Abstract: A partial explanation of the phenomenon of abnormal growth of Goss oriented grains in silicon iron is proposed. It follows from an observation that texture evolves toward Goss orientation if grains of low surface energy have high probability of growth. A corresponding simulation starts with primary recrystallization texture and results in strong Goss texture. To explain development of early secondaries in subsurface layers, an option is considered that besides capillary forces, the growth is caused by surface induced lattice instabilities

A Morawiec

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This is an electronic version of an article whose final and definitive form has been published
in Scripta Mater. 64, 466-9 (2011).
Scripta Materialia is available online at: http://www.sciencedirect.com/.
On abnormal growth of Goss grains
in grain oriented silicon steel
A.Morawiec
Polish Academy of Sciences
Institute of Metallurgy and Materials Science
Reymonta 25, 30–059 Kraków, Poland
E-mail: nmmorawi@cyf-kr.edu.pl
Tel.: ++48–122952854, Fax: ++48–122952804
Abstract: A partial explanation of the phenomenon of abnormal growth of Goss oriented
grains in silicon iron is proposed. It follows from an observation that texture evolves to-
ward Goss orientation if grains of low surface energy have high probability of growth. A
corresponding simulation starts with primary recrystallization texture and results in strong
Goss texture. To explain development of early secondaries in subsurface layers, an option
is considered that besides capillary forces, the growth is caused by surface induced lattice
instabilities.
Keywords: iron alloys, soft magnets, secondary recrystallization, recrystallization texture,
simulation
1
Grain oriented silicon iron (Fe ∼3%wt Si) is a conventional material for cores of electric
transformers [1, 2]. Although the use of amorphous and nanocrystalline soft magnetic al-
loys is increasing, the ’GO electrical steels’ still play a very significant role [3]. Among
factors determining their good magnetic properties is very sharp texture dominated by the
{110}<001> (Goss) component. There are a number of processing routes leading to the
Goss texture (two or one cold rolling + annealing cycles, MnS or AlN as inhibitors of normal
growth) [4]. Our considerations are limited to the phenomena that follow primary recrys-
tallization, and the mechanism of this last phase of processing is believed to be the same
for all routes. Briefly, this stage is a high temperature annealing of decarburized primary
recrystallized material (∼0.30mm thick sheet) with the initial grain size of about 20µm,
impurities in the form of dispersed–phase particles, and relatively weak texture consisting
mainly of the {111}<uvw> (γ) and {hk0}<001> (η) fibers. The texture is slightly inhomo-
geneous with the η fiber stronger in the subsurface layers than at the sheet center; e.g., [5].
The annealing leads to secondary recrystallization: normal grain growth is suppressed by
precipitates while some grains (secondaries) grow abnormally to large dimensions.
The control of the sharpness of the Goss texture is crucial for the quality of the final
product. Progress in optimizing the texture is possible through understanding of the pro-
cesses involved. Thus, questions arise about the mechanism of the exaggerated growth of
grains and the cause of the selection of the particular Goss orientation. These issues have
not been explained despite intense research in this area. We will not make any attempt
to review the vast literature of the subject. Let us only note that numerous theories have
been proposed and then abandoned (e.g., size advantage of Goss oriented grains, greater
perfection of their interior, suitably oriented strains directing orientation changes). One of
the lasting postulates is that the secondary recrystallization is driven by capillary forces, i.e.,
that interfacial (surface and/or boundary) tension provides energy for boundary migrations,
which change proportions of texture components. Philip and Lenhart [6] noticed that early
secondaries are located under the surface inside the specimen; this led to the conclusion
that grain boundary energy, and not surface energy, is the actual driving force. Currently, a
common view is that ”secondary recrystallization is a process of nucleation followed by grain
growth. The driving force for this process is the reduction of grain boundary energy” [7,8].
All recent theories (high mobility of CSL boundaries [9, 10] or high energy boundaries [11]
caused by preferential particle coarsening, solid state wetting [12]) are based on the above
principle, with special properties ascribed to boundaries characterized by particular grain
misorientations. However, the misorientation–based models do not explain the sharpness of
the final texture: in such models, the selection process is determined by the orientations
of the primary grains, and the spread of these orientations is large; cf. [13, 14]. Also the
principle that growth is driven by grain boundaries becomes problematic in confrontation
2
with the intricate early stage of secondary recrystallization, when Goss grains are in the
minority and not larger than other grains. To describe this phase, the principle must be
augmented, and an additional driving force needs to be indicated.
Below, an explanation of the rise of the Goss component is proposed. Briefly, it is
