FIELD OF THE INVENTION
The present invention relates to a cold-rolled continuously annealed
weldable dual phase steel with a tensile strength~ 650-800MPa for
automotive applications. The invention further relates to a process of
producing the steel grade.
BACKGROUDD OF THE INVENTION
In the recent years, high strength multiphase advanced high strength
(AHSS) steels are used in manufacturing of light-weightffuel-efficient
vehicles that ensure passenger safety. AHSS is a term used to describe
several families of steels (dual phase, transformationiinduced plasticity,
complex phase and martensitic steels) that contain ferrite, martensite,
bainite and/or retained austenite in sufficient quantities to produce desired
mechanical properties. All AHSS are produced by controlling the cooling rate
of the austenite -ferrtte phase, either on the run-out table of a hot rolling
mill or in the cooling section of a continuous annealing line to ensure high
strength levels through bainitic-martenstiic tran~formation. However, such
transformations can not be achieved unless the cooling rate is over a critical
threshold that depends on steel composition. The critical cooling rate
increases with decreasing alloying additions and as the cooling rates
attainable by industries are limited, without the addition of alloying
elements, achievement of high strength is not possible. However, increasing
alloying elements, such as C, Si and Mn enhances quenching hardenability
of steels that affects weldability adversely and also increases production
cost. Thus, AHSS grades though appear to meet tensile properties, the
weldability associated with these steels has been a matter of concern.
Further, the AHSS with tensile strength>590 MPa are quite difficult to weld.

Typically, weldability of steels depends upon carbon equivalent (CE) that
provides a measure of hardenability of the steel and is a function of alloying
elements present.
The following expression is widely used for expressing CE of AHSS:
CE = C + Si/30 +Mn/20+2P+4S (mass%)
Resistance spot welding is a favoured method for joining advanced high
strength steels and the nugget diameter criterion is critical in AHSS. The
welding lobes of AHSS are narrower than mild and high strength steels.
Further, AHSS tend to exhibit higher expulsion, higher electrode wear and
most importantly, they fail during pull button test resulting in interfacial
failure. In addition, detrimental martensite forms in resistance spot welds
even at low carbon levels as the cooling rate associated with spot welding is
very high and of the order of 103-105 eC/s. The hard martensite thus
formed due to high cooling rate, provides a path for a crack to propagate.
In addition, rapid cooling can cause porosities toward the edges of the weld
where stress concentration is highest in a peel test and also solidification
cracking enhancing the chances of interfacial failure of a spot weld.
Shrinkage, solidification cracks and liquid metal embrittlement in DP 600
are observed in some welds and weld current and hold time are found
responsible for formation of such imperfections. Interfacial failures in DP600
steel when the weld size exceeded a minimum threshold size was also
reported, in addition to interfacial fracture at weld button for thicker gages
of steel. Such failures occur due to solidification cracks and shrinkage voids.

Further, softening of heat affected zones (HAl) may occur in case of AHSS
with significant amount of martensite in the base metal that can be
attributed to martensite tempering at the points of weld where the
temperature approaches a subcritical HAZ. As DP steel essentially derive its
strength from a composite microstructure of ferrite and martensite, and the
martensite being a thermally unstable phase, it tends to decompose in the
heat affected zone (HAZ) resulting in softening of the zone and thus causing
hardness reduction. This phenomenon is termed as HAl softening that is
affected for example, by steel composition, base metal martensite,
prestrain, and heat input, which interalia imposes detrimental impact on the
weldment. It is to be mentioned that HAl softening can be detrimental for
DP980 weldability due to martensite tempering, however, presence of a low
amount of martensite in the base metal of lower strength DP450 does not
cause HAl softening.
Europian patent EP7923791 teaches a welded (10-50 kJ/cm) 15 mm hot
rolled dual phase steel comprising a ferrite phase and about 40-80 vol % of
a martensite/baintte phase of which bainite is no more than about 50 vol %,
the ferrite phase containing carbide or carbonitride precipitates of
vanadium, mobium, molybdenum and mixtures thereof. The
martensite/bainite phase containing retained films of austenite of less than
500 Angstroms thickness, and the sum of the vanadium and niobium
concentrations is not more than 0.27 wt %. The strength of HAl is reported
to be 95% of the base metal.

Canadian Patent CA 2182813C teaches a hot rolled dual phase steel grade
of lOOksi for line pipes, comprising in hot rolled condition ferrite and
martensite/bainite phases, the ferrite phase having vanadium and niobium
carbide or carbonitride precipitates, which is produced by a first rolling
above the austenite recrystallization temperature, a second rolling below
the austenite recrystallization temperature and cooling between Ar3
transformation point, and a temperature of 500 °C; and water cooling
below about 400 °C.
US patent US2009071575 (IN201100478-P4) discloses a hot rolled dual
phase steel comprising a martensite phase less than 35% by volume and a
ferrite phase more than 50% by volume with composition (wt%): C: 0.01%
- 0.2%, Mn: 0.3% - 3%, Si: 0.2% - 2%, Cr and Ni: 0.2-2% , AI: 0.01% -
0.10% and N :> 0.02%, Mo<0.2% and Ca: 0.0005% to 0.01% . Hot rolling
of the slab was conducted in a range between about Ar3-60 °c and -980 °C
and cooling was conducted at a mean rate of at least about 5 °c /s to a
temperature not higher than about 750°C . This resulted in tensile strength
of more than about 500 MPa and a hole expansion ratio more than about
50%. While, a mean impact energy recoreded was> 10,000 g-m on a V-
notch Charpy specimen of about 5 millimeters thickness.
EP 07428418 describes a high strength hot rolled Iinepipe steel having
composition 0.05-0.12C 0.01-0.50 Si 0. 40-2.0 Mn 0.03-0.12 Nb 0.05-
0.15V 0.2-0.8 Mo 0.015-0.03 Ti 0.01-0.03 AI,and comprising ferrite and
martensite/bainite phases, the ferrite phase having primarily vanadium and
niobium carbide or carbonitride precipitates prepared by a first rolling
above the austenite recrystallization temperature, a second rolling below
the austenite recrystallization temperature; cooling between the Ar3
transformation point and 500°C; and water cooling below about 400 A°e.

