Dual phase steels

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DUAL PHASE STEELS +8665

M. S. Rashid
by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only.

Metallurgy Department, General Motors Research Laboratories, Warren,
Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org

Michigan 48090

INTRODUCTION
Dual phase steels are characterized by a microstructure conslstmg of
a bout 75-85 vol% ferrite (a-iron) with the remainder being a mixture of
martensite, lower bainite, and retained austenite (Figure 1). The name
"dual phase" was coined in the mid 1970s to describe ferrite-martensite
microstructures (I), but dual phase steels usually contain more than the
two phases implied by their name. They are essentially low carbon steels
that are thermomechanically processed to have better formability than
ferrite-pearlite steels of similar tensile strength.
The stress-strain behavior of dual phase steels (2) is characteristically
different from that of ferrite-pearlite steels such as plain carbon steel or
the microalloyed, high strength low alloy (HSLA) steels (Figure 2). The
ferrite-pearlite steels have yield point elongation, a high ratio of yield
strength to ultimate tensile strength (YSIUTS), and their strength and
ductility (uniform elongation) are inversely related (Figure 3). Dual phase
steels have a continuous stress-strain curve with no yield point elonga
tion. They work-harden very rapidly at low strains, have a low YS, a high
UTS and hence a low YS/UTS ratio. They have better formability than
the ferrite-pearlite steels of equivalent tensile strength and their strength
ductility data fall on a separate curve (Figure 3) than that for ferrite
pearlite steels.
Ferrite-martensite steels were developed in the mid 1960s concurrently
at BISRA (British Iron and Steel Research Association) in the United
Kingdom (3-5) and Inland Steel Corporation in the US (6). Both efforts
were concerned mainly with producing steels for tinplate by drastically
quenching low carbon steels from temperatures close to the critical
temperature. The BISRA objective was to develop a compact annealing
process for making steels of about 500 MPa tensile strength; Inland Steel
effort was directed toward producing steels with about 1000 MPa tensile

245

0084-6600/81/080 1-0245$0 1.00

246 RASHID

Transformation Product

Ferrit 10 m

Figure J Scanning electron micrograph of a dual phase steel. The microstructure consists
of a fine-grained ferrite matrix with a uniform distribution of about 20 volume percent
transformation product , which consists of martensite, retained austenite and bainite (2).

980X
600
GM

----- -----� ,

Plain Carbon

I-+---Totul Elongotion, et ----I
i
o I
10
-11-0.2 Percent Stroin

Figure 2 Schematic stress-strain curves for plain carbon, HSLA, and dual phase steels.
SAE 950X and 980X are Society of Automotive Engineers designations for HSLA steels of
different s trength levels. GM 980X is a General Motors developed dual phase steel. GM
980X is more ductile than SAE 980 X although both steels have similar tensile strength.

DUAL PHASE STEELS 247

700
3'OX
o
...

� 600
of
III
C

�
f
1
R
Ferrite-pearlite

.!! 500
.
;;
c

,! SAE 950
.!!
o
� 400 ,
5
" " Figure 3 Strength-ductility relationship

Plain Carbon � of dual phase steels compared with that
for plain carbon and HSLA steels. The
dual phase steel curve is far above that
Uniform Elongation, '70 for ferrite-pearlite steels (40).

strength. The BISRA process involved intercritical heating followed by
quenching either into water with post tempering or into low temperature
salt or liquid metal baths, the bath serving as the tempering medium.
Although the first Inland Steel experiments were directed toward a fully
quenched martensitic product, subsequent work was also conducted on
intercritically heated and quenched sheet. Neither approach recognized
or researched the potential for improved formability in these steels.
Grange (7, 8) later investigated methods for producing fibered ferrite
martensite microstructures. This process requires cold rolling a pearlitic
structure to produce elongated patches of pearlite, which after a subse
quent intercritical heating and quenching treatment produces fibers of
martensite in a ferrite matrix. Steel having such microstructures has a
better combination of strength and toughness than the same steel before
the thermomechanical treatment. Although indications of better forma
bility were evident in these data, the importance of this aspect was not
elaborated.
Bailey researched ferrite-martensite steels (9, 10) in the mid 1970s and
reported that the strength of low carbon steels can be increased at the
expense of ductility by intercritical heating followed by quenching and
tempering. The ductility of these "dual phase" steels was comparable to
that of ferrite-pearlite steels of similar tensile strength. Like some of the

