through thickness cube texture

1. See discussions, stats, and author profiles for this publication at: https://www.researchgate.net/publication/331575032
Development of Through-Thickness Cube Recrystallization Texture in Non-
oriented Electrical Steels by Optimizing Nucleation Environment
Article in Metallurgical and Materials Transactions A · March 2019
DOI: 10.1007/s11661-019-05167-3
CITATIONS
2
READS
135
5 authors, including:
Some of the authors of this publication are also working on these related projects:
Ultrahigh strength materials View project
Heusler alloys View project
Ning Shan
Northeastern University (Shenyang, China)
4 PUBLICATIONS 7 CITATIONS
SEE PROFILE
Jin Long Liu
Northeastern University (Shenyang, China)
18 PUBLICATIONS 94 CITATIONS
SEE PROFILE
Fang Zhang
Northeastern University (Shenyang, China)
46 PUBLICATIONS 348 CITATIONS
SEE PROFILE
Liang Zuo
Northeastern University (Shenyang, China)
645 PUBLICATIONS 6,056 CITATIONS
SEE PROFILE
All content following this page was uploaded by Jin Long Liu on 26 August 2019.
The user has requested enhancement of the downloaded file.

2. 1 23
Metallurgical and Materials
Transactions A
ISSN 1073-5623
Volume 50
Number 5
Metall and Mat Trans A (2019)
50:2486-2494
DOI 10.1007/s11661-019-05167-3
Development of Through-Thickness Cube
Recrystallization Texture in Non-oriented
Electrical Steels by Optimizing Nucleation
Environment
Ning Shan, Jinlong Liu, Yuhui Sha, Fang
Zhang & Liang Zuo

3. 1 23
Your article is protected by copyright
and all rights are held exclusively by The
Minerals, Metals & Materials Society and ASM
International. This e-offprint is for personal
use only and shall not be self-archived in
electronic repositories. If you wish to self-
archive your article, please use the accepted
manuscript version for posting on your own
website. You may further deposit the accepted
manuscript version in any repository,
provided it is only made publicly available 12
months after official publication or later and
provided acknowledgement is given to the
original source of publication and a link is
inserted to the published article on Springer's
website. The link must be accompanied by
the following text: "The final publication is
available at link.springer.com”.

4. Development of Through-Thickness Cube
Recrystallization Texture in Non-oriented Electrical
Steels by Optimizing Nucleation Environment
NING SHAN, JINLONG LIU, YUHUI SHA, FANG ZHANG, and LIANG ZUO
Texture evolution of 2.1 wt pct Si non-oriented electrical steel sheets was investigated using
macro- and micro-texture analysis. The desirable cube ({001}h100i) component successfully
dominates the recrystallization texture through sheet thickness. The nucleation sites of cube
grains are mainly identified at the interfaces of {001}h230i-{001}h130i and
{223}h362i-{114}h481i oriented deformation bands. The formation of through-thickness cube
recrystallization texture can be attributed to the optimization of nucleation environment,
featuring quantitative advantage of cube nuclei at both strong plane strain and strong shear
strain layers under the superiority of locally low density of cube nuclei.
https://doi.org/10.1007/s11661-019-05167-3
The Minerals, Metals Materials Society and ASM International 2019
I. INTRODUCTION
NON-ORIENTED electrical steels (NOES) are
widely used in iron cores with alternating magnetic flux,
such as generators and motors. The various fabricating
methods of NOES have been reported, including con-
ventional rolling and annealing, powder metallurgy,
selective laser melting, CVD, PVD, rapid solidification,
spray forming, hot dipping, strip casting, etc.[1,2]
The
conventional rolling and annealing process is the most
common method for fabricating NOES sheets and
widely used for large-scale industrial production. The
magnetic properties of NOES are highly sensitive to the
texture, and h100i is the easy magnetization direction. In
order to achieve a high magnetic flux density at weak
magnetic field, it’s beneficial to form a cube texture
which contains two h100i along rolling direction (RD)
and transverse direction (TD), respectively. However,
through-thickness c fiber (h111i//normal direction
[ND]), which has no h100i in the rolling plane, usually
dominates the recrystallized sheets of NOES after final
annealing[3,4]
due to the orientation-dependent stored
strain energy.[5]
Currently, texture optimization in NOES mainly
focuses on two issues, reducing c texture and enhancing
k texture (h100i//ND) including cube texture. As c
recrystallized grains mainly nucleate at grain bound-
aries,[6,7]
modifying the nucleation environment at grain
boundaries is a considered way to control c texture.
Park et al.[8]
reduced the number of c nuclei at grain
boundaries in 2.0 wt pct Si NOES by increasing grain
sizes of hot bands from 115 to 460 lm. And Cunha
et al.[9]
weakened the strain-stored energy of grain
boundaries which provide activation energy of c recrys-
tallized grains using a two-stage cold rolling process. In
addition, temper rolling and pre-annealing was also
proved to effectively suppress c recrystallization texture
by Grégori et al.[10]
But it is difficult to completely
eliminate c texture through all these methods.
