Design of ULC Sheet Steels

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log K,P,fic = log {[Ti][C]} = 9.0- 10330/T’,
where T is the temperature in Kelvin. Using this volubility
product equation, the precipitation temperature in a Ti-
New Approaches to the Design
of ULC Sheet Steels
M. Hua, C. I. Garcia, and A. J. DeArdo
Basic Metals Processing Research Institute
Department of Materials Science and Engineering
University of Pittsburgh,
Pittsburg~ PA 15261, USA
(412) 624-8593
ABSTRACT
The use of a new stabilization map, based on the
observed precipitation behavior, allows for the rational
design of ULC steels. The complete or partial stabilization
of carbo~ and, flu-therrnore, the grain boundary
composition of highly segregating elements, such as B, C
and Nb, can be achieved through adjustment of the bulk
composition. This diagram can be used to either analyze
the behavior of existing steels or to design new high
performance steels.
1. INTRODUCTION
In recent decades, interstitial-free or ultra-low
carbon (ULC) sheet steels have been widely used in the
automobile industry, because of their good formability,
especially deep drawing. The ULC steels are very flexible
since they can be produced using either batch annealing or
continuous annealing lines without strain aging or yield
point phenomena. These steels are produced using vacuum
degassing in steelmaking to reduce the contents of the
interstitial elements C and N to about 30 ppm, which
would then be stabilized by the addition of Ti and/or Nb. 14
This paper reviews the dependence of the carbon
stabilization mechanism upon the composition of ULC
steels and to illustrate the application of a newly-
developed stabilization diagram for ULC steels.
2. STABILIZATION MAP
The traditional view of stabilization results in the
stabilization map, Figure 1, which may be applicable to
microalloyed steels and ULC steels with low S and/or high
Mn. The volubility product of TiC may be described as’
40TH
,...
only steel (Table I, steel Al) is estimated about 806 K
(533”C), which is very low compared to the temperature
for the formation of carbosulfide H phase. For steels
whose composition are above the stoichiometric line
(Ti*/C=4, Ti* = Tititil - 3 .42N+Nb/l .94), there is sufficient
Ti to stabilize C. However, whether complete stabilization
can actually be obtained is also dependent on having
adequate kinetics. Hence, similar steels might exhibit
different levels of stabilization when processed on different
strip mills.
Steels designed to fall into region 2 of Figure 1
would be bake hardenable, since the fi-ee carbon would be
available for strengthening through strain aging when the
steel is heated during the painting cycle.899 Furthermore,
steels designed to fall into region 2 of Figure 1 can be
fhrther stabilized through the addition of Nb. In this case,
the Nb would act as a particle former in a manner similar
to its behavior in conventional microalloyed steels.
The present work suggests a new precipitation
map to aid in the understanding of precipitation and
stabilization in S-containing (0.004-0.1 Clwt pet.), low P
(<0.02), low Mn (<0.2) ULC steels, Figure 2. The
experimental basis for Figure 2 is that the formation of H
phase is not by independent nucleation and growth, but
rather by the in-situ transformation of TiS. 10-14
The major difference between Figures 1 and 2 is
the zero point. In Figure 1, point Ti*=C=O represents,
theoretically, the completion of TiN precipitation and the
starting of TiC. In Figure 2, the “origin” for the formation
of TiC is pushed back to a new point (Ti*=TiEF3 S,
C=CH3 S/8). The slope of the stoichic}metric line is the
same as in Figure 1, but this stoichiometric line starts at
Ti*=l.5 S (C=O) since TiS forms before epitaxial carbides
(TiC). and free standing carbides (TiC)~. When S=0,
Figure 2 would degenerate into Figure 1. In the region I
of Figure 2, which is above the stoichiometric line and
C=CH, all the carbon can be stabilized by forming H-
TiqC2Sz,
Since H phase is much more stable than TiC (lower
free energy of formation AG\ plus large particle size), the
soluble C, which is in equilibrium with H, should be much
lower than that with TiC. Therefore, the new mechanism
indicates a better stabilization condition, or a real
“interstitial-free” matrix. This is supported by the
observation that high S levels can eliminate the bake
hardenability or aging index of ULC steels.8’9’15 A recent
experiment using internal friction also showed that a ULC
steel in region V has 0.5 ppm C and 0.2 ppm N remaining
in solution (not stabilized) in the hot rolled band material,
while the residual C in steels of regions I and 11 is below
the detection limit (about 0.1 ppm). 16
MWSP CONF. PROC., 1SS 1998 – 111
Ti
o
Fig
sto
the
stab
f. - volume fraction of preci
Al A2
i
r
1
I I
C (ppm) 30 30
1 I
N (ppm) I 30 I 30
S (ppm) I 80 I 80
Mn (ppm) I 15000 I 10000
Ti (ppm) 500 750
Nb (ppm) I I
fv (TiN) 1.94X104 1.94X104
fv (H) 0.715X 104
f. (MC) 2.41x 10A 2.12X104
Fpi.*, MPa 0.5913 0.5430
I 1
ml (PPm) 1 -
* using flexible boundary model, Fpt
r~c = 5 nm, rH = 200 nm,
112 – 40TH MWSP CONF. PROC., 1SS
pitates, ~] - Nb in solution
A3 B c
I I I
30
~
30 30
1 I
I 30 I 30 I 30
I 80 I 80 I 80
I 15000 I 1000 I 400
103 500 223
I 250 I 300 I 432
I.94X104 1.94X104 1.94X104
5.94X104 6.22x104
2.65x104
0.6308 0.004 0.004
I I I
200 200
= 3 G f,m / (rcr), (ref. 20), and assuming
TiN = 500 nm, and cr=O.8 J/mz.
