Effect of Zr content on the precipitation and dynamic softening behavior in Al–Fe–Zr alloys

Journal Pre-proof Effect of Zr content on the precipitation and dynamic softening behavior in Al–Fe–Zr alloys

Jieyun Ye, Renguo Guan, Hongjin Zhao, Anmin Yin PII:

S1044-5803(19)32738-X

DOI:

https://doi.org/10.1016/j.matchar.2020.110181

Reference:

MTL 110181

To appear in:

Materials Characterization

Received date:

8 October 2019

Revised date:

30 January 2020

Accepted date:

4 February 2020

Please cite this article as: J. Ye, R. Guan, H. Zhao, et al., Effect of Zr content on the precipitation and dynamic softening behavior in Al–Fe–Zr alloys, Materials Characterization (2020), https://doi.org/10.1016/j.matchar.2020.110181

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© 2020 Published by Elsevier.

Journal Pre-proof Effect of Zr content on the precipitation and dynamic softening behavior in Al– Fe–Zr alloys Jieyun Ye a , Renguo Guan a,b,* , Hongjin Zhao a,** , Anmin Yin c a

Faculty of Material Metallurgy and Chemistry, Jiangxi University of Science and

Technology, Ganzhou 341000, China b

School of Materials Science and Engineering, Northeastern University, Shenyang

110819, China c

Part Rolling Key Laboratory of Zhejiang Province, Ningbo University, Ningbo

Zhejiang 315211, China

of

Abstract

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In this paper, the precipitation and dynamic softening behavior of Al–Fe–Zr alloys were investigated by transmission electron microscopy (TEM), electron back scattering diffraction (EBSD), and the thermodynamics

was

studied

by

JmatPro

-p

precipitation

software.

Results

show

that

the

re

optimizing temperature of Al 3 Zr precipitates was 350℃ and the fraction of butterfly-shaped Al 3 Zr and rod-shaped Al 3 Fe reached a peak at ~350℃. With the increase of Zr content, the fraction of Al 3 Zr

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increased, whereas its size and morphology had almost no change. The fraction and morphology of Al 3 Fe showed less variation with Zr content. Dynamic recovery and dynamic recrystallization

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occurred during extrusion. The two kinds of precipitation particles in aluminum matrix pinned the dislocations and inhibited dynamic recrystallization effectively. As the Zr content went up from 0.1 to

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0.4 wt.%, the inhibition got stronger and the fraction of recrystallized grains decreased from 44.2% to 86.9%. As a consequence of the dynamic recrystallization suppression , dynamic recovery became the major softening mechanism. Keywords:

Al–Fe–Zr

alloy;

dynamic

softening;

precipitation;

dynamic

recovery;

dynamic

recrystallization; pinning force

Introduction The development of electric power and the comprehensive construction of smart grids has proposed high requirements for the performance of overhead transmission wires. Heat-resistant aluminum alloy has been used extensively in long-distance, large-span and ultra-high-voltage transmission because of its high conductivity, high transmission capacity and high specific strength [1-5]

. Alloy properties can be enhanced by rare-earth-element addition, such as scandium, erbium and

yttrium

[6-9]

. However, the high cost of these elements has limited their extensive application

[6,10-11]

.

Recently, aluminum (Al) alloys with zirconium (Zr) addition have been used extensively in electrical systems (including high-voltage transmission wires) because of their excellent combination of

1

Journal Pre-proof strength, conductivity and heat resistance

[6-7]

. Previous studies have found that Zr and Al can form

metastable Al 3 Zr, which is coherent with the matrix, can suppress recrystallization and improve the alloy heat resistance

[6-7, 12]

.

Ferrum (Fe) has long been regarded as an impurity element in aluminum alloys and is difficult to eliminate. Fe and Al can form Al 3 Fe at grain boundaries with coarse lamellar or acicular morphology, which deteriorate the alloy properties significantly

[13-14]

. However, it has been found in recent years

that Al–Fe alloy has an excellent heat resistance therefore is a promising candidate material for quality-sensitive equipment and products

[15-17]

The previous researches and findings

[18-20]

indicate

that the second phase particles can effectively pin dislocation and subgrain boundary, thus inhibit

of

dynamic recrystallization significantly. As for the heat-resistant aluminum alloy, the final

