the strain rate of 5/s at 300 °C, which shows 60% recrystallization. Kraner et al. [39] also reported that the 6 mm finished thickness of the hot-rolled band had 64% of the random texture components and 83% recrystallized grains, whereas the 3.25 mm thickness had 42% random texture components and 55% recrystallized grains. During the high-temperature and high-dwell-time annealing operation of the Al5052 sheet processed through high amounts of thickness reductions, Wu et al. [40] found that some grains take part in the static recrystallization mechanism, leading to grain refinement.

4. Conclusions

In this work, we tried to understand the effects of hot-rolling process conditions on the development of the recrystallization texture during the following annealing. Comparing the hot-rolling process conditions without annealing, it is noted that the 13% volume fraction of the Brass component is larger than the S and Copper components. The hot-rolled state shows the 10% volume fraction of the Cube component under the 10% recrystallization. Then, after annealing, the maximum fraction of the Cube component is 55% under the reduction of 85% with the strain rate of 5/s or 90/s at a high temperature of 350 °C for complete recrystallization. The maximum fraction of the Cube is about 50% under the reduction of 60% with the strain rate of 5/s at 300 °C for 60% recrystallization. This finding is of industrial importance in tailoring the texture components for different applications by varying the thermomechanical process conditions.