European patent EP104020S discloses an ultra high strength hot rolled dual
phase steel for cryogenic applications. The ultra-high strength, weldable,
low alloy, dual phase steel with high cryogenic temperature toughness
having a tensile strength greater than 830 MPa and a microstructure
comprising about 10 - 40 vol% ferrite and about 60 - 90 vol% of a second
phase of fine-grained lath martensite- lower bainite. The steel was produced
by heating a steel slab that contained carbon, manganese, nickel, nitrogen,
copper, chromium, molybdenum, silicon, niobium, vanadium, titanium,
aluminum and boron, a finish rolling of the plate between Ar3 and the Arl
transformation temperature, and quenching of the steel plate in a cooling
rate range of 10- 40°C per second, and a Quench Stop temperature below
about Ms transformation temperature plus 200°C , so as to achieve the
desired phase fraction, The resultant tensile strength was > 830 MPa with
DBTI of lower than about-73°C.
All the mentioned prior arts were largely in hot rolled conditions. Thus, it
was necessary to develop weldable and ductile high strength AHSS in cold-
rolled and annealed condition that are used in automotive.
OBJECTS OF THE INVENTION
It is therefore an object of the invention to propose a cold-rolled-
continuously annealed weldable dual phase stee! with a tensile strength-
6S0-800MPa for automotive applications, which adapts optimum selection of
alloying elements including microstructural engineering in selection of the
base metal.

Another object of the invention is to propose a cold-rolled-coniinuously
annealed weldable dual phase steel with a tensile strengthv 650-800MPa
for automotive applications, which allows defect-free welding.
A still another object of the invention is to porpose a process for producing
cold-rolled continuously annealed weldable dual phase steels with tensile
strength of 650-800 MPa.
A further object of the invention is to propose a process for producing cold-
rolled continuously annealed weldable dual phase steels with tensile
strength of 650-800 MPa ., in which the cooling rate is optimized from
intercritical annealing temperature.
A still further object of the invention is to porpose a process for producing
cold-rolled continuously annealed weldable dual phase steels with tensile
strength of 650-800 MPa, in which a predetermined amount of bainite is
incorporated in the ferritic/martensitic structure.
Yet another object of the invention is to propose a process for producing
cold-rolled continuously annealed weldable dual phase steels with tensile
strength of 650-800 MPa, in which a short cycle of low temperature
tempering in association with nano-scale microalloying precipitates is
conducted.

SUMMARY OF THE INVENTION
The present invention provides a process for producing in cold-rolled-
continuously annealed conditions a defect-free weldable dual phase steel
with ultimate tensile of strength of ~650-800MPa, spot welded nugget
diameter, 6Vt by optimally adapting alloying elements including an
innovative base metal microstructural engineering'.
Three trial heats were made in an air induction melting furnace with (wt%)
0.09-0.12-.C, 1.38-1.55:Mn,0.34-0.41:5i, 0.21-0.38:Mo, 0.007-0.031:Ti,
30-70ppm:Nb, 0.008:P, 0.008-0.013:S. The concept practiced in alloy
design was to incorporate low resistivity alloying elements in association
with low carbon equivalent to achieve a plug type failure in resistance spot
welding. Dilatometric studies for continuous heating and cooling
transformations were conducted on Gleeble 1500D to obtain'an intercritical
temperature regime and a transformation temperature for the second
phases (bainite and martensite). The steels were subsequently hot rolled
and cold rolled and were subsequently annealed in the inter critical region
using Gleeble 3500 to simulate a continuous annealing process. Three types
of annealing cycles were designed based on effect of i) cooling rate ii)
bainite and iii) tempered martensite, to produce crack-free weldments with
desired tensile properties. Plug type failure was achieved for all the samples
studied. The annealing cycles with < 30% martensite and cooling rate >
300C/s showed a poor weldability and elongation. HAZ softening was found
to be the main cause for poor weldability in some of the cycles. However,
the cycles that introduced bainite/tempered martensite in the ferritic-
martensitic dual phase structure in association with typical nano-sized
precipitates, such as (Ti-Nb-Mo)C, TiC, Mo:C, resulted in weldable steels
with desired tensile properties.