On the other hand, cube texture has won a high
degree of attention not only in NOES but in grain-ori-
ented electrical steels (GOES). Actually, cube is difficult
to develop as a dominant recrystallization texture in
virtue of the disadvantages of the size or quantity in the
early stage of recrystallization produced by the conven-
tional rolling method.[7]
It is easy to form strong Goss
and a (h110i//RD) textures in bands after conventional
hot rolling under heavy deformation at high tempera-
ture. These textures have an important effect on texture
evolution. In NOES, they convert to c or Goss recrys-
tallized grains after cold rolling and primary recrystal-
lization annealing, while in GOES, Goss grains in hot
bands are proved to be the origins of abnormal growth
of Goss grains.[7,11,12]
These evolutions prefer the
formation of the c or Goss texture, and make it difficult
to obtain favorable nucleation environment for cube
grains. Therefore, some special methods were adopted
for cube recrystallization texture, including using colum-
nar-grained polycrystals as starting material, cross
rolling, skew rolling, surface annealing, strip casting,
NING SHAN, JINLONG LIU, YUHUI SHA, FANG ZHANG,
and LIANG ZUO are with the Key Laboratory for Anisotropy and
Texture of Materials (Ministry of Education), School of Materials
Science and Engineering, Northeastern University, No. 3-11, Wenhua
Road, Heping District, Shenyang 110819, China; Contact e-mails:
liujl@atm.neu.edu.cn; yhsha@mail.neu.edu.cn
Manuscript submitted September 22, 2018.
Article published online March 6, 2019
2486—VOLUME 50A, MAY 2019 METALLURGICAL AND MATERIALS TRANSACTIONS A
Author's personal copy

5. and phase transformation.[13–19]
However, the complex-
ity of these methods induces high cost of production or
equipment investment and limits large-scale industrial
application.
Recently, warm rolling is proved to be an effective
way to control recrystallization texture of NOES, and it
has been found that the recrystallization texture is
closely related to the content of interstitial atoms (C, N)
and the process parameters of warm rolling.[20,21]
The
reports showed that the warm rolling temperature at
which dynamic strain aging (DSA) occurs was beneficial
to the formation of shear bands and further promoting
the preferential development of Goss recrystallization
texture in NOES and low C steel,[20–24]
based on the
widely observation of nucleation of Goss at shear
bands in the deformed grains with some special orien-
tations, such as {111}h112i, {111}h110i, and
{112}h110i.[7,20,25–28]
In contrast, when warm rolling at
higher temperatures, strongly positive rate sensitivity
effectively eliminates the shear banding tendency, thus
the nucleation can be promoted at the interface of
deformation bands or deformed grains.[21–24]
Moreover,
a relatively higher strain-stored energy was obtained for
the k grains when warm rolling above DSA tempera-
ture.[20]
In this way the competitiveness of cube grains is
effectively promoted at the early stage of recrystalliza-
tion, and then some weak cube and rotated cube
recrystallization textures remain after warm rolling at
700 C and annealing.[21]
In order to obtain strong cube
texture, more experiments and in-depth discussions are
needed for understanding the environment of formation
and development of cube recrystallization texture,
especially through-thickness cube texture due to the
remarkably inhomogeneous distribution of strain along
sheet thickness under rolling deformation condition.
In current work, through-thickness cube recrystalliza-
tion texture was successfully obtained under appropriate
warm rolling parameters in 2.1 wt pct Si NOES pro-
duced by conventional rolling and annealing method.
Microstructure and texture evolution were further
investigated to explore the origins and control of cube
recrystallization texture.
II. EXPERIMENTAL PROCEDURE
Fe-2.1 wt pct Si ingots, which contain 0.01 wt pct C,
2.1 wt pct Si, 0.2 wt pct Mn, 0.002 wt pct Al,
0.004 wt pct S, 0.017 wt pct P, and balance Fe, were
prepared by induction melting and homogenized at
1473 K (1200 C) for 120 minutes. The ingots were
forged into 140 mm 9 110 mm plates and then hot
rolled from 80 to 3.5 mm with the finishing temperature
of 1223 K (950 C). The microstructure and texture of
hot bands have been reported in the previous paper.[29]
Afterwards, hot bands were normalized at 1303 K
(1030 C) for 10 minutes and further rolled at 773 K
(500 C) with rolling reductions of 60, 70, 80, and 90
pct. Some sheets with different rolling reductions were
annealed in Ar atmosphere at 1123 K (850 C) for 10
minutes to investigate the recrystallization texture after
considerable grain growth. Other sheets of 70 pct rolling
reduction were further annealed at 973 K (700 C) for
1.7 minutes, 973 K (700 C) for 3 minutes, and 1273 K
(1000 C) for 10 minutes to obtain 7 pct recrystallized
fraction, 44 pct recrystallized fraction, and the stage of
grain growth for texture evolution analysis.
