Impact of two-phase region rolling on the microstructure and properties distribution in heavy gauge structural steel plate (INCROHSS) : final report.
Date: 2021-02-22
Institutions: Directorate-General for Research and Innovation
Subjects: applied research, innovation, metalworking, physical process, research report, steel, technological process

Impact of two-phase region 
rolling on the 
microstructure and 
properties distribution in 
heavy gauge structural 
steel plate 
(INCROHSS)

EUR [25330] EN 

EUR 30539 EN 

 
 
 
 
 
Impact of two-phase region rolling on the microstructure and properties distribution in 
heavy gauge structural steel plate (INCROHSS) 
European Commission 
Directorate-General for Research and Innovation 
Directorate D – Clean Planet 
Unit D.3 — Low Emission Future Industries 
Contact   Andrea Gentili 
E-mail   RTD-STEEL-COAL@ec.europa.eu 

RTD-PUBLICATIONS@ec.europa.eu 

European Commission 
B-1049 Brussels 

Manuscript completed in 2019. 

This  document  has  been  prepared  for  the  European  Commission  however  it  reflects  the  views  only  of  the  authors,  and  the 
Commission cannot be held responsible for any use which may be made of the information contained therein. 

More information on the European Union is available on the internet (http://europa.eu). 

Luxembourg: Publications Office of the European Union, 2020 

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ISBN 978-92-76-27847-4 

ISSN 1831-9424 

doi: 10.2777/639791 

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EUROPEAN COMMISSION 

Research Fund for Coal and Steel 

Impact of two-phase region rolling 
on the microstructure and 
properties distribution in heavy 
gauge structural steel plate 

(INCROHSS) 

T. Nguyen-Minh, M. Caruso, I. Tolleneer 
Centre de Recherches Metallurgiques (CRM) 
Technologiepark 903C B-9052 Zwijnaarde (Belgium) 

U. Mayo, N. Isasti, P. Uranga  
Asociación Centro Tecnológico CEIT-IK4 
Paseo Manuel De Lardizabal 15 ES-20018 San Sebastian (Spain) 

U. Lorenz, L. Duprez 
Onderzoekscentrum voor Aanwending van Staal (OCAS) 
President JF Kennedylaan 3 B-9060 Zelzate (Belgium) 

F. H. Akbary, R. Petrov, J. Sietsma 
Technische Universiteit Delft (TUDT) 
Stevinweg 1 NL-2628 Delft (Netherlands) 

Grant Agreement RFSR-CT-2015-000014 

01/07/2015 – 31/12/2018 

Final Report 

2020                      Directorate-General for Research and Innovation                     EUR 30539 EN 

 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
Table of contents 

Final summary .................................................................................................................... 5 

WP1: Coordination and reporting ........................................................................................... 9 

TASK 1.1 Technical project management ............................................................................ 9 

TASK 1.2 Full reporting ..................................................................................................... 9 

TASK 1.3 Dissemination and promotion of results/knowledge gained ...................................... 9 

WP2: Selection, production & determination of thermomechanical characteristic temperatures of 
material ............................................................................................................................ 12 

TASK 2.1 Chemistry selection and material workflow ........................................................... 12 

TASK 2.2 Production & collection of test material ................................................................ 12 

TASK 2.3 Determination of the critical temperatures ........................................................... 14 

WP3: Full mapping of the effect of the α/γ  balance for the mechanical properties of the reference 
alloy ................................................................................................................................. 23 

TASK 3.1 Laboratory rolling trials with single pass intercritical deformation ............................ 23 

TASK 3.2 Microstructural characterization .......................................................................... 25 

TASK 3.3 Mechanical properties ........................................................................................ 30 

TASK 3.4 Microstructural evolutions and correlation ............................................................ 33 

WP4: Interaction between intercritical rolling & microstructural evolution at high temperature: 
strain partitioning ............................................................................................................... 37 

TASK 4.1 MFS mapping of composite and “single phase” materials ........................................ 37 

TASK 4.2 High temperature microstructural evolution in the composite materials .................... 41 

TASK 4.3 Strain partitioning: law of mixtures ..................................................................... 58 

WP5: Effect of strain path in the intercritical temperature region on microstructure and mechanical 
properties ......................................................................................................................... 67 

TASK 5.1 Effect of strain .................................................................................................. 67 

TASK 5.2 Effect of strain distribution ................................................................................. 68 

TASK 5.3 Microstructural characterization .......................................................................... 71 

TASK 5.4 Characterization of mechanical properties ............................................................ 73 

TASK 5.5 Completion of the strength-toughness map versus intercritical deformation ............. 77 

WP6: Impact of chemical composition ................................................................................... 79 

TASK 6.1 Laboratory rolling trials ...................................................................................... 79 

TASK 6.2 Microstructural characterization .......................................................................... 81 

TASK 6.3 Determination of the mechanical properties .......................................................... 89 

TASK 6.4 Quantification of the impact of C, Nb and V on the microstructural and property 
relationship .................................................................................................................... 93 

WP7: Definition of the optimum processing conditions – Link between thin gauge homogeneous 
material behavior and real heavy gauge process .................................................................... 99 

TASK 7.1 Micromechanical modelling ................................................................................. 99 

TASK 7.2 Effect of temperature gradient .......................................................................... 103 

TASK 7.3 Microstructural characterization over thickness ................................................... 105 

TASK 7.4 Mechanical characterization over thickness and the correlation with microstructure . 112 

3 

 
 
 
TASK 7.5 Definition of the guidelines ............................................................................... 114 

Conclusions .................................................................................................................. 118 

Exploitation and impact of the research results ................................................................. 122 

Publications .............................................................................................................. 122 

Conference presentations ........................................................................................... 122 

List of figures ................................................................................................................... 124 

List of tables .................................................................................................................... 130 

List of acronyms and abbreviations ..................................................................................... 132 

List of references .............................................................................................................. 134 

Appendices ...................................................................................................................... 136 

4 

 
 
 
 
 
 
Final summary 

Heavy gauge line-pipe and structural steel plates are often rolled in the two-phase region for strength 
reasons. However, strength and toughness show opposite trends and the exact effect of each rolling 
process  parameter  remains  unclear.  A  stable  process  window  can  only  be  achieved  by  a  more 
profound  understanding  of  the  microstructure  development  during  the  intercritical  rolling  and  its 
relationship with final microstructures and mechanical properties. By means of recently developed 
microstructure  investigation  techniques  and  modeling,  the  relationship  between  the  temperature 
gradient, α-γ phase balance at high temperature, strain partitioning between phases and subsequent 
transformation was studied in detail to allow for wider process windows. 

To answer the research question on “the impact of two-phase region rolling on the microstructure 
and properties distribution in heavy gauge structural steel plate”, this project is divided into work 
tasks with several key objectives as given below: 

 

Identifying the relation between strain partitioning, microstructure evolution, and mechanical 
properties. 

 

  Determining  the  effect  of  process  parameters  on  the  high  temperature  microstructure 
evolution,  in  particular  the  strain  partitioning  between  phases  and  its  effect  on  final 
microstructure. 
Establishing the impact of strain, strain rate and additional micro-alloying elements Nb and 
V on the intercritical deformation. 
Investigating the impact of C during strain partitioning and how the C itself behaves during 
the complex transformations in the two-phase region. 

 

  Correlating the α-γ phase balance and therefore the strain partitioning to the thermal profile 

over thickness and width of the plates. 

  Describing  the  analysis  techniques  able  to  distinguish  the  two  differently  formed  ferrite 

colonies. 

  Clarifying  the  capabilities  of  recently  developed  characterization  techniques  and  data 

modeling to reconstruct the strain partitioning behavior. 

WP1 - Coordination & Reporting 

The project was carried out for 42 months (01/7/2015 – 31/12/2018). The consortium is formed by 
4 partners from 3 different European countries, consisting of one steel company (OCAS, R&D center 
of ArcelorMittal), two research institutes (CRM and CEIT) and one university (TU Delft). CRM is the 
coordinator of the project. Periodical meetings and conference calls were regularly organized to track 
the  technical  progress  of  all  research  activities.  Periodical  reports  were  carefully  prepared  and 
submitted to the Commission on-time. Results of the project were appropriately disseminated. One 
journal article and seven oral presentations at international conferences are the output of this project. 
On 27 September 2018, a seminar on “Intercritical rolling of high strength steel” was organized by 
CEIT at the MSE Congress 2018, Darmstadt, Germany to disseminate and promote results of this 
project. 

WP2 - Material selection, production and determination of characteristic temperature 

At the beginning of this project, all partners discussed and decided to choose three basic C-Mn steels 
as studied materials. The first grade of steel, called the CMn-0.06%C steel, contains 0.06wt.% C and 
1.5wt.% Mn. The second grade, CMn-0.12%C steel, has the C concentration of 0.12wt.%, while the 
Mn content is similar to the first one. The last grade, NbV steel, has the same C and Mn content as 
the first one. Additionally, the NbV steel contains 0.05wt.%Nb and 0.03wt.%V. All of these materials 
are lab produced in vacuum induction furnaces at CRM and OCAS. Although the material production 
was unexpectedly prolonged because of the large amount of materials, no delay were observed from 
other  activities  of  the  project.  These  materials,  afterward,  were  distributed  to  all  partners  for 
experiments of WP3, 4, 5 and 6. 

Critical  temperatures  of  all  three  studied  materials  (CMn-0.06%C,  CMn-0.12%C  and  NbV  steels) 
were intensively investigated by CEIT, CRM and OCAS. These temperatures were determined by both 

5 

 
 
 
dilatometry  and  hot  torsion  experiments.  Impacts  of  different  austenite  microstructures,  i.e. 
deformed or recrystallized, were also considered in this task. 

Besides the three main studied materials, a set of 10 different “single phase” materials corresponding 
to  intercritical  ferrite  and  austenite  compositions  was  also  casted  at  OCAS.  These  “single  phase” 
materials are mainly used for works of the WP4, in which mechanical behaviors of individual phases 
at intercritical rolling condition could be specifically studied. 

WP3 - Mapping of the γ/α balance on the mechanical properties of the CMn-0.06%C steel 

To produce samples of single pass intercritical deformation, laboratory rolling trials were performed 
at  OCAS  and  CRM.  Based  on  the  information  from  task  2.3,  three  intercritical  temperatures 
corresponding to 25%, 50% and 75% ferrite volume fraction were chosen. The intercritical rolling 
reduction was kept constant at 25% and the final thickness of rolled plates was set as 12 mm for the 
Charpy toughness samples. During these rolling trials, OCAS focused on direct quench and air cooled 
samples. On the other hand, CRM concentrated on producing samples of accelerated cooling. 

All single pass intercritical rolled samples were subjected to microstructure and mechanical property 
characterization.  Specifically,  optical  microscopy  and  mechanical  property  measurements  were 
carried out at OCAS and CRM. Additionally, the same set of samples were also delivered to TU Delft 
for microstructure and texture analysis by EBSD technique. All results collected by different partners 
were  finally  combined  and  correlated  as  relationships  among  strength,  toughness,  microstructure 
and rolling temperature. Mechanical properties of samples produced by laboratory intercritical rolling 
are in line with the ones produced by  industrial practice. The YS increase is of the same order of 
magnitude as observed industrially. With the reduction in FRT, the larger increase in YS than in TS 
results in a decrease in TS/YS ratio. The work-hardening in samples at room temperature therefore 
is reduced by lower intercritical rolling temperature. However, the strength of hot rolling plates is 
inversely  proportional  to  the  toughness,  especially  when  accelerated  cooling  or  direct  quench  is 
applied.  Crystallographic  texture  is  not  the  main  factor  to  define  absolute  values  of  toughness 
properties, because very different toughness values were measured from samples of quite similar 
crystallographic  textures.  Nevertheless,  because  of  the  crystallographic  textures,  toughness 
properties measured along different (longitude and transversal) directions of the same samples are 
not  similar.  On  the  other  hand,  a  clear  relationship  between  DBTT  and  the  ferrite  grain  size  was 
established. Crystal regions enclosed by boundaries having misorientation angles larger than 15° is 
found to be the effective grains governing ductility of material. 

WP4 - Strain partitioning between ferrite and austenite during intercritical deformation 

To  investigate  the  mechanical  behavior  of  intercritical  ferrite  and  austenite  separately,  CEIT 
performed hot torsion test on single  austenite and ferrite  materials.  These  single phase  materials 
have the same chemical compositions as intercritical ferrite and austenite of the CMn-0.06%C steel 
at three chosen temperatures. These intercritical temperatures are corresponding to 25%, 50% and 
75% of ferrite in the CMn-0.06%C steel. The approach of studying each component of the intercritical 
deformation  as “single phase  material”,  however,  was not entirely successful.  Because  the  γ to α 
phase  transformation  could  occur  well  below  the  equilibrium  temperature,  such  “single  phase 
materials”, in general, do not exist at practical hot rolling condition. Furthermore, intercritical ferrite 
and  austenite  could  transform  during  deformation.  Therefore,  the  plastic  flow  behavior  of  pure 
austenite and pure ferrite outside of the intercritical region was investigated. Flow behaviors of ferrite 
and austenite in the intercritical region, afterward, were extrapolated. This approach finally allowed 
the establishment of a strain partitioning law between γ and α phases during intercritical deformation. 

Although  a  more  complex  relationship  between  γ  and  α  phases  during  intercritical  deformation  is 
generally desired, practical limitation of intercritical rolling to rather low strains has suggested that 
a  composite model  under  the  iso-strain assumption  would give reasonable  predictions  of the flow 
behavior during intercritical deformation. 

Furthermore,  regarding  the  microstructural  evolution  during  intercritical  deformation,  a  basic 
phenomenological study was developed understanding the impact that process parameters such as 
phase fractions, holding times, effect of deformation have on the intercritical deformation and the 
final microstructure. A quantitative procedure of microstructural characterization was developed in 
6 

 
 
the frame of this project using EBSD technique. The quantification was made upon GOS parameter 
and  grain  boundary  maps,  and  a  successful  differentiation  between  deformed  and  non-deformed 
microstructures was achieved. 

WP5 - Effect of strain magnitude and strain path in the intercritical region 

To evaluate the effect of strain in the intercritical region, rolling trials as similar to the ones in WP3 
were performed at CRM. Laboratory rolling trials with single pass reductions of 10% and 40% in the 
intercritical region were implemented. Rolling trials of 40% thickness reduction with 2  intercritical 
rolling passes were initially proposed. However, intercritical rolling of high thickness reduction, for 
instance 20% and 40% per pass, is not technically feasible, particularly at low rolling temperatures. 
Therefore, intercritical rolling trials of  1, 2  and  3 rolling passes with 10% thickness reduction per 
pass were performed. Microstructures and mechanical properties of these samples, afterward, were 
measured and analyzed. 

For the study on effect of strain magnitude, microstructures and mechanical properties of the two 
intercritical rolled samples with 25% and 40% thickness reduction in single rolling passes at 745oC 
were compared. As the thickness reduction increases from 25% to 40%, ferrite grains in the CMn-
0.06%C  steel  become  more  elongated.  Sub-grain  structures  in  deformed  ferrite  grains  are  also 
visible, particularly in the sample of 40% thickness reduction. The yield strength and tensile strength 
of  the  sample  with  40%  thickness  reduction  are  both  higher  than  those  of  the  sample  with  25% 
thickness reduction. On the contrary, upper shelf energy measured by Charpy V-notched test of the 
later sample is higher than the corresponding value of the former sample. 

Strain  distribution  and  strain  accumulation  during  intercritical  rolling  although  expected  to  have 
significant  effect  on  microstructures  and  mechanical  properties  of  the  CMn-0.06%C  steel  did  not 
bring any special result. Not many different features  in microstructures of the  intercritically rolled 
samples  were  observed.  Ferrite grains in these  microstructures are  all  in the  range of  12-17 μm. 
Intercritical rolling of multi deformation passes, however, produces slightly finer microstructures than 
the single pass intercritical rolling. In-grain orientation gradients and low angle grain boundaries are 
visible  in  ferrite  grains  of  the  samples  deformed  by  multi  intercritical  rolling  passes  and  at  lower 
rolling temperature. The increase in number of intercritical rolling passes from 1 to 3 passes brings 
50 MPa to the tensile strength of the CMn-0.06%C material. However, the total elongation of this 
material reduces about 2%, and the upper shelf energy (USE) measured the Charpy V-notched tests 
decrease 70 J. 

WP6 - Impact of chemical composition 

To  understand  the  effects  of  chemical  composition  (C  and  micro-alloying  elements)  on 
microstructures  and  mechanical  properties  of  intercritically  rolled  samples,  OCAS  performed  lab 
rolling trials on CMn-0.12%C and NbV steels. The systematic approach as applied in WP3 an WP5 
was used. The number of rolling trials in this WP, however, were reduced as compared to the previous 
ones. Only two intercritical temperatures,  corresponding  to low and high ferrite  volume  fractions, 
were  selected  for  further  studies.  These  temperatures  of  the  CMn-0.12%C  and  NbV  steels  were 
selected based on the “relaxation like” torsion experiment, carried on in WP2. Besides the two cooling 
strategies  of  air  cooling  and  accelerated  cooling,  direct  quench  was  also  applied  in  this  WP  to 
investigate the state of materials before intercritical rolling. 

Microstructure  investigation  on  samples  of  hot  rolling  and  then  cooling  to  room  temperature  by 
various cooling rates shows different behaviors between CMn-0.12%C and NbV steels. For the NbV 
steel,  ferrite  volume  fractions  at  770°C  and  750°C  are  mostly  similar,  irrespective  to  cooling 
conditions. Just by intercritical rolling and then directly quenching, different ferrite volume fractions 
were  obtained  between  the  two  intercritical  temperatures.  For  CMn-0.12%C  steel,  equal  volume 
fraction of ferrite was obtained in samples of intercritical rolling at 750°C and 700°C and then air 
cooling to room temperature. To evaluate the effect of microalloying elements on precipitation in the 
NbV  steel,  TEM  analyses  were  carried  out  with  carbon  extraction  replicas  by  CEIT.  When  the 
transformation occurs from recrystallized austenite, a high density of fine precipitates  is obtained. 
These precipitates are considered to be formed in ferrite during or after the intercritical deformation. 

7 

 
 
In the sample corresponding to deformed austenite, strain induced precipitation in austenite are also 
observed. 

Irrespective  to  chemical  composition,  with  the  increase  in  cooling  rate  from  air  cooling  to  direct 
quench, the strength of intercritically rolled samples increases. Additionally, the DBTT increases, i.e. 
the toughness of materials decreases. These evolutions of mechanical properties  were linked to the 
volume  fraction  and  morphology  of  transformation  products  from  intercritical  austenite  (i.e. 
martensite,  bainite  and  pearlite).  With  lower  intercritical  temperature,  corresponding  to  higher 
amount of ferrite, the strength of steels slightly increases because of the higher strength of the 2nd 
phase causing by C enrichment. 

With  respect  to  composition,  it  was  observed  that  the  microalloying  of  Nb  and  V  increases  the 
strength  of  NbV  steel  at  least  100MPa,  thanks  to  the  precipitated  hardening,  independent  of  the 
cooling  rate.  The  addition  of  C  in  the  CMn-0.12%C  steel  increases  the  strength  significantly, 
particularly in direct quench samples. For air cooling samples, the strength increases by the larger 
fraction of pearlite microstructure in the CMn-0.12%C steel than in the CMn-0.06%C steel. However, 
the  strength  increasing  by  pearlite  is  lower  than  by  Nb-V  microalloying.  In  addition,  Nb  and  V 
significantly  reduce  ductile  to  brittle transition  temperature.  Results  of  mechanical  measurements 
indicate that toughness is less dependent on the cooling rate. The addition of C is deteriorating the 
toughness,  i.e.  ductile  to  brittle  temperature  increases  while  the  upper  shelf  energy  reduces. 
Toughness in CMn steels with the absence of Nb-V microalloying elements improves mainly by the 
reduction  of  the  cooling  rate.  Transformation  products  of  intercritical  austenite  changes  from 
martensite into ferrite and pearlite microstructure. However, appropriate combination of intercritical 
rolling and cooling indicates would bring good effects on toughness. Adding Nb and V to steels brings 
significant impacts on mechanical properties. Unexpectedly, the CMn-0.12%C steel after intercritical 
rolling  and  air  cooling  has  better  ductile  to  brittle  transition  temperatures  than  the  CMn-0.06%C 
steel. 

WP7 - Definition of the optimum processing conditions 

Objective of the WP7 is to link all research data collected in previous WPs, build a micromechanical 
model for microstructure and texture evolution during intercritical rolling and finally define guidelines 
for  future  steel  production.  From  previous  studies,  it  has  been  known  that  the  evolution  of 
microstructure and texture during intercritical rolling is not a simple phenomenon. Intercritical rolling 
is not just a hot rolling process for materials of dual phase ferrite-austenite microstructures. More 
than frequent, steels are rolled at high temperature, in the fully austenitic condition. Depending on 
the thickness reduction and the rolling temperature, after the rolling, austenite could be presence as 
deformed or recrystallized state. Even if the rolling temperature for fully  austenitic rolling is below 
the non-recrystallized temperature (Tnr), the holding or slow cooling from high temperatures could 
promote  recovery  and/or  static  recrystallization  of  austenite.  Apparently,  for  intercritical  rolling, 
when  the  steel  is  cooled  down  to  a  specific  intercritical  temperature  after  fully  austenitic  rolling, 
transformation of austenite to ferrite will occur. The rolling at intercritical temperatures, therefore, 
will  lead  to  deformation  of  both  ferrite  and  austenite.  Finally,  the  cooling  from  intercritical 
temperature to room temperature will define the final microstructure and texture of intercritical rolled 
samples.  Along  the  whole  complicate  process,  there  are  three  important  physical  phenomena 
occurring  in  the  materials:  plastic  deformation,  recrystallization  and  phase  transformation. 
Simulating all phenomena during this complete process is beyond the scope of this study. Therefore, 
the visco-plastic self-consistence  (VPSC)  crystal  plasticity model  has been chosen to simulate the 
plastic deformation and strain partitioning between ferrite and austenite during intercritical rolling. 