BRIEF DESCRIPTION OF THE ACCOMPANYING DRAWINGS
Fig. 1- shows a strip annealing device on a Gleeble 3500 thermomechanical
simulator.
Fig. 2- shows a CAL cycles conducted on the Gleeble 3500 for cooling rate
effect.
Fig. 3- shows a CAL cycles conducted on the Gleeble 3500 for bainite effect.
Fig. 4- shows the modified annealing device of the invention.
Fig. 5- shows a continuous cooling transformation diagram for trial Heat-9,
with micrographs exhibiting microhardness in VPN for corresponding
phases.
Fig. 6- shows a continuous cooling transformation diagram for trial Heat-10,
with micrographs exhibiting microhardness in VPN for corresponding
phases.
Fig. 7- shows a continuous cooling transformation diagram for trial Heat-14,
with micrographs exhibiting microhardness in VPN for corresponding
phases.
Fig. 8- shows a Ferritic-martensitic structure with 40% martensite for
annealing cycle 14-CR-l
Fig. 9: shows FEG-SEM and TEM micrographs indicating ferritic-martessitic
structure with 31% martensite for annealing cycle 14-CR-2.
Fig. 10: TEM micrographs showing (Ti-Nb-Mo) CN in association with TiC
nano-sized precipitates for annealing cycle 14-CR-2.

Fig. 11: Showing Top-Ieft)-photograph of pullout button failure, Bottom-
left) OM micrograph of predominantly martensitic structure in weld zone
along-with C-scan of weld zone, Bottom-right) OM micrograph showing
liquation crack in the fusion line [FEG-SEM magnified view in the top inset]
along with HAZ crack at the Mn-S stringers [bottom inset] and Top-right)
Macrograph of nugget with weldment hardness for annealing cycle 14-CR-1.
Fig.12: Showing Top-Ieft)-photograph of pullout button failure, Bottom-
left) OM micrograph of predominantly martensitic structure in weld zone
along-with the C-scan of weld zone, Bottom-right) OM micrograph
showing weld, heat affected zone and base metal and Top-right)
Macrograph of nugget with weldment hardness for annealing cycle 14-CR-2.
Fig. 13: FEG-SEM and TEM micrographs showing ferritic-martessitic
structure with 42% martensite for annealing cycle 9-CR-1.
Fig. 14: shows FEG-SEM and TEM micrographs showing ferritic-martessitic
structure with 35% martensite in association with nano-sized TiC precipitate
for annealing cycle 9-CR-2.
Fig. 15: Showing Top-Ieft)-photograph of pullout button failure with base
metal crack [inset], Bottom-left) OM micrograph of predominantly
martensitic structure in weld zone along-with weld crack at the MnS
stringers [inset] and C-scan of weld zone, Bottom-right) OM micrograph
showing liquation crack in the fusion line [FEG-SEM magnified view in the
top inset] along with HAZ crack at the Mn-S stringers [bottom inset] and
Top-right) Macrograph of nugget with weldment hardness for annealing
cycle 9-CR-1.

Fig. 16: Showing Top-left)-photograph of pullout button failure Bottom-
left) OM micrograph of predominantly martensitic structure in weld zone
along-with weld crack and C-scan of weld zonet Bottom-iigh)) OM
micrograph showing. HAZ crack at MnS stringers [FEG-SEM magnified view
in the bottom inset along with EDS line scan] and Top-righ)) Macrograph of
nugget with weldment hardness for annealing cycle 9-CR-2.
Fig. 17: FEG-SEM and TEM micrographs showing ferrite and martensite in
association with bainite (B+M-17%) for annealing cycle 14-B.
Fig. 18: Showing Top-left)-photograph of pullout button failure, Bottom-
left) OM micrograph of predominantly martensitic structure in weld zone
along-with the C-scan of weld zonet Bottom-iigh)) OM micrograph
showing weldt heat affected zone and base metal Top-righ)) Macrograph of
nugget with weldment hardness for annealing cycle 14-B.
Fig. 19 FEG-SEM and TEM micrographs showing ferritic-mattensitic structure
along with bainite (B+M-25%) for annealing cycle 9-B-1.
Fig. 20: Showing Top-left)-photograph of pullout button failure Bottom-
left) OM micrograph of predominantly martensitic structure in weld zone
along-with weld crack and C-scan of weld zonet Bottom-iigh)) OM
micrograph showing HAZ crack at MnS stringers [FEG-SEM magnified view
in the top inset along with EDS point analysis] and Top-righ)) Macrograph
of nugget with weldment hardness for annealing cycle 9-B-1.

Fig. 21: FEG-SEM and TEM micrographs showing ferritic-mattensitlc
structure along with bainite (B+M-37%) and intergranular Mo-C precipitate
for annealing cycle 9-B-2.
Fig. 22: Showing Top-Ieft)-hhotograph of pullout button failure, Bottom-
left) OM micrograph of predominantly martensitic structure in weld zone
along-with the C-scan of weld zone, Bottom-right) OM micrograph
showing weld and heat affected zone and Top-right) Macrograph of nugget
with weldment hardness for annealing cycle 9-B-2.
Fig. 23: FEG-SEM and TEM micrographs showing ferritic-tempered
martensitic structure (with Fe3C) for annealing cycle 10-T.
Fig. 24: TEM micrographs showing intergranular precipitation of nano-sized
Mo-C, (Ti-Mo)-C and Ti-C precipitates for annealing cycle 10-T.
Fig. 25: Showing Top-Ieft)-photographof pullout button failure, Bottom-
left) OM micrograph of predominantly martensitic structure in weld zone
along-with the C-scan of weld zone, Bottom-right) OM micrograph
showing weld and heat affected zone and Top-right) Macrograph of nugget
with weldment hardness for annealing cycle 10-T.
Fig. 26: FEG-SEM and TEM micrographs showing ferritic-partially tempered
martensitic structure in association with nano-sized TiC and TiN precipitates
for annealing cycle 9-T.

