J. FODER et al.: MEAN-FLOW-STRESS ANALYSIS OF LABORATORY HOT-ROLLED S1100QL STEEL ...
901–908
MEAN-FLOW-STRESS ANALYSIS OF LABORATORY
HOT-ROLLED S1100QL STEEL WITH MINOR Nb ADDITION
MFS ANALIZA LABORATORIJSKO VRO^E VALJANEGA JEKLA
S1100QL Z MANJ[IM DELE@EM Nb
Jan Foder
1
, Grega Klan~nik
1*
, Jaka Burja
2
, Samo Kokalj
1
, Bo{tjan Brada{kja
1
1
RCJ d.o.o., Cesta Franceta Pre{erna 61, 4270 Jesenice, Slovenia
2
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
Prejem rokopisa – received: 2020-07-06; sprejem za objavo – accepted for publication: 2020-10-05
doi:10.17222/mit.2020.125
Laboratory hot-rolling of ultra-high-strength martensitic steel S1100QL was carried out by measuring the material resistance
during hot-rolling using rolling force measurements. A simplified approach was used to calculate the mean flow stress as an in-
direct method to analyse the material’s response during hot rolling. It was shown that the method gives satisfactory results, re-
garding the rough detection of different temperature-deformation regions like crossing the region of considerable solute drag
and precipitation hardening, known as the non-recrystallization region. The transition from the recrystallization region, where
softening mainly by static or meta dynamic recrystallization is replaced with a non-recrystallization region having austenite
pancaking, is clearly distinguishable for given rolling parameters and final strip thickness. The grain size distribution of the
S1100QL was also evaluated from the as-cast state to the heat-treated condition.
Keywords: S1100QL, micro-alloying, hot-rolling, CALPHAD
Izvedeno je bilo laboratorijsko vro~e valjanje ultra-visoko-trdnostnega martenzitnega jekla S1100QL z analizo materialnega
odziva na preoblikovanje preko meritev sil valjanja. Izra~unana je bila srednja napetost te~enja (MFS), preko katere lahko
indirektno analiziramo odziv materiala med vro~im valjanjem. Dokazali smo, da z uporabo MFS analize dobimo zadovoljive
rezultate, vezano na detekcijo razli~nih temperaturno-deformacijskih podro~ij, kot je obmo~je kjer raztopljeni elementi
povzro~ajo dodaten upor in obmo~je izlo~evalnega utrjevanja, znano kot obmo~je ne-rekristalizacije. Za dane parametre valjanja
in kon~no debelino traku je jasno viden prehod iz podro~ja, kjer poteka meh~anje s stati~no ali metadinami~no rekristalizacijo, v
podro~je kjer je ta omejena oziroma ne poteka, s prehodom avstenitnega kristalnega zrna iz ekviaksialne v pala~inkasto obliko.
Dolo~ena je bila tudi razporeditev velikosti prvotnih avstenitnih kristalnih zrn od litega do kon~nega toplotno obdelanega stanja
jekla S1100QL.
Klju~ne besede: S1100QL, mikrolegiranje, vro~e valjanje, CALPHAD
1 INTRODUCTION
Ultra-high-strength steels (UHSS) and high-strength
low-alloyed steels (HSLA) have a high strength-to-
weight ratio and are often used in the automotive indus-
try, construction, mining and elsewhere. HSLA and
UHSS are both used for demanding welded structures,
with the emphasis on achieving a high yield strength and
ductility. Various strengthening mechanisms like grain
refinement, the formation of deformation substructures
(dislocations), solid-solution strengthening and precipita-
tion strengthening are used to achieve the ultra-high
strength. The strengthening mechanism depends on the
given alloy composition and the production route. In this
paper the influence of hot rolling is emphasized. Never-
theless, a combination of secondary metallurgy and cast-
ing practice with final heat treatment is also of great im-
portance for consistent production quality. Therefore,
similar chemical compositions and different technologi-
cal routes, such as different rolling schedules, will un-
doubtedly lead to different material properties, which
will also be measured in the final heat-treated condition.