8 

 
 
 
 
 
WP1: Coordination and reporting 
Objectives of the WP1 are: 

Ensure that the project is carried out according to the agreed time and cost schedules. 

- 
-  Organization of project meetings and provision of annual, mid-term and final reports. 
-  Dissemination and promotion of results/knowledge gained through publications in 

international journals and conferences. 

TASK 1.1 Technical project management 

After the kick-off of the project, (CRM Gent on 16/09/2015), physical meetings were held on a regular 
basis  of  semi-annual  (21/04/16  (Delft),  15/09/16  (Delft),  16/02/17  (Gent),  19/09/17  (Gent), 
16/03/18 (San Sebastian)). The last physical meeting was held on 20/09/18  at Darmstadt, on the 
occasion of the public dissemination of the project at the MSE Congress 2018. The main roadblocks 
were regularly discussed and shared among the partners. To facilitate document exchanges and allow 
all partners to have access to the latest research results, a Dropbox folder has been created to share 
information  in  a  secure  way:  login  and  password  are  required,  and  sharing  is  only  possible  upon 
invite of the coordinator. 

TASK 1.2 Full reporting 

All annual, the mid-term and the draft final reports of the project have been delivered as scheduled. 
Technical results and achievements of the project have been regularly reported at the TGS6 meetings 
by the project coordinator. 

TASK 1.3 Dissemination and promotion of results/knowledge gained 

A seminar on the intercritical rolling of high strength steel was organized at the MSE Congress 2018 
(https://www.mse-congress.de)  at  Darmstadt,  Germany  on  26th  –  28th  September  2018  to 
disseminate  the  project  results  to  public  audiences  from  industry  to  academia.  Presentation  and 
discussion  from  all  partners  of  the  INCROHSS  project  have  received  great  attention  and  positive 
reaction from audiences. 

Figure 1: Technical program on the INCROHSS seminar at the MSE Congress 2018 on 26th 
- 28th September 2018 in Darmstadt, Germany. 

9 

 
 
 
 
10 

 
 
 
 
11 

 
 
 
 
WP2: Selection, production & determination of thermomechanical characteristic 
temperatures of material 
The objectives of the WP2 are: 

Final selection of the material matrix 
Production & collection of material and its pre-processing 

- 
- 
-  Generation of essential data for the definition of the experimental set-up in WP3 to WP6. 

TASK 2.1 Chemistry selection and material workflow 

As indicated in the proposal, a basic CMn alloy (CMn-0.06%C) has been selected as the reference 
material. The contents of Si, Al, P and S have been chosen according to the current industrial practice, 
after discussion among all the partners. 

To  evaluate  the  impact  of  C  and  micro-alloying  elements  (Nb+V)  on  the  connection  among 
intercritical  rolling,  microstructural  evolution  and  mechanical  properties,  two  additional  material 
grades,  (CMn-0.12%C)  and  NbV),  are  also  studied  in  this  project.  Their  nominal  chemical 
compositions are given in Table 1. 

Table 1: Nominal chemical compositions measured in wt.% of the investigated materials. 

Grade 

CMn-0.06%C 

CMn-0.12%C 

NbV 

C 

0.06 

0.12 

0.06 

Mn 

1.5 

1.5 

1.5 

Si 

P 

S 

Al 

0.25 

0.012  0.002 

0.03 

0.25 

0.012  0.002 

0.03 

Nb 

- 

- 

V 

- 

- 

0.25 

0.012  0.002 

0.03 

0.05 

0.03 

TASK 2.2 Production & collection of test material 

Initially,  some  materials  were  proposed  to  be  collected  from  industrial  slabs  or  transfer  bars. 
However, there were difficulties to acquire desired materials at the beginning of this project. This 
obstacle  was  foreseen  causing  a  delay  on  other  activities  of  the  project.  Therefore,  lab  casted 
materials were decided to use through-out the project. The steel production which was supposed to 
be completed within the first semester of the project has been prolonged until the end of the 2nd year 
to provide enough materials for all partners. To fulfill this task, the large number of different material 
grades were produced by CRM and OCAS. 

CRM concentrated on producing studied materials, i.e. the CMn-0.06%C, the CMn-0.12%C and the 
NbV grades. 22 lab casted ingots having dimensions of 240 mm * 120 mm * 280 mm were produced 
in a laboratory vacuum induction melting furnace. Actual chemical compositions of all three material 
grades, measured by OES, are given in Table 2, Table 3 and Table 4, respectively. 

On the other hand, OCAS focused on “single-phase” materials which are used for the study on strain 
partitioning during intercritical rolling in WP4. Because three intercritical temperatures are desired 
to be investigated, six laboratory grades having the same equilibrium compositions as intercritical 
ferrite and austenite for each of the two CMn steels were expected at the beginning of the project. 
However, since the carbon content in intercritical ferrite is hardly different among various intercritical 
temperatures, only  two  low carbon grades of  0.000 wt.%C  and  0.007  wt.%C  for  each  of the  two 
studied  CMn  materials  were  casted.  The  total  number  of  “single  phase”  materials  are  ten, 
corresponding  to  4  ferritic  and  6  austenitic  chemical  compositions  at  3  intercritical  temperatures. 
Realized chemical compositions of the 10 casted blocks are given in Table 5 and Table 6. 

One lab casted ingot for each grade was initially pre-processed by laboratory hot rolling in order to 
produce  homogenized  material  for  the  dilatometry,  torsion  and  plane  strain  compression  tests  in 
Task 2.3 (determination critical temperature reference grades) as well as in WP4 (strain partitioning 
during  intercritical  deformation).  The  geometry  of  the  rolling  blocks  and  the  thermomechanical 

12 

 
 
 
 
schedule  were  applied  to  produce  16  mm-thick  plates,  as  requested  by  the  CEIT  and  TUDT.  In 
additional, 13 mm-thick plates were hot-rolled for OCAS and CRM in order to machine dilatometry 
(1 plate per grade, OCAS) and torsion samples (1 plate per grade/OCAS, 1 plate per grade/CRM) for 
the tests foreseen in the task 2.3. There are totally 9 hot rolled plates. The remaining heat blocks 
were shared between OCAS and CRM for the numerous rolling trials in WP3 and WP5. 

Table  2:  Chemical  composition  (OES)  of  the  CMn-0.06%C  grade,  measured  in  wt.%  on 
each individual casting blocks. 

Heat 

C 

Mn 

Si 

P 

S 

Al 

V 

Nb 

Cr 

N 

target  0.06 

1.5 

0.25  0.012  0.002  0.03 

2342 
2346 
2347 
2348 
2349 
2350 
2370 
2371 
2372 
2373 
2374 
2375 

0.063  1.535  0.250  0.0130  0.0030  0.035  0.005  0.002  0.0012  0.003 
0.064  1.546  0.249  0.0123  0.0026  0.037  0.002  0.001  0.0130  0.0044 
0.063  1.538  0.249  0.0115  0.0024  0.043  0.002  0.001  0.0123  0.0041 
0.063  1.536  0.251  0.0124  0.0025  0.040  0.002  0.001  0.0121  0.0044 
0.063  1.531  0.249  0.0117  0.0025  0.042  0.002  0.001  0.0124  0.0040 
0.063  1.534  0.249  0.0126  0.0026  0.042  0.002  0.001  0.0123  0.0046 
0.065  1.572  0.260  0.014 
0.003  0.025  0.005  0.001  0.017  0.0039 
0.060  1.592  0.268  0.0110  0.0021  0.028  0.004  0.001  0.0153  0.0038 
0.002  0.032  0.005  0.001  0.016  0.0031 
0.060  1.549  0.266  0.012 
0.002  0.030  0.005  0.001  0.016  0.0033 
0.059  1.553  0.264  0.013 
0.002  0.031  0.005  0.001  0.016  0.0033 
0.061  1.541  0.263  0.013 
0.002  0.019  0.005  0.001  0.017  0.0048 
0.061  1.534  0.263  0.013 

Table  3:  Chemical  composition  (OES)  of  the  CMn-0.12%C  grade,  measured  in  wt.%  on 
each individual casting blocks. 

Heat 

C 

Mn 

Si 

P 

S 

Al 

V 

Nb 

Cr 

N 

target  0.12 

1.5 

0.25  0.012  0.002 

0.03 

2343 
2351 
2352 

0.124  1.524  0.247  0.0130  0.0030  0.040  0.004  0.001  0.012  0.003 
0.123  1.524  0.255  0.0114  0.0024  0.0424  0.004  0.001  0.012  0.0045 
0.123  1.522  0.246  0.0119  0.0026  0.0400  0.004  0.001  0.012  0.0044 

Table  4 :  Chemical  composition  (OES)  of  the  NbV  grade,  measured  in  wt.%  on  each 
individual casting blocks. 

Heat 

C 

Mn 

Si 

P 

S 

Al 

V 

Nb 

target  0.06 

1.5 

0.25  0.012  0.002  0.03  0.03  0.05 

Cr 

- 

N 

- 

2344 
2353 
2354 
2376 
2377 
2378 
2379 

0.062  1.520  0.248  0.013 
0.003  0.038  0.030  0.056  0.012  0.004 
0.061  1.522  0.253  0.0115  0.0025  0.042  0.028  0.055  0.012  0.0043 
0.061  1.533  0.251  0.0119  0.0024  0.042  0.035  0.058  0.012  0.0045 
0.002  0.019  0.034  0.054  0.017  0.0049 
0.058  1.537  0.260  0.013 
0.002  0.023  0.035  0.056  0.017  0.0043 
0.060  1.532  0.262  0.012 
0.002  0.031  0.035  0.056  0.016  0.0035 
0.061  1.542  0.265  0.013 
0.002  0.028  0.035  0.054  0.016  0.0034 
0.060  1.517  0.262  0.014 

Table  5:  Chemical  compositions  measured  in  wt.%  of  the  single  phase  materials, 
corresponding to 25% - 50% - 75% ferrite and austenite for the CMn-0.06%C grade  

13 

 
 
 
 
 
 
 
 
 
 
 
 
 
Table  6:  Chemical  compositions  measured  in  wt.%  of  the  single  phase  materials, 
corresponding to 25% - 50% - 75% ferrite and austenite for the CMn-0.12%C grade  

TASK 2.3 Determination of the critical temperatures 

CEIT and OCAS are the two partners in charge of the determination of the critical temperatures like 
Ar3, Ar1 and Tnr by dilatometry and torsion testing. In order to check the reproducibility of the results 
on the thermomechanical simulators of among all partners, a common thermomechanical schedule 
was  agreed  to  perform  a  small  round-robin  test.  This  was  conducted  because  reliable  critical 
temperatures are identified as key issues for the next WPs. 
More  specifically,  CEIT  and  OCAS  carried  on  dilatometry  trials  on  all  3  studied  material  grades. 
Kinetics of the austenite-to-ferrite transformation, the starting (Ar3) and finishing (Ar1) temperatures 
of  the  phase  transformation  have  been  determined  with  different  austenite  conditions,  i.e.  non-
deformed and deformed austenite, as well as by the cooling with different cooling rates. 
Multi-pass  torsion  tests  were  carried  on  by  CEIT,  OCAS  and  CRM.  By  applying  a  common 
thermomechanical schedule (number of deformation passes and inter-pass durations), Tnr, Ar3, Ar1 
temperatures are determined from the work-hardening curves of torsion tests and morphology of 
prior austenite grains. 
Finally, the transformation kinetics of deformed austenite similar to conditions of rolling trials in WP3  
is investigated by CRM with the “relaxation-like” torsion tests. 

A. - Dilatometry tests 

Procedures and detailed analyses on determining critical temperatures of the three studied materials 
from  dilatometry trials are  given in  the  Appendix.  To  produce  different austenite  microstructures, 
two heat treatment cycles (Cycle A and Cycle B) as given in the Figure 2 were applied. In the Cycle 
A, samples are first soaked at 1250 ºC during 10 min, then deformed at 1150º C with the strain of 
0.3 by the deformation rate of 1 s-1, holding 10 s at the deformation temperature. Afterward, the 
first cooling step is applied to cool the samples to 900 oC with the cooling rate of 5 oC/s. Finally, the 
second cooling step is implemented to cool samples to room temperature with different cooling rates 
from 0.1 ºC/s to the maximum cooling rate of ~100 °C/s. In Cycle B, the specimens are treated as 
similar to samples of Cycle A. However, between the first and the second cooling steps, there is a 
deformation of 0.3 with the strain rate of 1 s-1 at 900 ºC. It is expected that with the Cycle A, the 

14 

target phaseRequestCMnSiPSAlNblock IDTarget0.00001.500.250.0120.00200.0300.0050Realised0.00081.500.250.0130.00140.0370.0033Target0.00701.500.250.0120.00200.0300.0050Realised0.00731.500.250.0120.00130.0390.0034Target0.07801.500.250.0120.00200.0300.0050Realised0.07701.500.230.0130.00090.0270.0037Target0.11001.500.250.0120.00200.0300.0050Realised0.10701.500.250.0110.00090.0290.0037Target0.21501.500.250.0120.00200.0300.0050Realised0.21701.500.250.0130.00120.0230.0036J571J57025% γ 0.06C50% γ 0.06C75% γ 0.06C75% α 0.06C25/50% α 0.06CJ575J573J572target phaseRequestCMnSiPSAlNblock IDTarget0.00001.500.250.0120.00200.0300.0050Realised0.00081.500.250.0130.00140.0370.0033Target0.00701.500.250.0120.00200.0300.0050Realised0.00731.500.250.0120.00130.0390.0034Target0.15501.500.250.0120.00200.0300.0050Realised0.14901.500.250.0140.00090.0270.0038Target0.22801.500.250.0120.00200.0300.0050Realised0.22701.500.250.0120.00110.0270.0042Target0.44801.500.250.0120.00200.0300.0050Realised0.44301.500.240.0120.00120.0300.003725% α 0.12CJ57050/75% α 0.12CJ57125% γ 0.12C50% γ 0.12C75% γ 0.12CJ577J576J574 
 
 
 
 
 
transformation of austenite to ferrite would occur with initially recrystallized austenite microstructure. 
On the other hand, with the Cycle B, the transformation would be from deformed austenite. 
Dilatometric  curves  from  all  heat  treatment  cycles  have  been  analyzed  in  the  conjunction  with 
microstructural characterization by optical and scanning electron microscopes. Results of this work 
are  summarized  in  Table  7.  Equilibrium  critical  temperatures  for  the  austenite  to  ferrite 
transformation  calculated  by  the  ThermoCalc®  software  are  also  included  in  the  same  table.  The 
results obtained by CEIT and OCAS are quite consistent. In most of the cases, the difference on the 
same critical temperatures determined by the two partners is less than 20 oC. 

Figure  2:  Schematics  of  the  thermomechanical  schedules  performed  at  the  dilatometer 
(non-deformed austenite in Cycle A and deformed austenite in Cycle B). 

In Figure 3, transformation start and finish temperatures are given as a function of cooling rate for 
both recrystallized and deformed austenite respectively. It can be seen that the addition of C (from 
0.06 wt.%C to 0.12 wt.%C) and the presence of microalloying elements (Nb and V) retard phase 
transformations,  and  consequently  decrease 
(mainly 
transformation start temperature). For a similar carbon content (0.06 wt.%C) the addition of Nb and 
V causes the decrease of transformation start and finish temperatures, leading to the formation of 
bainite microstructure and reducing the presence of polygonal ferrite microstructure. In terms of the 
cooling rate, the increase in cooling rate as expected leads to the formation of bainitic/martensitic 
microstructures, which occur at lower temperatures than the one of polygonal ferrite microstructure. 
Regarding to the initial austenite microstructure before the transformation, the stored strain energy 
in deformed austenite stimulates the phase transformation at higher temperatures (see Table 7 and 
Figure 3) as compared to the one of recrystallized austenite. 

transformation 

temperatures 

the 

Based on the dilatometry trials, continuous cooling transformation diagrams (CCT diagrams) for all 
three studied materials with the 2 different initial austenite microstructures have been built. These 
results  are  given  in  the  Appendix  and  will  be  used  for  the  selection  of  intercritical  rolling 
temperatures. 

Table  7:  Transformation  start  (Ar3)  and  finish  (Ar1)  temperatures  for  different  steels 
(CMn-0.06%C,  CMn-0.12%C  and  NbV),  selected  cooling  rates  (0.1,  1  ,10  and  100  ºC/s) 
and  different  austenite  conditions  prior  to  transformation  (recrystallized  and  deformed 
austenite) 

15 

1250 ºC, 10 minTEMPERATURE TIMECoolingrates: 0.1 ºC/s, 1 ºC/s, 10 ºC/s, maxcoolingCycleACycleB=0.3; =1 s-10 s=0.3; =1 s-110 sRefinement5ºC/s 
 
 
 
(a) 

(b) 

Figure 3: Transformation start and finish temperature as a function of cooling rate for (a) 
Non deformed and (b) Deformed austenite. 

B.- Torsion tests 

In order to determine the critical temperatures Tnr, Ar3, Ar1, each partner (CEIT, OCAS and CRM) has 
carried out multi-pass torsion tests following the thermomechanical schedule illustrated in Figure 4. 
After a soaking at 1250 ºC for 10 min, a deformation step was performed at 1150º C (deformation 
of 0.3, deformation rate of 1 s-1 and a holding time of 10 s), in order to obtain a fine recrystallized 
austenite. Afterwards, 26 deformation passes were applied at decreasing temperature in the 1150 ºC 
and 650 ºC range. The temperature reduction between torsion passes is 20 ºC. Different strains per 
pass (ɛ ) and interpass times (tip) has been used, as shown in Table 8. A strain rate of 1s -1 is defined 
in all the cases.  

16 

Ts (ºC)Tf (ºC)Ts (ºC)Tf (ºC)Ts (ºC)Tf (ºC)841682826692849681CEIT799636771606801689OCAS821646768625778623CEIT766614708582709556OCAS761628720583CEIT700559669513651546OCAS733579673495630538OCAS100 ºC/s549417489338539408CEIT824720783614804714OCAS818716796616818718CEIT769624756571774554OCAS779632751581768541CEIT753571725528704540OCAS753565744493721537OCAS100 ºC/s569375417330475325Equilibrium Temperatures for ferrite (Thermocalc)CycleLabCooling RateNon Deformed ɣCycle A0.1 ºC/s1 ºC/s10 ºC/sDeformed ɣCycle B 0.1 ºC/s1 ºC/s10 ºC/sCMn (0.06%C)CMn (0.12%C)NbV 
 
 
 
 
 
Figure 4: Schematics of the thermomechanical schedules for the MFS torsion trials 

Table 8: Deformation conditions defined in the multipass torsion tests. 

Soaking temperature (ºC) 
Deformation per pass 
Interpass time (s) 
Cooling rate (s-1) 

1250 
0.2, 0.3, 0.4 
5, 10, 20 
1, 2, 4 

CEIT has carried on the multi-pass torsion tests with all three studied materials. From the stress-
strain  curves  obtained  from  these  trials,  mean  flow  stress  (MFS)  corresponding  to  each  pass  has 
been extracted. This mean flow stress (MFS) is calculated by the area under each stress-strain curve 
divided by the pass strain. 

Figure  5  shows  an  example  for  the  change  of  the  MFS  with  deformation  temperature  in  the  NbV 
material. The strain applied to each torsion pass is 0.3, while the inter-pass times are 20 s, 10 s and 
5 s, respectively. 

(a) 

(b) 

(c) 
Figure 5: Mean flow stress as a function of temperature from the multipass torsion test on 
the NbV material. The strain of each torsion pass is 0.3, while the inter-pass times are 20 
s, 10 s and 5 s respectively. 

17 

1250 ºC, 10 minTEMPERATURE TIME=0.3; =1 s-110 s1; =1 s-1; =1 s-1RefinementMultipassDeformation () = 0.2, 0.3, 0.4Interpasstimes (tip) = 5 s, 10 s, 20 stiptip050100150200250300350600700800900100011001200MEAN FLOW STRESS (MPa)T (ºC)Tnr=1027ºCNbVtip=20sɛ=0.3Ar3=810ºCAr1=750ºC050100150200250300350600700800900100011001200MEAN FLOW STRESS (MPa)T (ºC)Tnr=1032ºCNbVtip=10sɛ=0.3Ar3=810ºCAr1=750ºC050100150200250300350600700800900100011001200MEAN FLOW STRESS (MPa)T (ºC)Tnr=1039ºCNbVtip=5sɛ=0.3Ar3=810ºCAr1=750ºC 
 
 
 
 
 
 
The value of Tnr is determined from the intersection between the two linear lines, corresponding to 
the  changes  of  MFS  with  temperatures,  as  showed  in  Figure  5.  Besides  the  Tnr  temperature, 
transformation  start  (Ar3)  and  finish  (Ar1)  temperatures  are  also  able  to  identify  from  the  same 
figures. 