Fig. 27: Showing Top-Ieft)-photorraph of pullout button failure, Bottom-
left) OM micrograph of predominantly martensitic structure in weld zone
along-with the C-scan of weld zone, Bottom-right) OM micrograph
showing weld and heat affected zone and Top-ri~ht) Macrograph of nugget
with weldment hardness for annealing cycle 9-T.
Detailed Description Of the Invention:
Experimental
Material
Three heats were made using a 25-kg in-house laboratory induction
furnace. Chemical composition (in wt%) of the heats (Heat Nos. 9, 10 and
14) along with the carbon equivalents are presented in Table I.
Table I: Chemical composition and carbon equivalent of the developed
steels


Continuous Annealing Simulation
The ingot samples were duly forged and hot rolled. Subsequently the hot
rolled samples were cold rolled to ~ 0.8-1 mm using the laboratory cold
rolling set-up at NML Jamshedpur for subsequent continuous annealing
(CAL) simulation. To simulate continuous annealing cycles in, a strip
annealing device was utilized (Fig. 1). The strip samples of 260mmX50mm
were used for the annealing simulation whi.e a K-type thermocouple (T1)
was welded to the samples at the centre as shown in Fig 4. Further, to
confirm uniform heating, two more thermocouples T2 and T3 were attached
on each side of the centre, at a distance of 25-mm as shown in Fig 4.
The annealing simulation was carried out within the intercritical temperature
range (u+y) that was determined by continuous heating transformation
tests on Gleeble 1500D. Three types of continuous annealing cycles were
designed for strip annealing simulation on Gleeble 3500 to examine the
effect of cooling rate, bainite and tempering on tensile properties and
weldability of ferritic-martessitic dual phase steels to obtain optimized
processing parameters for achieving weldable steel with the desired tensile
properties.
1) Effect of cooling rate
Cold rolled samples of heat-14 and heat -9 were used to examine the
effects of cooling rate in the simulated annealing cycles. The parameters
used in the annealing cycles along with their thermal cycles are given in
Table III and Fig 2. respectively.


Effect of bainite
Cold rolled samples of Heat-14 and Heat-9 were used to examine the effects
bainite in the simulated annealing cycles. The parameters used in the
annealing cycles along with their thermal cycles are given in Table IV and
Fig. 3 respectively.


Effect of tempering
Cold rolled samples of heat-10 and heat -9 were used to examine the
effects of tempering in the simulated annealing cycles. The parameters used
in the annealing cycles along with their thermal cycles are given in Table V
and Fig. 4 respectively.


Findings
ContInuous cooling transformation
Heat-9
From continuous cooling transformation curves in Fig 5, it can be seen that
below 5°C/s cooling rate, ferrite, pearlite in association with bainite were
obtained, while above 5°C/s and below 20°C/s, the co-existence of ferrite
and bainite were noticed. Further increase in cooling rate till 70°C/s showed
the presence of martensite in association with ferrite and bainite, while
above 70°C/s, bainite disappeared and only ferrite and martensite
remained. The phases were identified by morphological and microhardness
evidences that are depicted by the representative micrographs presented in
Fig 18.

Heat-10
In this steel, up to 10°C/s cooling rate, only polygonal and acicular ferrite
was observed and between 10-20°C/s, the presence of pearlite and bainite
was also evident (Fig. 6). However, at higher cooling rates (above 20°C/s till
70oC/s) pearlite disappeared and the microstructure consisted of ferrite,
bainite and martensite.
Heat- 14
In Heat-14, pearlite was observed in association with ferrite and bainite in
the cooling rate range of 10-20oC/s, while till 10°C/s ferrite structure was
dominant (Fig. 7). While from 20 to 100°C/s, ferrite, bainite and martensite
were present.
For designing continuous annealing cycles, fast cooling rate was maintained
in the pearlite formation regime to avoid pearlite formation as the presence
of pearlite develops yield point elongation and which in turn causes surface
problems. Thus, with the help of the CCT diagrams (Figs 5-7), the cooling
rate for Heat-9 was kept above 5°C/s, while for Heat-10 and 14, the cooling
rate maintained was above 20°C/s.
Continuously annealed samples, tensile properties and weldability
Effect of cooling rate
Annealing cycles 14-CR-l & 14-CR-2:
Samples that underwent cycles 14-CR-l.. 10oC/s-830oC (120s)-2°C/s-
800oC-100°C/s and 14-CR-2: 10oC/s-850oC (120s)-2°C/s-800°C-30°C/s,
experienced cooling rates of 100 and 30oC/s cooling rates, respectively. In
both the samples ferrite martensitic microstructure was obtained with 40%
and 31% martensite fraction, respectively (Figs. 8, 9). Due to fast cooling