Unfortunately, all strengthening mechanisms, with the
exception of grain refinement, have a negative effect on
the impact toughness and ductility. Hot rolling is there-
fore a key step in the production of UHSS and HSLA
with superior properties.
1.1 S1100QL
Niobium microalloyed steel S1100QL (Mat.No
1.8942) is a non-standardized fine-grained, fully marten-
sitic UHSS. S1100QL must meet the minimum yield
strength of 1100 MPa (with a tensile strength between
1200 MPa and 1500 MPa) and a Charpy impact tough-
ness of at least 27 J transverse and 30 J longitudinal to
the rolling direction at –40 °C.
1
Grain size control is
achieved with proper hot-rolling and selected reheating
temperatures. Low-temperature tempering is performed
after water quenching to maintain adequate yield
strength and increase in ductility. This paper focuses on
achieving the minimum yield strength of 1100 MPa by
quenching and low-temperature tempering. In the case of
on-line heat treatment, austenite conditioning prior to
quenching must be considered. Irrespective of the type of
Materiali in tehnologije / Materials and technology 54 (2020) 6, 901–908 901
UDK 621.77:620.1:669.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 54(6)901(2020)
*Corresponding author's e-mail:
klancnik.grega@gmail.com (Grega Klan~nik)

heat treatment, it is recognized that the martensite block
and packet size tend to scale linearly with the parent aus-
tenite grain size (PAGS), affecting the size of the
high-angle substructures (>15° misorientation).
2
This is
important if the packet size is considered as the "effec-
tive grain size" in relation to the yield strength and
toughness of the lath martensite.
3
The effective grain size
and distribution in combination with steel cleanliness,
which is the result of secondary metallurgy practice, de-
termines the functional properties such as the weldability
and formability of high-strength steels, therefore an opti-
mal rolling schedule is important and necessary.
4,5
Nowa-
days, the advanced hot-rolling technology for HSLA and
UHSS plate production is based not only on austenite
grain refinement and control over the precipitation of ni-
trides, carbides and carbo-nitrides, but also on transfor-
mation control during cooling after hot rolling with the
possibility of significantly reducing the alloying, such as
reducing the expensive molybdenum, and still obtaining
excellent or even improved mechanical properties. This
is possible if the process of controlled hot rolling is cou-
pled with on-line controlled cooling using accelerated
cooling (ACC), ultra-fast cooling (UFC) or even direct
quenching (DQ) technology.
6
This promising technology
of thermo-mechanical control processes (TMCP) with
controlled cooling, not only to improve the strength and
ductility properties for certain steel grades, but also for
steel mill time and energy efficiency, still needs to be
adapted for HSLA and UHSS production in the local re-
gion.
1.2 Strip/plate rolling
Phenomena occurring during the hot deformation
(hardening, softening etc.) of steels are usually studied
by hot compression, tension and torsion tests on
thermo-mechanical simulators.
7–9
Numerous specimens
and tests are required to obtain the relevant data on the
hot-deformation behaviour. The use of a laboratory roll-
ing mill for the evaluation of the mean flow stress (MFS)
is an alternative method for the investigation of the hot
deformation of steels. It is particularly useful for obtain-
ing important data during successive reductions in the
hot-rolling process. In addition, basic mechanical proper-
ties such as strength and toughness, bending, etc. are ob-
tained from the sample in a similar way to industrial
strip/plate production, as larger samples are produced
and tested.
The final grain size can be achieved through strain-
free austenite, formed by plastic deformation. Only lim-
ited grain refinement is expected after cooling and the
transformation from austenite to ferrite. This means that
during conventional hot rolling and recrystallization-con-
trolled rolling (RCR), the main deformations are carried
out as part of the roughing phase. The latter with an em-
phasis on higher per-pass reductions. When finishing
passes are introduced, they are mainly adjusted to the
target dimension tolerance and flatness in relation to the
selected finish rolling temperatures (FRT). A finer grain
size is expected when additional deformation at lower
temperatures is introduced to obtain deformed austenite
grains with deformation bands and strain-induced precip-
itates (SIP). Combined two-stage rolling of roughing and
finishing with a certain time delay between the roughing
and finishing phase is called controlled rolling.