shown that growth of grains with {hk0} planes parallel to the rolling plane, with particular
emphasis on {110}, leads to the Goss texture. An important characteristic of the {hk0}
planes, in particular {110} planes, is their low surface free energy. Therefore, the texture
evolution toward the Goss orientation follows, if the orientation changes are assumed to be
influenced by the surface of the sheet. It will only remain to clarify the mechanism of the
transmission of the changes to other grains. It is possible that driving force coming from the
reduction of interfacial energy may work in combination with other mechanisms of lattice
reorientations. With the involvement of surfaces, the explanation partly departs form the
keystone assumption of recent theories that the process is determined by grain boundaries.
The main aspects of the proposed growth scheme are considered below in more detail.
Let us begin with the issue of texture changes. It is known that the primary texture
initially evolves toward the η fiber, and only later the Goss component gets increasingly
sharper; e.g., [5]. In light of this, the following observations are of interest: If only grains with
{hk0} planes close to the rolling plane have high probability of growth, the resulting texture
will be dominated by the η fiber because there are no other primary texture components of
the type {hk0}<uvw>. Moreover, if only grains with {110} planes parallel to the rolling
plane have high probability of growth, the resulting texture will be dominated by the Goss
component because there are no other significant {110}<uvw> components in the primary
texture. Consequently, these two theoretical facts are sufficient for faithful modeling of the
actual texture evolution.
The above statements have been confirmed by a simple computer simulation. The sim-
ulation was based on the frequency of occurrence of particular orientations without spatial
relationships between grains. In effect, it involved only texture induced correlations be-
tween orientations of grains. In each simulation step, a randomly selected ’grain’ grew at
the expense of another ’grain’. The probability of growth depended on ’grain’ orientation
with respect to the rolling plane. For a given grain, it was, respectively, 50% or 45% if
one of its {110} or other {hk0} planes was within 7◦ to the rolling plane; otherwise, the
probability was only 10%. These probabilities agree with with the rules specified in the
previous paragraph. At the outset of the simulation, there were N = 3 × 106 ’grains’ of
equal size. Their orientations were generated in such a way that the orientation distribution
corresponded to an experimental primary recrystallization texture (taken from [15]). The
total number of ’grains’ declined with the progress of the simulation process. After some
steps, the texture contained mostly η fiber, and with a sufficiently large number of steps,
3
Goss became the only strong component; see Fig. 1. Thus, the texture changes in the simu-
lation are analogous to those occurring during secondary recrystallization, and these results
substantiate the theoretical claims made above. The simulation is stable with respect to
the numbers (probabilities of growth) used; different values give different texture intensities
but the trends are the same. We are not aware of any other simulations based on similarly
simple postulates, which would mimic the experimental texture evolution so closely.
The crystallographic planes pointed out in previous paragraphs correspond to low surface
free energy. Although experimental data on surface energy are strongly influenced by the
environment, impurities, adsorption and diffusion from the bulk, there are indications that
the energy has the lowest values at the {110} and {100} planes [16–18]. Results of theoretical
calculations also vary considerably but data for iron hint that surface energies corresponding
to {110} and other low index {hk0} planes are low compared to those of {111} and high
index planes; see, e.g., [19, 20].
The differences in the surface free energies of particular crystallographic planes are
proven to play a leading role in secondary recrystallization in thin–gauge specimens of high
purity Fe–Si [21, 22]. It is likely that the surface energy also plays a part in the secondary
recrystallization in the presence of second phase particles. Moreover, there is a general ar-
gument for the involvement of the surface: for Goss selection, the orientation changes must
be governed by a factor strongly linked to the sample coordinate system, and besides the
rolling direction, this system is determined by planes parallel to the surface.
Based on the above observations, it is clear that the texture evolution could be explained
by the growth driven by surface free energy. This idea, however, is inconsistent with afore-
mentioned conclusion of Philip and Lenhart [6] and with similar subsequent observations
that early secondaries occupy the subsurface layer, and the advancement of change along the
very surface is retarded, e.g., [23]. Therefore, one cannot rely on the conventional growth
model with low surface energy driving boundaries which come into contact with the surface,
and one must allow for other growth mechanisms to be involved. We are not at the position
to indicate a clear solution to this issue. Let us only mention that the orientation changes