rate of 100 °C/s, no new precipitates could be formed during the cooling
cycle of 14-CR-l, however, as the cooling rate in 14-CR-2 was slower
(30oC/s), some TiC and (Ti-Nb-Mo)C precipitates were also formed (Fig. 10.
As solubility increases with temperature, the Ti-containing precipitates some
diffusion of the solutes (Ti) from the periphery of Ti containing precipitates
to the matrix also possibly took place at t~e intercritical annealing
temperature, that finally resulted in complex precipitates of (Ti-Nb-Mo)C,
during the slower cooling cycles ~30°C/s.
The welded sample of 14-CR-l showed plug type nugget failure (Fig. 11,
Top-left) with predominantly martensitic in association with bainitic welded
structure (Fig. 11, Bottom-left). However, the nugget quality manifested by
the non-destructive method, C-scanning was found to be poor (Fig. 11,
Bottom-left). The weld zone was examined by FEG-SEM also and porosities
were observed thus supporting the assessment made by C-scanning (It is to
be noted that at the output of C-scanning, the data are presented as color
maps and colors from the bottom to the middle portion of the color scale
imply satisfactory weld quality). Further, cracks initiated in a zone where
the alloy remained in the solidus-liquidus temperature range during
welding. The dendrite structure inside open cracks (Fig. 11, Bottom-right)
proves that the liquid had to be present at the moment of crack formation
and therefore, it is liquation cracking according to the classification of
Hemsworth et -aI [38]. In other words, if a crack exists during Heating, it
may be easily filled because of high pressure inside the liquid nugget and
the healed structure appears like dendrites [39]. Liquation cracking usually
occurs while the high residual stresses that occur during weld cooling tend
to rupture the melted boundaries containing loW melting films. It is seen
that while the weld deposit is still liquid, compressive stresses tend to close
up liquation cracks; however, as the melt solidifies, the stresses in the HAZ

become tensile and these open up cracks [40]Itt is to be noted that
liquation cracking is usually associated with grain boundary segregation,
however, in this case no segregation was observed. In addition, cracking in
the HAl affected zone also was initiated by the presence of MnS stringers.
This can be attributed to the development of triaxial stresses at the sharp
corners of the stringers during the cooling cycle of the welding. This helped
lower the cohesive force at the MnS-matrix interface and thus initiated a
crack (Fig. 11, Bottom-right.. The nugget diameter is within the desired
limit for 14-CR-1 sample (6Vt=577 mm), 4.62 mm, as presented in the Fig
11, Top-right. While hardness values for the weld, HAl and base metal were
309, 256 and 290 VPN, respectively. Thus, hardness of the HAl was quite
lower than that of the weld region and base metal that is indicative of the
fact that the HAl softening was the possible reason for the HAl cracking
and poor weldability. It was therefore observed that HAl softening occurred
in this case while significant amount of martensite (40%) was present in
the base metal, which was in good agreement with the conclusion made in
the literature [8,9]. In HAZ softening, martensite tempering takes place at
the points of the weld where the temperate approached A1, subcritical HAZ.
This is because a hard weld metal can not deform plastically to relieve
welding stresses and they are transferred to the HAZ, causing it to crack
[41]. This is the reason why softer base metal consistently suffers a smaller
amount of HAZ cracking than the harder consisting of large amount of
martensite. This is consistent with the suggestion of Borland Younger [42]
that a weld metal with lower plastic resistance than the base metal is
beneficial since the amount of stress that is needed to be accommodated by
the base metal is reduced, which in turn reduces the amount of cracking.

However, the weldment of 14-CR-2 exhibited better welding performance
than 14-CR-1 (Fig. 12, Top-left)). A plug type nugget failure was obtained
and the microstructure of the welded zone was again bainitic martensitic.
However, no imperfections were observed in the welded region as per C-
scanning (Fig 12, Bottom-left) that was duly confirmed by FEG-SEM
examination. No cracks were found in the weldment ( Fig. 12, Bottom,
right) and the nugget diameter of 4.73 mm (Fig 12, Top-right) was also
within the desired limit (6Vt=622 mm). Further, hardness of the HAZ (390
VPN) was higher than both the welded (353 VPN) and base metal (250
VPN). Thus, there was no HAZ softening and subsequent cracking.
The YS and UTS of cold-rolled and continuously annealed 14-CR-1 were
registered as 405 MPa (Table VII) and 750 MPa, respectively, with % UEL
(%Uniform Elongation) of 13% and %TEL (%Total Elongation) of 19% for
100oC/s cooling rate and 40% martensite content. On the other hand, as
the cooling rate was reduced to 30oC/s in cycle 14-CR-2, the martensite
content reduced to 31%. In addition, nano-sized precipitates such as, TiC in
association with complex precipitates, (Ti-Nb-Mo)C were observed. The
precipitates helped increase the YS to 494 MPa compared to 405 MPa in 14-
CR-1, by grain refinement and restricting dislocation movement. However,
UTS was reduced to 690 MPa for 14-CR-2, as martensite content reduced,
while %UEL and TEL were increased to 17% and 27%, respectively. The
strain hardening exponents (n) for 14-CR-1 and 2 were satisfactory with
0.18 and 0.20, respectively.

Table VII: Microstructure and tensile properties of CAL simulated samples
along with weld nugget diameter.