248 RASHID

earlier work, this study also emphasized strength rather than formability
improvement but it dealt with steels thicker (up to 2 mm) than tinplate.
The development of dual phase steels that were tailored for improved
formability was triggered in the early 1970s by conflicting demands in
the US automotive industry for reduced weight to increase fuel economy,
and increased weight to satisfy newly imposed safety and ecological
standards. Steels were sought with strength-ductility combinations sub
stantially better than existing grades of high strength sheet steels to

enable the fabrication of complexly shaped automotive components using

techniques perfected for plain carbon steel.
Concurrent development of the desired steels occurred in Japan and
the US. Matsuoka & Yamamori ( 1 1) and Hayami & Furukawa ( 1)
reported the development of intercritically annealed, microalloy-free
(vanadium, niobium and titanium) Si-Mn steels, while Rashid (2, 12, 13)
reported on the development of intercritically annealed, microalloyed
dual phase steels. These steels had far better formability than any of the
previously reported high strength sheet steels and represented a
breakthrough in high strength steel development. Bucher & Hamburg
demonstrated that the research data reported by Rashid could be dupli
cated in commercially produced steels ( 14, 15). Later, Coldren & Tither
( 16) showed that dual phase steels can also be produced as-rolled directly
off the hot mill. Subsequently, numerous researchers have reported on
variations of these approaches using modified steel compositions ( l7-20).
The various developments made it clear that a dual phase microstructure
by itself did not automatically guarantee good formability, but that the
objective of good formability combined with high strength can be accom
plished by proper control of steel composition and process variables.
Low to high carbon steels with martensite-austenite microstructure,
with martensite being the major phase, have sometimes also been referred
to as "dual phase" steels (2 1 22) although they do not satisfy the
,

strength-ductility criteria shown in Figure 3. Martensite-austenite steels
have unique combinations of mechanical properties but do not fall into
the category of "dual phase" steels described earlier and should be dealt
with separately.
Publications and research on the highly formable dual phase steels
have increased exponentially ( 11-20) since their inception just a few
years ago, and a unified treatise of published literature is attempted in
this article. This review summarizes the historical development of the
steels and describes the present state of the art of the various approaches
that have been used to produce the steels. Current understanding of the
transformation and deformation mechanisms are discussed and the
strength ductility relationships in these new steels are briefly reviewed.


PROCESSING METHODS
Dual phase steels have been produced by continuous annealing, as-rolled
directly off the hot mill, and by batch annealing. Considerably more
research and production activity is reported in continuous annealing than
in the latter two approaches.

Continuous Annealed

All reported continuous annealing processes have three common salient

features (Figure 4), namely (a) rapid heating to above the critical
temperature A I, (b) a short time holding at temperature, and (c) cooling
below the martensite start ( Ms ) temperature. Some processes also include
a short time tempering below 5 00°C after cooling from above Al to
improve the ductility and toughness of the steel at the expense of tensile
strength. The microstructure of most steels prior to continuous annealing
consists of ferrite, pearlite, and grain boundary iron carbides (1, 2,
11-20). Some nominal compositions are listed in Table 1.
The rate of heating is far less critical ( 12) than heating temperature,
time, or cooling rate. Intercritical heating was preferred by most investi
gators although steels were also produced by heating supercritically, i.e.
above A3 (12, 14). Intercritical heating temperatures offer the inherent
control of volume fraction and composition of the ferrite and austenite as
dictated by the equilibrium phase diagram (Figure 5). The time at
temperature ranged from a few seconds to a few minutes and it is not
clear whether equilibrium conditions were attained. The kinetics of
austenite formation in these steels have not been investigated but the
short heating times suggest that the austenite is probably nonhomoge
neous and composition gradients may exist.

