ACICULAR FERRITE AND BAINITE ... - metal 2013

Metal 2005 24. – 26. 5. 2005, Hradec nad Moravicí
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ACICULAR FERRITE AND BAINITE MICROSTRUCTURE PROPERTIES AND
COMPARISON OF THEIR PHYSICAL METALLURGY RESPONSE
Eva Mazancová a
Zdeněk Jonšta a
Petr Wyslych a
Karel Mazanec a
a VŠB-TU Ostrava, Tř. 17. listopadu 15, 708 33 Ostrava-Poruba, ČR, eva.mazancova@vsb.cz
Abstract
The inclusions acting as preferential acicular ferrite nucleants are analysed. The
comparison with upper bainite nucleation parameter is performed. The acicular ferrite
nucleation is discussed from point of view of following parameters. Increase in chemical free
enthalpy change due to formation of the Mn-depleted zone in the neighbourhood of Ti2O3
inclusions. The role of decrease in interfacial energy between ferrite particles and nitrides is
presented. Further, the misorientation of acicular ferrite and upper bainite particles are
evaluated. The higher frequency of high-angle boundaries in acicular ferrite is described.
Simultaneously, the consequences resulting from the acicular ferrite microstructure beneficial
effect on the achievement of higher toughness in comparison with upper bainite
microstructure are clarified.
1. INTRODUCTION
Acicular ferrite (AF) is formed in the same temperature range as bainite (B) by the same
transformation mechanism. The ferrite plates in the bainitic microstructure are nucleated at
austenite (A) grain boundaries and/or at active interface allotriomorphic ferrite/A and form
packets consisting of parallel plates having similar crystallographic orientations [1]. On the
contrary, AF particles nucleate intragranularly. The AF transformation starts through a
nucleation of the primary plates at non-metallic inclusions and progresses during the
sympathetic nucleation of secondary ferrite grains nucleated at the A/primary AF plate
interface. This additional mechanism of AF formation represents a second stage in Adecomposition.
The term of AF is usually applied to describe intragranularly nucleated transformation
product by displacive mechanism characterised with a relatively wide range of morphologies
[2]. The AF plates usually exhibit the different spatial distribution and as a consequence of
this process the AF microstructures are generally make up a chaotic (interwoven) plate
arrangement. This type of microstructure is characterised as fine grained interlocking
morphology.
In medium and low carbon steels [3, 4], the AF formation is associated with beneficial
combination of strength and toughness properties. The achieved toughness values of the AF
microstructure can be related to the increased density of the high-angle boundaries found
among primary ferrite plates with high frequency [5].
The boundaries act as strong obstacles to propagation of cleavage cracks forcing these
cracks to change the microscopic propagation planes in order to accommodate the new local
crystallography. On the contrary, the low-angle boundaries are week obstacles which are not
effective hindrances in crack growth. For this reason, their influence on the toughness is
negligible. This leads to the conclusion, the crystallographic packets play an important role
and the present work is devoted to define the crystallographic parameters of AF and their
differences in comparison with B-microstructure.

Metal 2005 24. – 26. 5. 2005, Hradec nad Moravicí
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2. PROMOTION OF INTRAGRANULAR AF FORMATION DUE TO
NON_MRTALLIC NUCLEANTS
The most of the works devoted to the AF formation has been carried out on welds. The
high inclusion density which can be detected in weld materials favours AF nucleation in
comparison with B-microstructure formation. Recently, the application of this microstructure
to wrought steel types has been attempted by controlling the effect of selected non-metallic
inclusions acting as AF potential nucleants [6]. At the present time, a limited number of steels
containing nucleating non-metallic oxides particles is available. The attained results in Tideoxidized
steels are hopeful and demonstrate the perspective using of intragranular AF
nucleation process in structural steels. Under these conditions it is necessary to determine
which phases of the non-metallic inclusions, having heterogeneous character usually, are
more effective as intragranular nucleants and which nucleation mechanism is realized.
MnS
Al 2O 3.CaO
TixO
MnO.TiO 2
Fig.1a Extraction replica of Ti2O3 inclusion adhered with MnS and TiN.
Mn - depleted zone
MnS
Ti2O3
AF
TiN
Fig.1b Schematic illustration of inclusion
in steel deoxidized with Ti.
Although the mechanisms by which
inclusions nucleate AF are not fully
elucidated in detail, in following four
possible variants can be taken into
consideration: a) simple heterogeneous
nucleation on an inert nucleant; b) epitaxial
nucleation on the non-metallic inclusion
having a very good lattice registry with
ferrite matrix; c) nucleation arising from the
strain energy associated with the different
thermal expansion coefficients of matrix and
inclusion; d) nucleation accompanied with
deformation of depletion zone in the matrix
being in the close neighbourhood of
inclusion [6]. As it results from our last
investigation, the basic inclusion is also
bearing particles for further active nucleants having a higher nucleation capacity for AF
formation. Such basic inclusion is e.g. Ti2O3 formed in steel after secondary Ti-addition.
Figure 1a shows Ti2O3 extraction replica with adhered MnS and TiN. Figure 1b presents
schematic illustration of steel deoxidized with Ti-addition.
By inclusion investigation in the steel with different Ti concentration, it was found, Ti2O3
is more potent for AF nucleation than e.g. MnSiO3 (in steel without Ti-addition). In this case,
TiN