Critical temperatures (Tnr, Ar3 and Ar1) of all 3 different materials, including various interpass times 
are  listed  in  Table  9.  For  the  highest  deformation  of  0.4,  in  some  cases  the  thermocouple  was 
detached from the sample surface during the tests, due to too many deformation passes and high 
deformation per pass. Therefore, the Ar3 and Ar1 temperatures were not able obtained in these cases. 
It is evident that the  NbV material has higher Tnr temperatures than the other two (CMn-0.06%C 
and CMn-0.12%C) materials. The effect of micro-alloying elements on the Tnr temperature could be 
argued  by  the  solute  drag  effect  regarding  to  solute  atoms  and  the  pinning  effect  due  to  strain 
induced precipitates in the NbV material.  

Table 9: Critical temperatures (Tnr, Ar3 and Ar1) obtained from multipass torsion test for 
all the studied materials with various inter-pass times and deformation per torsion passes. 

In  terms  of  transformation  temperatures,  the  results  shown  in  Table  9  suggest  that  chemical 
composition has impacts on transformation start and finish temperature, although the temperature 
resolution of multipass torsion test (20 ℃ per pass) is rather coarse. Impacts of chemical composition 
on the Ar1 and Ar3, although are not as clear as on the Tnr temperature, show the addition of carbon 
from  0.06%  to  0.12%  results  in  the  reduction  of  transformation  start  (Ar3)  and  finish  (Ar1) 
temperatures.  

In  order  to  have  better  understanding  on  the  influence  of  microalloying  elements  on  the  critical 
temperatures, MFS curves of the CMn-0.06%C and the NbV steels are compared. For both of the two 
samples, the deformation of 0.2 per torsion pass and the inter-pass time of 20 s are applied. It can 
be observed from the Figure 6 that the presence of micro-alloying elements has significant influence 
on  the  non-recrystallization  temperature,  whereas  the  austenite  to  ferrite  phase  transformation 
kinetics are not affected.  

18 

Tnr(ºC)Ar3(ºC)Ar1(ºC)Interpasstimetip (s)5s10s20s5s10s20s5s10s20sɛ=0.2CMn(0.06%C)927922913810790790770750750CMn(0.12%C)915913914770770770730730730NbV103010371036790810790750770750ɛ=0.3CMn(0.06%C)910908903790790790750750750CMn(0.12%C)914902902770770770730730730NbV104210261019810810810750750750ɛ=0.4CMn(0.06%C)899895898-810810-750770CMn(0.12%C)895877859------NbV102110191004810810810750750750 
 
 
Figure 6: Influence of composition on critical temperatures (Tnr, Ar3 and Ar1): Mean flow 
stress curves corresponding to the steels CMn-0.06%C and NbV steel, using an interpass 
time of 20 s and a deformation per pass of 0.2.  

In  order  to  evaluate  the  effect  of  interpass  time  and  deformation  per  pass,  non-recrystallization 
temperature have been plotted as a function tip and deformation in Figure 7a and b, respectively. 
When the deformation of 0.3 is used (see Figure 7a), although influence of inter-pass time on Tnr of 
the CMn steels is not significant, higher Tnr temperatures of the NbV steel are measured with shorter 
interpass times. Similar trends are also observed for the deformation of 0.2 and 0.4. Concerning the 
influence  of  deformation  per  pass,  the  results  shown  in  Figure  7b  suggest  that  the  application  of 
higher deformation leads to a lower Tnr temperatures of all three material grades. 

(a) 

(b) 

Figure  7:  (a)  Influence  of  chemical  composition  and  interpass  time  in  the  non-
recrystallization temperature. (b) Effect of deformation per pass in the Tnr. 

At CRM, multi-pass torsion trials are also carried on. Results of all measurements are given in the 
Appendix. To check the consistency of critical temperatures determined by multi-pass torsion trials, 
a round-robin comparison has been implemented among CEIT, OCAS and CRM. 

The values of the critical temperature determined by all partners are given in Table 10. The Ar3 and 
Ar1 values determined by CEIT and CRM are in good agreement for all the investigated conditions 
and similar to those predicted by the empirical formulae. Large differences on Tnr, temperatures are 
due to the non-unique definition of this temperature. In any case, the temperatures determined by 
CRM are closer to those reported in literature for similar compositions.  For the rolling trials which 
were implemented in the later WPs, temperature evolutions in samples through-out all rolling passes 
were  recorded  either  by  pyrometer,  thermocouple  or  both  of  this  two  methods.  Afterward, 

19 

050100150200250300350600700800900100011001200MEAN FLOW STRESS (MPa)T (ºC)CMn-0.06%CNbVTnr=913ºCtip=20sɛ=0.2Ar3=790ºCAr1=750ºCTnr=1036ºC80085090095010001050110011500510152025Non-Recrystallization Temperature, Tnr (ºC)Interpass Time, tip (s)DEFORMATION 0.3CMn-0.06%CCMn-0.12%CNbV80085090095010001050110011500.10.20.30.40.5Non-Recrystallization Temperature, Tnr (ºC)Deformation per pass INTERPASS TIME 10sCMn-0.06%CCMn-0.12%CNbV 
 
 
 
 
microstructures and mechanical property analyses  for each specific plates  were always carried  on 
with recorded temperature history. 

Table 10 : Critical temperatures determined by torsion testing (ε=0.3) by three partners 
(CEIT,  CRM  and  OCAS),  in  a  comparison  with  results  of  dilatometry  measurements, 
calculated equilibrium temperatures and empirical formulae. 

C - “RELAXATION-LIKE” TORSION TRIALS 

The  austenite  to  ferrite  transformation  kinetics  is  strongly  influenced  by  the  initial  condition  of 
austenite (e.g. grain size, recrystallization/deformation, precipitation …). However, multi hot rolling 
passes  are  not  able  to  be  consistently  simulated  by  plasto-dilatometry  experiments.  Dedicated 
torsion  experiments,  therefore,  have  been  carried  on  by  CRM  to  simulate  a  multi-pass  rolling 
schedule including the subsequent transformation. Figure 8 shows the thermo-mechanical schedule 
for a “relaxation-like” torsion trial which is similar to the one applied in in WP3. A remarkable feature 
on  this  experiment  is  the  continuous  deformation  at  very  low  strain  rate  of  0.001  s-1  during  the 
cooling  after  inter-critical  deformation,  in  order  to  monitor  the  stress  evolution  in  samples.  The 
formation of “soft” ferrite is detected by a drop in the stress level. After having determined the stress-
temperature dependence of pure austenite and pure ferrite (nicely described of an Arrhenius-type 
law) by extrapolation, it is also possible to determine the entire transformation kinetics by a simple 
lever rule, as illustrated in Figure 9. 
The  transformation  kinetics  determined  for  the  same  austenite  conditioning  as  during  the  rolling 
trials in WP3 allow to determine  the temperature for 25%, 50% and 75% ferrite formation which 
will be used in WP3 (Figure 10). The NbV steel has the highest transformation temperatures because 
of austenite pancaking occurring during the last deformation passes. Strain accumulation does not 
occur in the reference CMn (0.06%C) grade and consequently the transformation takes place about 
20  °C  lower.  Increasing  the  carbon  content  (CMn  (0.12%C)  grade)  decreases  further  the 
transformation temperatures. 

Figure 8: Schedule “relaxation-like” torsion trials 

20 

 
 
 
 
 
 
 
Figure 9: Extrapolation of the stress temperature dependence for of the pure ferrite and 
austenite and determination of the ferrite transformation kinetics 

Figure  10:  Determination  of  the  temperatures  corresponding  to  25%,  50%  and  75% 
ferrite formed 

21 

 
 
 
 
 
 
 
 
22 

 
 
 
 
WP3: Full mapping of the effect of the α/γ  balance for the mechanical properties 

of the reference alloy 
The objective of the WP3 is: 

- 

To get the first systematic insight, for the reference alloy, on the impact of intercritical rolling 
on the mechanical properties. The reference alloy with the largest γ/α temperature range is 
selected and the effect of the γ/α balance is studied for two cooling strategies for plates: air 
cooling and accelerated cooling. A full mapping of the impact of γ/α balance on the strength-
toughness balance is the main goal of this WP. 

TASK 3.1 Laboratory rolling trials with single pass intercritical deformation 

The objective of WP3 is to get a first systematic insight on the impact of intercritical rolling on the 
mechanical properties of thin gauge plates (12 mm thickness). More particularly a  full mapping of 
the impact of γ/α ratio on the strength-toughness balance is realized. 

The CMn-0.06%C steel, which exhibits the largest γ/α temperature range, was selected. The effects 
of the γ/α ratio were studied for 3 cooling strategies after the last intercritical rolling pass: air cooling, 
accelerated cooling and direct quench. 

The  experimental plan  of the  rolling  trials  with  single  pass  IC  deformation  is  shown  in  Figure  11. 
Trials  were  shared  between  OCAS  (blue,  air  cooling  and  direct  quench  routes)  and  CRM  (green, 
accelerated cooling routes): 

  Set A is the schedule set for industrial processing via intercritical rolling. Three cooling 

paths were followed, i.e. air cooling, accelerated cooling and direct quench. 

  Set  B  is  the  reference  rolling  schedule  where  the  complete  rolling  is  done  in  austenitic 

region.  

  Set C is a set of trials with quenching just before the last rolling pass in the intercritical 

area. The objective of this set is to estimate the actual ferrite fraction before the last pass. 

Several finish rolling temperatures were investigated: the aim was to apply the last rolling pass on 
microstructures  containing  25%,  50%  and  75%  of  ferrite  respectively.  The  details  of  the  rolling 
schedules for the respective sets are presented in Table 11. 

Figure 11: Experimental plan of the first rolling trials of WP3. 

After a roughing simulation, a finishing simulation was done with a finish rolling temperature in the 
austenitic region. For set A, a last pass of 25% was given at various temperatures in the intercritical 

23 

 
 
 
 
 
area. Temperature evolutions during rolling was recorded by thermocouples inserted to the mid width 
of the slab and pyrometers attached to the rolling mill. 

Table 11: Detailed rolling schedule of the trials foreseen in WP3. 

The realization of the intercritical rolling temperatures corresponding to initial microstructures with 
25%, 50% and 75% ferrite is not easy because ferrite transformation occurs far from equilibrium. 
The actual amount of ferrite formed during continuous cooling depends not only on temperature, as 
it would be the case under equilibrium conditions, but also on the transformation kinetics. The latter 
is strongly influenced by the  thermomechanical schedule (prior austenite grain size and austenite 
work-hardening) and by the Mn content. 

In  the  first  months  of  the  project  (WP2),  a  significant  effort  was  made  to  investigate  the  ferrite 
transformation kinetics by means of different techniques (torsion and plasto-dilatometry). The results 
are summarized in Figure 12. 

Figure 12: Summary of the ferrite fraction formed as function of temperature as estimated 
by  torsion  and  dilatometry,  and  comparison  with  the  fraction  determined  by 
metallography. 

OCAS based the selection of intercritical rolling temperatures on the first results of CEIT, while CRM 
based their selection on the “relaxation-like” torsion trials. As a consequence, the target intercritical 
rolling temperatures used by CRM were lower compared to OCAS, although both of the two partners 
used the same techniques (thermocouple  and  pyrometer) to measure temperature during rolling. 
CRM has used a lower FRT in the fully austenitic temperature regime as well, which has also an effect 
on microstructure and properties. In Table 12 a comparison of the used target rolling temperatures 
are given. In any case, the actual fraction of ferrite is verified by quenching into water the plate prior 
to the last intercritical rolling pass (set C). 

24 

 
 
 
 
Table 12: Comparison of the rolling process temperatures targeted by OCAS and CRM 

TASK 3.2 Microstructural characterization 

Microstructural characterization on DQ samples before intercritical deformation 

Figure  13  shows  the  microstructures  of  samples  quenched  just  before  and  after  the  intercritical 
deformation for the various temperatures selected by OCAS, which targeted to correspond to 25%, 
50% and 75% ferrite respectively. Before intercritical rolling the actual fractions of ferrite at ~790°C 
and  ~750°C  are  quite  close  to  the  targeted  amounts  of  75%  and  25%.  However,  a  mismatch  at 
~770°C is observed. Instead of 50% martensite, only ~30% was measured. Again, a mismatch in 
ferrite / martensite fraction is also observed when intercritical rolling is applied.  For the other two 
target temperatures, 790°C and 750°C, the values after intercritical rolling are in accordance to the 
ones before deformation, see Figure 14. The differences in ferrite / martensite fraction are due to 
the applied deformation and the lower start cooling temperature. 

Set C – DQ from intercritical target temperature 

Before deformation 

After deformation 

64% MA 

(a) Target 25% ferrite. Realized: 36 %F 

(d) Target 25% ferrite. Realized: 40% F 

26% MA 

(b) Target 50% ferrite. Realized: 74 %F  

(e) Target 50% ferrite. Realized: 50% F 

25 

OCASCRMFRT820°C795°CIC-target ~25% ferrite790°C755°CIC-target ~50% ferrite770°C745°CIC-target ~75% ferrite750°C735°CAir coolingAccelerated cooling (~30K/s)Direct QuenchDirect QuenchCooling 
 
 
 
 
 
  
  
 
29% MA 

(c) Target 75% ferrite. Realized: 71 %F 

(f) Target 75% ferrite. Realized: 75% F  

Figure 13: Microstructures at quarter thickness as obtained in the set C after etching with 
Nital 2%, i.e. quenched just before (a-c) and after (d-f) the last intercritical pass (OCAS 
trials) 

Figure  14:  Comparison  of  the  martensite  fraction  after  direct  quench  from  intercritical 
temperature  range  before  (full  symbols)  and  after  intercritical  rolling  (cross  symbol), 
determined from optical micrographs 

Microstructural characterization on ACC samples before and after intercritical deformation 

The as-quenched microstructures prior to intercritical rolling and ACC produced by CRM are presented 
in Figure 15. The microstructures obtained after quenching at 795°C and 785°C consist mainly of 
martensite, bainite and fine/acicular ferrite. Such phases are likely to be formed upon cooling at 70 
– 80 °C/s, i.e. the cooling rate measured at mid-thickness, from partially non recrystallized austenite. 
The microstructure at T≥785°C is supposed to be fully austenitic. At temperature equals or below 
765°C, some large ferrite grains are observed. They are likely to be formed during the slow cooling 
prior to water quench. The actual ferrite fraction ranges from 50% at 765°C to about 70% at 735 
°C.  It can be concluded form this trial set that the intercritical process window ranges from about 
770°C to 735 °C, i.e. 20°C lower than the temperature range estimated by OCAS. 

Set C – DQ from intercritical target temperatures 

26 

 
 
  
 
(a) Target 795°C, realised: 80% MA, 20% ferrite 

(b) Target 785°C, realised: 85% MA, 15% ferrite 

(c) Target 765°C, realised: 47 %MA, 53% ferrite 

(d) Target 755°C, realised: 42 %MA, 58% ferrite 

(e) Target 745°C, realised: 32 %MA, 68% ferrite 

(f) Target 735°C, realised: 34% MA, 66% ferrite 

Figure  15:  Microstructures  as  obtained  in  the  set  C  after  etching  with  Nital  2%,  i.e. 
quenched just before the last intercritical pass (CRM trials) 

The  microstructure  obtained  after  fully  austenitic  rolling  (FRT=795  °C)  and  accelerated  cooled 
(quench in water at 60 °C for about 10 s) is presented in Figure 16. It consists mainly of ferrite and 
pearlite,  with  some  bainite/martensite.  No  evidences  of  deformed  ferrite  (as  for  instance  ferrite 
grains exhibiting a recovery substructure) are observed. 

The  microstructures  resulting  from  intercritical  rolling  followed  by  accelerated  cooling  are  quite 
complex  (Figure  16).  A  plate  rolled  at  710  °C,  i.e.  in  the  almost  fully  ferritic region,  is  added  for 
comparison,  although  it  was  not  one  of  the  intercritical  conditions  for  investigation.  All  the 

27 

 
 
 
 
 
 
 
 
microstructures  obtained  after  intercritical  rolling  contain  some  bainite/martensite  (dark  phases) 
resulting  from  the  decomposition  of  C  enriched  γ.  This  confirms  the  fact  that  deformation  is 
performed  in  the  intercritical  region.  Estimating  the  bainite/martensite  fraction  from  LOM 
micrographs is nevertheless a tricky issue. A significant fraction of ferrite is also present after all the 
intercritical rolling temperature considered. Small ferrite grains (about 10 µm) and coarser (40 µm) 
grains coexist. A clear and pronounced recovery substructure is never observed into the grains by 
LOM,  probably  because  of  the  small  deformation  applied  in  the  intercritical  region  (ε=0.33).  As 
various  metallurgical  phenomena  are  likely  to  take  place  during  and  after  intercritical  rolling 
(recovery  of  deformed  ferrite,  ferrite  recrystallization,  decomposition  of  deformed  γ),  the  specific 
origin of each grain population is still unclear and needs more detailed investigations (EBSD, rolling 
trials  with  higher  deformation  levels).  However,  the  fact  that  the  sample  deformed  at  the  lowest 
temperature (710 °C) contains mainly coarse ferrite grains, indicates that possibly the coarse grains 
are  formed  before  deformation  while  the  small  ones  result  from  the  decomposition  of  deformed 
austenite during cooling. Indeed, the fraction of coarse grains seems to increase as the deformation 
temperature decreases. 

Set A and B – ACC after last rolling pass (cooling in hot water at 60 °C) 

(a) Target 100% A. Pass at 795°C. HV=140 

(b) Target 25% F. Pass at 755°C. HV = 173 

(c) Target 50% F. Pass at 745°C. HV = 188 

(d) Target 75% F. Pass at 735°C. HV = 191 

28 

 
 
 
 
 
 
(e) Target 100% F. Pass at 710°C.HV = 177 

Figure 16: Microstructures obtained after intercritical rolling at various temperatures and 
accelerated cooling (simulated by quench in hot water at 60 °C) (set A and set B-ACC) 

Microstructural characterization on Air cooled samples before and after intercritical deformation 

The microstructures obtained after  austenitic (FRT=820°C) and intercritical rolling followed by air 
cooling,  are  presented  in  Figure  17.  As  expected,  the  microstructures  are  ferritic-pearlitic.  The 
pearlite  fraction  is  about  10%  in  all  samples,  independent  on  the  deformation.  However,  without 
intercritical deformation the pearlite seems to have a more pronounced band structure. The ferrite 
grain  structure  changes  with  the  application  of  intercritical  rolling.  Based  on  the  light  optical 
micrographs,  the  ferrite  grains  are  somewhat  finer  and  more  elongated,  especially  at  the  lower 
temperature of intercritical deformation.  

Set A and B – Air cooling after last rolling pass 

(a) Target 100% A. Pass at 820°C. 

(b) Target 25% F. Pass at 790°C.  

(c) Target 50% F. Pass at 770°C.  

   (d) Target 75% F. Pass at 750°C.  

Figure  17:  Microstructures  obtained  after  intercritical  rolling  at  various  temperatures 
followed by air cooling 

29 

 
 
 
 
 
 
 
TASK 3.3 Mechanical properties 

The  tensile  properties  have  been  measured  by  full  thickness  A50  geometry  samples  taken  in 
transverse and rolling direction of the plate. The final plate thickness was deliberately set to 12mm 
thickness making it possible to test full size Charpy V-notch. Therefore, from the plate top and bottom 
1mm  thickness  was removed. Charpy samples were  also machined in both rolling and transverse 
direction. 

Mechanical properties after intercritical deformation and air cooling 

The  tensile  properties  after  intercritical  rolling  followed  by  air  cooling  are  shown  in  Figure  18.  An 
increase  in  strength  is  observed  when  decreasing  the  intercritical  rolling  temperature.  The  yield 
strength  is  increasing  faster  (76-92MPa)  than  the  tensile  strength  (60-53MPa),  leading  to  higher 
yield  to  tensile  ratios  with  lowering  temperature.  The  elongation  is,  however,  decreasing  with 
decreasing temperature, dropping from 20% to 13% in both test directions.  

(a) transversal direction (TD) 

(b) longitudinal direction (LD) 

Figure 18: Strength as a function of start cooling temperature for plates rolled intercritical 
followed by air cooling. 

Figure 19 shows the Charpy impact toughness for the air-cooled plates as a function of finish rolling 
temperature.  When  finish  rolled  in  the  austenitic  temperature  zone  (820°C)  or  the  ferritic  zone 
(730°C)  the  toughness  is  better  than  intercritically  rolled.  This  means  that  intercritical  rolling 
deteriorates the toughness, particularly ductile to brittle transition temperature. The worst toughness 
is observed when rolled and start cooling at around 790°C.  

IC 

(a) transversal direction (TD) 

(b) longitudinal direction (LD) 

Figure 19: Charpy impact toughness of the plates air cooled after intercritical rolling, as a 
function of finish rolling / start cooling temperature 

Mechanical properties after intercritical deformation and accelerated cooling (ACC) 

The tensile properties after intercritical rolling followed by accelerated cooling are presented in Figure 
20. The main tendencies are the following: 

 

Increasing ferrite fractions (during the lasts rolling pass) lead to an increase in both yield 
and tensile strength. The increase in YS is more pronounced than the increase in TS, which 
seems to approach a steady state. 

30 

 
 
 
 
 
 
  Decreasing FRT (i.e. increasing deformed ferrite  fraction)  decreases the uniform and total 

elongation. 