10
Accord-
ing to Dutta et al.,
11
the recrystallization limit tempera-
ture (RLT) is the lowest temperature for the roughing
phase to prevent the formation of a mixed microstructure
by partial recrystallization. A mixed microstructure
causes the scattering of the final mechanical proper-
ties.
5,10
The RLT temperature can be interpreted as the
non-recrystallisation temperature T
nr
, due to limited
recrystallisation.
Based on the recrystallization behaviour, plastic de-
formation can be divided into three main characteristic
regions (Figure 3):
5,10
• Region I – typical for deformation above the RLT, the
temperature where approx. 95 % recrystallization still
occurs. The deformation per pass and the interpass
times control the softening ratio achieved by static
recrystallization (SRX) or meta dynamic recrystal-
lization (MDRX) during rough rolling or RCR.
• Region II – is below the recrystallization stop tem-
perature (RST), the temperature below which less
than 5 % recrystallization still occurs.
• Region III – is between region I and II (between RLT
and RST) and should be avoided due to the limited
control of recrystallization, grain growth and precipi-
tation. It is also referred to as the partial recrys-
tallization region (95 % to5%ofSRX).
Below the RST, where recrystallisation of less than
5 % is expected, both finish rolling and hot levelling are
usually carried out for controlled rolling. In some cases,
due to the sufficient number of passes and associated
temperature drop of the final passes, this is also done by
conventional hot rolling of thin plates or strips. Softening
at low temperatures can be influenced by the introduc-
tion of microalloying elements such as niobium for SIP.
According to Dutta and Sellars,
11
the retarding effect of
niobium on austenite recrystallization is observed to a
lesser extent for niobium in a solid solution and more
pronounced with the SIP of Nb(C,N) from the austenite
matrix, which means that sufficient accumulation of the
deformation for SIP can effectively stop recrystallization
recognized as Region II (below the RST). It has also
been recognized that RST occurs when the time for 5 %
recrystallization (or less) is equal to the time for5%of
precipitation of Nb(C,N).
11
Inhibition of recrystallization
is effective until the end of precipitation and loss of the
effective pinning force (Zenner pinning). Austenite
grains are therefore elongated (pancaked) in rolling di-
rection if sufficient strain is accumulated and that mini-
mum incubation time is respected for the successful SIP
of Nb(C,N) (see Figure 1c).
12
It is also recognized that
deformation accelerates precipitation, providing prefer-
J. FODER et al.: MEAN-FLOW-STRESS ANALYSIS OF LABORATORY HOT-ROLLED S1100QL STEEL ...
902 Materiali in tehnologije / Materials and technology 54 (2020) 6, 901–908

ential sites for the nucleation and precipitation of
Nb(C,N) at lower supersaturation, as in the case of the
undeformed austenite
11
and with that expected change in
flow stress behaviour (sometimes written as k
f
), as it is
influenced by the introduction of the dislocation density
variation,  , inside the grains:
11,13
  = Gb (1)
Where  is a constant, G is the shear modulus and b
is the Burgers vector. The dislocation density is not so
trivial to measure,
14
so instead of observing metallurgical
phenomena associated with the change in dislocation
density, a recognized change of grain-boundary migra-
tion due to SIP is possible with various methods, but also
by measuring the change in the average flow stress.
15
1.3 Alloy design
Alloy design is crucial for the rolling schedule design
and final microstructure that is achieved after heat treat-
ment. The chemical composition of S1100QL with mi-
nor niobium and titanium additions is summarised in Ta-
ble 1.