Annealing cycles 9-CR-1 & 9-CR-2:
Cold rolled samples of Heat-9 underwent the cycles, 9-CR-1: 10oC/s-780oC
(120s)-2°C/s-730oC-100oC/s and 9-CR-2: 10oC/s-780oC (120s)-2°C/s-
730oC-50°C/s, to examine the cooling rate effect of 100oC/s and 50oC/s on
the tensile properties and voldability of the samples.
The annealed microstructures of both the cycles, 9-CR-1 (Fig. 13) and 9-
CR-2 (Fig. 14), resulted in a ferritic microstructure along with 42% and 35%
martensite fractions, respectively. In addition, like cycle 14-CR-2 with slow
cooling rate of 30°C/s, TiC precipitates (Fig. 14) were also observed in 9-
CR-2 sample that experienced a slower cooling rate of 50oC/s.
However, in this case, both the cycles resulted in samples that manifested
poor weldability. 9-CR-1 sample though showed pull-out fracture (Fig. 15,
Top-left), the same had predominantly martensitic weld microstructure with
weld crack at MnS stringers and porosities that were observed using C-
scanning and FEG-SEM (Fig. 15, Bottom-left). In addition, liquation cracking
in the partially melted zone was also found (Fig. 15, Bottom-righ)) [39,40..
The weldment had full of cracks at the crack initiation sites of MnS stringers
in the base metal (Fig. 15 Top-left) as well as in the HAZ (Fig. 15, Bottom-
right) and weld metal (Fig. 15, Bottom-left).
While the nugget diameter (4.53 mm) (Fig. 15, Top-right) was within the
desired limit (6Vt=5.6 mm), the hardness data for all the zones in the
weldment, weld (326 VPN), HAZ (330 VPN) and base metal (320 VPN) were
almost the same. This can be possibly due to the presence of very high
amount of martensite (42%) in the base metal. Thus, as all the zones in the
weldment were hard, the complex stresses that generate during the cooling

of the weldment, could not be accommodated by plastic deformation and
therefore, all the zones crack to dissipate the stresses developed.
The cycle 9-CR-2 also resulted in a poor weldment with a plug type failure
(Fig. 16, Top-left.. The welded region had solidification crack without any
low melting phases in the crack (Fig. 16, Bottom-left.. This was examined
using C-scanning and FEG-SEM. The presence of solidification crack can be
attributed to the stress generated during weld cooling in association with
the presence of predominantly hard martensitic structure. In this case as
well, HAl cracking was observed. However, no liquation cracking was
noticed apart from the MnS stringer related cracks ( Fig. 16, Bottom-right).
Further, HAl cracking occurred on account of HAl softening [8, 9] as the
hardness value registered for the HAl was lower with 319 VPN than that of
the weld region hardness (345 VPN) (Fig. 16, Top-right). However, the
nugget diameter of 4.64 mm satisfied the desired range of <6^t=5.9 mm.
For cycle 9-CR-1 with 100°C/s cooling rate, YS & UTS were 417 and 827
MPa (Table VIII) respectively while, %UEL was a bit low with 12% and
%TEL was 20%, while martensite in the ferrite matrix was 42%. Cycle 9-
CR-2 with 50°Cjs cooling rate, on the other hand, could register a YS of 546
MPa that is much higher than the YS obtained from 9-CR-1 cycle. This can
be attributed to the presence of TiC precipitates in ferrite matrix, for 9-CR-2
cycle. While UTS, %UEL and %TEL for the cycle ~ere recorded as, 806 MPa,
10% and 17% respectively. For both the cycles, strain hardening exponent
was also low, 0.14 (9-CR-1) and 0.15 (9-CR-2).

Effect of Bainite
Annealing cycle 14-B:
Cold rolled sample of Heat-14 was used to see the effect of bainite in the
annealed microstructure in association with ferrite and martensite on tensile
properties and weldability of the steel. Thus, the sample was submitted to
the annealing cycle: 10oC/s-760°C (60s)-100oC/s-450°C (60s)-100oC/s.
The resultant microstructure obtained was ferrite with 17% martensite and
bainite [Fig. 17].
The base metal microstrucrure resulted in a good weld with plug type
failure (Fig 18, Top-left) with ferritic-mattensitic-bainitic weld
microstructure without any weld imperfection as shown by C-scanning data
(Fig. 18, Bottom-left.. Further, no evidence of HAl cracking was noticed (
Fig. 18, Bottom-right) that was also supported by the hardness
measurements of the weldment, as there was no HAl softening (the HAl
hardness, with 309 VPN was more than the weld ( 294 VPN) as well as the
base metal (219 VPN)). This can be attributed to the replacement of a part
of martensite by bainite that rendered beneficial effect on weldability. In
addition, the weld nugget diameter was 4.62 mm that was also within the
acceptable range of 6.2mm (6Vt)t
With the introduction of bainite in the base metal matrix along with ferrite
and martensite (17% B+M), the % UEL (17%) % TEL (25%) and n (0.2)
values improved (Table VII), while YS, UTS were recorded as 529 and 664
MPa, respectively.

Annealing cycles 9-8-1 and 9-8-2:
The base metal microstructure of the cold rolled and continuously annealed
sample of cycle 9-B-1 (10oC/s-780°C (120s)-2°C/s-760oC-40oC/s-470°C
(60s)-70oC/s) was ferrite with.25% martensite and bainite (Fig. 19). In this
case, bainite was quite less compared to martensite content. While the
sample was welded a plug type failure (Fig. 33, Top-left) was observed after
peeling. The welded zone with bainitic-martensitic structure though had no
porosities or cracks ( Fig. 33, Bottom-left), the HAl was not devoid of
cracks (Fig. 20, Bottom-right). Short cracks in the HAl were again found to
have initiated and propagated at the MnS stringers. Crack generation in the
HAZ can be attributed by HAl softening [8,9]. The hardness of the HAl was
recorded to be 250 VPN while the hardness values of the weld and base
metal were 330 and 300 VPN, respectively (Fig. 20, Top-right). The nugget
diameter although was within the desired limit (4.33 mm) and less than 6Vt
(6 mm).
Another annealing cycle, 9-B-2: 10oC/s-780oC (120s)-2°C/s-760oC-
20oC/s-470oC (60s)-10oC/s showed ferritic-bainitic-martensitic
microstructure with bainite and martensite contents of 37%, in which the
amount of bainite was more than martensite (Fig. 21). Further,
intergranular precipitation of Mo-C was also observed because of slower
cooling rate of 20oC/s from quenching temperature to the bainitic
temperature. The sample was welded successfully with pull-out nugget
failure (Fig. 22, Top-left), without any weld (Fig. 22, Bottom-lef))
imperfection, as indicated by C-scanning, and HAl cracking (Fig. 22,
Bottom-right). The successful welding can be attributed to the presence of
substantial amount of bainite in the base metal microstructure that helped