(a) transversal direction (TD) 

(b) longitudinal direction (LD) 

Figure  20:  Tensile  properties  as  function  of  the  last  pass  temperature  (FRT=795  °C/ 
austenitic rolling), (FRT= 755, 745 and 735 °C – intercritical rolling). Accelerated cooling 
conditions. 

The results of the Charpy V-notch impact tests are shown in Figure 21. The impact toughness values 
of intercritical rolled plates are  compared to  those  corresponding to plates rolled in the austenitic 
region (FRT=795 °C) and in almost full ferritic region (FRT=710 °C).  The best toughness properties, 
both in term of upper shelf energy as DBTT, are obtained for the material rolled in the austenitic 
region. The toughness properties become worse as the FRT decreases. No improvement is observed 
if  the  last  pass  is  performed  in  the  fully  ferritic  region  (710  °C).  Here,  the  worst  properties  are 
measured, contrarily to what observed when ferritic rolling is followed by air cooling. A significant 
scattering  of  the  Charpy  values  is  observed  in  the  intercritical  rolled  plates.  This  is  probably  a 
consequence of the thermal gradient present on the plate during rolling and cooling: small differences 
in  temperature  in  the  intercritical  region  can  result  in  very  different  ferrite  fractions  and 
microstructures. Indeed, scattering is less important  in plates produced by fully austenitic or fully 
ferritic rolling. 

(a) transversal direction (TD) 

(b) longitudinal direction(LD) 

Figure  21: Charpy impact toughness as function of the  last pass temperature (FRT=795 
°C/ austenitic rolling), (FRT= 755, 745 and 735 °C – intercritical rolling) 

Mechanical properties before intercritical deformation  

In  Figure  22  the tensile strength levels before and  after  the  intercritical deformation, followed by 
direct quench, are shown. The strength appears to be independent on the start cooling temperature 
with a standard  deviation between 20  and  40MPa.  The same  behavior is  observed when applying 
intercritical deformation before direct quench, just at higher strength (Figure 22). When comparing 
the values of before and after intercritical rolling, it can be noticed, that the deformation leads to an 
increase in strength of about 100MPa in average, for both YS and TS. The values obtained at 770°C 
with direct quench before deformation  appear to  be higher than expected (black circle),  probably 
due to the higher cooling rate. 

31 

 
 
 
 
 
 
(a) transversal direction (TD) 

(b) longitudinal direction (LD) 

Figure 22: Comparison of tensile results before (full) and after (empty) intercritical rolling 
reduction followed by direct quench in both, transverse and longitudinal direction 

Figure  23  compares  the  Charpy  impact  toughness  with  and  without  intercritical  rolling  for  direct 
quenched  plates.  As  can  be  seen,  the  general  trend  is  the  same  observed  in  air-cooled  and 
accelerated  cooling  plates:  independent  of  intercritical  rolling,  the  toughness  deteriorates  with 
decreasing  start  cooling  temperature  (black  arrows).  It  is  assumed  that  the  toughness  does  not 
significantly change if intercritical rolling is applied or not when the plates are direct quenched, see 
Figure 24. 

DQ before IC 

DQ after IC 

Figure 23: Comparison of the Charpy impact toughness with direct quench before or after 
intercritical rolling as a function of start cooling temperature 

32 

90° Rp0.2Rt0.5Rm790_DQ341362583770_DQ402434664750_DQ324351600730_DQ346375645IC_790_DQ413443686IC_770_DQ404428665IC_750_DQ469496729IC_730_DQ462484700avg_DQ353380623avg_IC_DQ437463695stdev_DQ343738stdev_IC_DQ3332270°Rp0.2Rt0.5Rm790_DQ348375605770_DQ369402636750_DQ304329600730_DQ339364642IC_790_DQ421447686IC_770_DQ428442700IC_750_DQ424457711IC_730_DQ424445693avg_DQ340367620avg_IC_DQ424448698stdev_DQ273021stdev_IC_DQ3611 
 
 
 
 
 
 
 
 
 
Figure 24: Direct comparison of Charpy impact toughness for a selection of start cooling 
temperature - 770°C and 750°C 

TASK 3.4 Microstructural evolutions and correlation 

Comparison cooling strategies 

In  Figure  25,  the  tensile  properties  resulting  from  the  3  different  cooling  strategies  (i.e.  air, 
accelerated (ACC) and direct quench cooling) after intercritical rolling are compared. The following 
trends can be observed: 

 

 

 

For the same FRT, the strength of DQ samples are the highest while the strength of the air 
cooled samples are the lowest among all samples of the 3 cooling strategies. 
The  strength  of  intercritically  rolled  samples  increase  with  reduction  of  the  finish  rolling 
temperature. 
The strength of air cooled and direct quenched samples are less sensitive to the change of 
FRT than that of the ACC samples. Yield strength of air cooled and direct quenched samples 
increases about 1.00 MPa/K when the finish rolling temperature reduces from 790℃ to 730℃. 
Tensile strength of these samples increases by 0.50-0.60 MPa/K with the reduction of the 
FRT in the same range of temperature. For the ACC samples, the yield strength increases 
1.85 MPa/K and the tensile strength increases 1.25 MPa/K, when the FRT reduces from 795℃ 
to 710℃.  

Figure  25:  Comparison  of  the  tensile  properties  obtained  after  intercritical  rolling  for 
different cooling strategies 

The toughness properties for all the cooling strategies are summarized in Figure 26. The upper shelf 
energy (USE) and the ductile brittle transition temperature (DBTT) are plotted as a function of the 
finishing  rolling  temperature.  The  general  trend  is  clear:  in  the  intercritical  range,  toughness 

33 

 
 
 
 
 
 
deteriorates as the FRT decreases, both in term of USE as DBTT. The exception is the behavior  of 
samples after the air cooling. With the reduction of the FRT, the USE decreases but the DBTT also 
reduces. 

Figure  26:  Comparison  of  the  impact  toughness  properties  obtained  after  intercritical 
rolling for different cooling strategies 

Strength/toughness maps 

Figure 27 present the toughness yield strength (more relevant than TS for end user plate application) 
balance for the intercritical rolled plates. The opposite trends of toughness (DBTT, USE) and yield 
strength,  despite  the  cooling  strategy  applied,  is  further  emphasized  in  these  strength/toughness 
maps. The only exception is the behavior after ferritic rolling and air cooling, which seems to improve 
YS  without  affecting  toughness,  maybe  because  of  ferrite  recrystallization  occurring  during  slow 
cooling. Further investigation is needed to confirm and explain this finding. 

Figure  27:  Strength  vs.  toughness  maps:  Upper  Shelf  Energy  (USE)  and  Ductile  Brittle 
Transition Temperature (DBTT) as function of the Yield Strength (YS) level 

Link microstructure – strength/toughness balance 

For the link between microstructure and mechanical properties of the material, EBSD measurements 
were implemented on  intercritically  rolled samples. Two features of ferrite, which have significant 
impacts on mechanical properties, are the average ferrite grain size and the amount of deformation 
energy  stored  in  ferrite  grains.  Figure  28  shows  the  change  ferrite  grain  size  and  the  fraction  of 
ferrite grains having the grain orientation spread larger than 2.0°. According to the results in WP4, 
the grain orientation spread (GOS) is a good indicator for the amount of strain energy accumulated 
in ferrite grains. And thus the grains which have the GOS value higher than 2.0° could be considered 
as deformed, while the grains having the GOS  lower than 2.0° are non-deformed ones. From the 
Figure 28, it could be observed that the reduction of FRT results in the increases of average ferrite 
grain size and the fraction of deformed ferrite. However, by linking these microstructure features to 
corresponding  mechanical  properties,  it  could  be  concluded  that  the  dominant  strengthening 

34 

 
 
 
 
mechanism  in  intercritically  rolled  and  ACC  samples  is  not  by  the  Hall-Petch  relationship.  The 
reduction of finishing rolling temperature results in both an increase in tensile/yield strength and a 
development  of  the  ferrite  grain  size.  On  the  contrary,  strengthening  by  the  presence  of  second 
phase is expected to play a major role on both strength and toughness of the material. Reducing the 
finish rolling temperature causes an increase in the area fraction of deformed ferrite. For the samples 
of higher fraction of deformed ferrite, the strength is higher while the toughness is lower than those 
of the samples having lower fraction of deformed ferrite. The strengthening in intercritically rolled 
samples could eventually link to the carbon partitioning between ferrite and austenite before the last 
rolling  pass,  as  well  as  the  volume  fraction  of  transformation  products  from  austenite  during  the 
cooling to room temperature. For the lower FRT, the volume fraction of intercritical austenite is lower, 
but the carbon concentration in this phase will be higher. Depending on the carbon concentration in 
austenite and the cooling condition after the intercritical rolling, various transformation products of 
austenite could be obtained. For the air cooled samples, the transformation product from austenite 
is  generally  expected  as  pearlite.  For  the  quenched  samples,  it  is  the  martensite  and/or  bainite. 
Because  of  its  intermediate  cooling  rate,  accelerated  cooling  may  produce  a  mixture  of  different 
products,  for  instance  martensite,  bainite  and  even  pearlite.  Different  products  with  their  volume 
fractions should have various impacts on mechanical properties of samples. A deep microstructure 
investigation  on  samples  of  various  cooling  conditions  could  identify  transformation  products  of 
austenite and link their volume fractions to the strength and toughness of materials. However, such 
an analysis is not always straight forward to be applied for all samples of this WP.  In most of the 
samples, there is no clear distinguish among ferrite, bainite and martensite. Statistical distributions 
of various parameters, commonly used to differentiate morphologies and strain states of ferrite (i.e. 
KAM, GOS, …) are usually unimodal, rather than bi- or multi-modal. The GOS value of 2.0 degrees 
which  is  used  to  differentiate  between  deformed  and  undeformed  ferrite  is  just  a  representative 
parameter. Further microstructure analysis based on this rough estimation would bring more errors 
and uncertainties than the simple one. 

Figure 28: The change of average ferrite grain size and area fraction of deformed ferrite 
grains which have GOS > 2.0° with finish rolling temperatures in ACC samples. 

Conclusively, the link between microstructure and mechanical properties of intercritically rolled and 
ACC samples could be summarized by Figure 29. By applying single intercritical rolling pass at lower 
temperature and accelerated cooling, the strength of CMn0.06%C steel could be enhanced while the 
toughness would be reduced. The strengthening of intercritically rolled samples is due to the increase 
in volume fraction of deformed ferrite or the second hard phase (bainite, martensite), rather than by 
the  reduction  of  ferrite  grain  size.  Approximately,  for  every  10oC  reduction  of  the  finish  rolling 
temperature  into  the  intercritical  region,  the  area  fraction  of  deformed  ferrite  in  ACC  samples 
increases  0.03.  This  increase  in  area  fraction  of  deformed  ferrite  would  bring  13-19  MPa  to  the 
strength of material, but reduce the toughness measured by the USE of 2.7 J. 

35 

 
 
 
 
Figure  29:  The  link  between  microstructure  and  mechanical  properties  of  intercritically 
rolled and ACC samples. 

36 

 
 
 
 
WP4: Interaction between intercritical rolling & microstructural evolution at high 
temperature: strain partitioning 
Objectives of the WP4 are: 

-  Monitor and map the strain partitioning between austenite and ferrite by combining MFS 

measurements on the single phase and the composite materials with detailed 
microstructural analysis. 
Evaluation the influence of α/γ balance on hot ductility. 

- 
-  Determine how the strain partitioning evolves with: 

o  Deformation temperature (γ/α balance) 
o  C-content 
o  The amount of deformation given 

- 

Link between the strain partitioning and the final microstructure 

o  Microstructural quantification: grain size, distribution and angle boundary nature 
o  Dislocation density quantification by means of EBSD in the final microstructure 

TASK 4.1 MFS mapping of composite and “single phase” materials 

Torsion tests have been carried out by CEIT to investigate the hot deformation behavior of the “single 
phase  materials”  (100%  ferrite  and  100%  austenite).  These  materials  have  the  same  chemical 
compositions as the equilibrium austenite and ferrite in the CMn(0.06%C) and CMn(0.12%C) grades 
at  three  desired  intercritical  temperatures.  The  three  considered  α/γ  intercritical  microstructures 
contain 25%, 50 % and 75% ferrite. Chemical compositions of these “single phase materials” have 
been given in WP2. At the same time, the flow behavior of the reference grades CMn(0.06%C) and 
CMn(0.12%C)  in  the  intercritical  region  (which  can  be  referred  to  as  “composite  materials”)  is 
investigated.  These  torsion  tests  will  allow  analyzing  the  effect  of  phase  balance  on  hot  ductility 
behavior. 

Composite materials 

To obtain different ferrite contents in the final microstructures, different thermo-mechanical cycles 
are designed. Torsion tests have been carried out in both composite materials (CMn(0.06%C) and 
CMn(0.12%C)). After a soaking at 900 ºC for 3 min, the samples were cooled down slowly (1 ºC/s) 
to  various  intercritical  temperatures,  in  order  to  obtain  different  intercritical  ferrite  contents.  A 
holding  at  the  intercritical  temperatures  during  15  min  and  then  direct  quenching  to  room 
temperature are desired to evaluate the ferrite content prior to deformation. Single hit tests until 
rupture  were  performed  and  stress-strain  curves  have  been  obtained  for  all  compositions  (single 
phase materials and composite materials) and entire range of ferrite contents. 

Figure 30: Thermomechanical cycle applied in torsion. 

In  both  CMn(0.06%C)  and  CMn(0.12%C)  grades,  ferrite  fractions  at  different  intercritical 
temperatures before deformation have been measured on the quenched samples. Results of these 
measurements are given in Table 13, while the corresponding microstructures are shown in Figure 

37 

Tdef-low-ɑ(25% ɑ)Tdef-med-ɑ(50% ɑ)Tdef-high-ɑ(75% ɑ)1ºC/sDirectQuench1ºC/s900ºC, 3min15minTemperature(ºC)TimeSingle hit testsuntilrupture 
 
 
31.  As  can  be  observed  from  these  results,  volume  fraction  of  intercritical  ferrite  in  both  of  the 
materials gradually increases with the reduction of deformation temperature. 

Table 13: Measured ferrite fraction in each deformation temperature (torsion tests) 

CMn(0.06%C) (Direct quench) 

CMn(0.12%C) (Direct quench) 

e
t
i
r
r
e
f

w
o
L

e
t
i
r
r
e
f

i

m
u
d
e
M

e
t
i
r
r
e
f

h
g
H

i

(a) 

(d) 

(b) 

(e) 

(c) 

(f) 

Figure  31:  Microstructures  obtained  after  a  quenching  at  different  deformation 
temperatures  and  both  compositions  (torsion  tests):  (a,d)  low  ferrite  content,  (b,e) 
medium ferrite content and (c,f) high ferrite content. 

For  each  deformation  temperature  and  ferrite  content,  MFS  behavior  of  composite  materials 
(CMn(0.06%C) and CMn(0.12%C)) has been determined. In Figure 32a and b stress-strain curves 
are shown for CMn(0.06%C) and CMn(0.12%C) steels, respectively. The curves obtained in the entire 
range  of  deformation  temperature  are  presented  in  each  case,  showing  the  effect  of  both  ferrite 
content  and  deformation  temperature.  The  comparisons  presented  in  Figure  32  suggest  that  the 
hardening caused by the application of lower deformation temperatures is counterbalanced by the 
presence of a higher fraction of ferrite (softer phase than austenite). Therefore, similar tensile curves 
are achieved at the different deformation temperatures and ferrite contents.  

38 

Objective: 25% ɑObjective: 50% ɑObjective:75% ɑSteelTemperature, ºClowɑMeasuredferriteFraction(%)Temperature, ºCmediumɑMeasuredferriteFraction(%)Temperature, ºChighɑMeasuredferriteFraction(%)CMn(0.06%C)78024.17705076070.3CMn(0.12%C)76021.975039.674065.3MFMF 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
(a) 

(b) 

Figure 32: Mean Flow Stress curves (single hit tests until rupture) obtained at the different 
deformation temperatures for the (a) CMn(0.06%C) and (b) CMn(0.12%C) steels. 

Single phase materials 

To confirm the presence of completely ferritic and austenitic microstructures at desired intercritical 
temperatures  (740oC  -780oC),  microstructures  in  quenched  samples  of  all  single  phase  materials 
have been analyzed. The J577, J576 and J575 steels, indeed, contain fully austenitic microstructure 
at the desired temperatures. Nevertheless, the quench microstructures of the J574, J573 and J572 
materials still show a significant fraction of intercritical ferrite. For instance, microstructures obtained 
after  direct  quenching  of  the  J577,  J572  and  J571  steels  are  given  in  Figure  33,  respectively.  At 
780ºC, the microstructure of the J572 material, which is supposed to be fully austenitic, still contains 
a  noticeable  fraction  of  ferrite.  However,  single  phase  microstructures  of  austenite  and  ferrite  at 
desired temperatures are obtained in the J577 and J571 materials.. 

J577 (100% austenite) 

J572 (austenite+ferrite) 

J571 (100% ferrite) 

(a) 

(b) 
Figure  33:  Optical  micrographs  obtained  after  quenching  at  different  deformation 
temperatures and compositions: (a) quenching at 740ºC in J577 steel (100% austenite), 
(b) quenching at 780ºC in J572 steel (austenite + ferrite) and (c) quenching at 760ºC in 
J571 steel (100% ferrite). 

(c) 

For  single  phase  materials,  mechanical  behaviors  at  high  temperatures  also  are  investigated  as 
similar to the composite materials. Stress-strain curves from single hit torsion tests of single phase 
materials and the corresponding composite materials are plotted in the same graphs for comparison 
(Figure  34). In all cases,  the  MFS curves  of composite  materials (CMn-0.06%C  and  CMn-0.12%C 
steels) are  in between the curves of  fully  austenitic and  ferritic materials.  All  stress-strain curves 
show the work hardening behaviors of materials until a peak stress is reached. The gradual reduction 
of stress beyond  the peak value  is due to the softening by dynamic recrystallization  in materials. 
Most of the single phase materials (either ferritic or austenitic steels) rupture at higher strain levels 
than the corresponding composite materials. However, these are not the cases of the J572, J573 and 
J574 steels. At the experimental temperatures, these materials are expected to be pure austenite. 
Nevertheless, as observed previously, the presence of intercritical ferrite in these materials leads to 
early rupture during torsion deformation. As observed previously, the presence of intercritical ferrite 
at the austenite grain boundaries promotes the premature rupture during the torsion deformation. 

39 

05010015020025001234Stress (MPa)Strain780ºC (25% ferrite)770ºC (50% ferrite)760ºC (75% ferrite)CMn-0.06%C05010015020025001234Stress (MPa)Strain760ºC (25% ferrite)750ºC (50% ferrite)740ºC (75% ferrite)CMn-0.12%CMMFF 
 
 
 
 
 
 
(a) 

(d) 

(b) 

(e) 

(c) 

(f) 

Figure 34: Mean Flow Stress curves (single hit tests until rupture) obtained for  the CMn 
steels and pure phase materials (100% austenite and 100% ferrite): (a,b,c) CMn(0.06%C) 
and  (d,e,f)  CMn(0.12%C).  Different  deformation  temperatures  are  applied,  obtaining 
different ferrite content: (a,d) low ferrite, (b,e) medium ferrite and (c,f) high ferrite. 

To  complete  the  study  on  mechanical  behaviors  of  both  single  and  composite  materials  at  high 
temperatures, microstructural  characterization  has  been carried out for the rupture samples after 
the torsion experiments. Figure 35 shows optical micrographs corresponding to the austenitic (J577), 
ferritic (J571) and composite (CMn(0.12%C)) materials after the torsion trials at 740ºC. For both of 
the two single phase materials, microstructural characterization confirms the occurrence of dynamic 
recrystallization during the torsion experiments. In the composite material, softening mechanisms 

40 

05010015020025001020304050Stress (MPa)StrainTemperature 770ºCJ573 (100% AUSTENITE)CMn-0.06%CJ570(100%FERRITE)Lowferrite05010015020025001020304050Stress (MPa)StrainTemperature 760ºCJ574 (100% AUSTENITE)CMn-0.12%CJ570(100%FERRITE)Lowferrite050100150200250010203040Stress (MPa)StrainTemperature 780ºCJ572 (100% AUSTENITE)CMn-0.06%CJ570(100%FERRITE)Medium ferrite0501001502002500102030Stress (MPa)StrainTemperature 750ºCJ576 (100% AUSTENITE)CMn-0.12%CJ571(100%FERRITE)Medium ferrite0501001502002500102030Stress (MPa)StrainTemperature 760ºCJ575 (100% AUSTENITE)CMn-0.06%CJ571(100%FERRITE)High ferrite0501001502002500102030Stress (MPa)StrainTemperature 740ºCJ577 (100% AUSTENITE)CMn-0.12%CJ571(100%FERRITE)High ferrite 
 
 
 
 
 
 
 
also  occur  and  different  microstructures  are  observed  before  (Figure  31)  and  after  deformation 
(Figure 35). 

Figure 36 shows microstructures of the CMn(0.12%C) materials after single hit torsion until rupture 
and  then  water  quenched  to  room  temperature.  No  significant  differences  are  observed  between 
these microstructures, although deformation temperatures for these samples are not similar. When 
a high deformation temperature is applied (low ferrite content), higher ferrite fractions are clearly 
observed  after  rupture  compared  to  the  sample  obtained  before  deformation  (see  Figure  31d). 
Therefore,  it  can  be  stated  that  the  applied  deformation  activates  and  accelerates  phase 
transformation. For example, in the highest deformation temperature, the ferrite fraction is increased 
from 22 to 69% with the application of deformation passes. 