In the case of niobium and titanium carbo-nitride
forming, they act as a retarding force on grain-boundary
migration. This is also important for limiting the grain
growth during cooling. In addition to Zenner pinning
with rather moderate niobium and titanium contents,
grain-boundary mobility can be influenced by the solute
drag of niobium, titanium and also molybdenum.
2,16
Tita-
nium and niobium additions are minor, which excludes
the possibility of coarse nitride and carbonitride forma-
tion already from the molten steel, see Figure 1 a and
Figure 4.
17
Niobium has an important effect on RST and RLT, ei-
ther in solution or even more so as strain-induced
Nb(C,N) precipitates.
11,15–21
As mentioned above, nio-
bium in solid solution segregates to austenite grain
boundaries, which causes drag and thus retards the
recrystallization by the solute-drag effect.
16,18
Nb(C,N)
precipitates are more effective in retarding recrystal-
lization. They nucleate first at the austenite grain bound-
aries (RLT) and later at the dislocations with matrix pre-
cipitation (RST), pinning them and thus retard
recrystallization by preventing grain-boundary migra-
tion.
19–21
The non-recrystallization temperature (T
nr
)i sn o r -
mally used to determine the critical temperature at which
strain accumulation occurs during hot rolling. T
nr
is nor-
mally assumed to be similar or close to the RLT, al-
though the exact position depends on other variables
such as strain, strain rate, experimental method used, etc.
This means that T
nr
can be identified between RLT and
RST, depending on the experimental method and the
given parameters. According to literature
5,10
, the MFS
range for Region III in standard HSLA grades during hot
rolling with typical per pass reductions and strain rates
for thin plate or strip rolling is between 150 MPa and
200 MPa. To correlate the measured MFS change in cor-
respondence to T
nr
, we have used the widely accepted
Boratto’s equation for HSLA grades:
22
T
nr
= 887 + 464C + 6445Nb – 644 Nb + 732V – 230
V+ 890Ti + 363Al – 357Si (2)
where all the elements are given in mass percent.
1.4 Rolling load model
According to Sims
23
, the calculation of rolling force
F can be easily calculated by:
F = k
fm
bl’
d
Q
f
(3)
where b is the plate width, l’
d
is the horizontal projec-
tion length of the contact arc between the squashed
roller and plate, Q
f
is the geometric factor and k
fm
is the
mean metal deformation resistance or the mean value of
constrained yield stress in the roll gap. It is also possible
to introduce the effect of the rolling force by tensile
stress (i.e., Steckel).
24
It is important that the equations
can also be used for strain-hardening materials, not only
for ideal plastic-rigid material. If plane deformation is
assumed, some precautions must be taken if the
width/thickness ratio per pass is less than 5. In this pa-
per an additional simplification has been made by set-
ting Q
f
to unity.
2 EXPERIMENTAL PART
The experimental S1100QL steel was produced in a
vacuum induction furnace and cast as a 12-kg ingot
(80 × 80) mm
2
. The ingot was heated to 1150 °C and
soaked for 2 h before hot forging into a (60 × 60) mm
2
billet. The billet was cut into two parts. Individual billets
intended for strip hot rolling were then reheated to
1200 °C for 30 min (a total 2.5 h for temperature homo-
geneity) to ensure that the niobium was completely in the
austenite solution in order to achieve an effective and
uniform recrystallization retardation by solute drag or
SIP during hot rolling. The solubility temperature for
Nb(C,N) was determined to be 1122 °C using the
Irvine
25
equation. An additional infrared camera Optris
PI 1ML with a working temperature range between
450 °C and 1800 °C was used to pre-define the tempera-
ture distribution measurement before and during rolling
at interpass times. The laboratory billet was then re-
J. FODER et al.: MEAN-FLOW-STRESS ANALYSIS OF LABORATORY HOT-ROLLED S1100QL STEEL ...
Materiali in tehnologije / Materials and technology 54 (2020) 6, 901–908 903
Table 1: Chemical composition of S1100QL billet, in mass fractions, (w/%)
CS iM nC rN iM oA lT iN bBNF e
0.18 0.23 0.87 0.52 1.3 0.41 0.06 0.013 0.014 0.0011 0.0031 bal.