The macro-texture of the specimens was measured
and studied by an X’Pert PRO X-ray diffractometer
with a radiation source of Co Ka operated at 35 kV and
40 mA through sheet thickness layer. Here, the thick-
ness layer is defined as the parameters S = 2a/d, where
a represents the distance away from the center layer and
d is the whole sheet thickness. The micro-texture of the
specimens was investigated and analyzed via electron
backscattered diffraction (EBSD) technique. EBSD
specimens were electro-polished under the voltage of
20 V for 15 s in an 8 pct perchloric acid/alcohol
solution. EBSD tests were performed in a field emission
gun scanning electron microscope (FEG-SEM, JEOL
JSM-7001F) with an electron accelerating voltage of
15 kV at a working distance of 15 mm, and the selected
step size was marked in the corresponding maps. The
EBSD data were analyzed with the HKL Channel 5
software. Referring to the previous reports,[18,30–32]
the
recrystallized grain based on EBSD data is defined as an
area that simultaneously satisfies the following rules: (a)
surrounded by grain boundaries having misorientation
greater than 5 deg; (b) having internal average misori-
entation less than 1 deg; and (c) its equivalent size larger
than three times of the selected step size. According to
the reports,[18,30–32]
the texture components in EBSD
analysis are defined within 15 deg around the ideal
orientations. The individual orientation of the selected
region is determined by the statistical result of the ODFs
calculated at each point in the interior of this region.
The magnetic flux density at 5000 A/m (B50) was
measured using a single sheet tester (IWATSU, B-H
analyzer SY8232) along the rolling direction of the
sheets annealed at 1273 K (1000 C) for 10 minutes with
100 mm (RD) 9 30 mm (TD) size.
III. RESULTS
A. Texture and Magnetic Properties After Warm Rolling
and Annealing
Figure 1 shows warm rolling microstructure and
texture of 2.1 wt pct Si NOES. The microstructure is
characterized with grains extending along rolling direc-
tion. And a small number of elongated grains contain
shear bands with an inclination angle of ~30 deg. Before
warm rolling, Goss and cube components dominate
S = 0.5 and S = 0 layers, respectively. And after warm
rolling, textures are mainly composed of strong a and c
fibers and weak cube component under various rolling
reductions. As rolling reductions increase, {001}h110i
and {111}h112i components increase, while cube com-
ponent decreases. Furthermore, all warm rolling com-
ponents show obvious texture gradient along sheet
thickness. Compared with the S = 0.5 layer, the S = 0
layer has much strong {001}h110i and cube components
which belong to plane strain texture, and relatively weak
METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 50A, MAY 2019—2487
Author's personal copy

6. {111}h112i component which locates at the intersection
of plane strain texture and shear texture in Euler space.
Hereinafter, the vicinity of the S = 0.5 and S = 0
layers are referred to as a region of strong shear strain
(SSS) and strong plane strain (SPS), respectively.
After annealing at 1123 K (850 C) for 10 minutes,
complete recrystallization occurs in all of the warm
rolled sheets under different rolling reductions. The
most striking feature of annealing texture is that
through-thickness cube component dominates the
recrystallization texture under 70 pct rolling reduction
(Figures 2(a) and (b)). Moreover, cube texture slightly
enhances with the value of S decreasing and has the
highest orientation density at S = 0 layer.
It should be noted that the symmetrically equivalent
positions of {111}h112i and {111}h110i orientations are
twice as large as those of cube and Goss orientations due
to the cubic crystal symmetry and orthogonal sample
symmetry in rolling, thus the observed intensities of
{111}h112i and {111}h110i components should be
increased by a factor of 2 in order to be comparable
to those of cube and Goss components.[33]
After
considering this symmetry, cube component still dom-
inates the recrystallization texture in the sample for 70
pct reduction due to the extremely weak {111}h112i and
{111}h110i components as shown in Figure 2(b).
The characteristic of recrystallization texture in pre-
sent study is evidently different from that of previously
reported NOES produced by conventional rolling meth-
ods, in which through-thickness c fiber usually domi-
nates the primary recrystallization texture.[3,4]
As the
warm rolling reductions increase to 80 pct, Goss
significantly enhances and has the similar intensity to
the cube at S = 0 layer. With the further increase of the
rolling reductions to 90 pct, the intensity of Goss
continues to increase and cube component shows
distinct weakness.