800

� 600
P!.
.2
e 400
!.
E
GI
... 200

Figure 4 Schematic representation of
oL-------�-- various steps on the continuous an
Time- nealing process.

250 RASHID

Table 1 Nominal compositions of some dual phase steels

Maximum
Reference temperature" Cooling Composition, wt.%

Continuous Annealed
Hayami (I) Air cool 0.09 C, 0.92 Si, 0.97 Mn, 0.32 Cr
Matsuoka and
Yamomouri (J I) I Water quench 0.07 C, 0.39 Si, 2.96 Mn,

Rashid (2) I Air cool 0.12 C, 0.51 Si, 1.46 Mn, 0.11 V
Bucher (14) I&S Fast air cool 0.11 C, 0.5 Si, 1.4 Mn, 0.06 V

As-Rolled
Coldren (16) S Air cool 0.06 C, 0.9 Si, 1.2 Mn, 0.6 Cr, 0.4 Mo

Batch-Annealed
Parker (23) Furnace cool (14°Cjhr) 0.08 C, 0.3 Si, 3.0 Mn

a 1= intercriticaJ, S= supercriticaJ.

Cooling rates in the range between air cooling and water quenching
have been used to produce dual phase microstructures. Faster cooling
rates are required for steels with lower hardenability. The nonmicroal
loyed Si-Mn steels are usually produced by water quenching, while slower
cooling rates have been used for the microalloyed compositions. Slower
cooling rates produce better strength-ductility combinations and are
generally preferred. Higher cooling rates induce a larger number of
lattice defects and residual stresses into the matrix and may reduce

Typical Carbon Content
I' of Duol Phose Steels
I
Austenite, Y

v
I
o
I
i
:>
'0700
OJ
Q.
E
'"
I- 600

500

Percent Carbon

Figure 5 Schematic representation of a portion of the iron-carbon phase diagram.


ductility slightly. However, ductility can be improved by tempering the
steel.
Some researchers have attempted ( 19, 20) to use existing carbon
equivalent formulas to identify compositions and cooling rates that will
produce the desired mechanical properties. Such formulae are, at best,
only rough indicators of the expected mechanical properties, but they
provide a good screening mechanism to test steel compositions and
processing parameters. Many unknown parameters exist, however, and

reliable predictive techniques for steel composition and cooling rates that

would produce the optimum microstructure have yet to be perfected.

As-Rolled
In the as-rolled process (16), the steel composition is chosen (Table I)
such that 80- 90% of the steel transforms to ferrite after the final roll pass
in normal conventional hot rolling and before entering the coiler. The
remaining 10-20% does not transform until much later, during slow
cooling in the coil. This is possible with steel compositions that exhibit
certain special characteristics in their continuous cooling transformation
(CCT) diagram (Figure 6): namely (a) an elongated ferrite C-curve, i.e.
the ability to form very large amounts of ferrite over a wide range of
cooling rates on the run-out table; (b) a suppressed (delayed) pearlite
nose and high pearlite finish temperature to ensure avoidance of pearlite
formation during cooling to the coiling temperature; and (c) a gap
between pearlite and bainitic regions to provide a temperature range
within which no transformation occurs, permitting sufficient time for the
steel to be coiled. For the composition listed in Table I, the range of

1000

To I
AC3 ACI
u 800
°
A
€ 600
Z
2
� 400 Ms
E -
III
M
.... 200 coiling
window
0
10 0 101 102 103 104 105
Time to Cool from 960°C

Figure 6 Continuous cooling transformation diagram of an as-rolled dual phase steel
(A = austenite, PF=poJygonal ferrite, P-pearlite, BF=bainitic ferrite, M=martensite of
average C content, M'=martensite from carbon-enriched austenite, Ta =austenitization
temperature) (16).

252 RASHID

cooling rates through the ferrite region appears to be rather wide and
produces the desired microstructure.