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microstructure has an significant effect on strength, heat resistance and other properties, hence, studying precipitation, the dynamic softening behavior and mechanism is an effective technique to

-p

control the ultimate microstructure. This research focuses mainly on the effect of Zr content on the

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and microstructure in Al–Fe–Zr alloys. We prepared a series of pre-drawing Al–Fe–Zr semi-finished wire products with different Zr contents by backward extrusion. Precipitation and dynamic softening

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behavior were investigated by transmission electron microscopy(TEM), electron back scattering diffraction(EBSD) and JmatPro software. The formation mechanism of Al–Fe and Al–Zr precipitates

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and its effect on dynamic softening behavior during deformation is discussed. We hope this study will provide a theoretical basis for the development of low-cost, high-performance Al–Fe–Zr wires.

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Materials and methods

Four Al–Fe–Zr alloys with different Zr contents were prepared by using pure Al ingot (99.99 wt%), Al–Fe master alloy and Al–10 Zr master alloys. The alloy composition was measured by an Iris Advantage 1000 inductively-coupled plasma spectrometer (ICP) as shown in Table 1. Backward extrusion was performed with a 500T oil press. The blank and extrusion cylinde rs were heated to 400℃ and 350℃, respectively. The extrusion ratio was 22 and the extrusion deformation was ~95%. The extrusion speed was 5 m/min. Graphite emulsion was used for lubrication. A schematic diagram of the extrusion process, concrete size and sample (SEM, EBSD, TEM) positions was given in Figure 1.

2

Journal Pre-proof

ED

ED Load

ED

φ8.6 8

6

Sample for

φ22

microstructure observation

Fig. 1 schematic diagram of extrusion process 、concrete sizes and sample (SEM, EBSD, TEM)

of

positions

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The EBSD samples were ion polished by a Gatan 697 Ilion II polishing machine. An Ultra Plus scanning electron microscope was used for grain-size and orientation measurements. The minimum

-p

angle and scanning step were set as 2° and 0.2 μm, respectively, to reduce orientation noise at the

re

grain boundaries.

A Tecnai G2 20 transmission electron microscope was used to investigate the precipitates, grain

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boundaries, dislocations and other substructures. Thin foils for TEM were prepared by elecrospark wire electrode cutting sample pieces from the extruded materials and by polishing them mechanically

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to ~50 μm. Discs with a 3 mm diameter were punched and dimpled to 60 μm, followed by twin -jet electropolishing at 18 V (Struers Tenupol5). A solution of 1/3 (volume fra ction) nitric acid and 2/3

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methanol at –25 °C was applied during electropolishing. The volume fraction and average radius of precipitation particles were quantified by image analysis of TEM images. Five different areas (1.5×1.5μm 2 for each) and more than 200 particles were measured for each alloy. Table 1: Chemical composition of the four alloys (wt.%)

Samples

Fe

Zr

Al

Al–0.35Fe–0.1Zr

0.352

0.122

Bal.

Al–0.35Fe–0.2Zr

0.341

0.247

Bal.

Al–0.35Fe–0.3Zr

0.357

0.298

Bal.

Al–0.35Fe–0.4Zr

0.354

0.380

Bal.

3. Results 3.1 Effect of Zr content on precipitation Figure 2 and 3 show the morphology and distribution of the precipitates in the Al–Fe–Zr alloy. Al 3 Zr and Al 3 Fe could be identified from the diffraction pattern. Rod-shaped phase is Al 3 Fe and the

3

Journal Pre-proof fine spherical phase is Al 3 Zr, with diffraction patterns shown in Figure 2 (g) and 2(h), respectively. After magnification, the Al 3 Zr precipitates had a typical butterfly-shaped morphology (as shown in Fig.2(e) and 2(f)), and were located in the grains and at grain boundaries and subgrain boundaries. In the Al–0.35Fe–0.1Zr alloy, a small amount of Al 3 Zr was observed (marked with red arrows) with an uneven distribution. The average radius of Al 3 Zr was ~（16.8± 0.2）nm and its volume fraction was ~(4.79 ± 0.11) × 10 -3 . In the Al–0.35Fe–0.2Zr alloy, the average radius was ~（17.2± 0.3）nm. Its volume fraction increased to ~(7.56 ± 0.22) × 10 -3 . In the Al–0.35Fe–0.3Zr alloy, the Al 3 Zr was distributed evenly and its amount increased significantly to a volume fraction of ~(1.065 ± 0.02) × 10 -1 . While the average radius had changed little . As shown in Figure 2d, in the Al–0.35Fe–0.4Zr