heated as described above and hot rolled in a predefined
11 passes to a final thickness of 12 mm, using a single
stand mill with subsequent continuous air-cooling. The
soaking temperature and the number of passes with the
specified per-pass reduction made it possible to achieve a
low-temperature finish. Hot rolling was performed with-
out interruption or time delay between the roughing and
finishing to observe the MFS behaviour in the region of
the RLT and/or RST (defined as T
nr
in this paper) for the
given chemical composition and related rolling parame-
ters. The rolling loads were continuously measured for
indirect MFS evaluation by stress measurements on the
rolling mill. The last step in the production process was a
laboratory heat treatment by re-austenitization of the
samples in a temperature-controlled furnace using addi-
tional thermocouples (Type-K) for the samples’ austenit-
ization control. Thus, after austenitization at 900 °C and
a holding time of 20 min, the steel was first water
quenched, followed by low-temperature tempering at
200 °C and a holding time of 60 min.
The general chemical composition was determined
by optical emission spectroscopy (OES), ARL MA-310.
Carbon and sulphur were determined by the combustion
method using a LECO CS-600 and nitrogen using a
LECO TC-500, Table 1. The base composition also cor-
responds to other grades with a lower yield strength ac-
cording to standardized steel qualities defined in
reference
26
.
The microstructure and PAGS evolution of S1100QL
in the as-cast, hot-forged, hot-rolled and heat-treated
conditions was characterised, using scanning electron
microscope with field emission gun (FE-SEM), Zeiss
Supra VP55. Samples were prepared following a stan-
dard metallographic procedure of grinding, polishing and
etching using 2 vol. % Nital solution. PAGS and distribu-
tion were determined using imageJ software by analys-
ing an average of 400 grains.
Thermodynamic prediction, using Thermo-Calc
2020a, was also conducted using TCFE7 and TCFE9 da-
tabases for the estimation of the thermodynamically sta-
ble characteristic temperatures for carbo-nitride precipi-
tation. A simplified isopleth equilibrium Fe-C phase
diagram (Figure 4) was calculated to provide a better vi-
sual representation of the temperature dependence of the
titanium- and niobium-rich precipitates with respect to
the carbon variation.
To evaluate the mechanical properties in the final
tempered condition, longitudinal tensile specimens were
prepared and tests were carried out according to EN ISO
6892-1:2017 B60 using Zwick/Roell Z600.
27
Standard
V-notch specimens for the Charpy pendulum impact test
were machined according to EN ISO 148-1:2016.
28
Im-
pact toughness should be higher than 27 J for transverse
test pieces, tested at –40 °C, as specified in the standard
EN 10025-6:2019 for similar HSLA grades.
26
J. FODER et al.: MEAN-FLOW-STRESS ANALYSIS OF LABORATORY HOT-ROLLED S1100QL STEEL ...
904 Materiali in tehnologije / Materials and technology 54 (2020) 6, 901–908
Figure 1: Secondary electron image of microstructure of: a) as-cast, b) hot-forged, c) hot-rolled and d) QT state of S1100QL steel

3 RESULTS AND DISCUSSION
3.1 Temperature evolution during hot rolling
In order to determine a semi-empirical equation for
continuous cooling during the hot rolling of S1100QL,
several laboratory rolled heats at the same re-heating
temperature (1200 °C) and final thickness (12 mm) were
required. No water de-scaling was used. Temperature
evolution during the hot rolling of S1100QL can be cal-
culated using the obtained semi-empirical equation:
T = 1194.6 exp(–0.003t) (4)
where T is the rolling temperature in °C and t is the time
in seconds elapsed since the reheated billets were taken
from reheating furnace. The temperature profile is
shown in Table 2. From the rolling schedule, 7 roughing
passes lie between 1158 °C and 971 °C followed by 4
finishing passes down to 882 °C. Achieved grains size
and an example of the final as-rolled microstructure af-
ter cooling from 882 °C in air is shown in Figure 1c and
Figure 2.