may not necessarily start at (rare) Goss nuclei, and proceed by the reduction of interface
energy. They may, for instance, be caused by lattice instabilities similar to those preced-
ing diffusionless structural transformations, with mechanical energy necessary for atomic
displacements provided by thermal vibrations; see, e.g., [24]. Their amplitudes in ”soft” di-
rections can be sufficiently large to cause displacements resulting in a more stable reoriented
lattice. Presence of such directions is implied by the high elastic anisotropy, which drasti-
cally grows with temperature (Fig. 2). On the phase diagram, the stability of the α phase
is related to the distance to the γ loop. The actual distance may be small at some locations
in the material due to solute segregation or dissolution of precipitates, and the orientation
4
changes can be triggered by local instabilities1. In the considered case, the first destabi-
lizing factor is believed to be the sheet surface. The scenario involving instability–caused
reorientations seems to be confirmed by the kinetics of orientation changes: Experiments
demonstrate that secondary recrystallization is not a smooth process but the steady growth
is accompanied by burst motions [4, 28]. It is also clearly visible that some grains become
engulfed, and growth fronts are very different from those corresponding to minimization of
boundary area.
With the above observations in mind, one can sketch a possible mechanism of Goss
selection. It is believed that orientation changes toward the η fiber and Goss orientation are
initiated at some properly oriented grains on the surface, and then the changes penetrate
to the subsurface layer and spread out there. First, crystals with low surface energy {hk0}
planes appear or grow on the surface, and this enhances the η fiber. With the {110} cusp of
the surface energy deeper than the other {hk0} cusps, the Goss grains ultimately dominate
the texture via the reduction of the surface energy. Particular aspects of the process are
subject of speculation. Besides the conventional growth driven by capillarity forces, the
changes of orientations may be caused by lattice instabilities induced at the surface or
instigated at grain boundaries by newly reoriented neighbors.
Summarizing, the key conclusion of this note is that the texture evolution and the
sharpness of final texture in grain oriented silicon iron can be explained if grains with
low surface free energy have high probability of growth. The proposed description of the
phenomenon of abnormal growth of Goss grains is intended to circumvent deficiencies of
the theories based purely on properties of grain boundaries. However, only some concepts
were outlined, and many gaps need to be filled. The main questions concern: – actual
roles played by the ’reduction of surface energy’, ’reduction of boundary energy’ and ’lattice
instability’ mechanisms, – the character of the transmission of the change to subsequent
grains, – the type of atomic displacements responsible for the orientation changes, – the
state of the surface and the impact of surface segregation on surface energy, – the impact
of coatings and annealing atmosphere, – the influence of the sheet thickness, – the kinetics
of the progress of growth toward the sheet center. Properly designed experiments may test
the validity of the proposed hypothesis and clarify the above issues.
1The underlying reasons for such instabilities are anisotropy and anharmonic behavior of the lattice [27].
5
References
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[16] B. Mills, M. McLean, E.D. Hondros, Phil. Mag. 27 (1973) 361.
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[23] N.C. Pease, D.W. Jones, M.H.L. Wise, W.B. Hutchinson, Met. Sci. 15 (1981) 203.
6
[24] T.R. Finlayson, Aust. J. Phys. 36 (1983) 553.
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[28] H. Homma, B. Hutchinson, Acta Mater. 51 (2003) 3795.
7
Captions
Figure 1: (a) Experimental distribution of orientations used as the starting point for the
simulation. The φ1–projection is used. (b) Distribution of orientations after 20 × N steps
(1154119 ’grains’); mainly η fiber. (c) Distribution of orientations after 1000 × N steps
(65733 ’grains’); mainly Goss component. (d) Values of the orientation distribution at Goss
orientation and representatives of η and γ fibers versus (number of steps)/N .
Figure 2: Schematic illustration of the dependence of Zener’s anisotropy index A on temper-
ature for Fe 2.1 wt% Si. The curve is based on data reported in [25]. Since the anisotropy
grows with Si content [26], the values of A for the considered material are higher.
8
Isolines:
1.0
2.0
3.0
4.0
5.0
Fig. 1a
A.Morawiec, ... growth of Goss grains ...
9
Isolines:
1.0
5.0
10.0
15.0
Fig. 1b
A.Morawiec, ... growth of Goss grains ...
10
Isolines:
1.0
20.0
60.0
120.0
Fig. 1c
A.Morawiec, ... growth of Goss grains ...
11
{110}<001>
{310}<001>
{111}<112>
1 10 30 100 300 1000
5
10
15
20
25
30
35
Fig. 1d
A.Morawiec, ... growth of Goss grains ...
12
100 300 500 700 900
3
4
5
6
7
A
T [C]
Fig. 2
A.Morawiec, ... growth of Goss grains ...
13


On abnormal growth of Goss grains in grain oriented silicon steel

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On abnormal growth of Goss grains in grain oriented silicon steel

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