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alloy, the volume fraction of Al 3 Zr increased to (5.233 ± 0.13) × 10 -1 and the average radius almost

ro

kept unchanged. In summary, the fraction of Al 3 Zr increased significantly and tended to be distributed evenly with an increase in Zr content, whereas its size and morphology showed almost no

-p

change.

a

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Al 3 Zr

c

Al 3 Fe Al 3 Zr

na

Al 3 Fe

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re

b

0.5μm

0.5μm

d

Dislocation

Al 3 Zr Al 3 Zr

Al 3 Fe

0.5μm

0.5μm

4

Journal Pre-proof e

f

Dislocation Al 3 Zr Al 3 Zr

Subgrain

dislocation wall 0.2μm

0.2μm

g

of

h

ro

（ -440）

（ -3-31）

-p

（ -13-2）

lP

re

（ -201）

Fig. 2 Al 3 Zr morphology and distribution map of extruded Al-Fe-Zr alloys

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(a) Al–0.35Fe–0.1Zr (b) Al–0.35Fe–0.2Zr (c) (e) Al–0.35Fe–0.3Zr (d) (f) Al–0.35Fe–0.4Zr (g) electron diffraction pattern corresponding to phase marked by black circle in Fig. 1(c)

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(h) electron diffraction pattern corresponding to phase marked by red circle in Fig. 1(d)

Figure 3 shows the morphology and distribution of Al 3 Fe in the four alloys. Al 3 Fe tended to precipitate at the grain boundaries along the elongated direction. Some Al 3 Fe had been broken into small pieces during the extrusion. These pieces were unseparated and maintained their original morphology. Some Al 3 Fe was elliptic because of tip melting, as indicated by the red line i n Figure 3. The average size of the Al 3 Fe was ~350 nm. With an increase in Zr content from 0.1% to 0.4%, the Al 3 Fe size and morphology did not change significantly.

5

Journal Pre-proof b

a

Al 3 F

Al 3 Fe 1μm

d

re

-p

ro

of

c

1μm

Al 3 Fe

1μm

lP

1μm

Al 3 Fe

Fig. 3 Al 3 Fe morphology and distribution map of extruded Al–Fe–Zr alloys

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(a) Al–0.35Fe–0.1Zr (b) Al–0.35Fe–0.2Zr (c) Al–0.35Fe–0.3Zr (d) Al–0.35Fe–0.4Zr

In addition, the equilibrium phase diagram was given using JmatPro software as shown in Figure

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4 in order to investigate the precipitates in the Al-Fe-Zr alloys. Results showed that Al 3 Fe and Al 3 Zr precipitates were found in the alloys which were consistent with the TEM results. The optimizing temperature of Al 3 Zr precipitates was 350℃, and it is approximately the extrusion temperature. When the temperature was above 350 °C, the mole fraction of Al 3 Zr increased with the decrease of temperature and reached the peak at 350℃, without any change thereafter. By comparing four alloys, with the increase of Zr content, the mole fraction of Al 3 Zr increased rapidly. As to Al 3 Fe, the mole fraction was almost unchanged with the variation of Zr content and temperature of below 600℃. This is unanimous to the results given in Figure 2 and 3.

6

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Fig. 4. Equilibrium phase diagram of Al-Fe-Zr alloys fro m Jmatpro software

3.2 Effect of Zr content on dynamic softening

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Figure 5 shows the grain-boundary distribution of the grains in the Al–Fe–Zr alloys with

ro

different Zr contents. Grain boundaries with a misorientation larger than 15° were marked in black, whereas those from 2° to 15° were marked in white. The grains were elongated significantly into a

-p

fibrous morphology along the extrusion direction. Large amounts of low-angle equiaxial subgrains

re

that resulted from dynamic recovery were found in the elongated grains. Fine recrystallized grains were located near the band structure. With an increase in Zr content from 0.1% to 0.4%, the average

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length and width of the band structure increased and the fraction of low-angle grain boundaries in the

b

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a

na

band structure increased too.