avoid HAZ softening ( HAZ-360 VPN, Weld 355 VPN, base metal-444 VPN)
( Fig. 22, Top-right) by providing a formable base metal to accommodate
the plastic residual stresses during weld cooling. The nugget diameter was
measured to be 3.89 mm that is less than 6Vt (5.1 mm).
The YS and TS with 25% M+B in the ferrite matrix of 9-B-1 were obtained
as 600 and 702 MPa, respectively (Table VII). However, in this case bainite
amount was much lower than the martensite. The %El in terms of %UEL
and %TEl, was found to be 16 and 23%, respectively, while V value was
quite satisfactory with 0.23. The cycle 9-B-2 offered one of the best
annealing cycles in the present study with YS and UTS of 449 and 753 MPa
with % UEl, %TEL and V as 16%, 22% and 0.18 (Table VII), respectively.
A substantial amount of the second phase bainite in association with nano-
sized Mo-C and TiN precipitates helped achieve the above properties.
Effect of Tempering
Annealing cycles 10-T &9-T:
The cold rolled sample of Heat-10 was submitted to the following annealing
cycle consisting of a tempering step: 10oC/s-780oC (120s)-10oC/s-700oC-
22°C/s-200°C(300s)-3°C/s (10-T). The sample manifested the presence of
ferrite along with 23% tempered martensite structure (Fig. 23) in
association with nano-sized precipitates of Mo-C, (Ti-Mo)C (heterogeneous
nucleation on AIN) due to slower cooling (22°C/s) (Fig. 24) from the quench
temperature, and manifested good weldability (Fig. 24). A favourable plug
type failure along with porosity- and crack-free weld and HAZ were
observed (Fig. 25, Top-left, Bottom-left and Bottom-right)Thss can be
attributed to the absence of HAZ softening (HAZ-375, Weld-313 and Base

Metal-240 VPN) (Fig. 25, Top-right) due to the presence of formable base
metal in the presence of temperared martensite. The nugget size of 3.92
mm was also within the desired limit of 6Vt =6 mm.
An introduction of tempering step in the 10-T cycle resulted in improvement
ofweldability and tensile properties.
On the other hand, the cold rolled sample of Heat-9 with annealing cycle 9-
T, 10oC/s-7500C(120s)-2oC/s-730oC-100oC/s-200°C (120s)-60oC/s
resulted in ferrite and tempered martensite (30%) microstructure along
with TiC precipitates (Fig 26). The presence of nano-sized TiN precipitates
that precipitated during solidification process of heat making is also
presented in Fig. 26. A favourable plug type failure (Fig. 27, Top-left) was
noticed that was associated with porosity- and crack-free weld zone and
HAZ. In this case also the presence of good weld can be accounted for the
absence of HAl softening (HAZ-367, Weld-343 and Base metal-280
VPN). This can be due to the introduction of tempering step and reduction of
solute carbon in ferrite matrix. The tempered martensite in the base metal
helped accommodate the plastic strain generated during weld cooling and
thus avoided HAl cracking [41,42]. The nugget diameter size was 4.93 mm
that was within the desired limit of 5.9 mm (6Vt ) (Fig. 27, Top-right).
The 9-T samples resulted in YS of 491 MPa and a high UTS of 812 MPa
(Table VII) along with 12%, %UEL and 19% %TEL The microstructure
observed was ferrite with partially temper~d 31% M in association with
nano-sized TiC and TiN (formed during solidification) precipitates.

The data for all the annealing cycles that depict good (marked as OK) and
poor welds in the developed steels, are provided. From the characterization,
it is clear that the base metal microstructure plays an important role in
affecting the HAZ hardness. It was inferred that HAl hardness needed to be
higher than the weld metal (i. e.HHAz > H WM.. while the same requires to be
substantially stronger than the base metal to avoid HAl softening and thus
HAl cracking.
The steels developed with current process and given chemistry have
ultimate tensile strength in the range of -660-810 MPa and total elongation
in the range of 19-27%. In all the annealing conditions the desirable plug
type nugget failure was observed with nugget diameter <6Vt.
Further, Microstructure containing martensite >~ 30% in the base metal
and cooling rate of >30oC/s from intercritical annealing temp., (without
bainite/tempeiing effect) adversely affected weldability. In weldments
where HHAzz HBM (HAl softening) liquation cracks, and cracks at MnS
stringers in HAl/HAZ and WM were observed. While, cracks in base metal
were also noticed when HHAzz HWELD~ HBM
Introduction of optimum amount of bainite in the ferrite and martensite
structure resulted in satisfactory weldability behavior. Tempering of
martensite in the ferrite-martensite matrix registered favourable weldability.
Nano-sized precipitates like TiC, TiN, and (Mo-Ti) C, (Ti-Nb-Mo) CN etc
helped in contributing to strengthening. Cycle 9-6-2 (10oC/s-70o°C (120s)-
2°C/s-760oC-20oC/s-470oC (60s)-10°C/s) with ferrite, 37% bainite and
martensite along with nano-sized precipitates of Mo-C and TiN and 10-T-2
(10oC/s-780oC (120s)-10oC/s-700oC-22°C/s-200oC(300s)-3°C/s) with
ferrite and 23% tempered martensite in association nano-sized (Mo-Ti)C
and TiC precipitates resulted in the best combination of tensile properties