Batch Annealed
Dual phase steels have also been produced by batch annealing techniques
(23) modified for heating in the intercritical temperature range. The very
slow cooling rates inherent in this approach (several days to cool to room
temperature) necessitate the use of steels with very high alloy content

(Mn and Mo) and high hardenability. This approach is presently the least

researched of the three.

TRANSFORMATION MECHANISM
Continuous Annealed
The mechanisms by which dual phase microstructures are formed are
reasonably well understood and can be explained with reference to steel
microstructure prior to heating above the critical temperature. As men
tioned earlier, the starting steel consists of a ferrite matrix with grain
boundary iron carbides and small islands of pearlite (Figure 7). Microal
loyed steels also contain microalloy carbonitrides uniformly distributed
throughout the ferrite matrix (Figure 8).
Ferrite Pearlite

Iron Carbide Pearlite 5 m

Figure 7 Scanning electron micrograph of a high strength, low alloy steel. The microstruc
ture consists of a fine-grained ferrite matrix, grain boundary iron carbides, and islands of
pearlite (2).


Figure 8 Bright-field micrograph revealing carbonitride precipitates distributed in the
. ferrite. Arrows I, 2 and 3 point to precipitates of different sizes. Arrows 2 point to
medium-sized carbonitrides precipitated at low angle boundaries, for example, separating
regions L and N. Arrows P reveal dislocations pinned by precipitates (26).

Upon heating the steel above the critical temperature (Figure 5),
islands of carbon-rich, nonequilibrium austenite form at the carbide
locations. The heating temperature determines the carbon content and
volume fraction of austenite that can exist in equilibrium with ferrite.
Given sufficient time at temperature the austenite nuclei grow until this
criterion is fulfilled. In steels heated to just above the critical tempera
ture, the proportion of austenite formed is at a minimum and it has a
high carbon content because carbon is more soluble in austenite than in
ferrite (Figure 5). At higher annealing temperatures, the volume fraction
of austenite is larger and it has a lower carbon content and hence lower
hardenability. Steels heated supercritically transform entirely to austenite
of the carbon content of the steel.
The composition of the austenite is also influenced by other alloying
elements in the steel. The presence of Si in the ferrite promotes carbon
migration from ferrite to the austenite (24) thereby adding to the carbon
content of the austenite, while Mn partitions preferentially to the austenite
and increases its hardenability. The austenite composition is usually
nonhomogeneous and concentration gradients exist because of the short
heating times involved. The substructure and composition of the untrans
formed ferrite that coexists with the austenite at elevated temperature (2,
26) are also modified. Carbon partitions out of the ferrite, when present
microalloy precipitates coarsen, and dislocations rearrange themselves
into low energy configurations.
The transformation product that forms upon cooling the steel back to
room temperature depends on the austenite composition and cooling

254 RASHID

rate. Si suppresses the pearlite transformation (25) while C and Mn
stabilize the austenite and lower the Ms temperature. At rapid cooling
rates all the austenite transforms to martensite. At slower rates, depend
ing on austenite hardenability, various proportions transform to marten
site, bainite, and ferrite, with some austenite remaining untransformed
(retained austenite). High carbon austenite transforms to twinned
martensite while low carbon austenite transforms to lath martensite; the
former martensite is stronger than the latter. The volume change and

shear accompanying the austenite� martensite transformation generates

numerous new mobile dislocations ( 1, 2) in the surrounding ferrite
matrix. The ferrite matrix consists of transformed and untransformed
ferrite. As mentioned previously, the untransformed ferrite has low
strength and is relatively free of interstitials such as carbon and nitrogen,
which have either diffused out of the ferrite or remain in the ferrite and
contribute to precipitate coarsening. The transformed ferrite formed
when the austenite is cooled below the critical temperature is very similar
to the ferrite of the starting steel (26) in dislocation substructure, carbide
distribution, and strength. Untransformed ferrite is not expected in
supercritically heated steels but very short heating times sometimes
preclude complete austenitization and some untransformed ferrite might
be observed.
As-Rolled
Most of the hot rolling is done in the austenite range of the Fe-C phase
diagram and the finish rolling temperature is usually in the intercritical
region. As mentioned previously, 80-90% of the austenite transforms to
ferrite on the run-out table and the remainder transforms in the coil to
transformation products similar to those observed after continuous an
nealing. Some autotempering of the martensite and decomposition of
retained austenite may be expected during cooling in the coil.