3.2 Microstructure evolution
Microstructures of as-cast, hot-forged, hot-rolled and
heat-treated (QT) conditions are shown in Figure 1a to
1d, respectively. In the as-cast to hot-rolled condition, a
mixed martensitic-bainitic microstructure can be ob-
served. Martensite occurs predominantly in the segre-
gated regions with solute enrichment. The mixed
microstructure is related to relatively slow cooling rates
achieved under natural cooling in air. The grain, except
in the case of the as-rolled state, is equiaxial. In Fig-
ure 1c the grains are elongated or "pancaked", indicating
that the final rolling temperature (FRT) corresponds to
the desired low FRT. The average PAGS and the distribu-
tion of 400 grains evaluated for each condition are pre-
sented in Figure 2.
By combining the PAGS distribution for a separate
process step, the average grain refinement from coarse
320 μm to rather fine 18 μm from as-cast to heat-treated
condition is obvious. It is interesting to note that by con-
tinuous hot rolling the maximum PAGS is still quite
coarse in the range of 130 μm. Similar results were ob-
tained after forging with a distribution factor of G
avg
/G
max
= 0.28 for rolled condition, where G stands for grain size
in μm. Comparing the factor after re-austenitization with
G
avg
/G
max
= 0.38, it becomes clear that re-austenitization
of the forged-rolled condition has considerably improved
the material isotropy by reducing the maximum grain
size and narrowing the grain size distribution. This ex-
plains why the toughness is sometimes improved by
off-line heat treatment compared to on-line, with respect
to the achieved austenite conditioning before on-line
quenching. Nevertheless, isotropy is also improved when
preferential grain orientation is changed from pan-
cake-shaped to equiaxed by re-austenitization.
The results of the PAGS for similar chemical compo-
sitions but different processing routes, according to liter-
ature, were from 30 μm to 7 μm.
4,29
PAGS in the
hot-rolled condition is comparable to PAGS that Muckle-
roy et al.
4
reported for the direct quenched, controlled
rolled condition with a 40 % reduction below T
nr
, for
steel of similar composition. They emphasised that
martensitic blocks, as discussed before, determine the ef-
fective grain size of martensitic steels and not PAGS,
which is in accordance with earlier reports.
30,31
3.3 MFS analysis
The results of the MFS analysis of laboratory
hot-rolled S1100QL steel are presented in Figure 3.
Rolling schedules and obtained MFS are shown in Ta-
ble 2. Rolling consisted of 7 roughing passes (R1–R7),
followed by 4 finishing passes (F1–F4). As can be seen
in Figure 3, the MFS increases continuously up to R7 in
a moderate manner, which is a result of a temperature
drop and SRX. Limited secondary recrystallization or
so-called grain growth appears in the case of R2, as esti-
mated from the MFS curve. Rolling in the SRX region
leads to grain refinement and improving the through-sec-
J. FODER et al.: MEAN-FLOW-STRESS ANALYSIS OF LABORATORY HOT-ROLLED S1100QL STEEL ...
Materiali in tehnologije / Materials and technology 54 (2020) 6, 901–908 905
Figure 3: MFS analysis for laboratory hot rolling of S1100QL steel.