Figure 35: Microstructures observed after rupture in the austenitic (J577), ferritic (J571) 
and composite material (CMn(0.12%C)). Deformation temperature of 740ºC is applied. 

CMn(0.12%C) (Direct quench) 

(a) 

(b) 
Figure  36:  Optical  micrographs  obtained  after  rupture  and  quenched  at  different 
deformation temperatures and CMn(0.12%C) steel (torsion tests): (a) low ferrite content, 
(b) medium ferrite content and (c) high ferrite content. 

(c) 

TASK 4.2 High temperature microstructural evolution in the composite materials 

Isothermal austenite-to-ferrite transformation under para-equilibrium conditions 

The austenite-to-ferrite transformation kinetics in the reference CMn(0.06%C) grade was calculated 
by  ThermoCalc  and  then  experimentally  assessed  by  dilatometry.  The  volume  fraction  of  ferrite 
formed during isothermal holding was calculated under para-equilibrium condition. Para-equilibrium 
assumes that phase transformation occurs without any redistribution of the substitutional alloying 
element M (M = Mn, Ni, Cr, etc.) and the chemical potential of carbon across the interface is constant 
[Kirk58,  Coat72].  Figure  37  compares  the  kinetics  of  austenite  decomposition  measured  by 
dilatometry during isothermal holding and the ones calculated assuming para-equilibrium conditions. 

41 

0501001502002500102030Stress (MPa)StrainTemperature 740ºCJ577 (100% AUSTENITE)CMn-0.12%CJ571(100%FERRITE) 
 
 
 
 
 
Figure  37:  Experimental  (N  refers  to  normal  and  C  refers  to  compression  mode  of 
dilatometer) and theoretical fractions of ferrite formed during isothermal holding. 

The volume fractions of ferrite measured at the end of each isothermal holding step are summarized 
in Table 14. The equilibrium carbon concentrations of ferrite at isothermal temperature, as calculated 
by ThermoCalc, are also added to Table 14. The carbon concentration of austenite was determined 
from following relation: 

𝐶𝛾 =

(𝐶0−𝐶𝛼𝑓𝛼)
𝑓𝛾

 (1) 

where 𝐶𝑖 and 𝑓𝑖 are carbon concentration and volume fraction of phase i (i=α and γ), respectively. 𝐶0 
is carbon concentration of the steel (0.06 wt.%). The values which have uncertainty of ±0.05 wt.% 
are given in Table 14. 

Table  14:  Ferrite  fraction  formed  and  carbon  concentration  of  the  phases  at  isothermal 
temperature. 

Temperature 
(˚C) 
780 

Fraction of 
ferrite 
0.28 

Carbon in ferrite 
(wt.%) 
0.006 

Carbon in austenite 
(wt.%) 
0.080 

770 

760 

0.46 

0.53 

0.007 

0.008 

0.10 

0.15 

As  it  can  be  seen  in  Figure  37,  the  austenite-to-ferrite  transformation  stops  before  reaching 
equilibrium. This can be explained based on the fact that, while the calculation was done assuming 
para-equilibrium condition, the transformation occurs under local-equilibrium conditions. In the local-
equilibrium, it is assumed that the interface migrates under full local-equilibrium with partitioning of 
both  carbon  and  Mn.  Due  to  the  large  difference  in  diffusivity  between  Carbon  and  Mn,  the 
transformation  is  subdivided  into  two  stages.  Within  the  first  stage,  both  C  and  Mn  partition  into 
austenite. However, due to the low diffusion coefficient of Mn, the concentration of Mn in ferrite is 
the  same  as  that  in  austenite  and  a  Mn  spike  in  front  of  the  migrating  interface  is  formed.  The 
transformation rate within this stage is effectively controlled by carbon diffusion and is relatively fast. 
By increasing the Mn concentration in the interface to level that stabilizes austenite in that area, the 
second stage of the transformation starts. Within this stage, the formation of new ferrite requires Mn 
diffusion toward  inside  austenite grains.  The  transformation kinetics  is therefore controlled  by  Mn 
diffusion which is extremely sluggish. The point at which the transformation mode switches from C-
controlled    to  Mn-controlled  mode  is  regarded  as  the  termination  of  the  transformation  [Zuro08, 
Chen14]. 

Effect of intercritical deformation on the  austenite-to-ferrite transformation kinetics 

The effect of intercritical deformation on the austenite-to-ferrite transformation was investigated by 
means of a second set of dilatometry trials. 

The specimens were austenitized at 900 ˚C for 180 s and then cooled to 760 ˚C, 770 ˚C or 780 ˚C 
in which isothermally hold for different times ranging from 0 sec till 5000 s. Then, the specimens 
were subjected to 0.4 deformation under a strain rate of 0.002 sec-1. The process was finished by 
direct quench or followed by keeping for 10 sec, 100 sec or 600 sec at deformation temperature prior 

42 

 
 
 
 
    
 
 
 
 
 
 
 
 
 
quenching. Here,  XXX  –YYY-q refers to a  non-deformed specimen which was  isothermally hold  at 
XXX ˚C for YYY sec and then quenched. Furthermore, XXX-YYY-d-ZZZ-q represents a specimen which 
was isothermally hold at XXX  ˚C for YYY sec, subjected to compression deformation, isothermally 
hold for ZZZ sec and finally quenched. Specimens which were directly quenched after deformation is 
shown with XXX-YYY-d-q. 

Figure 38 represents dilatometry plots of the 770-10-q and 770-3000-q specimens during quenching 
from isothermal holding temperature. Four transformations can be identified during quenching of the 
specimens from 900 ˚C. The first one is formation of ferrite during cooling from 800 ˚C till (isothermal 
temperature)  770  ˚C.  The  second  phase  transformation  happens  within  isothermal  holding 
temperature.  The  fraction  of  ferrite  formed  during  isothermal  temperature  in  the  770-3000-q 
specimen  is  higher  than  770-10-q  specimen.  The  third  phase  transformation,  austenite  to  ferrite 
decomposition, happens as the quenching from isothermal temperature starts. Regarding formation 
higher fraction of ferrite during isothermal temperature in the 770-3000-q specimen compared to 
770-10-q specimen, carbon concentrating of austenite and consequently stability of austenite in the 
former  specimen  is  higher  than  later  specimen.  The  same  phenomenon  explains  the  fact  that 
significant bainite formation (transformation number 4 in Figure 38) is not detected in the 770-3000-
q specimen. The last phase transformation, indicated as transformation number 5 in Fig.2, which is 
austenite  to  martensite  transformation  happens  only  in  the  770-3000-q  specimen.  Accordingly, 
nearly all  of  the  austenite transforms to ferrite  in  the 770-10-q specimen before martensite start 
temperature. Considering that dilatation of the 770-3000-q specimen is less than 770-10-q indicates 
presence of retained austenite in the 770-3000-q specimen. 

Figure 38: Dilatometer plots of specimens 770-10-q and 770-3000-q. 

Figure 39.a and Figure 39.b show optical micrographs of the 770-10-q and 770-3000-q specimens, 
respectively.  Microstructure  of  the  770-10-q  composed  of  a  ferrite  matrix  which  formed  during 
isothermal holding. There are bainite grains which formed during quenching, as indicted in  Figure 
39.a.  Furthermore,  dark  areas  which  represent  bainite  ferrite  grain  appear  in  the  micrograph. 
Microstructure  of  the  770-3000-q  specimen  consists  of  ferritic  grains  with  a  bigger  grain  size 
compared to the 770-10-q specimen. Etched areas in Figure 39.b are related to martensite. 

43 

 
 
 
 
 
 
Figure 39: OpticalOptical micrograph of specimens (a) 770-10-q and (b) 770-3000-q. 

Figure 40: Compression stress-strain curve of specimens 770-10-q and 770-3000-q 

Compression strain-stress plots of specimens are given in Figure 40. Figure 41 shows microstructures 
of  specimens  after  deformation  with  direct  quenching  and  after  10  sec  of  isothermal  holding  and 
after 600 sec of isothermal holding. 

44 

 
 
 
 
 
 
 
 
 
 
Figure  41:  Optical  micrograph  of  specimens  (a)  770-10-0.4d-q,  (b)  770-3000-0.4d-q, 
(c)  770-10-0.4d-10-q,  (d)  770-3000-0.4d-10-q,  (e)  770-10-0.4d-600-q  and  (f)  770-
3000-0.4d-600-q. 

The  ferrite  formation  prior  and  after  deformation  at  different  isothermal  holding  temperatures  is 
shown in Figure 42. This figure also includes equilibrium fractions of ferrite which were determined 
from ThermoCalc. This figure shows that applying compression deformation stimulate austenite to 
ferrite transformation however total fraction of ferrite remains less than equilibrium fraction.  

Applying compression deformation during isothermal transformation makes austenite less stable and 
therefore driving force for austenite decomposition increases. This will lead to further progress of the 
transformation; however, for the amount of applied deformation the total fraction of formed ferrite 
is lower than equilibrium. 

Figure  42:  Isothermal  formation  of  ferrite  before  and  after  deformation  at  different 
intercritical temperatures. 

At CEIT, uniaxial compression tests were carried out using a Bähr DIL805D deformation dilatometer 
in order to analyze the high temperature microstructure evolution. Different thermomechanical cycles 
were designed to obtain various ferrite contents in the final microstructures. After an austenitization 
at  900ºC  for  3  minutes,  the  samples  were  cooled  down  slowly  (1ºC/s)  to  different  deformation 
temperatures, in order to obtain different ferrite contents before deformation. Different deformation 
temperatures are defined depending on the chemical composition. With the purpose of generating 
different ferrite contents of 25, 50 and 75%, in CMn(0.06%C) deformation temperatures of 760, 750 
and 740ºC are selected. However, for CMn(0.12%C), the formation of different fractions of ferrite is 
achieved by deformation temperatures of 730, 720 and 710ºC (25, 50 and 75%, respectively). After 
a  holding  time  of  5s,  the  samples  were  quenched  (Cycle  4B)  at  entire  range  of  deformation 
temperatures, in order to measure ferrite content prior to deformation. Furthermore, a deformation 
step was performed, using different deformation per pass (0.1, 0.25 and 0.4) and a deformation rate 
of 1s-1. Finally, different cooling strategies have been applied. Cycle 4A is followed by a cooling rate 
of 1 ºC/s (air cooling), whereas in Cycle 4C a maximum cooling rate is applied.  

45 

 
 
 
 
 
 
Figure 43: Thermomechanical cycle applied in dilatometer. 

The  ferrite  fractions  formed  in  each  deformation  temperature  and  each  composition  have  been 
measured. The measured ferrite fractions are summarized in Table 15. Ferrite contents are measured 
in the samples quenched from each deformation temperature before deformation (after Cycle 4B). 
In  Figure  44  the  microstructures  observed  in  the  quenched  specimens  are  shown  in  both 
compositions and entire range of deformation temperature. It is clearly observed that the reduction 
of the deformation temperature promotes the increment of the ferrite content gradually (see Figure 
43 and Table 15). 

Table 15: Measured ferrite fraction in each deformation temperature  (dilatometry tests) 

In addition to the characterization of the samples quenched before deformation, the microstructural 
characterization of the microstructures generated after deformation is required. Both quenched and 
slowly cooled (Cycle C and Cycle A, respectively) samples have been analyzed by means of different 
characterization  techniques  such  as  optical  microscopy  and  FEG-SEM.  As  an  example,  the 
microstructures observed in the samples obtained after Cycle A (air cooled samples) are shown in 
Figure  45,  for  both  chemical  compositions  and  deformation  temperatures.  Combinations  between 
non-deformed ferrite and deformed ferrite are clearly distinguished in all the cases. Moreover, the 
formation of pearlite is also observed in the resulting microstructure, mainly in the steel containing 
the highest C content (see measured fractions in Table 16). Figure 45 shows that the deformed ferrite 
is characterized by a significant presence of substructure. It is evident that the content of deformed 
ferrite is increased as the deformation temperature is decreased, due to a higher presence of ferrite 
in  the  microstructure  previous  to  deformation  (high  content  of  ferrite).  Furthermore,  the  detailed 
microstructural characterization performed suggests that the deformed ferrite formed in CMn-012%C 
contains a more relevant substructure compared to CMn-0.06%. 

Table 16: Pearlite fractions measured in each deformation and chemical composition after 
deformation and air cooled. 

Low ferrite 
Medium ferrite  9.5% 
High ferrite 

10.2% 

CMn(0.06%C) 
def0.1 
11.3% 

def0.25 
11.1% 
8.2% 
8.8% 

def0.4 
13% 
10.7% 
12.6% 

def0.1 
18.4% 
16.7% 
17.7% 

CMn(0.12%C) 
def0.25 
14.6% 
16.5% 
17% 

def0.4 
19% 
18.8% 
17.3% 

46 

Tdef-low-ɑ(25% ɑ)Tdef-med-ɑ(50% ɑ)Tdef-high-ɑ(75% ɑ)Air cooled1ºC/s1ºC/sDirectQuenchQuenchCycle4BCycle4CCycle4A1ºC/s900ºC, 3min5s1sTemperature(ºC)Timeɛ, 1s-1Deformation0.1, 0.25, 0.4Objective: 25% ɑObjective: 50% ɑObjective:75% ɑSteelTemperature, ºClowɑMeasuredferriteFraction(%)Temperature, ºCmediumɑMeasuredferriteFraction(%)Temperature, ºChighɑMeasuredferriteFraction(%)CMn(0.06%C)76025.475052.474075.1CMn(0.12%C)73030.472060.171074.1 
 
 
 
 
 
 
 
 
CMn(0.06%C) (Cycle 4B) 

CMn(0.12%C) (Cycle 4B) 

(a) 

(d) 

(b) 

(e) 

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(c) 

(f) 

Figure  44:  Optical  micrographs  obtained  after  a  quenching  at  different  deformation 
temperatures and both compositions (dilatometry tests): (a,d) low ferrite content, (b,e) 
medium ferrite content and (c,f) high ferrite content. 

CMn(0.06%C) (Cycle 4A) 

CMn(0.12%C) (Cycle 4A) 

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(a) 

(d) 

47 

MFMFNon-deformedferrite 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
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(b) 

(e) 

(c) 

(f) 

Figure  45:  Optical  micrographs  obtained  after  air  cooling  at  different  deformation 
temperatures and both compositions (dilatometry tests): (a,d) low ferrite content, (b,e) 
medium ferrite content and (c,f) high ferrite content. 

In order to gain a better understanding of the influence of composition and deformation temperature 
on the final microstructures, the characterization has been completed by means of scanning electron 
microscopy (FEG-SEM). Figure 46 illustrates the morphology of the phases formed in CMn-012%C 
at different deformation temperatures and both air cooling and maximum cooling steps. Concerning 
the air cooled microstructures, the effect of deformation temperature is significant. When the highest 
deformation  temperature  of  760ºC  is  applied  (low  ferrite  content  before  deformation),  polygonal 
ferrite  or  non-deformed  ferrite  is  noticed,  which  is  formed  during  the  final  air  cooling  step. 
Conversely, the microstructure formed at the lowest deformation temperature of 740ºC, in which a 
high  fraction  of  ferrite  is  formed  previous  to  deformation,  is  characterized  by  the  presence  of  a 
considerable  fraction  of  deformed  ferrite  and  low  fraction  of  non-deformed  ferrite.  In  addition  to 
these  phases,  the  formation  of  low  content  of  pearlite  is  also  noticed.  In  reference  to  the 
microstructures  obtained  after  Cycle  C  (quenched  microstructures  after  deformation),  a  mixture 
between martensite and deformed ferrite is observed in both deformation temperatures, as shown 
in Figure 46b and d. As expected, the reduction of the deformation temperature leads to a higher 
presence of deformed ferrite and a lower content of martensite.  

Low ferrite 

High ferrite 

48 

Non-deformedferriteDeformedferrite 
 
 
 
 
 
 
 
 
 
 
)
s
/
C
º
1
(

A
4

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C

)
g
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(a) 

(c) 

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(
C
4

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C

(b) 

(d) 

Figure  46:  FEG-SEM  micrographs  related  to  the  samples  obtained  after  different 
deformation temperatures for CMn(0.12%C) steel: (a,b) 760ºC (low ferrite content) and 
(c,d) 740ºC (high ferrite content). Different cooling strategies have been used: (a,c) air 
cooling and (b,d) quenching. 

Additionally,  the  morphology  of  the  resulting  microstructure  varies  significantly  depending  on  the 
applied  deformation  per  pass.  The  microstructural  characterization  has  been  extended  for  the 
samples  obtained  after  the  different  deformation  values  (0.1,  0.25  and  0.4)  and  both  chemical 
composition. Figure 47a, b and c clearly show the modification of the microstructure caused by the 
application  of  different  deformation  per  pass  in  CMn(0.12%C)  steel    and  high  ferrite  content 
(deformation of 0.1, 0.25 and 0.4, respectively). It is evident that the increment of the deformation 
promotes the formation of a more intense substructure within the deformed ferrite grains. When a 
deformation of 0.1 is applied (see Figure 47a), the substructure is almost negligible, whereas as the 
deformation increases a higher fraction of low angle boundaries is noticed in the final microstructures 
(see Figure 47b and c). Furthermore, the microstructures shown in  Figure 47 suggest that as the 
deformation increases a grain size refinement could be achieved. Similar trends are observed for the 
other chemical composition and ferrite contents (low and medium content). In  Figure 48 FEG-SEM 
micrographs  corresponding  to  CMn(0.06%C)  steel  and  an  intermediate  deformation  temperature 
(medium ferrite content) are presented for different deformation values (0.25 and 0.4 in Figure 48a 
and b, respectively).   

(a) 

(b) 

(c) 

49 

MFDeformedferriteMartensite 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
Figure 47: Influence of deformation. Optical micrographs corresponding to 0.12%C-CMn 
and high ferrite content (75%): (a) ε=0.1, (b) ε=0.25 and (c) ε=0.4. 

Figure 48: Influence of deformation. FEG-SEM micrographs corresponding to 0.06%C-CMn 
and medium ferrite content (50%): (a) ε=0.25 and (b) ε=0.4. 

(a) 

(b) 

In  the  current  study,  the  microstructural  analysis  has  been  extended  to  Electron  Backscattered 
Diffraction (EBSD) technique, in order to further the understanding of the effect of different process 
parameters (chemical composition, deformation temperature, balance between ferrite and austenite 
before deformation and deformation per pass) when intercritical rolling is considered. EBSD scans 
for all the samples quenched before deformation (Cycle B) have been obtained for both steels and 
entire range of deformation temperatures (low, medium and high content of ferrite). In  Figure 50 
Image  Quality  maps  corresponding  to  the  quenched  samples  before  deformation  are  shown,  for 
different deformation temperatures and ferrite contents. The optical micrographs corresponding to 
these cases have already been exposed in Figure 49. 

NbV (Recrystallized austenite) 

NbV (Deformed austenite) 

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(a)  Tdef 750ºC 

(d) Tdef 770ºC 

(b) Tdef 730ºC 

(e) Tdef 750ºC 

50 

Real ferritefraction26% Real ferritefraction31% Real ferritefraction49% Real ferritefraction61%  
 
 
 
 
 
 
 
 
 
 
 
 
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(c) Tdef 720ºC 

(f) Tdef 740ºC 

Figure  49:  Optical  micrographs  obtained  after  a  quenching  at  different  deformation 
temperatures  and  both  austenite  conditions  (Recrystallized  and  Deformed):  (a,d)  low 
ferrite content, (b,e) medium ferrite content and (c,f) high ferrite content. 

Low ferrite 

Medium ferrite 

High ferrite 

(a) 

(b) 

(c) 

(d) 

(e) 
Figure  50:  IQ  maps  corresponding  to  the  samples  quenched  before  deformation  for 
CMn(0.12%C) and different ferrite content: (a) low ferrite, (b) medium ferrite, (c) high 
ferrite. Image Quality distributions corresponding to: (d) low ferrite, (e) medium ferrite, 
(f) high ferrite contents. 