with plug type nugget failure and weldabllity-(B-B-2: YS-449 MPa , UTS-
753 MPa, %UEL-16%, %TEL-22% and n-0.18 and 10-T-2: YS-415 MPa ,
UTS-750 MPa, %UEL-18%, %TEL-23% and n-0.17.. Strong HAZ and
formable base metal resulted in favorable combinaiion of strength-ductility-
weld ability.

We Claim:
1. A cold-rolled continuously annealed, high strength dual phase steel with
a tensile strength >650 MPa, comprising:
Carbon (C) in an amount from 0.09 weight percent to 0.12 weight
percent;
Manganese (Mn) in an amount from 1.38 weight percent to 1.55
weight percent;
Silicon (Si) in an amount from 0.34 weight percent to 0.41 weight
percent;
Molybdenum in an amount from 0.21 weight percent to 0.38 weight
percent;
Titanium (Ti) in an amount from 0.007 weight percent to 0.031 weight
percent;
Niobium (Nb) in an amount from 30 ppm to 70 ppm; Phosphorous (P)
in an amount of 0.008 weight percent; and Sulfur (5) in an amount
from 0.008 weight percent to 0.013 weight percent;
balance being iron,
wherein the steel comprises from about 63 to about 77 percent by
volume of a first phase, the first phase comprising a ferrite mean
grain size of about 5 microns or less;
and wherein , the steel further comprises 23 to 37 percent by volume
of a second phase, the second phase comprising fine-grained
martensite/mattensite

bainite/tempered martensite in association with nano-sized Mo-C, Ti-C
and (Ti-Nb-Mo)C or any mixture thereof.
2. The steel as claimed in claim 1, wherein ultimate tensile strength of the
steel is in the range of 660-810 MPa.
3. The steel as claimed in claim 1, wherein total elongation of the steel is
in the range of 19-27%.
4. The steel as claimed in claim 1, wherein spot welded nugget diameter is
less than 6vt in cold-rolled continuously annealed conditions.
5. A method of preparing cold-rolled continuously annealed, high strength
dual phase steel with a tensile strength >650 MPa, the method
comprising:
- hot and cold rolling a steel composition comprising carbon (C) in an
amount from 0.09 weight percent to 0.12 weight percent; manganese (Mn)
in an amount from 1.38 weight percent to 1.55 weight percent; silicon (5i)
in an amount from 0.34 weight percent to 0,41 weight percent;
molybdenum in an amount from 0.21 weight percent to 0.38 weight
percent; titanium (Ti) in an amount from 0.007 weight percent to 0.031
weight percent; niobium (Nb)in an amount from 30 ppm to 70 ppm;
phosphorous (P) in an amount of 0.008 weight percent; sulfur (5) in an an
amount from 0.008 weight percent to 0.013 weight percent and balance
being iron; and inducing said steel composition on to a plurality of annealing
cycles consisting of:

6. The method as claimed in claim 5, wherein the steel composition has the
following characteristics:

ABSTRACT

The present invention is related to developing defect-free weldable dual
phase steels in laboratory of strength ~650-800MPa with spot welded
nugget diameter, 6Vt in cold-rolled-continuously annealed conditions by
judicious selection of alloying elements and base metal microstructural
engineering. The method involves hot and cold rolling a steel composition
comprising of carbon (C) in an amount from 0.09 weight percent to 0.12
weight percent; manganese (Mn) in an amount from 1.38 weight percent to
1.55 weight percent; silicon (5i) in an amount from 0.34 weight percent to
0,41 weight percent; molybdenum in an amount from 0.21 weight percent
to 0.38 weight percent; titanium (Ti) in an amount from 0.007 weight
percent to 0.031 weight percent; niobium (Nb) in an amount from 30 ppm
to 70 ppm; phosphorous (P) in an amount of 0.008 weight percent; sulfur
(5) in an amount from 0.008 weight percent to 0.013 weight percent and
balance being iron; and inducing said steel compositing to three types of
annealing cycles
Three types of annealing cycles were designed: effect of i) cooling rate ii)
bainite and iii) tempered martensite, to obtain crack free weldments with
desired tensile properties. Plug type failure was obtained for all the samples
studied. The annealing cycles with < 30% martensite and cooling rate >
30°C/s cooling rate showed poor weldability and elongation. HAZ softening
was found to be the main cause for poor weldability in some of the cycles.
However, the cycles that introduced bainite/tempered martensite in the
ferritic-martensitic dual phase structure in association with typical nano-
sized precipitates, such as (Ti-Nb-Mo)C, TiC, Mo-C resulted in weldable
steels with desired tensile properties.