Metal 2005 24. – 26. 5. 2005, Hradec nad Moravicí
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it is useful to accept the realisation of additional nucleation mechanism to a simple
heterogeneous nucleation process. Concerning Ti2O3 it has been proposed, the local depletion
of A-stabilizing element (e.g. Mn by way of MnS formation adhered at TI-oxide) around
inclusions promotes AF nucleation due to increase of chemical free enthalpy. The discussed
depletion process is not bound with Ti2O3 (in all examples) because Mn-depletion is also
detected in steel with very low sulphur content (20ppm) accompanied with very low MnS
volume fraction. Further, a very important is the finding that increased Ti-addition leads to a
higher AF-formation probability in comparison with intergranular B-formation. It was
observed in steel without Ti-addition. For this reason it is useful to accept the proposition, the
Mn-depleted zone formation may be also directly connected with Mn diffusion into Ti2O3
inclusions. The source of this mechanism is a higher density of cation vacancies in Ti2O3 what
leads to the enrichment of this inclusion with Mn due to increased diffusion of this element
(Mn-depletion zone is found near this inclusion). It is known, Ti-oxide inclusions are
characterized as particles with higher density of cation vacancies (Al2O3 inclusions are rich on
anion vacancies and are not acting as intragranular nucleants) being contributing to Mn
diffusion into Ti-oxide inclusion [7].
3. EFFECT OF HIGH LATTICE REGISTRY OF SOME NITRIDES AND FERRITE
PARTICLES
TiN and VN nucleation activity represent a very interesting parameter influencing the AFformation.
In this case, a decrease in interfacial energy accompanied the nucleation process of
alpha phase on TiN can be held for an effective parameter influencing A-decomposition. The
effect of TiN should be also discussed because AF-formation is markedly suppressed in the
TiN absence on Ti2O3. The TiN formation is favoured by an increase in the amount of free
nitrogen in steel being available for precipitation of given nitride phase. The Ti-addition to
steel by deoxidation process can be bound as Ti2O3 right after solidification [8]. The TiNformation
is a result of an exchange reaction with Mn. At low and intermediate nitorogen
levels in steel, the MnOSiO2 particles can be formed as it results from the following reaction
[9]:
Ti2O3 + Mn + N → MnOTiO2 + TiN (1)
The same effect on AF nucleation can be found in the case of VN [10]. The effects of these
inclusions are coincident as it results from the crystallographic similarity. The interfacial
energy of the ferrite/nitride boundary is less than the interfacial energy of the A/nitride
interface (γF/nit./γA/nit.< 1). Closed matching of important planes of AF and nitride lattices
would promote this effect. The energy of A/nitride interface markedly depends on the
formation mode of such particles. The TiN particles may compose those formed in melt prior
the solidification and those precipitated from A-phase. In the case of TiN particles formed
prior to solidification the interface A/TiN is incoherent (high value of γA/nit interface).
However, since V-nitride is also important from point of view AF formation, factors
influencing nitride precipitation in A must be considered (according Ishikawa et al. [10]), the
interface properties are considered to be the same as in case of A grain boundaries.
The crystal structures and lattice parameters of discussed nitrides are as follows: TiN
(NaCl, cubic) a = 0.424nm and VN (NaCl cubic) a = 0.407nm. The corresponding lattice
dimensions of these potential AF nucleants are presented in Fig.2. The Al nitride has h.c.p.
lattice, which is very different from TiN and VN. The both considered nitrides precipitate in
austenitic and ferritic matrix as plates on {100} planes. These planes are favourable since
modulus E is low in direction. As these particles grow in size the interface A/nitride