In view of the dominant cube texture obtained in
Figure 2(b), the warm rolled sheets under 70 pct
reduction were further annealed at 1273 K (1000 C)
Fig. 1—Warm rolling microstructure and texture of Fe-2.1 wt pct Si sheets: (a) microstructure under 70 pct rolling reduction, (b) constant
u2 = 45 deg section of ODFs at different layers under 70 pct rolling reduction, and (c) orientation densities of main texture components at
different rolling reductions.
Fig. 2—(a) Constant u2 = 0 and 45 deg sections of ODFs at S = 0.5 layer under 70 pct rolling reduction, (b) orientation densities of main
texture components at different thickness layers under 70 pct rolling reduction, and (c) orientation densities of main texture components under
different rolling reductions in Fe-2.1 wt pct Si sheets after annealing at 1123 K (850 C) for 10 min.
2488—VOLUME 50A, MAY 2019 METALLURGICAL AND MATERIALS TRANSACTIONS A
Author's personal copy

7. for 10 minutes and the texture is measured in OIM areas
of 29 9 106
lm2
and represented in Figure 3. Recrys-
tallization texture is also dominated by cube component.
Moreover, cube grains have the largest quantity fraction
and relatively larger average grain size among the grains
with different orientations, leading to the clearly
increased magnetic induction (B50) of 1.794 T, which
are clearly higher than reported samples between 1.736
and 1.776 T in ~2 wt pct Si NOES.[34–36]
B. Nucleation of Cube Grains
In order to elucidate the origin of cube recrystalliza-
tion texture, the 7 pct recrystallized sheets under 70 pct
rolling reduction were analyzed by EBSD. Goss grains
are found to nucleate at the shear bands in c deformed
matrix at surface (Figure 4(a)). Cube grains A, C, and D
nucleate at the grain boundaries of {001}h120i deformed
matrix at S = 0.3-0.4 layers, {113}h361i deformed
matrix at S = 0.2-0.3 layers, and {114}h481i deformed
matrix at S = 0.5-0.6 layers, respectively, and cube
grain B nucleates in the interior of {113}h361i deformed
matrix as shown in Figure 4(b).
Since a large number of grain boundaries have not
nucleated in the stage of 7 pct recrystallization, the
microstructure and texture in 44 pct recrystallized
samples were further observed in 3.23 9 106
lm2
OIM. In this stage, the texture of recrystallized grains
is composed of strong cube as well as weak c and Goss
at both SSS and SPS layers (Figure 5). Figure 6(a)
shows more nucleation sites of cube grains. It is clearly
observed that the nucleation environment of cube grains
is closely related with the deformed matrices with some
special orientations, including {001}h230i, {001}h130i,
{113}h361i, {112}h241i, and {223}h362i, although cube
grains have a considerable growth. These sites can be
divided into two categories (Figure 6(b)): Region I from
{001}h230i to {001}h130i which belongs to the {100}
texture and Region II from {223}h362i to {114}h481i
which belongs to the R texture.[37]
It should be noted
that the density of cube nuclei is remarkably low in the
specially oriented deformation bands (Figures 4 and 6)
in virtue of relatively small strain-stored energy.
Fig. 3—(a) Orientation image maps, (b) constant u2 = 0 and 45 deg sections of ODFs, and (c) number fraction and average grain size of main
texture components at S = 0 layer in Fe-2.1 wt pct Si sheets after annealing at 1273 K (1000 C) for 10 min.
Fig. 4—(a) Orientation image maps and (b) local enlarged maps in Fig. 4(a) and corresponding u2 = 45 deg section of ODFs of cube
recrystallized grains and surrounded deformed matrices in 7 pct recrystallized Fe-2.1 wt pct Si sheets.
METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 50A, MAY 2019—2489
Author's personal copy

8. In addition, the cube texture has an evidently larger
area fraction than Goss, {111}h112i, or {111}h110i
texture (Figure 7(a)). Figure 7(b) further gives the
number fractions of specific textures in various layers.
It can be seen that the number fraction of cube grains
becomes the dominant one at all layers. This demon-
strates distinctly that cube grains have a prominent
number advantage for the formation of dominant cube
recrystallization texture during subsequent grain
growth.
Fig. 5—(a) Orientation image maps and (b) constant u2 = 0 and 45 deg sections of ODFs of recrystallized grains within different thickness
layers in 44 pct recrystallized Fe-2.1 wt pct Si sheets.
Fig. 6—(a) Orientation image maps and corresponding constant u2 = 45 deg section of ODFs of cube recrystallized grains and surrounded
deformed matrices and (b) orientations of deformed matrices in Fig. 6(a) marked in the u2 = 45 deg section in 44 pct recrystallized
Fe-2.1 wt pct Si sheets.
Fig. 7—(a) Area fraction and (b) number fraction of main texture components of recrystallized grains in 44 pct recrystallized Fe-2.1 wt pct Si
sheets.
2490—VOLUME 50A, MAY 2019 METALLURGICAL AND MATERIALS TRANSACTIONS A
Author's personal copy

9. IV. DISCUSSION
Then some interest questions are submitted, for
example, where are these cube nuclei originating from?