Batch Annealed
The transformation mechanisms here are similar to those observed dur
ing continuous annealing but the grain size and substructure are char
acteristic of the much slower cooling rates involved.
The various phase transformations discussed produce a microstructure
consisting of untransformed ferrite, transformed ferrite, martensite,
bainite, retained austenite, and carbide precipitates. The relative volume
fraction, morphology, distribution, composition, and mechanical prop
erty of each constituent is governed by steel composition and processing
parameters, and determines the deformation behavior of the steel. Defor
mation of the steel itself can also induce some phase transformations.


DEFORMATION BEHAVIOR
The deformation behavior of dual phase steels is quite complex. A
thorough understanding of the interactions between the various micro
constituents discussed and their influence on mechanical properties is
lacking, but steel deformation behavior can be explained in generalized
terms.
The deformation behavior of most metals, especially plain carbon steel,

may be described (27) in terms of a simple empirical relationship between

the stress-strain data obtained in a tension test, namely

a=Ke; 1.

or
log a=log K+n log€p' 2.

where a is the true stress, K is the strength constant or a at lOp = 1.0, lOp is
the true plastic strain, and n is the strain hardening exponent, a measure
of the ability of the metal to distribute strain. The true strain at
maximum load, fu' also called the true uniform strain, will numerically
equal n when Equation I is satisfied (27). When Equation I is satisfied, a
plot of log a vs log lOp will be a straight line with slope n. Such behavior is
attributed to truly uniform plastic deformation behavior.
Equation I is not satisfied for dual phase steels (28). A plot of log (J VS
log fp (Figure 9) deviates substantially from linearity, which suggests that

5.2,------ ---,
• Ferrite-Pearlite

Cooling Rate.o IC
5.1
• 5
09

5.0 -12
014

Figure 9 Variation in plots of log
true stress vs. log true plastic strain for
steels with the same composition but
4.7
with different microstructures. The
ferrite-pearlite steel was heated at

4.6 '--_--'-__=-'-="__.J..-_-:-'o--__ ..
....
788°C for 3 min and cooled to room
-3.0 -1.0 temperature at the rates shown to
Log True Plastic Strain produce dual phase steels.

256 RASHID

several deformation processes are operative in this steel. At least two
approaches have been used for detecting changes in deformation behav
ior. In one approach (28) an incremental value of n, ni' is calculated for
'
each segment of the stress-strain curve, with n i being defined as
log OJ -log 0j_ .
n;(j)=J for /=lto/. 3.
og Epu> - 1og Epu_1) •

Here I is the number of small segments of the stress-strain curve. In the
other approach (29) the stress-strain curve is represented by the equation

4.
where a i s the true stress, 00 i s the true yield stress, Ep is the true plastic
strain, and Band m are constants (m is not strain rate sensitivity).
Equation 4 can be differentiated (30, 31), logarithms taken on both sides
and written as
do
In- =In Bm+ ( m-l ) ln Ep• 5.
d lOp

Plots of n i vs lOp (Figure 10) or In dojdEp vs In lOp (Figure 11) are straight
lines for ferrite-pearlite steels but delineate several different stages of

0.30

!'
en /
c:
'c
� 0.20
o
:z:
c
SAE 9 50X /
~
"e
.;;
g
'E
;
f
0.10 7SA' .. OX

///
u

.5
'p"'ni

0.2
True Plostic 5troin, Ep

Figure 10 Variation of incremental strain-hardening rate, ni with increasing true plastic
strain in various steels: ni is relatively constant in the ferrite-pearlite steels but not in the
dual phase steel (28).


Annealed 4min. @ 810°C

�,
�
,

" '" 53"e/sec.