Recrystallized grain Recrystallized grain

10μm

10μm

c

d Recrystallized grain

Recrystallized grain

10μm

10μm

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Journal Pre-proof Fig. 5 EBSD diagrams of extruded Al–Fe–Zr alloys with different Zr contents (a) Al–0.35Fe–0.1Zr (b) Al–0.35Fe–0.2Zr (c) Al–0.35Fe–0.3Zr (d) Al–0.35Fe–0.4Zr

Figure 6 shows the change in mean grain misorientation and the low-angle grain boundary fraction. Large amounts of low-angle grain boundaries were found in the extruded alloys. With an increase in Zr content from 0.1% to 0.4%, the Zr fraction increased gradually from 57.84% to 78.36%. The average misorientation among the grains decreased, and was measured to be 20.1°, 16.9°, 16.35° and 11.97°.

0.14 0.12

0.08 0.06 0.04

FLAGBS=60.48 1.3

0.08

ro

Frequency/%

0.10

FLAGBS=57.84 2.1%

0.06 0.04

15

-p

15

0.02

0.02

0.00

0.00

0

10

20

30

40

50

lP

c

0.14 0.12 0.10

FLAGBS=70.28 2.4

na

0.08 0.06 0.04

0.00 0

Jo ur

15

0.02

10

0

20

30

40

10

re

Misprientation/degree

60

0.14

20

30

40

50

60

Misprientation/degree

d

0.12 0.10

Frequency/%

Frequency/%

of

0.12

0.10

Frequency/%

b

0.14

a

FLAGBS=78.36 1.7

0.08 0.06 0.04

15

0.02 0.00 50

0

60

10

20

30

40

50

60

Misprientation/degree

Misprientation/degree

Fig. 6 Distribution of grain-boundary misorientation in Al–Fe–Zr alloys with extrusion (a) Al–0.35Fe–0.1Zr (b) Al–0.35Fe–0.2Zr (c) Al–0.35Fe–0.3Zr (d) Al–0.35Fe–0.4Zr

Figures 7 and 8 show the submicrostructure distribution and its fraction change as a function of Zr content after extrusion. For the Al–0.35Fe–0.1Zr alloy, the recrystallized grains accounted for 44.2% and the submicrostructures accounted for 46.3%, which were marked blue and yellow, respectively, in Fig. 5(a). The recrystallized grains were surrounded by high -angle grain boundaries (shown as black lines), whereas the low-angle grain boundaries (shown as white lines) existed mainly in the submicrostructures. Based on Figures 8, it can be concluded that the extrusion process induced dynamic recovery and dynamic recrystallization. For the Al–0.35Fe–0.2Zr alloy, the fraction of recrystallized grains decreased to 33.4%, whereas the fraction of submicrostructure increased to

8

Journal Pre-proof 56.1%. With an increase in Zr concentration, the fraction of recrystallized grains and submicrostructure continued to decrease and increase and reached 86.9% and 5.2%, respectively, in the Al–0.35Fe–0.4Zr alloy.

b

ED

of

a

c

d

10μm

ED

na

lP

re

-p

ED

ro

10μm

10μm

10μm

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Fig. 7 Substructure distribution and content change of different Zr contents in Al–Fe–Zr alloys (a) Al–0.35Fe–0.1Zr (b) Al–0.35Fe–0.2Zr (c) Al–0.35Fe–0.3Zr (d) Al–0.35Fe–0.4Zr

90

Fraction/%

Recrystallized Substructure Deformed

60

30

0

Al-0.35Fe-0.1Zr Al-0.35Fe-0.2Zr Al-0.35Fe-0.3Zr Al-0.35Fe-0.4Zr

Fig. 8 Substructure content change of different Zr contents in extruded Al–Fe–Zr alloys

9

Journal Pre-proof 3.3 Effect of Zr content on dislocation density To quantitatively study the dislocation density, the kernel average misorientation （KAM） was adopted to determine the local misorientation using the EBSD. Where KAM value is high, plastic deformation degree or defect density is high [21] . In general, to extrapolate the density information, a simple method according to the strain gradient theory was used as following [22-23] ： ρGND =

2KAMave μb

where ρGND is the mean geometrically necessary dislocations, μ is the unit length (200nm) and

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b= 0.2863nm [24] is the magnitude of the Burgers vector. KAM ave represents the average misorientation

ro

of the selected area.

Figure 9 shows the KAM diagrams in the Al–Fe–Zr alloys with different Zr contents. With an

and 0.619387, respectively.