998
1~.o1 Stabilizedpi] + TiC 41~c$$~*66,@o o2 Unstabilized*LO TiC + [C]
c, Wt’%o
ure 1. Traditional precipitation map. Above the
ichiomettic line (Ti*/C=4, Ti*=Titil-3 .42N+0.5 lNb)
re is sufficient Ti and the steel is generally considered as
ilized by forming TiC.
?
I II III
Ti*
Ti~
Ti*
ok====
CH c, w pet
Figure 2. Proposed precipitation map for S-containing, Ti
stabilized ULC steels. Ti*=Tihtil - 3.42N (@us Nb/1 .94
only when TiS formed). The region II has epitaxial
carbides (TiC)., whereas regions HI and IV have both
(TiC)e and free-standing carbides (TiC)~. Ti*=l .5S is the
amount of Ti needed to form TiS; Ti~3 S and Cm3S/8
stand for the maximum amount of Ti and C in H phase
(TiiC2S2), respectively.
Table I. Examples of Designed ULC Steel Compositions and Mass Balances
Usually, the minimum bulk C content obtainable in
a commercial ULC is fixed by steelmaking and the amount
of S can be varied (e.g., to balance the C in forming the H
phase). However, there should be an optimum
(iv) ~] on cxferrite grain boundaries for amelioration
of out bursting phenomena,
T
combination of S and Ti, even in the “stabilization zone”.
If Ti is too high, some surface defects may result. I*
It should be understood that, in order to make use
of the stabilization mechanism proposed in Figure 2, the
effects of other elements should also be taken into
account. For example, Mn should be controlled to avoid
the formation of MnS. Also, the addition of other
elements must be viewed fi-om the standpoint of their
effects on the stabilization diagram. For the Ti+Nb steel,
Nb may be counted as ATi*=Nb/1 .94, but a minimum
amount of Ti (e.g., Ti*=TiA) is needed such that TiS
formation is insured, since Figure 2 is not valid for Nb-
only ULC steel, as fhrther discussed below.
Although Nb-only, Ti-only and Ti+Nb ULC steels
are ~allcalled “interstitial free” steels, the mechanisms by
which the stabilization of C occurs are often quite
different. Stabilization in the Nb-only steel behaves like
what is expected in a HSLA steel, that is, largely
independent of the Mn and S contents of the steel, while in
the Ti-containhg steel the stabilization of C depends upon
the amount of Mn and S in the steel.
Previous work revealed Nb-containing precipitates
(in Ti+Nb ULC steel) only in the following conditions: 1)
as-cast, which were produced by continuous casting,2)
reheated below 1000”C and 3) TMP below 1000”C.
Reheating or TMP above 1000”C resulted in the absence
of Nb in the precipitates. 13
It follows that Nb does not have strong influence
on the precipitation in filly Ti-stabilized steel. Therefore,
the ‘effects of Nb on the behavior of the ULC steels, such
as yield strength and ductility, lZ could then be explained
by the amount of Nb in solution, rather than arguments
based on precipitation. Furthermore, the major portion of
the total Nb, found to be present in solution on the ferrite
grain and subgrain boundaries, appears to be responsible
for the stronger (557)[~ T9 5] textures and deep
drawability found in the Nb-bearing stabilized steels.12 In
addition, the Nb in solution has been confirmed to be
responsible for the enhanced resistance to powdering,
outbursting and cold work embrittlement ofien observed in
single-stabilized Ti ULC steels. I*
Therefore, the possible roles of Nb in ULC steels
include:
(i) intercalated TiS to (TiO.sNbO.s)dCzS*,
(ii) ~] on cx ferrite grain boundaries for texture
control,
(iii) ~] on cx ferrite grain boundaries for CWE (cold
work embrittlement) control,
40
(v) ~] on a ferrite grain boundaries for improved
powdering resistance after galvannealing.