'�'"
...-
1//

20"C/sec
.�'


�Uniform Wt. % &.
Strain C 0.15
Mn 1.37
��
Si 0.27
Figure II Plot of Inda/dEp vs In Ep for a
N 0.0048
dual phase steel of indicated composition
and annealing conditions for different
True Plastic Strain cooling rates (29).

strain hardening in the deformation behavior of dual phase steels,
confirming earlier indications that multistage, nonhomogeneous deforma
tion occurs in these steels in contrast with homogeneous deformation in
ferrite-pearlite steels.
Some of these deformation processes can be identified by correlating
certain features of the tensile stress-strain curve with the deformation
behavior of particular microconstituents in the steel. As mentioned
earlier the yield strength in dual phase steels is only 0.5-0.6 of the tensile
strength and the yield point elongation is absent. The work-hardening
rate is very high at low strains and decreases with increasing strain. The
total elongation is higher than that in ferrite-pearlite steels of similar
tensile strength. Each of these characteristics can be traced to the
deformation behavior of one or more microconstituents described previ
ously.
The shear and volume change accompanying the austenite _ martensite
transformation on cooling from above the critical temperature generate
numerous free mobile dislocations in the surrounding ferrite matrix.
Upon application of a load, the free dislocations move at stresses much
lower than required to move restrained dislocations, such as those
commonly observed in ferrite-pearlite steels. Hence, dual phase steels
yield or commence plastic flow at much lower stresses compared to the
ferrite-pearlite steels of equivalent tensile strength. Furthermore, since
interstitial solutes such as carbon and nitrogen have either diffused out of
the untransformed ferrite or are in combined form, solute-dislocation
interactions are severely reduced; an initial threshold stress is not needed
to break the dislocation away from the solutes and, hence, no yield point
elongation is observed. Upon continued application of tensile load,

258 RASHID

plastic flow continues in the microconstituent with the lowest yield
strength, this being the untransformed ferrite. After the untransformed
ferrite has work-hardened to the yield strength of the transformed ferrite,
both phases presumably deform and work-harden simultaneously (2, 26,
28).
Metals work-harden because mobile dislocations interact with other
dislocations, solutes, precipitates and other microconstituents. This is
also true in dual phase steels. But the magnitude of the work-hardening

rate observed at low strains is too large to be explained by dislocation

interactions alone, and could be caused by the added contribution due to
deformation of retained austenite. The strain-induced transformation of
austenite to martensite increases the ductility and work hardening rate in
-

several material systems (32, 33). This phenomenon, called transforma
tion-induced plasticity or TRIP, also occurs in dual phase steels (34, 35).
Dual phase steels contain as much as 10 vol% retained austenite, which
transforms to martensite in direct rel�tion to increasing strain in a
tension test (34). Most of the transformation occurs at low strains, nears
completion at about 10% strain and could enhance the work-hardening
,

rate because of dislocations interacting with other dislocations and with
the strain-induced martensite.
Deformation obviously continues in the untransformed and trans
formed ferrites beyond 10% strain, but because of differences in the
strengths of the two phases, the work-hardening rate is not constant but
continues to decrease (Figure 10) until a constant incremental work
hardening rate is approached at strains just prior to uniform strain. The
deformation characteristics of the ferrite phases are, of course, influenced
by composition and process variables. These effects are treated in detail
in the next section. Martensite is usually regarded as a nondeformable
hard phase that contributes primarily to steel strength. However,
martensite deformation has been observed (Figure 12) in dual phase
steels at very high strains (28). Concentration gradients in the austenite
during continuous annealing evidently produce martensite of varied
composition and strength, some of which deforms and contributes differ
ently to work-hardening behavior than nondeformable martensite. Fur
thermore, the strain-induced martensite would have yet different defor
mation characteristics from the original martensite and also contribute to
the observed deformation behavior.
The higher total elongation of dual phase steels is a consequence of the
multistage deformation behavior just described. When dual phase steel is
strained, slip leading to deformation occurs first in the constituent with
the lowest yield strength. When this constituent work-hardens to the
yield strength of the second constituent, plastic flow begins to occur in it.
This continues until all constituents are involved in the deformation