Why through-thickness cube recrystallization texture
forms under the conditions of 70 pct rolling reduction?
A. Origin of Through-Thickness Cube Texture
It is well known that the main convergent orientations
during the rolling process are a and c fibers in bcc
alloys.[38]
Supposing no initial cube orientation, other
orientations are difficult to rotate to the cube orienta-
tion. Therefore, the cube nuclei in Figures 4 and 6
mainly originate from the retaining of the cube regions
of hot bands. Stojakovic et al.[39]
also found that cube
nuclei can be largely restored from hot bands based on
the calculation of crystal plasticity finite element method
(CPFEM), but they believed that this retention operated
under small rolling reduction and recrystallization by
the strain induced boundary migration (SIBM)
mechanism.
These retained cube micro-regions can exist near the
deformed matrices in the Regions I and II (Figure 6(b)).
And the deformed matrices in these regions have been
proved to be closely related to the strong initial cube
texture. Among them, deformed matrices in Region I
rotate from cube orientations with small deviation angle,
leading to the formation of {001}h230i-{001}h130i ori-
ented deformation bands.[18,40]
And the deformed matri-
ces in Region II are thought to rotate from near-cube
orientations deviated beyond 10 deg,[18,41]
resulting in the
formation of {223}h362i-{114}h481i oriented deforma-
tion bands.
Zhang et al.[13]
considered that the formation of above
deformation bands can be promoted by the adjacent
hard grains with high Taylor factor to cube grains,
because of more heterogeneous distribution of reaction
stress along the grain boundaries. Moreover, the reduc-
ing of strain hardening due to recovery during warm
rolling is beneficial to retard the appearance of shear
bands, which transfer the strain away from the matri-
ces,[22]
further driving the formation of these special
deformation bands.
There are two kinds of boundaries of deformation
bands: grain boundaries and transition bands in the
interior of grains.[42]
Some retained cube micro-regions
are found in transition bands of {113}h361i and
{223}h362i oriented deformation bands, and near grain
boundary of {001}h230i oriented deformation bands, as
shown in Figure 8. Obviously, the accommodation slip
at grain boundaries and transition bands promotes the
retaining and formation of these cube deformed
micro-regions when the adjacent regions rotate towards
other orientations as shown in Figures 8(b) through (d).
Because of the large misorientation within rather small
regions, grain boundaries and transition bands have
higher strain-stored energy than those in the interior of
deformation bands and the nucleation would be
enhanced at these sites.[42,43]
In the early stage of
recrystallization, the retained cube micro-regions at
these boundaries preferentially nucleate and further
Fig. 8—(a) Orientation image maps and (b, c, d) corresponding local enlarged maps in Fig. 8(a) and constant u2 = 45 deg section of ODFs of
cube micro-regions retained at (b, c) transition bands and (d) grain boundaries and surrounded deformation bands at S = 0 layer of warm
rolled Fe-2.1 wt pct Si sheets under 70 pct reduction.
METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 50A, MAY 2019—2491
Author's personal copy

10. grow into the interior of neighboring deformation
bands.
B. Development of Through-Thickness Cube Texture
Nucleation of recrystallization is considered as no
more than discontinuous subgrain growth at sites of
high strain energy. In current work, both
{001}h230i-{001}h130i and {223}h362i-{114}h481i
deformed matrices, where cube grains mainly nucleated,
have low strain-stored energy due to the relatively small
Taylor factors,[5,39,44]
so cube grains have a superiority
of locally low nucleation rate, as shown in Figure 4. If
cube nuclei are excessively dense, cube grains are prone
to clusters, resulting in orientation pinning[45,46]
and
inevitable reduction during grain growth. Furthermore,
cube grains obtain a prominent number advantage at
both SSS and SPS layers in the partially recrystallized
state, and finally develop the dominant texture through
sheet thickness after primary recrystallization. There-
fore, the formation of through-thickness cube texture is
derived from the quantitative advantage of cube nuclei
at both SPS and SSS layers under the superiority of
locally low density of cube nuclei. Based on the above
analysis, the formation process of the cube texture is
identified, and the schematic diagram is shown in
Figure 9.
C. On the Nucleation Environment at Shear Bands
Based on extensive experiments, it is agreed that cube
grains nucleate mainly at {001}h230i-{001}h130i and
{223}h362i-{114}h481i oriented deformation bands, as
well as shear bands within {111}h110i-{111}h112i and
{112}h110i-{110}h110i deformed matrices.[7,13,17,18,46,47]
It should be pointed out that there are several problems
of nucleation environment for promoting through-thick-
ness cube texture by means of controlling shear bands.
Table I briefly summarizes the nucleation of cube and
Goss grains at shear bands in different deformed
matrices through sheet thickness.