From above discussion, the alloy design of fiture
ULC steels may follow a different path, as will be
discussed in sections 3B and 3C. Through better
understanding of the physical metallurgy, fiu-ther
improvement of these steels should be possible.
3. APPLICATION OF THE POLYPLOT
As a result of new understanding of C stabilization
in ULC steels, a computer simulation program (polyplot)
has been developed to predict and analyze, for a given
ULC steel composition, the stabilization condition in the
hot band material and the microstructural behavior (such
as pinning force for recrystallization and grain boundary
composition) during hot rolling.
3.1 Traditional Approach to Alloy Design for
Stabilization
As previously mentioned, early alloy design was
based on the stoichiometry of the particles TiN and TiC.
Later, the carbosulfide TidC& was incorporated into the
stabilization equation:
Ti(total) = Ti(ex) + 3.43N + 4C + 1.5S.
There had been several approaches based on these
concepts as described below.
Case I. Ti addition is intended for the formation of
TiN and TiC through the reactions
Ti + N = TiN, and
Ti + C = TiC.
This approach assumes that the nucleation and
growth of precipitates take place independently and the
alloy is considered as a dilute microalloyed steel. An
example (Al) is shown in Table I. It should be pointed
out that, in this case, the small precipitates (<5 nm, for
example) may exert a pinning force sufficient to retard the
recrystallization of the ferrite after cold rolling.
Case II. Ti addition is intended for the formation
of TiN, H and TiC through the assumed reactions
Ti + N = TiN,
4Ti + 2S + 2C = TidC*Sz (H), and
Ti + C = TiC.
This approach assumes that the carbosulfide
TidCN* (H-phase) was formed directly from the austenite,
H MWSP CONF. PROC., 1SS 1998 – 113
4. suMMARYbut the ULC steel is still considered as a dilute
microalloyed steel. An example (A2) is shown in Table I.
Case III. Ti addition is intended for the formation
of TiN and the Nb addition for NbC assuming the
98
It has been shown that there are two forms of
following reactions
Ti + N = TiN, and
Nb+c=Nbc.
This approach assumes that the nucleation and
growth of precipitates takes place independently and the
alloy was again considered to behave like a microalloyed
steel. An example (A3) is shown in Table I.
3.2 New Approach to Alloy Design for Stabilization
and Control of Grain Boundary Compositions
Ti+Nb dual stabilized ULC steels, with the Ti
addition intended for the formation of TiN and TiS and also
to intercalated along with C the TiS to H. The Nb addition is
intended for remaining in solution on the a ferrite grain
bourldaries. The reactions involved are
Ti + N = TiN,
Ti + S = TiS,
TiS + Ti + C = %Ti4CzSz, and
ml= ml.
It should be noted that in these ULC steels the
nucleation and growth of H phase is not an independent
process but an in-situ transformation and the alloy does not
behave like a microalloyed steel. An example (B) is shown
in Table 1. It should be pointed out that, in this case, the
large precipitates (> 200 nm, for example) should not be
enough to insert a pinning force to interact with the
recrystallization process of the material. About 70°/0 of the
total Nb is available for ferrite grain boundary segregation,
which may lead to improvements in CWE and other
properties.
3.3 Future Approach to Alloy Design for Stabilization
and of Grain Boundary Compositions
Ti addition is intended for the formation of TiN and
TiS while Nb is intended for both the transformation from
TiS to H and for being in solution on the ct ferrite grain
boundaries. The reactions are
Ti + N = TiN,
Ti + S = TiS,
TiS + Nb + C = %(Ti,Nb)dCzSz, and
ml= ml.
An example (C)is shown in Table I.
114 – 40TH MWSP CONF. PROC., 1SS 19
carbon stabilization in ULC steels. These are through the
formation of MC carbides or by the formation of
carbosulfides M.4C& The two approaches obviously
differ in the role played by sulfhr. The stabilization
reactions that actually occurs is dictated by the bulk
composition of the steel. This behavior can be predicted
through the use of a new stabilization diagram.
The use of the stabilization diagram allows for the
rational design of ULC steels. For example, the complete
or partial stabilization of carbon can be achieved through
adjustment of the bulk composition. Furthermore, this
diagram can be used to control the grain boundary
composition, which has been shown to be important in
controlling the texture, galvannealing, powdering, and cold
work embrittlement behavior in these steels.
Titanium is usually considered as a precipitate
former or C-stabilizer (when forming TiC and/or H).
Niobium can act as stabilizer (when substituting Ti in
forming NbC, MC and H) and/or grain boundary
strengthener.
ACKNOWLEDGEMENT
We thank the University of Pittsburgh Interstitial
Free Steel Consortium and the Basic Metals Research
Institute, Department of Materials Science and Engineering,
University of Pittsburgh, for sponsoring this work.
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0TH
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4
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MWSP CONF. PROC., 1SS 1998 – 115
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