Metal 2005 24. – 26. 5. 2005, Hradec nad Moravicí
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will become incoherent what is beneficial situation from point of view of AF formation. The
{100} planes in AF and nitride particles match as it is observed in case of thin TiN and VN
having a close matching between {100} planes. The mismatch parameters are very low δ =
0.0045 (TiN) and/or 0.007 (VN). The δ-value corresponds to the differences in nitride and
ferritic lattice parameters divided by ferritic lattice parameter.
In the case of AlN particles,
close matching between the
{100}, {110} and {111} planes
of ferrite and important planes
in the h.c.p., AlN lattice
({0001}, {1010} and {1020})
can not bee achieved. As a
result, AlN particles in A would
be unlikely sites for
intragranular AF plate
nucleation. In the case of ZrN
particles, the conditions for
matching between {100} planes
of nitride and ferritic lattice are
not so much favourable as it is
observed by TiN and VN
evaluation [11]. The TiN effect
should also be considered since
AF formation is partially
limited in the absence of TiN
nitride on Ti2O3 particles
Fig.2 Ferrite plates nucleation at intragranular nitride
particles in austenite.
(Fig.1a). This effect results
from decrease interfacial energy
for AF nucleation. Increase in
interfacial energy with
nucleation of AF is only 0.20Jm -2 at F/A semicoherent interface and reaches to 0.85Jm -2 at
A/A incoherent interface. The lattice registry between AF and TiN is very good and the
interface energy is considered to be only 0.15Jm -2 aproximatelly [12].
4. EVALUATION OF MISORIENTATION IN AF AND B PARTICLES
The nucleation conditions and resulting relationship between intragranular and
intergranular processes play a very important role by the arrangement of ferritic plates in AF
In upper B (parallel ferritic sheaves vs. interwoven AF plates directly nucleated on nonmetallic
inclusions).
The resistance to cleavage fracture in AF microstructure depends on misorientation level
of its components (particles). The successful results can be obtained by the EBSD technique
application which makes possible to study the local crystallographic properties of AF
microstructure [5, 13]. The level of misorientation angles between neighbouring AF plates
can be determined in great number of measuring. The evaluation is aimed at the finding the
occurrence frequency of high angle boundaries of interfaces representing a very important
hindrance from easy cleavage crack propagation in given microstructure. The deviation in
crack propagation can be observed when the misorientation between neighbouring particles is
higher than 15° [5, 13].

Metal 2005 24. – 26. 5. 2005, Hradec nad Moravicí
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The set of determined misoreintation angles observed within AF plates is plotted in Fig.3
which follows from the misorientation map of AF microstructure. The most boundaries
correspond to high angle boundaries between AF plates, while small misorientation
boundaries (< 15°) can be found within high angle boundaries. The majority of measured high
angle boundaries correspond to the misoreintation angle higher than 45° [13]. Concerning B
microstructure (upper B) the set of misorientation angles is different from AF microstructure.
The occurrence of misorientation peaks is scarce in B microstructure. The distance between
the high-angle boundaries is in AF microstructure 3-5µm approximately. The evaluated
distance is in the B microstructure 15µm approximately. The high angle misorientations can
be detected at grain boundaries of B-packets. In AF microstructure, the density of highly
crystallographic misoriented plates is
enhanced by a profuse direct nucleation
misorientation angle
[degree]
80
60
40
20
0
Point to point
0 5 10
Distance to point A [µm]
Fig.3 Misorientation angle between AF plates.
on inclusion particles. In this case it
does not seem possible to relate the
morphological packet to the
microstructural unit which controls the
cleavage crack propagation in B
microstructure. The microstructure unit
can be directly related to the set of AF
plates enclosed by high-angle
boundaries.
In the case of B, a morphological
packet size (dB) has been previously
defined and related to the unit crack
path (UCP) [13, 14]. This distance is
defined as the region in which the crack
propagates in a straight line and really
corresponds to the free path between
two neighbouring high-angle
boundaries. This demonstrates, the
misorientation higher than about 10-15° between {100} cleavage planes corresponding to two
adjacent B-packets are able to markedly deflect the brittle crack propagation [15].
Fig.4 Histogram of misorientation angle distribution found in AF and B microstructure