So

the

dislocation density

re

0.578176,

-p

increase in Zr content from 0.1% to 0.4%，the average KAM was measured to be 0.45319, 0.547626 ，

1.913×10 12 /cm 2 、 2.019×10 12 /cm 2 、 2.163×10 12 /cm 2 ,

respectively.

was

1.583×10 12 /cm 2 、

To

summarize,

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na

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the increase of Zr content, the dislocation density increased gradually.

Fig. 9 KAM diagrams in the Al–Fe–Zr alloys with different Zr contents

10

with

Journal Pre-proof (a) Al–0.35Fe–0.1Zr (b) Al–0.35Fe–0.2Zr (c) Al–0.35Fe–0.3Zr (d) Al–0.35Fe–0.4Zr

4. Discussion Because the solubility of Fe is extremely low in the aluminum matrix, the Al–Fe phase is formed mainly during solidification

[25]

. According to the Al–Zr binary phase diagram, Zr and Al form a

peritectic phase diagram with a limited solubility and the temperature of the peritectic reaction is 0.5°C higher than the melting point of Al. For equilibrium solidification, the maximum solubility of Zr is 0.28% and its solubility decreases as the temperature decreases. However, for the actual non-equilibrium solidification, the solubility o f Zr in Al can reach ~0.5%

[26]

. Therefore, for the

experimental alloys with varied Zr content from 0.1% to 0.4%, a Zr-rich precipitate was indicated to

of

form during extrusion instead of solidification. Elemental Zr existed mainly in the solid – solution state

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before extrusion.

According to the TEM observation and diffraction identification, two kinds of precipitates,

-p

Al 3 Zr and Al 3 Fe, formed in the experimental alloys. The Al 3 Fe was relatively coarse as it was

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precipitated from the liquid at a high temperature. Because of its high hardness, Al 3 Fe was broken along the direction of metal flow during extrusion and therefore behaved as clusters of ~350 nm

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rod-like or spherical particles. For the Al 3 Zr precipitate, large amounts of defects, such as dislocations and deformation bands were formed because of the three-dimensional compressive stress

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during the hot extrusion. As a result, the diffusion of the solute atoms was accelerated and more nucleation sites for Al 3 Zr were provided, which promoted the precipitation of the butterfly-shaped [27]

. Based on Figure 2, the average radius of Al 3 Zr was ~17 nm. As the Zr

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Al 3 Zr significantly

concentration increased from 0.1% to 0.4%, the volume fraction of Al 3 Zr increased from (4.791 ± 0.11) × 10 -4 to (5.233 ± 0.13) × 10 -2 and its distribution gradually became uniform. During the extrusion of the Al–Fe–Zr alloys, two kinds of softening processes, dynamic recovery and dynamic recrystallization, were likely to occur with the increase of strain. According to the serrated and wavy grain boundaries in Figure 5, it could be inferred that softening behavior existed throughout the extrusion process

[28]

. Figure 10 shows the longitudinal section microstructure of the

extruded alloy observed by TEM. Large amounts of substructures formed by interacted dislocations could be found. Those twists and steps on the subgrain boundaries composed of two groups of parallel dislocations indicated that the subcrystals were migrating

[29]

(as shown in Figure 10a).

During migration, the subgrain boundaries kept absorbing dislocations so that the dislocation density increased. Because of the repulsion among the congeneric dislocations, the dislocations started to increase and were arranged vertically on different slip planes. As a result, the total strain energy was reduced and the stain was offset by the congeneric dislocations, which lead to the formation of

11

Journal Pre-proof polygonal grains. It is in good coincidence with the results of Figures 9. The polygonization of grains indicates the occurrence of dynamic recovery during extrusion. Those subgrain boundaries in the deformed grains were curved with a high curvature interface, and thus, they were easy to migrate and merge [30] . Traces of subgrain merging existed in the cycled area in Figures 10(a) and 10(d). Because of the increase in subgrain size and dislocation density at the boundaries, the misorientations between the adjacent subgrains tended to increase. These subgrain boundaries transformed gradually into grain boundaries and thereby became preferred nucleation sites for recrystallization

a

[27]

.