Fradured Martensite

F rrite

Void

Martensite

Figure J2 Scanning electron micrograph of a dual phase steel at fracture initiation. Both
ferrite and martensite grains are elongated. Some martensite islands are fractured and voids
are formed at the ferrite-martensite interfaces (28).

process and are work-hardened to their maximum extent. Since deforma
tion is distributed among several constituents, strain is distributed more
uniformly and localized necking is delayed, resulting in far better forma
bility than ferrite-pearlite steels of similar tensile strength.
Fracture initiation has also been contrasted in dual phase and ferrite
pearlite steels. In the latter steels, deformation is restricted to the ferrite
phase with no obvious deformation of the iron carbides. When the ferrite
is work-hardened to its limit, voids form preferentially at the ferrite-iron
carbide interfaces (28) and failure is initiated. In dual phase steels,
deformation occurs in the microconstituents in a sequence related to their
yield strength. Martensite, being the strongest constituent, does not
deform until all other consitituents have deformed and are highly strained.
Voids leading to failure form at the ferrite martensite interface but are
not nucleated until more extensive deformation has occurred and the
martensite is also highly strained (Figure 12).

COMPOSITION-STRUCTURE-PROPERTY
RELATIONSHIPS

Dual phase steels have been modeled as two-phase composites: the ferrite
is treated as a homogeneous ductile matrix phase and the "martensite"
(or transformation product consisting of martensite, retained austenite,

260 RASHID

and bainite) is treated as a high strength reinforcing component with
homogeneous mechanical properties. In spite of these simplifying as
sumptions, the strength of the composite (dual phase steel) has been
predicted with reasonable success (36- 38) using the simple rule of
mixtures, namely 0c =(1- Vr)om + V.o" where 0c' om' and Or are the
strengths of the composite, the matrix and reinforcement, and v,. is the
volume fraction of the reinforcement.
The good agreement with experimental data suggests either an insensi

tivity in dual phase steels to nonhomogeneous mechanical properties of

the matrix and reinforcement, a cancellation of errors or just a fortunate
coincidence of the assumptions made in the model. Furthermore, these
models suggest the rather obvious: optimum properties can be developed
in the dual phase steel when the matrix and the reinforcement individu
ally have maximum strength as well as ductility (38).
Since "martensite" is the principal load-bearing constituent, various
attempts have been made to correlate volume percent of "martensite"
with steel strength. The two are linearly related (Figure 13), independent
of "martensite" carbon content (39, 40). Further work showed (41),
however, that carbon content is important and that separate linear
relationships exist (Figure 14) between yield and tensile strength and
percent "martensite" of constant carbon content. A smaller volume of
high carbon martensite produces the same strength as a larger volume of
martensite of lower carbon content, the high carbon martensite, of
course, being stronger than the low carbon martensite. The reason for
this apparent contradiction is not clear, but is probably due to scatter in
experimental data (Figure 13) and because the "martensite" reported in
the earlier work (39, 40) actually contained varying amounts of retained
austenite depending on intercritical annealing temperature. The strength
could also be affected by variations in the p r ecipitation in the ferrite due
to supersaturation. These variations result from the different annealing
temperatures used to produce different volumes of martensite.
Besides strengthening with carbon or alloy elements, martensite strength
can also be increased by decreasing its particle size. This is ensured, in
part, by continuously annealing steels with a fine grain size. With a small
grain size, the grain boundary iron carbides are proportionately small
and this produces correspondingly small martensite islands after trans
formation. Several types of "martensite" distributions have been ob
served but the one most conducive to homogeneous deformation is a
uniform distribution of very small, disconnected "martensite" islands
located at ferrite grain boundary intersections. Sometimes, larger islands
are located further apart and some steels are partially banded, with the
"martensite" content being higher in the bands. Continuous martensite

Dual phase steels

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Dual phase steels