To the SPS layer, the insufficient shear strain at this
layer during rolling seriously impedes the formation of
shear bands. Shear bands tend to form in {110}h110i,
{111}h112i, and {111}h110i deformed matrices with
relatively large Taylor factor above 3.4[48]
at SSS layer
for bearing much SSS. Xu et al.[17]
and Nguyen-Minh
et al.[47]
found that cube grains can nucleate at shear
bands in {110}h110i deformed matrices by both the
experiments and the calculation based on full-con-
straints Taylor model, but the warm and cold rolled
sheets of NOES usually have a low volume fraction of
{110}h110i deformation texture.[7,18]
In the case of
{111}h112i deformed grains, Goss shear bands are
easier to form in bcc alloys according to the previous
reports[7,20,25–27]
rather than cube shear bands. More-
over, there is no clear way during rolling and annealing
to ensure the formation of cube shear bands rather than
Goss shear bands in {111}h112i, {111}h110i, and
{112}h110i deformed matrices,[7,18,20,25–28,46]
although
shear bands in {111}h110i deformed grains have been
the mainly nucleation sites of cube recrystallized grains
observed in 2.8 wt pct Si NOES.[18]
Therefore, opti-
mization of nucleation environment at deformation
bands is a more stable way for the control of
through-thickness cube texture in the process of NOES
production.
D. Effect of Warm Rolling on the Nucleation
Environment of Cube Grains
Three critical factors need to be satisfied to optimize
the nucleation environment of cube grains in warm
rolled NOES sheets: (i) enough initial cube or near-cube
grains in hot bands as the origins, (ii) the high warm
rolling temperature to prevent the occurrence of DSA
which promotes the shear banding tendency favorable
to the Goss texture evolution, and (iii) optimum warm
Fig. 9—Schematic representation of the formation of through-thickness cube texture during primary recrystallization: (a) warm rolling
microstructure, (b) formation of cube grains at the boundaries of deformation bands with orientation in Regions I and II in the early stage of
recrystallization, and (c) formation of cube texture after primary recrystallization.
2492—VOLUME 50A, MAY 2019 METALLURGICAL AND MATERIALS TRANSACTIONS A
Author's personal copy

11. rolling reductions to balance the quantity and strain-s-
tored energy of deformation bands for their effective
nucleation at an early stage of recrystallization. Because
of low strain-stored energy, the sites of the cube nuclei
need relatively large strain to supply sufficient activation
energy for recrystallization. Meanwhile, the number of
remained cube micro-regions originated from hot bands
will gradually decrease under the heavy strain due to the
characteristics of metastability of cube orientation under
deformation conditions in BCC alloys.[43]
In the present work, the facilitating role of warm
rolling on the formation and development of cube
texture is mainly achieved by inhibiting the nucleation of
c and Goss grains. Compared with rolling at room
temperature,[20]
it can be seen from Figure 5 that the
nucleation of c grains is obviously inhibited during
recrystallization in warm rolled sheets due to the
insufficient strain-stored energy at c deformed grain
boundaries. Compared with warm rolling in the tem-
perature range where DSA phenomenon occurs,[20–22]
the shear bands are effectively inhibited during warm
rolling at high temperature as shown in Figure 1.
Therefore, cube grains can obtain the competitive
advantage of nucleation environment through sheet
thickness in the early stage of recrystallization as shown
in Figures 4 and 5.
The optimizing nucleation environment of cube grains
also requires a reasonable warm rolling reduction. When
rolling at 773 K (500 C) under 60 pct rolling reduction,
the nucleation of cube deformed micro-regions is
retarded by relatively low strain-stored energy, which
is further weakened by dynamic recovery during rolling.
As rolling reductions reach up to 70 pct, a considerable
number of {001}h230i-{001}h130i and
{223}h362i-{114}h481i oriented deformation bands
form at different thickness layers in warm rolled sheets.
These deformation bands obtain enough activation
energy by accumulated plastic strain, while the quantity
of remained cube micro-regions gradually decreases.
With increasing of rolling reductions to 90 pct, the
strain-stored energy of deformation bands increases
with the further accumulated plastic strain, whereas the
quantity of remained cube micro-regions is insufficient.
Consequently, it can be deduced that the applied
processing parameters are responsible for both the
sufficient quantity of remained cube micro-regions and
enough strain-stored energy of deformation bands at
different thickness layers in warm rolled sheets, which
provide the favorable nucleation environment for cube
grains, and thus for the development of through-thick-
ness cube recrystallization texture.
V. CONCLUSIONS
A through-thickness cube recrystallization texture is
successfully obtained in 2.1 wt pct Si NOES after final
annealing by designing the initial texture and the applied
processing parameters. Cube grains mainly nucleate at
the boundaries of {001}h230i-{001}h130i and
{223}h362i-{114}h481i oriented deformation bands.