Regions of different recrystallization behaviour are marked in accor-
dance with the literature
5,10
Figure 2: Prior austenite grain size distribution and average grain size
for as-cast, hot-forged, hot-rolled and QT state of S1100QL steel

tion grain size distribution by successive reductions with
limited grain refinement. From R7 to F4 a distinct and
almost linear MFS deflection is detected. The slope
change was taken as the critical temperature defined as
T
nr
and determined as an average between R7 to F1 due
to the limited temperature raster taken by an individual
test. If the limits are taken according to references
5,10
the
MFS curve is interpreted as R1–R4 belong to Region I,
which means that SRX refines PAGS during successive
passes, given that the interpass time and deformation are
sufficient. Nevertheless, according to the MFS curve,
R1–R6 are taken as part of the roughing phase and
F1–F4 is the finishing phase with R7–F1 being the pass
in the region where Type III is expected and should be
avoided. This could explain the rather coarse maximum
PAGS found in the as-rolled state, knowing that during
precipitation of Nb(C,N) niobium depletion appears in
the matrix, affecting the grain-boundary mobility and en-
hancing the recrystallisation, thus promoting grain
growth.
32
If we apply temperatures calculated from equation
(4), strain accumulation between R7 and F1 occurs at an
average temperature of 958 °C. T
nr
, calculated according
to the Boratto’s equation, is 936 °C, i.e., 22 °C lower
than experimentally predicted, which seems reasonable.
The results of the MFS analysis, i.e., the change in
the slope after R7 to F4, can be associated with SEM mi-
crographs of the hot-rolled condition, where elongated
prior austenite grains can be seen.
In the industrial practice, MFS analysis can easily be
used to optimize rolling schedules based on the material
response during hot rolling, even in real-case rolling pro-
cesses. Due to the different starting conditions, such as
slab geometry, size and reheating time and temperature,
MFS values for industrial rolling are typically lower than
the values determined during laboratory rolling. Also,
RLT and RST, if assumed to be 150 MPa and 200 MPa,
are considered as orientational values only, but according
to the available data, this approximation still allows a
fairly good estimation of where Region I and Region II
are expected to be found.
5,10
Table 2: Per-pass reduction, calculated temperature, interpass time
and MFS for laboratory hot rolling of S1100QL steel
Pass no.
Reduc-
tion/%
Tempera-
ture/°C
Interpass
time/s
MFS/MPa
R1 5,0 1158 / 127
R2 11,4 1128 10 114
R3 13,6 1089 10 127
R4 15,8 1056 10 135
R5 17,4 1022 10 161
R6 16,8 993 9 174
R7 18,7 971 8 175
F1 17,1 945 8 231
F2 17,6 922 8 257
F3 14,3 903 8 309
F4 12,5 882 8 346
3.4 Thermodynamic analysis
The predicted equilibrium phase diagram Fe-C has
been simplified by excluding boron, sulphur and phos-
phorous for easier visualisation. The minor addition of
titanium shows no risk of primary (and coarse) TiN for-
mation under the equilibrium condition of solidification
due to the stability of nitrides below the solidus tempera-
ture (presented as Liquid, see Figure 4) for the given
chemical composition. The temperature stability of TiN
in the carbon region of interest is practically unchanged
due to the constant nitrogen content and covers the entire
temperature region for austenite soaking and multi-pass
hot-rolling for grain size control, including cooling after
rolling. The solidification path for S1100QL is presented
with a vertical line at 0.18 mass percent carbon. The
phase field stability of TiN is presented by the prototype
FCC_A1#3. The predicted isopleth diagram also shows
temperature-carbon concentration relation of Nb(CN)
precipitation based on the equilibrium thermodynamics
of a multicomponent system according to the chemical
composition in Table 1.