b

of

0.2μm

ro

Polygon

Al 3 Fe

-p

Dislocation Wall

re

Subgrain boundaries

1μm

lP

Dislocation entanglement

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Dislocation wall

Dislocation entanglement

1μm

Al 3 Zr

d Dislocation wall

na

c

Al 3 Fe

Dislocation entanglement

Dislocation wall

Dislocation entanglement

Subgrain boundaries 1μm

0.5μm

Fig. 10 TEM structure of longitudinal section of extruded Al –Fe–Zr alloys (a) Al–0.35Fe–0.1Zr (wt.%) (b) Al–0.35Fe–0.2Zr (wt.%) (c) Al–0.35Fe–0.3Zr (wt.%) (d) Al–0.35Fe–0.4Zr (wt.%)

The recrystallization process includes the formation of unstrained subgrains and their growth by consumption of the nearby strained matrix. If a critical recrystallized grain size is reached, its grain boundary can migrate into the nearby matrix. The critical radius Rc can be calculated from the Gibbs– Thomson equation

[31]

: 𝑅c =

4𝛾GB 𝑃D −𝑃Z

(1)

in which γ GB is the interface energy per unit area of the alloy, which is estimated to be a constant

12

Journal Pre-proof value

[32-34]

, P D is the nucleation driving force because of the stored strain energy and P Z is the

pinning force that suppresses the migration of the low- and high-angle grain boundaries. The interfacial velocity of the recrystallized grain boundaries can be calculated from 𝑉 = 𝑀(𝑃𝐷 − 𝑃𝑍 )

[31]

: (2)

in which M is the interface mobility. Because the extrusion condition of the four experimental alloys was constant, P D can be regarded as a constant. According to Equations (1) and (2), the critical subgrain radius R C and the interface velocity of the recrystallized grain boundary V are dependent on the pinning force of the precipitates. Therefore, the precipitates affect the recovery and recrystallization process significantly.

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Figure 10(b) shows that the precipitates were located at grain boundaries and impede the interface

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migration during deformation, which, to some extent, inhibits dynamic softening. Precipitates can also drag and pin dislocations. As shown in the red circle in Figure 2(f), the dislocations were bent

-p

and twisted through the particles.

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With an increase in pinning force, the critical size R c that enables the subgrain boundaries to migrate into the adjacent deformed matrix increases, whereas the grain boundary velocity V decreases,

lP

and recrystallization inhibits. It is generally acknowledged that the pinning force P Z by the precipitates is related to its size, fraction, morphology and orientation relationship with the matrix [27]

. For the butterfly-shaped Al 3 Zr, which maintains a coherent orientation relationship with

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interface

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the matrix, its maximum grain-boundary pinning force can be estimated from 𝑃𝑍 =

[34-36]

:

6𝑉𝑓 𝛾𝐺𝐵

(3)

𝑟

For the Al 3 Fe, which maintains an incoherent orientation relationship with the matrix, its maximum [34-36]

grain-boundary pinning force can be estimated from

𝑃𝑍 =

:

3𝑉𝑓 𝛾𝐺𝐵

(4)

2𝑟

where V f and r are the volume fraction and mean radius of the precipitates, respectively. γ GB is the interface energy per unit area of the alloy, which is estimated to be a constant value

[32-34]

.

By comparing Equations (3) and (4), it can be estimated that the Al 3 Zr pinning force is approximately four times larger than Al 3 Fe if the V f /r value of particle is same. So 4*V f /r value was given in the table in order to compare the pinning force of Al 3 Fe and Al 3 Zr. 4*V f /r value of Al 3 Zr increased gradually with increasing Zr content and V f/ r value of Al 3 Zr kept almost unchanged. So the pinning force of Al 3 Zr increased gradually and pinning force of Al 3 Fe were little changed in the four alloys. In Al-0.35Fe-0.1Zr alloy and Al-0.35Fe-0.2Zr alloy, 4*V f /r value of Al 3 Zr was slightly smaller than V f /r value of Al 3 Fe, so the pinning force of Al 3 Zr was smaller than the pinning force of Al 3 Zr. Nevertheless in the Al-0.35Fe-0.3Zr alloy and Al-0.35Fe-0.4Zr alloy, 4*V f /r value of Al 3 Zr increased

13

Journal Pre-proof significantly. The pinning force came mainly from Al 3 Zr and increased rapidly. So the critical size for subcrystals to grow towards the surrounding deformed matrix had to increase. As a result, the high-angle grain boundaries were difficult to form and grain -boundary migration slowed down, which suppressed the occurrence of dynamic recrystallization significantly. Table 2 Statistics of average radius of precipitates ( r ) and volume fraction of precipitates ( V f )