The formation of through-thickness cube recrystalliza-
tion texture is attributed to the optimization of nucle-
ation environment, featuring quantitative advantage of
cube nuclei at both SPS and SSS layers under the
superiority of locally low density of cube nuclei. No
special requirements for the initial columnar grains or
twin-roll casting, as well as the simple processing route
of rolling and annealing, are expected to provide an
efficient and significantly cheap way to optimize texture
of NOES for large-scale industrial applications.
ACKNOWLEDGMENTS
This work was supported by the National Key
Research and Development Program of China
(2016YFB0300305), the National Natural Science
Foundation of China (51671049), the Fundamental
Research Funds for the Central Universities
(N170213019), and the China Scholarship Council
(CSC) (201806085006).
REFERENCES
1. N.R. Overman, X.J. Jiang, R.K. Kukkadapu, T. Clark, T.J.
Roosendaal, G. Coffey, J.E. Shield, and S.N. Mathaudhu: Mater.
Charact., 2018, vol. 136, pp. 212–20.
2. M. Garibaldi, I. Ashcroft, M. Simonelli, and R. Hague: Acta
Mater., 2016, vol. 110, pp. 207–16.
3. S.K. Chang: J. Mater. Sci., 2006, vol. 41 (22), pp. 7380–86.
4. Y. Wan, W.Q. Chen, and S.J. Wu: J. Rare Earths, 2013, vol. 31
(7), pp. 727–33.
5. N. Rajmohan, Y. Hayakawa, J.A. Szpunar, and J.H. Root: Acta
Mater., 1997, vol. 45 (6), pp. 2485–94.
6. H. Inagaki: Trans. JIM, 1987, vol. 28 (4), pp. 251–63.
Table I. Nucleation of Cube and Goss at Shear Bands Through Sheet Thickness
Thickness Layer S
Deformed Matrices
Major Orientation of
Nuclei at Shear Bands
Orientation
Taylor
Factor[5]
The Formation
of Shear Bands
S = 0.25–1
SSS Layer
{110}h110i 4.3 easy cube[17,47]
{111}h112i 3.6 easy cube,[7]
goss[7,20,25–27]
{111}h110i 3.79 easy cube,[18]
goss[7]
{112}h110i 3.17 difficult near-cube,[46]
Goss[7,28]
S = 0–0.25
SPS Layer
insufficient shear strain difficult —
METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 50A, MAY 2019—2493
Author's personal copy

12. 7. J.T. Park and J.A. Szpunar: Acta Mater., 2003, vol. 51 (11),
pp. 3037–51.
8. J.T. Park and J.A. Szpunar: J. Magn. Magn. Mater., 2009, vol. 321
(13), pp. 1928–32.
9. M.A. Cunha and S.C. Paolinelli: J. Magn. Magn. Mater., 2008,
vol. 320 (20), pp. 2485–89.
10. F. Grégori, K. Murakami, and B. Bacroix: J. Mater. Sci., 2014,
vol. 49 (4), pp. 1764–75.
11. S.K. Nam, G.H. Kim, D.N. Lee, and I. Kim: Metall. Mater.
Trans. A, 2018, vol. 49A (5), pp. 1841–50.
12. M. Matsuo, T. Sakai, and Y. Suga: Metall. Trans. A, 1986,
vol. 17A (8), pp. 1313–22.
13. N. Zhang, P. Yang, and W.M. Mao: Mater. Lett., 2013, vol. 93,
pp. 363–65.
14. A. Sonboli, M.R. Toroghinejad, H. Edris, and J.A. Szpunar: J.
Magn. Magn. Mater., 2015, vol. 385, pp. 331–38.
15. M. Sanjari, Y. He, E.J. Hilinski, S. Yue, and L.A.I. Kestens: J.
Mater. Sci., 2017, vol. 52 (6), pp. 3281–3300.
16. R.Y. Liang, P. Yang, and W.M. Mao: J. Magn. Magn. Mater.,
2018, vol. 457, pp. 38–45.
17. Y.B. Xu, Y.X. Zhang, Y. Wang, C.G. Li, G.M. Cao, Z.Y. Liu,
and G.D. Wang: Scripta Mater., 2014, vol. 87, pp. 17–20.
18. Y.H. Sha, C. Sun, F. Zhang, D. Patel, X. Chen, S.R. Kalidindi,
and L. Zuo: Acta Mater., 2014, vol. 76, pp. 106–17.
19. T. Tomida, N. Sano, K. Ueda, K. Fujiwara, and N. Takahashi: J.
Magn. Magn. Mater., 2003, vols. 254–255, pp. 315–17.
20. S. Lee and B.C. De Cooman: ISIJ Int., 2011, vol. 51 (9),
pp. 1545–52.
21. M. Atake, M. Barnett, B. Hutchinson, and K. Ushioda: Acta
Mater., 2015, vol. 96, pp. 410–19.