The precipitation of Nb(C,N) in undeformed austen-
ite is presented with the prototype of FCC_A1#2. With
constant niobium and nitrogen concentration and in-
creasing carbon, the temperature for the homogeneous
nucleation of Nb(C,N) is increasing. This goes well with
the nucleation theory, which refers to the solubility prod-
ucts of niobium, carbon and nitrogen in austenite. Ac-
cording to the prediction, the stability of Nb(C,N) with
homogenic precipitation starts already under 1092 °C or
1104 °C using the TCFE7 or TCFE9 database, respec-
tively. When using a thermodynamic approach for a
multicomponent system, the temperatures are somewhat
lower when calculated by the reference
25
, which only
takes into account niobium, carbon and nitrogen, see Ta-
ble 3. Taking into consideration the thermo-mechanical
processing and dislocation density generated for poten-
tial nucleation sites (based on precipitation kinetics and
not included into the equilibrium phase stability calcula-
J. FODER et al.: MEAN-FLOW-STRESS ANALYSIS OF LABORATORY HOT-ROLLED S1100QL STEEL ...
906 Materiali in tehnologije / Materials and technology 54 (2020) 6, 901–908
Figure 4: Isopletic equilibrium phase diagram for Fe-C for S1100QL

tion) the total amount of precipitates at a given tempera-
ture can vary significantly due to heterogenic nucleation.
In this case it is also interesting to note that the experi-
mentally determined T
nr
and the calculated one according
to Boratto
22
are in a rather good agreement and both are
lower than the predicted homogenic nucleation start for
the precipitation of niobium-rich carbo-nitrides by
Thermo-Calc. The determined starts of the homogeneous
(equilibrium) and experimental precipitation of Nb(CN)
are gathered in Table 3. The variation between the ther-
modynamic calculations and the experimentally deter-
mined T
nr
is expected, because the MFS change is related
to the actual volume fraction and size of the precipitates
formed by rolling, in order to effectively delay the
recrystallisation kinetics and therefore to the rolling pa-
rameters.
Table 3: Calculated temperatures for Nb(C,N) precipitation and mea-
sured T
nr
Thermo-Calc – Nb(C,N) precipitation
TCFE7 1092,9 °C
TCFE9 1104,4 °C
Nb(C,N) precipitation according to
Irvine
25 1122 °C
T
nr
according to Boratto
22
936 °C
MFS exp. – T
nr
958 °C
3.5 Mechanical properties
The mechanical properties of S1100QL steel in the
QT condition are shown in Table 4. The yield strength
after quenching and low-temperature tempering exceeds
1100 MPa, indicating that a pre-rolling strategy with ap-
propriate pass numbers and re-heating temperatures be-
fore hot rolling provides a good starting point for the QT
process with off-line heat treatment. The ductility, ex-
pressed in A
5,65
, also exceeds 10 %, which is required for
this steel grade. Overall, the needed tensile properties are
met. Charpy impact toughness KV
2
at –40 °C is 49 J,
which is above 27 J, prescribed in the standard.
26
Table 4: Mechanical properties of S1100QL steel in QT state
R
p0,2
/MPa R
m
/MPa A
5,65
/% KV
2
/J
min. 1100 1250–1500 min. 10 min. 27
1160 1377 13.5 49
3 CONCLUSIONS
Laboratory hot rolling of S1100QL plate was suc-
cessfully performed, and MFS analysis as an indirect
method shows that the position of T
nr
is determined in
the range of 150 MPa and 200 MPa. This agrees well for
similar grades as the limitation for RLT and RST. A rela-
tively good agreement was obtained by comparing the
measured T
nr
and predicted T
nr
according to Boratto, with
958 °C and 936 °C, respectively. The tested rolling
schedule reveals the presence of a locally coarse grain
size and indicates that the optimization of rolling sched-
ules is necessary for laboratory single stand rolling for
the desired overall mechanical properties based on the
obtained MFS curve presented in this paper. Significant
improvements in the material isotropy were achieved by
off-line heat treatment by reducing the maximum grain
size and narrowing the grain size distribution.
The data obtained with the given parameters can be
used for rolling optimization with respect to the actual
rolling mill capability in terms of the maximum allow-
able force and torque as also for achieving the appropri-
ate dimensional and flatness tolerance.
Acknowledgments
This research was made as a part of ^MRLJ research
project co-financed by the Republic of Slovenia and the
European Union under the European Regional Develop-
ment Fund.
The authors also want to acknowledge Douglas
Stalheim, from DGS Metallurgical Solutions, the consul-
tant of CBMM, for overall cooperation and the experi-
ence exchange on the production of HSLA grades.
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