Alloy

Average radius r（ nm）

Volume fraction V f （ %）

V f /r （ nm -1 ）

4*V f /r （ nm -1 ）

Average radius R（nm）

-6

-5

350±2.2

2.21±0.15

6.31×10 -5

Al3Zr -3

Volume fraction V f （ %）

V f /r （ nm -1 ）

Al3Fe

16.8±0.2

(4.791±0.11)×10

1.14×10

Al–0.35Fe–0.2Zr

17.2±0.3

(7.561±0.22)×10 -3

4.40×10 -6

1.76×10 -5

352±3.7

2.18±0.13

6.19×10 -5

Al–0.35Fe–0.3Zr

17.4±0.3

(1.065±0.02)×10 -1

6.12×10 -5

2.448×10 -4

347±3.5

2.25±0.22

6.48×10 -5

Al–0.35Fe–0.4Zr

17.5±0.4

(5.233±0.13)×10 -1

2.99×10 -4

1.200×10 -3

2.23±0.27

6.32×10 -5

353±2.5

ro

2.85×10

of

Al–0.35Fe–0.1Zr

-p

As shown in Figure 10, the subgrain is ~880 nm in the Al–0.35Fe–0.1Zr alloy (subgrain boundary marked with a red arrow in Figure 10 (a)). As the Zr content increases, the subgrains are

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refined significantly. The average subgrains were 820 nm, 750 nm and 460 nm for the Al–0.35Fe– 0.2Zr, Al–0.35Fe–0.3Zr and Al–0.35Fe–0.3Zr alloys, respectively. According to the EBSD results in

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Figure 5, Figure 6, and Figure 7, dynamic recovery and dynamic recrystallization occurred during hot extrusion. With the increase in Zr concentration, the the fraction of recrystallization decreased

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gradually, whereas the dynamic recovery increased. The formation of a fine submicrostructure

delayed.

5. Conclusions

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because of dynamic recovery reduced the strain energy significantly, so that the recrystallization was

(1) The optimizing temperature of Al 3 Zr precipitates was 350℃ and it fitted right in with the extrusion temperature. Butterfly-shaped Al 3 Zr and rod-shaped Al 3 Fe were observed in the extruded Al–Fe–Zr alloys. With an increase in Zr content from 0.1 wt.% to 0.4 wt.%, the volume fraction of Al 3 Zr increased significantly from (4.791±0.11)×10 -3 to (5.233±0.13)×10 -1 . The volume fraction of Al 3 Fe changed little. The size and morphology of the two phases performed almost no change. (2) The Al 3 Zr and Al 3 Fe particles precipitated in aluminum matrix pinned the dislocations effectively. The pinning force of Al 3 Zr is smaller than the pinning force of Al 3 Fe at the low Zr content （0.1~0.2 wt.%）alloys. When the content of Zr increased to 0.3and 0.4 wt.%, the pinning force of Al 3 Zr increased remarkably. The total pinning force came mainly from Al 3 Zr and increased rapidly. (3) Dynamic recovery and dynamic recrystallization occurred during extrusion. As a

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Journal Pre-proof consequence of the pinning force got stronger, the dynamic recrystallization behavior was inhibited more intensively. With the Zr content increased from 0.1 to 0.4 wt.%, the fraction of recrystallized grains decreased from 44.2% to 86.9%. Dynamic recovery became the major softening mechanism.

6. Acknowledgements This work was supported by the National key research and development program (grant numbers 2018YFB2001800), the National Natural Science Foundation of China (grant numbers 51871184) and the High-Level Talent Support Program of Liaoning (grant number XLYC1 802128).

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Journal Pre-proof Conflict of interest We declare that we do not have any commercial or associative interest that

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represents a conflict of interest in connection with the work submitted.

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Journal Pre-proof Highlights

1、Al-Fe-Zr alloys with various Zr content were prepared using backward extrusion. 2、 With the increase of Zr content, the fraction of Al 3 Zr gradually increased. 3、In

Al-Fe-Zr alloys the dynamic softening mechanism was mainly due to dynamic recovery.

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4、Dynamic recrystallization was suppressed by the dispersive Al 3 Zr precipitates.

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