(f) 

In the recent years, the Electron Back Scattered Diffraction (EBSD) technique has been proposed as 
an appropriate tool to discriminate different phases. In the literature, different criterions have been 
applied  to  distinguish  different  type  of  phases:  Image  Quality  (IQ)  parameters  and  grain 
misorientation based methods. The criteria based on Image Quality (IQ), which describes the quality 
of  an  electron  back  scattering  diffraction  pattern,  has  been  taken  into  account  for  the  aim  of 
distinguishing ferrite and martensite in the microstructures obtained after Cycle 4B. The distinction 
of the phases  allows the characterization of  the  formed  ferrite previous to deformation. In  Figure 
50d, e and f, IQ distributions related to the corresponding microstructures shown in Figure 50a, b 
and c are plotted. The consideration of the IQ parameter for the distinction of different type of phases 
is  based  on  the  fact  that  the  diffraction  pattern  given  by  a  deformed  crystal  structure  (such  as 
martensite) has lower quality than the one obtained in a phase free of defects. Therefore, it must be 
considered that the IQ of martensite is noticeably lower than those obtained in ferrite. Figure 50d 
shows a bimodal distribution, reflecting the presence of different type of phases. In this case, a value 

51 

Real ferritefraction76% Real ferritefraction79% 00.010.020.030.040.050.0605001000150020002500Area FractionGrain Average IQCMn-0.12%C-LOW FERRITE00.010.020.030.040.050.0605001000150020002500Area FractionGrain Average IQCMn-0.12%C-MEDIUM FERRITE00.010.020.030.040.050.0605001000150020002500Area FractionGrain Average IQCMn-0.12%C-HIGH FERRITE 
 
 
 
 
 
 
 
 
 
 
 
 
of 1180 was defined for the distinction of ferrite and martensite. This methodology has been applied 
in all the cases, in order to analyze the ferrite previous to deformation. 
On  the  other  hand,  the  crystallographic  characterization  of  the  microstructures  generated  after 
deformation is required. The analysis has been focused on the slowly cooled (Cycle 4A) specimens. 
In  Figure  51,  different  EBSD  maps  corresponding  to  CMn(0.06%C),  deformation  of  0.4  and  high 
ferrite  content  (Image  Quality,  Inverse  Pole  Figure  and  Grain  Boundary  maps)  are  presented, 
illustrating the differences between both ferrite populations. Figure 51a clearly shows that a mixture 
between  deformed  and  non-deformed  ferrite  grains,  in  which  non-deformed  ferrite  grains  are 
characterized by better diffraction patterns (higher IQ). The differences between both ferrite type 
are also evident in the grain boundary map shown in Figure 51c. In all the Grain Boundary Maps, the 
low  and  high  angle  boundary  criteria  are  the  same.  High  angle  boundaries  are  considered  for  a 
misorientation bigger than 15º and low angle boundaries for a misorientation between 4 and 15º. 

(a) 

(b) 
Figure 51: (a) Image Quality, (b) IPF and (c) grain boundary  maps (black lines for high 
angle misorientation criteria, bigger than 15º,and red lines for low angle misorientation 
criteria,  between  4º  and  15º)  corresponding  to  CMn(0.06%C)  steel  and  high  ferrite 
content (deformation of 0.4 and deformation temperature of 740ºC). 

(c) 

Given  that  a  combination  between  different  ferrite  populations  is  observed  in  the  resulting 
microstructures,  a  methodology  able  to  distinguish  both  type  of  ferrite  has  been  developed.  The 
differentiation  of  the  different  type  of  ferrites  allows  the  microstructural  characterization  of  both 
deformed and non-deformed ferrite. In the current study, Grain Orientation Spread (GOS) parameter 
has been considered for the distinction of deformed ferrite and non-deformed ferrite. The present 
analysis  develops  a  methodology  for  quantifying  microstructural  features  in  a  microstructure 
obtained  after  intercritical  rolling,  which  cannot  be  correctly  shown  by  other  microstructural 
techniques (optical microscopy and FEGSEM). The procedure is based on the following steps. Once 
the EBSD scans have been performed, a cleanup should be applied to assimilate any non-indexed 
points into the surrounding neighborhood grains. Afterwards, the elimination of the formed pearlite 
is needed in order to only have a combination between non-deformed and deformed ferrite. For that 
purpose,  the  criteria  considering  IQ  (previously  explained)  has  been  taken  into  account.  The  IQ 
threshold  to  distinguish  pearlite  is  defined  using  the  pearlite  fraction  measured  in  the  optical 
microscope. The value for this threshold varies from samples to sample depending on the quality of 
the scan. In Figure 52a and b, an example related to CMn(0.06%C), deformation of 0.25 and slow 
cooling rate are shown, in which Image Quality maps before and after the elimination of pearlite are 
presented, respectively. 

52 

Non-deformedferriteDeformedferriteNon-deformedferriteDeformedferrite 
 
 
 
 
 
(a) 

(b) 

Figure  52:  Image  Quality  maps  corresponding  to  CMn(0.06%C)  Steel  and  low  ferrite 
content (deformation of 0.25 and cooling rate of 1ºC/s): (a) with pearlite and (b) without 
pearlite. 

The influence of the ferrite content has also been evaluated, as shown in Figure 53, in which Kernel 
maps are presented for both low and high ferrite content (CMn(0.06%C) and deformation of 0.25). 
In Figure 53a and b the images concerning to both ferrite populations (non-deformed and deformed 
ferrite) after pearlite elimination are shown. Looking at these maps, the differences between different 
deformation temperature and ferrite content before deformation is clearly noticed. Moreover, higher 
Kernel Average Misorientation values are measured for the lowest deformation temperature of 740ºC 
and highest ferrite content before deformation, attributed to a higher deformed ferrite fraction. In 
this  case,  KAM  value  of  0.76º  and  1.14º  have  been  obtained,  for  the  lowest  and  highest  ferrite 
content, respectively. 

(a) 

(b) 

Figure  53:  Kernel  maps  (pearlite  is  deleted)  corresponding  to  CMn(0.06%C)  steel  and 
deformation  of  0.25  and  different  ferrite  contents:  (a)  low  ferrite  (deformation 
temperature of 760ºC) and (b) high ferrite (deformation temperature of 740ºC). 

As mentioned previously, for the distinction of deformed and non-deformed ferrite Grain Orientation 
Spread  (GOS)  parameter  has  to  be  considered,  which  is  the  average  deviation  between  the 
orientation of each point in the grain and the average orientation of the grain. In Figure 54a and b 
GOS parameter distributions have been plotted for different ferrite contents, deformation of 0.4 and 
both steels (CMn(0.06%C) and CMn(0.12%C), respectively). During the deformation, a distortion in 
the crystal lattice is introduced, reflected in a higher GOS parameter. Conversely, after deformation 
the formation of new ferrite grains free of deformation occurs, which are characterized by low GOS 
values.  Therefore,  as  shown  in  Figure  54  the  fraction  of  ferrite  previous  to  deformation  affects 
significantly in the GOS parameter distribution. In CMn(0.06%C) steel, when a low fraction of ferrite 
is formed before deformation, a narrow GOS distribution is noticed, with a considerable fraction of 
grains with low GOS values. This could be attributed to the presence of high fraction of non-deformed 
grains (transformed after deformation) in the resulting microstructure. As the ferrite content before 
53 

Pearlite0º2º 
 
 
 
 
 
 
deformation  increases,  GOS  parameter  distributions  become  wider  and  most  of  the  grains  are 
characterized by higher misorientation values, associated with the presence of a higher fraction of 
deformed ferrite. Similar trends could be noticed for CMn(0.12%C) steel (see Figure 54b). This study 
has been performed in each EBSD scan for different compositions, ferrite contents and deformation 
per pass. In each GOS distribution, a minimum GOS value for the definition of a deformed ferrite 
grain  has  been  defined.  Similar  tolerance  angles  have  been  determined  in  all  the  cases,  ranging 
between 0.8º and 3.8º. A common value of 2º was finally defined for the distinction of non-deformed 
and deformed ferrite grains, which is the average of all the tolerance angles estimated in each scan. 
This minimum GOS value is able to distinguish both non-deformed and deformed ferrite in the final 
microstructure for different compositions, ferrite contents before deformation and deformation per 
pass. The measured deformed ferrite fractions by means of EBSD are shown in Table 17. 

(a) 

(b) 

Figure  54:  GOS  parameter  distributions  for  different  ferrite  contents  (low,  medium  and 
high)  and  both  chemical  compositions  (deformation  of  0.4):  (a)  CMn(0.06%C)  and  (b) 
CMn(0.12%C). 

Table  17:  Measured  fractions  of  deformed  ferrite  by  means  of  EBSD  (microstructures 
obtained after Cycle 4A). GOS values higher than 2º has taken into account. 

Deformation of 0.1 

Deformation of 0.25 

Deformation of 0.4 

Low 
Ferrite 

Medium 
Ferrite 

High 
Ferrite 

Low 
Ferrite 

Medium 
Ferrite 

High 
Ferrite 

Low 
Ferrite 

Medium 
Ferrite 

High 
Ferrite 

26.1% 

39.5% 

61.2% 

33% 

63.8% 

79% 

10.6% 

30.8% 

60.5% 

35.1% 

46.9% 

53.7% 

59.4% 

54.5% 

70.7% 

35.1% 

41.7% 

63.2% 

CMn-
0.06%C 

CMn-
0.12%C 

Taking  into  account  the  misorientation  criteria  based  on  Grain  Orientation  Spread,  the  different 
populations have been distinguished in the case shown in  Figure 51.  Figure 55. Figure 55a and b 
show the maps corresponding to non-deformed ferrite population, whilst Figure 55c and d are related 
to the population composed of deformed ferrite. It is clearly observed that the applied deformation 
promotes the modification of the formed ferrite, reflected in the presence of a higher substructure 
(see red lines in the grain boundary map shown in  Figure 55c). Moreover, the differences in both 
populations are clearly observed in Kernel maps. Figure 55b and d show that higher Kernel Average 
Misorientation value is achieved for deformed ferrite compared to non-deformed ferrite. KAM values 
of 1º and 1.21º have been measured, for the deformed and non-deformed grains, respectively. Even 
if these differences are not very important, it seems that the discretization made using GOS values 
is valid for this purpose. 

Non-deformed ferrite 

Deformed ferrite 

54 

00.050.10.150246810Area Fraction Grain Orientation Spread (º)CMn-0.06%CLow ferrite (def0.4)Medium ferrite (def0.4)High ferrite (def0.4)00.050.10.150246810Area Fraction Grain Orientation Spread (ºC)Low ferrite (def0.4)Medium ferrite (def0.4)High ferrite (def0.4)CMn-0.12%C 
 
 
 
 
 
 
 
(a) 

(c) 

(b) 

(d) 

Figure 55: EBSD maps (grain boundary and Kernel maps) corresponding to CMn(0.06%C) 
steel and high ferrite content (deformation of 0.4): (a,b) Non-deformed ferrite population 
(GOS<2º) and (c,d) deformed ferrite population (GOS>2º). 

A quantification of low and high angle misorientation unit sizes was performed for all the samples 
obtained after Cycle A, considering 4º and 15º misorientation criteria. In Figure 56 mean unit sizes 
considering low (a, b and c) and high (d, e and f) misorientation criteria have been plotted for the 
different chemical compositions, ferrite populations (deformed and non-deformed ferrite) and ferrite 
contents (low, medium and high). The evaluation of the mentioned parameters has been performed 
for each deformation value (0.1, 0.25 and 0.4). Concerning the effect of composition, as a general 
trend, it is observed that the increase of C leads to a significant refinement of resulting microstructure 
in most of the cases and both ferrite populations. This trend is clearly evident in both low and high 
misorientation  criteria.  In  terms  of  the  influence  of  deformation,  the  results  shown  in  Figure  56 
suggest that finer microstructures are achieved as the deformation increases. For example, in CMn 
(0.06%C) steel, when medium content of ferrite is formed, mean unit sizes (D15º) of 9.7, 12.5 and 
14.5 µm are measured in deformed ferrite, for the deformation of 0.4, 0.25 and 0.1, respectively. 
Regarding  ferrite  population  corresponding  to  deformed  grains,  the  EBSD  measurements  indicate 
that  slightly  coarser  microstructures  are  obtained  as  the  ferrite  content  prior  to  deformation 
increases, for both misorientation criteria. Conversely, looking at the results related to non-deformed 
population, the mean grain size decreases as ferrite fraction before deformation is increased. This 
trend is more important when the highest deformation values are applied (0.25 and 0.4).  

In order to gain a better understanding of the influence of deformation per pass, in Figure 57 D4º and  
D15º values are shown as a function of ferrite content for both steels and different deformation values. 
With the aim of evaluating the effect of deformation on each ferrite population, deformed and non-
deformed grains have been distinguished (deformed ferrite in Figure 57a and c, and non-deformed 
ferrite in Figure 57b and d). As mentioned previously, deformation value affects noticeably in the 

55 

0º2º 
 
 
 
 
 
resulting ferrite grain size. For both ferrite types (deformed and non-deformed) and misorientation 
criteria (4º and 15º), higher deformation values results in finer microstructures (see Figure 57).  

Low angle misorientation (4º) 

High angle misorientation (15º) 

4
.
0

f
o

n
o
i
t
a
m
r
o
f
e
D

5
2
.
0

f
o

n
o
i
t
a
m
r
o
f
e
D

1
.
0

f
o

n
o
i
t
a
m
r
o
f
e
D

(a) 

(d) 

(b) 

(e) 

(c) 

(f) 

Figure 56: Influence of chemical composition and deformation on mean unit size for both 
ferrite populations (deformed and non-deformed ferrite) and different deformations (0.1, 
0.25 and 0.4). Different misorientation criterion are considered (4º and 15º). 

Low angle misorientation (4º) 

High angle misorientation (15º) 

56 

02468101214Low FerriteMedium FerriteHigh Ferrite4º Mean Unit Size (μm)CMn (0.06%C) Deformed FerriteCMn (0.06%C) Non-deformed FerriteCMn (0.12%C) Deformed FerriteCMn (0.12%C) Non-deformed FerriteDeformation 0.40481216Low FerriteMedium FerriteHigh Ferrite15º Mean Unit Size (μm)CMn (0.06%C) Deformed FerriteCMn (0.06%C) Non-deformed FerriteCMn (0.12%C) Deformed FerriteCMn (0.12%C) Non-deformed FerriteDeformation 0.402468101214Low FerriteMedium FerriteHigh Ferrite4º Mean Unit Size (μm)CMn (0.06%C) Deformed FerriteCMn (0.06%C) Non-deformed FerriteCMn (0.12%C) Deformed FerriteCMn (0.12%C) Non-deformed FerriteDeformation 0.250481216Low FerriteMedium FerriteHigh Ferrite15º Mean Unit Size (μm)CMn (0.06%C) Deformed FerriteCMn (0.06%C) Non-deformed FerriteCMn (0.12%C) Deformed FerriteCMn (0.12%C) Non-deformed FerriteDeformation 0.2502468101214Low FerriteMedium FerriteHigh Ferrite4º Mean Unit Size (μm)CMn (0.06%C) Deformed FerriteCMn (0.06%C) Non-deformed FerriteCMn (0.12%C) Deformed FerriteCMn (0.12%C) Non-deformed FerriteDeformation 0.10481216Low FerriteMedium FerriteHigh Ferrite15º Mean Unit Size (μm)CMn (0.06%C) Deformed FerriteCMn (0.06%C) Non-deformed FerriteCMn (0.12%C) Deformed FerriteCMn (0.12%C) Non-deformed FerriteDeformation 0.1 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
e
t
i
r
r
e
f

d
e
m
r
o
f
e
D

e
t
i
r
r
e
f

d
e
m
r
o
f
e
d
-
n
o
N

(a) 

(c) 

(b) 

(d) 

Figure 57: Influence of deformation per pass on mean unit size for both ferrite populations: 
(a,  c)  deformed  and  (b,d)  non-deformed  ferrite.  Different  misorientation  criterion  are 
considered: (a,b) 4º and (c,d)15º. 

In Figure 58 Kernel Average Misorientation (KAM) values have been plotted for the different chemical 
compositions  and  ferrite  contents  (low,  medium  and  high).  The  analysis  of  these  parameters  has 
been evaluated for entire range of deformation value (see Figure 58a, b and c, for deformation of 
0.1,  0.25  and  0.4,  respectively).  On  the  one  hand,  the  addition  of  C  (from  0.06%C  to  0.12%C) 
promotes  the  increment  of  KAM  value,  showing  the  formation  of  a  microstructure  with  higher 
dislocation density. This effect is more significant as the deformation increases. On the other hand, 
and as expected, higher Kernel values are measured in the deformed ferrite population compared to 
non-deformed population. Finally, the results shown in  Figure 58 confirms that dislocation density 
increases considerably as deformation is increased, associated with the measurement of higher KAM 
angles when the highest deformation of 0.4 is applied.   

Deformation of 0.1 

Deformation of 0.25 

(a) 

(b) 

Deformation of 0.4 

57 

02468101214Low FerriteMedium FerriteHigh Ferrite4º Mean Unit Size (μm)CMn (0.06%C) def0.1CMn (0.12%C) def0.1CMn (0.06%C) def0.25CMn (0.12%C) def0.25CMn (0.06%C) def0.4CMn (0.12%C) def0.4Deformed ferrite0481216Low FerriteMedium FerriteHigh Ferrite15º Mean Unit Size (μm)CMn (0.06%C) def0.1CMn (0.12%C) def0.1CMn (0.06%C) def0.25CMn (0.12%C) def0.25CMn (0.06%C) def0.4CMn (0.12%C) def0.4Deformed ferrite02468101214Low FerriteMedium FerriteHigh Ferrite4º Mean Unit Size (μm)CMn (0.06%C) def0.1CMn (0.12%C) def0.1CMn (0.06%C) def0.25CMn (0.12%C) def0.25CMn (0.06%C) def0.4CMn (0.12%C) def0.4Non-Deformed ferrite0481216Low FerriteMedium FerriteHigh Ferrite15º Mean Unit Size (μm)CMn (0.06%C) def0.1CMn (0.12%C) def0.1CMn (0.06%C) def0.25CMn (0.12%C) def0.25CMn (0.06%C) def0.4CMn (0.12%C) def0.4Non-Deformed ferrite0.511.52Low FerriteMedium FerriteHigh FerriteKernel Average Misorientation (º)CMn (0.06%C) - Deformed FerriteCMn (0.12%C) Deformed Ferrite0.511.52Low FerriteMedium FerriteHigh FerriteKernel Average Misorientation (º)CMn (0.06%C) - Deformed FerriteCMn (0.12%C) Deformed Ferrite 
 
 
 
 
 
 
 
 
 
 
 
(c) 
Figure  58:  Influence  of  chemical  composition  and  deformation  on  Kernel  Average 
Misorientation  for  both  ferrite  populations  (deformed  and  non-deformed  ferrite)  and 
different deformations (0.1, 0.25 and 0.4). 

TASK 4.3 Strain partitioning: law of mixtures 

In order to gain a better understanding concerning the deformation behavior and strain partitioning 
between  phases,  flow  behavior  have  been  analyzed  in  austenitic  and  ferritic  region  for  both 
CMn(0.06%C) and CMn(0.12%C). For that purpose, the stress-strain curves have been obtained in 
both austenitic and ferritic region, by means of single hit tests until rupture in torsion tests. These 
torsion  tests  will  allow  evaluating  the  effect of  different  austenite-ferrite balances  on  hot  ductility 
behaviour. In Figure 59a and b, the thermomechanical schedules applied in the austenitic and ferritic 
region are shown, respectively. In Figure 59 the cycle corresponding to two-phase microstructure is 
observed.  In  order  to  obtain  different  austenite  structures,  three  different  austenitization 
temperatures of 900, 950 and 1000ºC were used and after a holding time of 3min, single hit tests 
until rupture were carried out (see Figure 59a). Conversely, in the ferritic region, the samples were 
heated up to four different deformation temperatures of 700, 650, 600 and 550ºC (see Figure 59b). 
Afterwards, as in ferritic region, single hit tests until rupture were performed. Figure 59c shows the 
cycle followed for the intercritically deformed samples. The deformation temperatures in this case 
were defined by the ferrite fraction sought. They depend on the chemical composition and for ferrite 
fraction of 25, 50 and 75% are of 780, 770 and 760ºC for the CMn(0.06%C) steel and of 760, 750 
and 740ºC for the CMn(0.12%C) steel, respectively. 

(a) 

(b) 

58 

0.511.52Low FerriteMedium FerriteHigh FerriteKernel Average Misorientation (º)CMn (0.06%C) - Deformed FerriteCMn (0.12%C) Deformed FerriteDirectQuench1ºC/sTdef, 3minTemperature(ºC)TimeSingle hit testsuntilruptureTdef900, 950, 1000ºCDirectQuench1ºC/sTdefTemperature(ºC)TimeSingle hit testsuntilruptureTdef700, 650, 600, 550ºC 
 
 
 
 
Figure 59: Thermomechanical cycles applied in torsion for (a) austenitic, (b) ferritic and 
(c) intercritical region. 

(c) 

The flow stress behaviour within the intercritical region is considered a competition between thermal 
softening and work hardening associated with austenite. Figure 60a and b show stress-strain curves 
obtained  in  the  intercritical  region,  for  CMn(0.06%C)  and  CMn(0.12%C)  steels,  respectively.  The 
curves  measured  in  the  entire  range  of  deformation  temperature  can  be  compared  in  each  case, 
showing the effect of both ferrite content and deformation temperature. The results shown in Figure 
60  indicate  that  the  hardening  caused  by  the  application  of  lower  deformation  temperatures  is 
counterbalanced  by  the  presence  of  a  higher  fraction  of  ferrite  (softer  phase  than  austenite). 
Therefore, similar stress-strain curves are obtained at the different deformation temperatures and 
ferrite contents.  

Figure 60: Mean Flow Stress curves (single hit tests until rupture) obtained at the different 
deformation temperatures for the (a) CMn(0.06%C) and (b) CMn(0.12%C) steels. 

In Figure 61 the evolution of saturation stress values is plotted for both CMn steels as a function of 
deformation temperature in the different regions (ferritic, austenitic and intercritical). In both ferritic 
and  austenitic  microstructures,  saturation  stress  decreases  as  the  deformation  temperature 
increases. Furthermore, higher stress values are measured for CMn (0.12%C), due to the higher C 
content. 