Metal 2005 24. – 26. 5. 2005, Hradec nad Moravicí
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adjacent B-packets are able to markedly deflect the brittle crack propagation [15]. Figure 4
shows the distributions of plate misorientation level determined in AF and B microstructure
(evaluated frequency is plotted in dependence on misorientation angle) in histogram form.
A very interesting crystallographic finding resulting from these dependences is the absence of
misorientation across plate boundaries within the range of 20-47° approximately [1, 5]. These
results are consequence of A-decomposition by the displacive mechanism. The formed AF
plates exhibit misorientation angles lying between 47-60° as it corresponds to near N/W
relationships with the parent (austenitic) phase. The same relationships were found in upper B
microstructure. In the case of the AF, the microstructure is found to be very close to the self
accommodating microstructure arrangement of martensite which is favourable for the
realisation of shears stress accommodation process [5, 13].
The ductile - brittle transition temperature Tx is usually expressed as being inversely
proportional to the root square of distance between high-angle grain boundaries. The crack
propagation can be arrested at high-angle boundaries. The parameters of this process depend
on the microstructural characteristics of given steel type. In AF microstructure, the density of
crystallographic misorientation planes is enhanced in comparison with upper B (Fig.4). The
observed AF microstructure is beneficial due to higher probability in realisation of
ascertainable brittle crack deflection than can be detected in upper B microstructure [5, 13,
15].
The following type of equation has been used to describe the ductile-brittle transition
temperature in steel [13]:
Tx = T0 – K.d -1/2 (2)
Where K is a constant, T0 only depends on the tensile properties of the material and d
represents the mean linear intercept between high-angle boundaries. The value of d in steel
having AF microstructure is smaller than the distance d between boundaries of B packets.
This comparison confirms a beneficial influence of interlocked AF microstructure [5].
5. CONCLUSIONS
The study has been carried out to clear up the basic physical metallurgy principles
controlling the AF and B formation. The effect of some inclusions, known as AF potential
nucleants, is analysed. The attention has been devoted to the study of nucleating behaviour of
Ti-oxides (Ti2O3) and the additional influence of MnS and TiN particles. Such particles form
together the heterogeneous type of nucleants (formation of Mn-depleted zone in the
neighbourhood of Ti2O3 oxides and the influence of low interfacial energy between TiN
and/or VN and ferritic matrix – AF).
The AF particles are intragranularly nucleated, while the intergranular nucleation
corresponds to B formation. The AF microstructure is characterised with increased frequency
of a high-angle boundary. On the contrary, B microstructure is characterised with higher
occurrence of low-angle boundaries what leads to lower resistance to cleavage crack
propagation and to larger unit crack path (UCP). The mechanical metallurgy consequences of
these parameters (AF vs. B) are elucidated.
Acknowledgements
The authors acknowledge the Grant agency of Czech Republic for financial support of
project No. 106/03/0264.

Metal 2005 24. – 26. 5. 2005, Hradec nad Moravicí
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REFERENCES
[1] BADESHIA, H.K.D.H. Bainite in steels, 2 nd edit., London, The Inst. of Materials,
Cambridge, 2001, 451p.
[2] MADARIAGA, I. et al. Met. Mater. Trans., 32A, 2001, pp. 2187-2197.
[3] LINAZA, M.A. et al. Scr. Mater., 32, 1995, pp. 395-400.
[4] DIAZ-FUENTES, M. et al. Mater. Sci. Forum, 284-286, 1998, pp. 245-252.
[5] GOURGUES, A.F. et al. Mater. Sci., Technol., 16, 2000, pp. 26-40.
[6] BYUN, J.S. et al. Acta Mater., 51, (2003), pp. 1593-1606
[7] TAKAMURA, S. et al. Roles of oxides in steels performance-metallurgy of oxides in
steels. In Proceedings of the 6 th International Iron and Steel Congress. ISIJ Nagoya,
1990, pp. 551-597.
[8] JUNG, I.H. et al. ISIJ International, 44, (2004), pp. 527-536.
[9] VAN DER Eijk,C. et al. Material Sci. Technol., 16, 2000, pp. 55-64.
[10] ISHIKAWA, F., TOKAHASHI, T. ISIJ International, 35, (1995), pp. 1128-1134.
[11] NORTH, T.H. et al. Welding Journal, Welding Research Supplement, 57, (1979), pp.
343s-354s.
[12] YAMAMOTO,K. et al. ISIJ Inter., 36,1996, pp. 80-86.
[13] DIAZ-FUENTES, M. et al. Met. Mater. Trans., 34A, 2003, pp. 2505-2516.
[14] BROZZO, P. Met. Science, 11, 1977, pp. 123-129.
[15] MAZANCOVÁ, E. et al. Zeszyty Naukowe Politechniky Opolskej, Mechanika, 78,
2004, pp. 113-118.