22. M.R. Barnett and J.J. Jonas: ISIJ Int., 1997, vol. 37 (7),
pp. 697–705.
23. B. Hutchinson: Philos. Trans. R. Soc. Lond. A, 1999, vol. 357
(1756), pp. 1471–85.
24. M.R. Barnett and J.J. Jonas: ISIJ Int., 1997, vol. 37 (7),
pp. 706–14.
25. T. Haratani, W.B. Hutchinson, I.L. Dillamore, and P. Bate: Met.
Sci., 1984, vol. 18 (2), pp. 57–66.
26. K. Ushioda and W.B. Hutchinson: ISIJ Int., 1989, vol. 29 (10),
pp. 862–67.
27. D. Dorner, S. Zaefferer, and D. Raabe: Acta Mater., 2007, vol. 55
(7), pp. 2519–30.
28. H.T. Liu, J. Schneider, H.L. Li, Y. Sun, F. Gao, H.H. Lu, H.Y.
Song, L. Li, D.Q. Geng, Z.Y. Liu, and G.D. Wang: J. Magn.
Magn. Mater., 2015, vol. 374, pp. 577–86.
29. N. Shan, Y.H. Sha, F. Zhang, J.L. Liu, and L. Zuo: Metall. Mater.
Trans. A, 2016, vol. 47A (12), pp. 5777–82.
30. H. Nakamichi, F.J. Humphreys, and I. Brough: J. Microsc., 2008,
vol. 230 (3), pp. 464–71.
31. D.I. Kim, J.S. Kim, J.H. Kim, and S.H. Choi: Acta Mater., 2014,
vol. 68, pp. 9–18.
32. W.C. Hsu, L.W. Chang, P.W. Kao, and I.C. Hsiao: ISIJ Int.,
2018, vol. 58 (5), pp. 958–64.
33. S. Mishra, C. Därmann, and K. Lücke: Acta Metall., 1984, vol. 32
(12), pp. 2185–2201.
34. C.G. Pei, P.K. Bai, and Z.X. Guo: Appl. Mech. Mater., 2013,
vols. 423–426, pp. 286–89.
35. S.C. Paolinelli, M.A. Cunha, and A.B. Cota: Mater. Sci. Forum,
2007, vols. 558–559, pp. 787–92.
36. K.M. Lee, M.Y. Huh, H.J. Lee, J.T. Park, J.S. Kim, E.J. Shin, and
O. Engler: J. Magn. Magn. Mater., 2015, vol. 396, pp. 53–64.
37. P. Gobernado, R.H. Petrov, and L.A.I. Kestens: Scripta Mater.,
2012, vol. 66 (9), pp. 623–26.
38. H. Inagaki: ISIJ Int., 1994, vol. 34 (4), pp. 313–21.
39. D. Stojakovic, R.D. Doherty, S.R. Kalidindi, and F.J.G.
Landgraf: Metall. Mater. Trans. A, 2008, vol. 39A (7),
pp. 1738–46.
40. H. Abe, M. Matsuo, and K. Ito: Trans. JIM, 1963, vol. 4 (1),
pp. 28–32.
41. I. Samajdar and R.D. Doherty: Acta Mater., 1998, vol. 46 (9),
pp. 3145–58.
42. Y.Y. Tse, G.L. Liu, and B.J. Duggan: Scripta Mater., 1999, vol. 42
(1), pp. 25–30.
43. I.L. Dillamore, P.L. Morris, C.J.E. Smith, and W.B. Hutchinson:
Proc. R. Soc. Lond. A, 1972, vol. 329 (1579), pp. 405–20.
44. S.F. Castro, J. Gallego, F.J.G. Landgraf, and H.J. Kestenbach:
Mater. Sci. Eng. A, 2006, vol. 427 (1–2), pp. 301–05.
45. O. Engler: Acta Mater., 1998, vol. 46 (5), pp. 1555–68.
46. H.T. Liu, Z.Y. Liu, Y. Sun, Y.Q. Qiu, C.G. Li, G.M. Cao, B.D.
Hong, S.H. Kim, and G.D. Wang: Mater. Lett., 2012, vol. 81,
pp. 65–68.
47. T. Nguyen-Minh, J.J. Sidor, R.H. Petrov, and L.A.I. Kestens:
Scripta Mater., 2012, vol. 67 (12), pp. 935–38.
48. L.A.I. Kestens and J.J. Jonas: Metall. Mater. Trans. A, 1996,
vol. 27A (1), pp. 155–64.
Publisher’s Note Springer Nature remains neutral with regard to
jurisdictional claims in published maps and institutional affiliations.
2494—VOLUME 50A, MAY 2019 METALLURGICAL AND MATERIALS TRANSACTIONS A
Author's personal copy
View publication stats
View publication stats