59 

Tdef-low-ɑ(25% ɑ)Tdef-med-ɑ(50% ɑ)Tdef-high-ɑ(75% ɑ)1ºC/sDirectQuench1ºC/s900ºC, 3min15minTemperature(ºC)TimeSingle hit testsuntilrupture05010015020025001234Stress (MPa)Strain780ºC (25% ferrite)770ºC (50% ferrite)760ºC (75% ferrite)CMn-0.06%C05010015020025001234Stress (MPa)Strain760ºC (25% ferrite)750ºC (50% ferrite)740ºC (75% ferrite)CMn-0.12%C 
 
 
 
 
(a) 

(b) 

Figure  61:  The  evolution  of  saturation  stress  values  as  a  function  of  deformation 
temperature strain rate for both steels: (a) CMn(0.06%C) and (b) CMn(0.12%C 

In  order  to  understand  the  evolution  of  the  microstructure  during  the  rupture  test,  the 
microstructures  quenched  at  different  stages  were  analyzed.  In  Figure  62a,b  and  c,  the  initial 
microstructures corresponding to ε=0 are shown, for 760, 750 and 740ºC deformation temperatures, 
respectively (ferrite fractions of 22, 40 and 65% are measured prior to deformation). In addition, 
the quenched samples obtained after single hit rupture test are also presented in Figure 62d, e and 
f. Comparing the microstructures before and after rupture, a higher fraction of ferrite is observed in 
the final microstructure, reflecting that strain induced ferrite transformation is occurring during the 
deformation. 

Quenched before rupture test 

Quenched after rupture test 

(a)  Tdef 760ºC 

(d) Tdef 760ºC 

(b) Tdef 750ºC 

(e) Tdef 750ºC 

60 

010020030040050050055060065070075080085090095010001050Saturation Stress (MPa)Deformation Temperature (ºC)CMn(0.06%C)-Ferritic RegionCMn(0.06%C)-Intercritical regionCMn(0.06%C)-Austenitic Region010020030040050060050055060065070075080085090095010001050Saturation Stress (MPa)Deformation Temperature (ºC)CMn(0.12%C)-Ferritic RegionCMn(0.12%C)-Intercritical regionCMn(0.12%C)-Austenitic Region 
 
 
 
 
 
 
 
 
(c) Tdef 740ºC 

(f) Tdef 740ºC 

Figure 62: Quenched microstructures for CMn(0.12%C) obtained before and after single 
hit rupture tests and different deformation temperatures (760, 750 and 740ºC): (a,d) low, 
(b,e) medium and (c,f) high ferrite content. 

In  addition,  in  order  to  quantify  the  evolution  of  transformed  ferrite  during  deformation,  an 
intermediate quenching have been performed for selected samples. The microstructure have been 
characterized after a strain of 0.65 and 0.75 (corresponding to peak stress), for CMn(0.06%C) and 
CMn(0.12%C),  respectively.  Figure  63  shows  the  quenched  microstructures  obtained  for  both 
chemistries and after a strain of 0.65 and 0.75. As expected, higher ferrite fractions are noticed at 
low  strain  values,  compared  to  the  ferrite  contents  prior  to  rupture  tests.  For  example,  in 
CMn(0.06%C) steel, when deformation temperature is 780ºC, the ferrite fraction increases from 24% 
to  61%  during  deformation.  Therefore,  it  can  be  claimed  that  strain  induced  transformation 
mechanism is activated at early stages of deformation. This fact would also minimize the differences 
between the intercritical stress-strain curves as shown in Figure 60. 

CMn(0.06%C) 

CMn(0.12%C) 

(a)  Tdef 780ºC (ε=0.65) 

(c) Tdef 760ºC (ε=0.75) 

(b) Tdef 770ºC (ε=0.65) 

(d) Tdef 750ºC (ε=0.75) 

Figure  63:  Quenched  microstructures  for  (a,b)  CMn(0.06%C)  and  (c,d)  CMn(0.12%C) 
obtained  after  a  strain of  0.65  and  0.75.  Different  deformation  temperatures  have  been 
included (760 and 750ºC): (a,c) low and (b,d) medium ferrite content. 

61 

Ferritefraction61%Ferritefraction51%Ferritefraction73%Ferritefraction55% 
 
 
 
 
 
 
 
With the aim of describing the hot deformation behaviour in all the regions (ferritic, austenitic and 
two-phase), constitutive equations listed in Table 18 have been adopted. The total flow stress can 
be is considered the sum of contributions corresponding to work hardening (WH), dynamic recovery 
(DRV)  and  dynamic  recrystallization  (DRX):  In  the  current  work,  dynamic  recrystallization 
mechanisms are not taken into account. From experimental stress-strain curves, the material fitting 
parameters have been determined in each region and for both steels (shown in Table 19). It is worth 
noting that values for some of the parameters (especially C and Q) differ significantly depending on 
the region of analysis. That is in the intercritical region, they are orders of magnitude lower. This is 
well reflected in the saturation stress charts plotted in Figure 61. Similar trends are also reported in 
the literature [Moha15]. 

Table 18 Constitutive equations obtained for describing hot deformation behavior. 

0.5

                                 …………    ……… …(2) 

𝜎𝑡𝑜𝑡𝑎𝑙 = 𝜎𝑊𝐻+𝐷𝑅𝑉 + 𝜎𝐷𝑅𝑋                                                                                          (1) 
where σWH, σDRV and σDRX is the contribution related to work hardening, dynamic 
recovery and dynamic recrystallization 
𝜎 = 𝜎0 + (𝜎𝑠𝑎𝑡 − 𝜎0) ∙ [1 − 𝑒(−𝛽∙𝜀)]
where σ0 is the strain-independent term resulting from the extrapolation of the flow stress 
at ε=0. σsat and β, are the saturation stress and fitting parameter depending on strain rate 
and temperature, respectively. 
𝑄
𝑅∙𝑇⁄
𝑍 = 𝜀 ∙̇ 𝑒(
                                                                                                            (3) 
Where Z is the Zener Hollomon parameter, 𝜀̇ is the strain rate, Q is the apparent activation 
energy for deformation and R universal gas constant. 
𝑍 = 𝐶 ∙ [sinh(𝛼 ∙ 𝜎𝑠𝑎𝑡)]𝑛                                                                                            (4) 
C, α and n are the material constants. 

)

Table 19 Material constants obtained in each region (ferritic, austenitic and intercritical) 
for both chemistries. 

CMn(0.06%C) 

CMn(0.12%C) 

Ferritic 
Region 

Austenitic 
Region 

Intercritic
al Region 

Ferritic 
Region 

Austenitic 
Region 

Intercritical 
Region 

30 

12  

40 

30 

12  

40 

1.89043E+
11 

3.98532E+
13 

57.03886 

8.84076E
+13 

7.8299E+13 

59.1294 

4.9 

6.4 

5 

4.9 

6.4 

5 

0.0054  

0.0054  

0.0054  

0.0054  

0.0054  

0.0054 

232738 

308095 

39871 

275240 

320625 

48842 

25-
0.7log(Z) 

12-
0.59log(Z) 

6-
0.59log(Z
) 

30-
0.7log(Z) 

12.5-
0.59log(Z) 

6-
0.59log(Z) 

σ0  
(MPa) 

C 

n 

α  
(MPa-1) 

Q 

β 

Figure  64a  and  b  shows  the  comparison  between  experimental  and  predicted  flow  stress  curves 
corresponding to fully ferritic samples, for CMn(0.06%C) and CMn(0.12%C) steels, respectively. The 
curves  obtained  in  the  entire  deformation  temperatures  (700,  650,  600  and  550ºC)  have  been 
included in the graphs. Moreover, as shown in Figure 64c and d, the prediction ability of the developed 
constitutive equations has been evaluated for the stress-strain curves achieved in austenitic region. 
Finally, the experimental and modeled stress-strain curves obtained in the intercritical region have 
also been compared in Figure 64e and f. Comparing experimental and predicted curves, a reasonable 

62 

 
 
 
 
 
 
agreement  is  achieved  for  both  type  of  microstructures  (ferritic,  austenitic  and  austenitic/ferritic) 
and both low carbon steels with different carbon contents. 

(a) 

(b) 

(c) 

(d) 

(e) 

(f) 

Figure  64:  Comparison  between  experimental  curves  and  predicted  curves  for  (a,b) 
ferritic, (c,d) austenitic and (e,f) intercritical regions: (a,c) CMn (0.06%C) and (b,d) CMn 
(0.12%C). 

Considering the developed constitutive equations for both fully ferritic and austenitic microstructures, 
the  stress-strain  curves  corresponding  to  the  deformation  temperatures  defined  within  the 
intercritical  region  can  be  estimated.  From  the  flow  curves  obtained  in  both  microstructures  and 
taking into account law of mixtures, the mechanical behavior for different ferrite-austenite balances 
could be achieved. 

In the current study, as a first approach, uniform strain has been considered and the iso-strain model 
was used to determine the composite materials mechanical behavior. The hypothesis of this model 

63 

010020030040050060000.10.20.30.40.50.60.7Stress (MPa)StrainCMn(0.06%C)-550ºC-ExperimentalCMn(0.06%C)-550ºC-ModelCMn(0.06%C)-600ºC-ExperimentalCMn(0.06%C)-600ºC-ModelCMn(0.06%C)-650ºC-ExperimentalCMn(0.06%C)-650ºC-ModelCMn(0.06%C)-700ºC-ExperimentalCMn(0.06%C)-700ºC-Model010020030040050060000.10.20.30.40.50.60.7Stress (MPa)StrainCMn(0.12%C)-550ºC-ExperimentalCMn(0.12%C)-550ºC-ModelCMn(0.12%C)-600ºC-ExperimentalCMn(0.12%C)-600ºC-ModelCMn(0.12%C)-650ºC-ExperimentalCMn(0.12%C)-650ºC-ModelCMn(0.12%C)-700ºC-ExperimentalCMn(0.12%C)-700ºC-Model05010015020000.10.20.30.40.50.60.7Stress (MPa)StrainCMn(0.06%C)-900ºC-ExperimentalCMn(0.06%C)-900ºC-ModelCMn(0.06%C)-950ºC-ExperimentalCMn(0.06%C)-950ºC-ModelCMn(0.06%C)-1000ºC-ExperimentalCMn(0.06%C)-1000ºC-Model05010015020000.10.20.30.40.50.60.7Stress (MPa)StrainCMn(0.12%C)-900ºC-ExperimentalCMn(0.12%C)-900ºC-ModelCMn(0.12%C)-950ºC-ExperimentalCMn(0.12%C)-950ºC-ModelCMn(0.12%C)-1000ºC-ExperimentalCMn(0.12%C)-1000ºC-Model05010015020025000.10.20.30.40.50.60.7Stress (MPa)StrainCMn(0.06%C)-760ºC-ExperimentalCMn(0.06%C)-760ºC-ModelCMn(0.06%C)-770ºC-ExperimentalCMn(0.06%C)-770ºC)ModelCMn(0.06%C)-780ºC-ExperimentalCMn(0.06%C)-780ºC-Model05010015020025000.10.20.30.40.50.60.7Stress (MPa)StrainCMn(0.12%C)-740ºC-ExperimentalCMn(0.12%C)-740ºC-ModelCMn(0.12%C)-750ºC-ExperimentalCMn(0.12%C)-750ºC-ModelCMn(0.12%C)-760ºC-ExperimentalCMn(0.12%C)-760ºC-Model 
 
 
 
 
 
 
 
is that the two constituents (austenite and ferrite) suffer the same local strain in respond to a global 
strain applied in the composite material.  
εMix = εγ = εα                                                                                                               (5) 

σMix = fγ ∙ σγ + (1 − fγ) ∙ σα                                                                                           (6) 

Using  the  Equation  (6),  the  experimentally  obtained  results  (measured  ferrite  fractions  from 
quantitative  metallography)  and  the  calculated  stress-strain  curves  obtained  for  one  phase 
component,  the  flow  curve  of  the  composite  material  for  both  compositions  can  be  calculated.  In 
Figure  65  predicted  flow  curves  corresponding  to  both  austenitic  and  ferritic  region,  as  well  as 
calculated flow curves obtained by the assumption of iso-strain are shown. Additionally, the stress-
strain  curves  considering  constitutive  equations  have  been  included  in  the  entire  deformation 
temperature range. Analyzing the comparisons plotted in Figure 65, it is observed that the predicted 
flow curves (assuming law of mixtures) are close to the modeled flow curves (considering constitutive 
equations). 

(a) 

(d) 

(b) 

(e) 

64 

05010015020025030000.511.52Stress (MPa)StrainCMn 0.06%C-780ºCAustenite ModelFerrite ModelMix Law of MixturesMix Model05010015020025030000.511.52Stress (MPa)StrainCMn 0.06%C-770ºCAustenite ModelFerrite ModelMix Law of MixturesMix Model05010015020025030035000.511.52Stress (MPa)StrainCMn 0.12%C-750ºCAustenite ModelFerrite ModelMix Law of MixturesMix Model 
 
 
 
 
 
(c) 

(f)  

Figure 65 Comparison between flow intercritical curves obtained by the law of mixtures 
(assumption of iso-strain) and the stress-strain curves considering constitutive equations 
for  both  steels  (a,b,c  CMn(0.06%C)  and  d,e,f  CMn(0.12%C))  and  different  deformation 
temperatures. 

65 

05010015020025030000.511.52Stress(MPa)StrainCMn 0.06%C-760ºCAustenite ModelFerrite ModelMix Law of MixturesMix Model05010015020025030035000.511.52Stress (MPa)StrainCMn 0.12%C-740ºCAustenite modelFerrite modelMix Law of mixturesMix model 
 
 
 
 
 
 
66 

 
 
 
 
WP5:  Effect  of  strain  path 
microstructure and mechanical properties 
Objectives of the WP5 are: 

in  the 

intercritical  temperature  region  on 

- 
- 

The effect of the amount of strain in the intercritical region on the strength-toughness map. 
The effect of the strain distribution in the intercritical region over one or more deformation 
passes. 

TASK 5.1 Effect of strain 

To  investigate  the  effect  of  strain  in  the  intercritical  region  on  microstructures  and  properties  of 
materials, laboratory rolling trials with reductions of 10% and 40% in the intercritical region of the 
CMn(0.06%C) reference materials were carried on. In the first phase, single intercritical deformations 
of  10%  and  40%  thickness  reductions  are  implemented  in  the  same  manner  as  for  the  single 
intercritical deformation of 25% thickness reduction in WP3. Details of these rolling schedules are 
given in the Table 20 and Table 21. To obtain the same final thickness of 12 mm, initial thicknesses 
of the plates before intercritical rolling with different strain levels are adjusted. In the second phase, 
multi rolling passes with the same total deformation of 40% thickness reduction in the intercritical 
region was proposed to be carried on. Nevertheless, because of the larger fraction of ferrite at lower 
intercritical temperature, single rolling passes of 40% thickness reduction could not be completely 
implemented.  Specifically,  single  rolling  pass  of  40%  thickness  reduction  at  745oC  to  have  50% 
intercritical ferrite was successfully  done.  However,  rolling force  and torque measured during this 
trial  were  already  exceeding  the  upper  limits  of  the  laboratory  rolling  mill.  Additionally,  a  single 
intercritical rolling pass of 40% thickness reduction is not practical for industrial rolling. Therefore, 
instead of splitting the total thickness reduction of 40% into multi-rolling passes, 1, 2 and 3 rolling 
passes  of  10%  thickness  reduction  per  pass  were  implemented.  The  rolling  schedule  for  3  pass 
intercritical rolling with 10% thickness reduction per pass is given in  Table 22. The total thickness 
reduction for 3 intercritical rolling passes is equivalent to 27%, and thus mechanical properties of 
these samples are better compared with the ones of 25% single intercritical rolling.  

Table 20. Rolling schedule for single intercritical pass of 10% thickness reduction. 

Table 21. Rolling schedule for single intercritical pass of 40% thickness reduction. 

67 

Th(in)Th(out)WidthLength(mm)(mm)(mm)(mm)reheating180.080.00%0.00120170280.068.015%0.191200120200368.057.116%0.201180120238457.148.016%0.201160238143548.040.316%0.201140238170640.333.116%0.231130238207733.127.816%0.20890238247827.823.615%0.19870238291923.619.816%0.208502383461019.816.517%0.228352384161116.513.319%0.24820238514intercritical pass13.31210%0.12795-735238571PassReductionTrue strainTarget T (°C)1250 °C - 30 min 
 
 
 
 
Table 22. Rolling schedule for 3 intercritical passes of 10% thickness reduction per pass. 

TASK 5.2 Effect of strain distribution 

The redistribution of the total strain of 40% thickness reduction into multi rolling passes is not feasible 
to  be  carried  on  by  lab  rolling  trials.  However,  the  effect  strain  distribution  on  microstructure 
evolution during these intercritical deformations could be simulated by applying thermomechanical 
cycles  on  plasto-dilatometry  or  hot  torsion  experiments.  Figure  66  shows  two  thermomechanical 
cycles for intercritical deformations of 0.4 in single and two deformation passes of the CMn(0.06%C) 
reference material which have been implemented at CEIT. All samples are reheated to 900oC and 
holding in 3 min to form fully austenitic initial microstructure. Subsequently, samples are cooled to 
750oC with the cooling rate of 1oC/s for the microstructure of 50% intercritical ferrite. Deformations 
of 0.4 and two passes of 0.2 are applied, before the air cooling to room temperature. Interpass time 
of 1 and 5 s are also considered in the case of multi intercritical deformation pass. 

68 

Th(in)Th(out)WidthLength(mm)(mm)(mm)(mm)reheating180.0080.00%0.00120170280.0068.015%0.191200120200368.0057.116%0.201180120238457.1250.216%0.151160238137650.1942.615%0.19890238161742.6335.816%0.20870238191835.8429.717%0.22850238230929.7424.119%0.248352382841024.1020.017%0.22820238343intercritical pass20.001240%0.59795-7352385711250 °C - 30 minPassReductionTrue strainTarget T (°C)Th(in)Th(out)WidthLength(mm)(mm)(mm)(mm)reheating180.080.00%0.00120170280.068.015%0.191200120200368.057.116%0.201180120238457.148.016%0.201160238143548.040.816%0.191140238168640.834.316%0.20890238200734.329.115%0.19870238235829.124.516%0.20850238280924.520.317%0.228352383371020.316.519%0.24820238416intercritical passes16.514.810%0.12795-735238463intercritical passes14.813.310%0.12795-735238514intercritical passes13.312.010%0.12795-735238571PassReductionTrue strainTarget T (°C)1250 °C - 30 min 
 
 
 
 
(a) 

(b) 

Figure  66:  Thermomechanical  schedules  applied  in  the  deformation  dilatometer:  (a) 
Deformation pass of 0.4 and (b) two deformation passes of 0.2 (interpass time of 0, 1 and 
5s). 

Figure  67  presents  the  microstructures  obtained  after  applying  only  one  deformation  pass  of  0.4 
(Figure 67 a) and after using two deformation passes and different interpass times (figure b, c and 
b, interpass time of 0, 1 and 5s, respectively). It is shown that modifying the deformation schedule 
and  interpass  time  (between  deformation  passes)  considerably  affects  the  resulting  air  cooled 
microstructures. A higher fraction of deformed ferrite is clearly distinguished when two deformation 
passes are applied compared to the microstructure formed after a deformation of 0.4. This could be 
attributed  to  the  transformation  taking  place  after  the  first  deformation  pass  of  0.2  and  during 
interpass time.  In order to confirm this fact,  two interrupted quenching were performed after the 
interpass  time  of  1  and  5s.  Figure  68a  and  b  illustrates  the  differences  between  the  quenched 
microstructures  obtained  after  an  interpass  time  of  1  and  5s,  respectively.  A  significantly  higher 
fraction of ferrite was quantified for the highest interpass time (58 and 78% of ferrite, in 1s and 5s, 
respectively), reflecting that austenite/ferrite phase transformation is occurring during holding time 
after deformation pass of 0.2. 

The obtained samples were also characterized by EBSD technique. Microstructures were analysed by 
the  same  method  as  for  the  ones  in  Task  4.2.  The  Kernel  maps  corresponding  to  different  strain 
distributions and interpass times are shown in Figure 69. The fractions of deformed ferrite determined 
are  also  indicated.  Lower  Kernel  values  are  measured  when  the  applied  thermomechanical  cycle 
consists of one deformation pass of 0.4, due to the presence of a lower fraction of deformed ferrite. 
Higher  Kernel  values  are  measured  when  two  0.2  deformation  passes  are  performed,  due  to  the 
formation of additional ferrite grains during the interpass time. The Kernel values increase further as 
the holding time increases, because of a higher deformed ferrite fraction in the final microstructure. 

The  EBSD  characterization  procedure  developed  allows  distinguish  non-deformed  ferrite  and 
deformed  ferrite.  Ferrite  grains  with  GOS  values  lower  than  2º  are  considered  as  non-deformed 
ferrite (transformed during final air cooling step). The ferrite grains located between 2 and 4º are 
assumed to be ferrite formed during the interpass time and then deformed in the last deformation 
pass  of  0.2.  GOS  values  higher  than  4º  correspond  to  heavily  deformed  ferrite  grains  (two 
deformation passes of 0.2). 

Deformation of 0.4 (one deformation pass) 

Deformation of 0.2 (two deformation pass) 
and no interpass time 

69 

Tdef-med-ɑ(50% ɑ)= 750ºCAir cooled1ºC/s1ºC/s1ºC/s900ºC, 3minTemperature(ºC)Timeɛ=0.4 (1s-1)750ºCAir cooled1ºC/s1ºC/s1ºC/s900ºC, 3minTemperature(ºC)Timeɛ=0.2 (1s-1)ɛ=0.2 (1s-1)tip(0,1,5 s)Quenching 
 
 
 
(a) 
Deformation of 0.2 (two deformation pass) 
and interpass time of 1s 

(b) 
Deformation of 0.2 (two deformation pass) 
and interpass time of 5s 

(c) 

(d) 

Figure  67:  Optical  micrographs  obtained  in  intercritically  deformed  and  air  cooled 
microstructures: (a) Deformation pass of 0.4, (b,c,d) two deformation passes of 0.2 and 
interpass times of 0, 1 and 5s, respectively. 

(a) 

(b) 
Figure  68:  Optical  micrographs  obtained  in  intercritically  deformed  and  quenched 
microstructures obtained after an interpass time of (a) 1 and (b) 5s. 

Deformation of 0.4 (one deformation pass) 

Deformation of 0.2 (two deformation pass) 
and no interpass time 

70 

58% deformedferrite78% deformedferrite 
 
 
 
 
 
 
 
 
(a) 
Deformation of 0.2 (two deformation pass) 
and interpass time of 1s 

(b) 
Deformation of 0.2 (two deformation pass) 
and interpass time of 5s 

(c) 

(d) 

Figure  69:    Kernel    maps  obtained  in 
intercritically  deformed  and  air  cooled 
microstructures: (a) Deformation pass of 0.4, (b,c,d) two deformation passes of 0.2 and 
interpass times of 0, 1 and 5s, respectively. 

TASK 5.3 Microstructural characterization 

To evaluate the effect of strain on microstructure of intercritical rolled samples, optical micrographs 
of the two samples which were intercritically rolled to reduce the thickness of 25% and 40% at the 
same temperature of 745oC and then cooled in warm water at 60oC have been analysed. Figure 70 
shows microstructures captured at the mid-thickness regions of the two corresponding samples. As 
the  thickness  reduction  increases  from  25%  to  40%,  ferrite  grains  become  more  elongated. 
Additionally,  sub-grain  structure  in  deformed  ferrite  grains  is  more  pronounced  and  easier  to 
recognizable  in  the  sample  of  40%  thickness  reduction  than  in  the  sample  of  25%  thickness 
reduction. 

8952g - Set A ACC. 
Intercritical pass 25% at 745°C followed by 
simulation of ACC (laminar cooling) 

8953g - Set A ACC. 
Intercritical pass 40% at 745°C followed by 
simulation of ACC (laminar cooling) 

HV=178 

HV=192 

Figure 70: Effect of reduction rate (25% vs 40%) on microstructure of intercritical rolled 
plates 

Because the strain distribution and strain accumulation during intercritical rolling also have effects 
on microstructures and mechanical properties of the material at room temperature, microstructures 
of  intercritically  rolled  samples  by  1,  2  and  3  rolling  passes  are  analysed.  Figure  71  shows 
microstructures at the mid thickness regions in samples of intercritical rolling at 770oC and 745oC 

71 

44% deformedferrite66% deformedferrite66% deformedferrite90% deformedferrite 
 
 
 
 
 
 
 
after 1, 2 and 3 rolling passes of 10% thickness reduction and finally cooling to room temperature 
by air. Not many different features could be observed among these microstructures. Nevertheless, 
deformation of multi rolling  passes produces  slightly finer  microstructures than the ones of  single 
pass.  Elongated  ferrite  grains  are  more  frequently  observed  in  samples  of  lower  deformation 
temperatures. 

Figure  72  shows  IPF  maps  measured  by  EBSD  technique  at  the  mid-thickness  regions  of  the  3 
samples  which  were  intercritically  rolled  at  745oC  by  1,  2  and  3  rolling  passes  and  subsequently 
cooling to room temperature by air. All of the three microstructures contain homogeneous and fine 
ferrite  grains.  With  the  high  angle  grain  boundary  of  15°  disorientation,  the  averaged  grain  size 
calculated for these microstructures are in the range of 12-17 μm. In-grain orientation gradients, 
although are not pronounce, are also visible in these microstructures, particularly in the sample of 3 
intercritical rolling passes. The observed in-grain orientation gradient indicates the presence of low 
angle  grain  boundaries  in  ferrite  grains  which  are  induced  by  plastic  deformation  during  the 
intercritical rolling. 

Crystallographic textures measure at  the  mid-thickness regions of  the  corresponding samples are 
given in Figure 73. As can be seen in the ODF sections φ2 = 45o, textures of all three samples are 
rather weak. The maximum intensities in these textures are about 2.5 multi random times. However, 
the highest intensity peaks in different textures do not located at the same texture component. As 
the number of intercritical rolling passes increases from 1 to 3 passes, typical texture components 
corresponding to deformed and subsequent transformed austenite are less pronounced. At the same 
time, intensities of the texture components corresponding to deformed ferrite become stronger. 

FRT = 770oC 

FRT = 745oC 

s
s
a
p
-
1

s
e
s
s
a
p
-
2

(a) 

(b) 

(c) 

(d) 

72 

 
 
 
 
 
 
 
 
 
s
e
s
s
a
p
-
3

(e) 

(f) 

Figure 71: Microstructures of intercritically rolled samples, after deformation by 1 (a, b), 
2 (c, d) and 3 (e, f) rolling passes of 10% thickness reduction per pass at 770oC (a, c, e) 
and 745oC (b, d, f) and air cooling to room temperature. 

Figure 72: Inverse Pole Figure (IPF) maps measured at the mid-thickness regions of the 
intercritically  rolled  samples  at  745oC  after  1,  2  and  3  rolling  passes  of  10%  thickness 
reduction. 

Figure  73:  Crystallographic  textures  measured  at  the  mid-thickness  regions  of  the 
intercritically  rolled  samples  at  745°C  after  1,  2  and  3  rolling  passes  of  10%  thickness 
reduction. 

TASK 5.4 Characterization of mechanical properties 

Characterization of mechanical properties were implemented on intercritically rolled samples which 
were produced in task 5.1.  Specimen geometries for  tensile and toughness measurements in this 
task is similar to the ones in WP3.  Specifically, tensile and toughness properties of these samples 
were measured, and the results are given as below. 

To  evaluate  the  effect  of  strain  magnitudes  on  mechanical  properties,  tensile  and  toughness 
measurements  have  been  carried  on  with  the  two  intercritically  rolled  samples  of  25%  and  40% 
thickness  reduction  at  the  same  intercritical  temperature  of  745oC.  Tensile  properties  of  the  two 
samples are shown in the Figure 74, while results of toughness measurements are given in Figure 
73 

 
 
 
 
 
 
 
75. As the thickness reduction for a single intercritical rolling pass increases from 25% to 40%, both 
yield strength and tensile strength of the material increase. Specifically, the yield strength increases 
40 MPa and the tensile strength increases 50 MPa when the higher thickness reduction of 40% is 
applied. Additionally, the total elongations of the two samples are rather similar. Both have the total 
elongation  of  ~12%.  However,  the  toughness  of  samples  which  was  intercritically  40%  thickness 
reduced  is  lower  than  that  of  20%  thickness  reduced.  For  the  same  samples,  fracture  energy 
measured in the longitudinal direction is always higher than that for the transversal direction. Non-
equiaxial microstructure and crystallographic texture are the main two reasons for the anisotropy of 
mechanical properties in these samples. 

Figure 74: Effect of reduction rate (25% vs. 40%) on the tensile properties of intercritical 
rolled plates  

Figure  75:  Effect  of  reduction  rate  (25%  vs  40%)  on  the  toughness  properties  of 
intercritical rolled plates. 

Impacts of the strain accumulation during intercritical rolling on mechanical properties of the CMn-
0.06%C steel were further investigated by mechanical property characterization on samples of multi 
intercritical rolling passes. Figure 76 shows tensile properties of all samples produced by intercritical 
rolling with 1, 2 and 3 passes of 10% thickness reduction at different temperatures and then cooling 
to room temperature by air. Regardless to the number of intercritical rolling passes, tensile strength 
and yield strength of the studied material increase with the reduction of the rolling temperature. On 
the contrary, total elongations of the corresponding samples reduce proportionally with the rolling 
temperature. However, at the same rolling temperatures, the number of intercritical rolling passes 
has  significantly  impacts  on  tensile  properties  of  material.  Specifically,  when  the  number  of 
intercritical  rolling  passes  increases  from  1  to  3  passes,  tensile  strengths  of  the  material  gain  50 
MPa. Additionally, the total elongation reduces about 2% by the increase in the number of intercritical 
rolling passes. Complete transition curves for all toughness measurements are given in Figure 77. 
For all material processing conditions, the upper shelf energy (USE) of the Charpy V-notched tests 
are measured in the range of 235 to 350 J, while the ductile-brittle transition temperatures (DBTT) 
are between -35°C to -90°C. Toughness properties measured in longitudinal direction of all samples 
74 

 
 
 
 
 
irrespective to the number of intercritical rolling passes are rather similar. On the contrary, there are 
differences in USEs and DBTTs measured in the transversal direction among samples produced 1, 2 
and 3 intercritical rolling passes. The highest USE among samples produced by 1 intercritical rolling 
pass is 350J, while the highest USE of samples produced by 3 intercritical rolling pass is 280J. The 
lowest DBTT measured in the former set of samples is -90oC, while the corresponding value in the 
later set of samples is -72oC. The differences in toughness properties measured for longitudinal and 
transversal directions is due to the texture strengthening by increasing intercritical rolling passes. 
Additionally,  the  larger  number  of  rolling  passes  applied  at  lower  intercritical  rolling  temperature 
introduces the higher amount of stored strain energy in material. As a result, the lowest USE of 235J 
and the highest DBTT of -35oC are obtained in the sample of 3 intercritical rolling passes at 710oC. 

(a) 

(b) 

(c) 
Figure 76: Tensile properties of the 3 samples which were intercritical rolled by (a) 1 pass, 
(b)  2  passes  and  (c)  3  passes  of  10%  thickness  reduction  at  different  temperatures 
between 710oC and 800oC. 

75 

 
 
 
 
 
 
 
 
Figure 77: Toughness measured in longitudinal and transversal directions of the samples 
which were intercritically rolled at different temperatures between 710oC and 800oC in 1, 
2 and 3 rolling passes of 10% thickness reduction. 

The effect of cooling rates on mechanical properties of intercritically rolled samples of 10% thickness 
reduction in a single rolling pass was studied. Tensile properties of these samples are given in Figure 
78. Tensile strengths of the ACC samples are always higher than those of the air cooled samples. At 
the lowest rolling temperature of 710°C, the ACC sample has the tensile strength of 570 MPa, while 
the  tensile  strength  of  the  air  cooled  sample  is  542  MPa.  However,  by  applying  the  ACC  after 
intercritical rolling, both yield strength and total elongation of the CMn-0.06%C steel reduce with the 
rolling temperature. The  highest ratio of TS/YS for all ACC samples is reached when the rolling is 
implemented at 755°C. Tensile properties of ACC samples, therefore, are not linear functions of the 
rolling temperature. Toughness properties in longitudinal and transversal directions of the air cooled 
and ACC samples are given in Figure 79. The USE measured in longitudinal directions of the air cooled 
and ACC samples are equivalence. For both set of samples, the USE are in the range of 250-350 J. 
The range of DBTT for ACC samples, however, are broader than that for air cooled samples. A broad 
range of DBTT for ACC samples is also observed in the toughness curves for the transversal direction. 
After rolling at lower temperatures (710°C – 745°C), ACC samples have the USE of 150-250 J and 
the DBTT of -17°C to -35°C. 

(a) 

(b) 

Figure 78: Tensile properties of the samples which were intercritically rolled to reduce the 
thickness 10% at different temperatures between 710oC and 800oC and cooled in (a) air 
and (b) warm water (ACC) to room temperature. 

76 

 
 
 
 
 
 
 
 
(a) 

(b) 

(c) 

(d) 

Figure 79: Toughness properties of the samples which were intercritically rolled to reduce 
the thickness 10% at different temperatures between 710oC and 800oC and cooled in (a, 
b) air and (c, d) warm water (ACC) to room temperature. 

TASK  5.5  Completion  of  the  strength-toughness  map  versus  intercritical 
deformation 

To  complete  the  strength-toughness  map  for  all  intercritical  rolling  trials  on  the  CMn-0.06%C,  6 
mechanical parameters measured in the Task 5.4 have been given in Figure 80. With the main aim 
to investigate the effect of strain accumulation by intercritical rolling on mechanical properties, rolling 
temperatures  for  each  samples  have  been  omitted  from  the  Figure  80.  However,  processing 
schedules of these samples are still systematically recorded for later studies. 

As can be seen in Figure 80a and Figure 80b, the increase in number of intercritical rolling passes 
from  1 to 3 brings higher YS  and TS to the CMn-0.06%C  steel. In  general,  the  YS and TS of the 
material gain at least 25-35 MPa, when the total intercritical thickness reduction increases from 10% 
to  27%.  By  applying  multi-intercritical  rolling  passes  instead  of  single  intercritical  rolling  pass, 
hardening  behavior  of  the  material  also  changed.  The  TS/YS  ratios  for  all  samples  of  single 
intercritical rolling spread in the range of 1.11 to 1.27. On the other hand, these ratios for samples 
of 3 intercritical rolling passes are all in between 1.11 and 1.17. This result, in combination with the 
reduction of total elongation as observed in Figure 80d, indicates the reduction of strain hardening 
when  the  number  of  intercritical  rolling  passes  increases.  The  same  trend  is  also  observed  in 
toughness  behaviors  of  the  material.  In  Figure  80e  and  Figure  80f,  as  the  number  of  intercritical 
rolling passes increase from 1 to 3 passes, the lowest DBTT increases from  -90°C to -75°C, while 
the highest USE reduces from 325J to 275J. 

77 

 
 
 
 
 
 
 
(a) 

(b) 

(c) 

(d) 

Figure 80: The strength-toughness map versus the number of intercritical rolling passes 
on all samples which were intercritically rolled and air cooled to room temperature. 

(e) 

(f) 

78 

 
 
 
 
 
 
 
 
 
 
 
WP6: Impact of chemical composition 

- 

The objective of WP6 is to understand and quantify the effect of the chemical composition 
(C and Nb-V micro-alloying) content on the microstructural and property relationship. 
Some selected TMCP paths are investigated to check if the processing window is drastically 
influenced by the presence of higher C or micro-alloying content as these are often the 
elements varied in this type of materials. The specific pinning effects of the small Nb and V 
carbides could strongly influence the strain partitioning behavior. 

TASK 6.1 Laboratory rolling trials 

The same systematic approach as before is used to assess the impact of the chemical composition 
on the microstructure and property relation. 

  C-effect 
  Microalloying (NbV) effect 

New cast ingots had to be produced for this WP. The actual chemical composition is shown in the 
Table 23 and Table 24. As compared to the one in the first casted ingots produced by CRM in the 
WP2, the carbon content in the new ingots is slightly higher. 

Table 23: Chemical composition of the CMn (0.12%C) grade heats cast 

Table 24: Chemical composition of the NbV grade heats cast 

To limit the number of rolling trials as compared to the ones of WP3 and WP4, only two intercritical 
temperatures have been selected for further studies in this WP6. The two intercritical temperatures 
are chosen based on the “relaxation-like” torsion trials, performed by CRM in WP2. They are desired 
to produce different intercritical ferrite volume fractions and thus sharply impacts on final mechanical 
properties of materials: 

- 

- 

For the CMn (0.12%C) grade, the intercritical rolling temperatures are 750°C to target low 
ferrite fraction and 700°C for high ferrite fraction. 
For grade NbV the intercritical temperatures are 770°C and 750°C, respectively. 

Knowing  the  difficulties  to  obtain  the  right  fraction  of  intercritical  ferrite  in  WP3,  interrupted 
quenching experiments have been carefully implemented right from the beginning of the rolling trials. 

79 

CMnSiPSAlNbNV0.121.50.250.0120.0020.03RealisedL5910.1281.5390.2390.01270.00100.0240.0020.00230.001RealisedL5920.1291.5430.2440.01390.00100.0260.0020.00330.000TargetCMn (0.12%C)CMnSiPSAlNbNV0.061.50.250.0120.0020.030.050.03RealisedL5930.0641.5330.2490.01310.00100.0240.0530.00560.034RealisedL5940.0661.5430.2480.01270.00110.0250.0520.00380.034TargetNbV 
 
 
 
 
Figure 81: Schematic of the experimental plan of the rolling trials of WP6. 

Table 25: Key rolling and cooling parameters of the rolling trials in WP6. 

In the same way as carried out in WP3, the accelerated cooling (ACC) was performed by cooling the 
plates in a warm (60°C) water tank while direct quench was performed in a cold water tank. 

As can be seen in Table 26, target rolling and cooling temperatures for both grades in DQ mode and 
ACC mode have been well achieved. However, cooling rates in DQ mode are broadly variant, between 
31  oC/s and 94  oC/s. This is likely to be related  to the  temperature  increase of water  in the  tank 
during the successive cooling trials. For the same reason, cooling rates of ACC mode are determined 
between 10  oC/s and 48  oC/s. The cooling rates (CR) have been calculated within the temperature 
range of 800°C and 500°C (t8/5). Quenching was stopped when the core plate temperature dropped 
below 100°C and accelerated cooling was stopped when the temperature dropped below 400°C, since 
phase transformations are almost completed above 500°C. The thermal profiles of each rolled plates 
are given in Figure 82. 

Similar to WP3, all plates were sampled for tensile and Charpy impact toughness in transverse and 
longitudinal  direction.  A  sample  per  plate  is  also  foreseen  for  microstructural  investigations.  To 
minimize the temperature inhomogeneity, the area close to the plate edges was not used. 

Table  26:  Realized  rolling  and  cooling  parameters,  temperatures  measured  by 
thermocouples at plate mid thickness 

80 

FRTcoolingTempReductionCoolingSet B820DQ12mm820Air7500820Air7000820Air75025%820Air70025%820Air75025%820Air70025%820Air75025%820Air70025%CMn (0.12C%C)Above Ar3γ-α-phase regionFinal thicknessSet CDQ16mmSet AAir12mmDQACCFRTcoolingTempReductionCoolingSet B820DQ12mm820Air7700820Air7500820Air77025%820Air75025%820Air77025%820Air75025%820Air77025%820Air75025%NbVAbove Ar3γ-α-phase regionFinal thickness16mmSet AAir12mmDQACCSet CDQ 
 
 
 
 
(a) 

(b) 

(c) 

(d) 

Figure 82: Thermal profiles during rolling and cooling: (a-b) for grade NbV and (c-d) for 
grade CMn (0.12%C). Graph (a & c) with direct quench before and after IC-rolling; Graph 
(b & d) with ACC. 

TASK 6.2 Microstructural characterization 

The  microstructure  of  the  rolled  plates  were  investigated  by  LOM.  An  overview  of  different 
microstructures  is given  in  the  Appendix  for  the  CMn  (0.12%C)  grade  and  for  the  NbV  grade.  All 
microstructures are captured at the quarter thickness region on ND-RD section of each samples. The 
fraction  of  ferrite  was  determined  by  an  in-house  built  image  analysis  software  at  OCAS  and 
summarized in Figure 83. The desired volume fractions of intercritical ferrite are the best achieved 
by fast cooling. However, it is worthy to mention that the difference in intercritical ferrite volume 
fraction for the NbV grade at 770°C and 750°C is very small and thus their effects on the mechanical 
properties might be negligible. In samples which were slowly cooled by air after intercritical rolling, 
the austenite still transforms into ferrite and pearlite during cooling to the room temperature. Effect 
of cooling rate on final microstructures of the CMn (0.12%C) grade is stronger than that of the NbV 
grade. 

81 

CMn (0.12%C)Settarget αFRTreal[°C]SCTreal[°C]FCTreal[°C]CR (t8/5)[K/s]CoolingL591AB-8117927231DQL591BClow8177347049DQL591CChigh8157006861DQL591FAlow7517397461DQL592AAhigh7027007094DQL592BAlow75174236010ACCL592CAhigh70068124248ACCNbVSettarget αFRTreal[°C]SCTreal[°C]FCTreal[°C]CR (t8/5)[K/s]CoolingL593AB-8198058145DQL593BClow8127677562DQL593CChigh8207457560DQL593FAlow7697636936DQL594AAhigh7507477662DQL594BAlow76976436720ACCL594CAhigh75074637215ACC 
 
 
 
 
 
 
 
 
Figure 83: Fraction of Ferrite as a function of realized start cooling temperature for (left) 
CMn (0.12%C) grade and (right) NbV grade at ¼ thickness 

In general, microalloying elements like Nb and V are added to steels to stimulate the formation of 
carbides, nitrides and carbonitrides. Based on the solubility of these elements in austenite and ferrite, 
the fractions of free Nb and V can be calculated as a function of temperature, shown in Figure 84. As 
can  be  seen,  when  entering  into  the  intercritical  temperature  regime  all  Nb  should  already  be 
precipitated, while V precipitation is delayed due to the  different solubilities of V in austenite and 
ferrite. 

ICP-MS  (Inductively  coupled  plasma  mass  spectrometry)  measurements  were  performed  on  the 
intercritically rolled and directly quenched plates. When a sample of NbV grade is rolled at 770°C 
and quenched, only 40% of Nb is precipitated. However, when a sample of the same grade is rolled 
at 750°C, 65% of Nb  is precipitated.  That means, that during the slow air cooling from  770°C to 
750°C another 25% of Nb is precipitated. In comparison, V is even less precipitated with 13% and 
18% at 770°C and 750°C, respectively. The slow cooling rate by air indeed promotes the precipitation 
of (Nb,V)C in the NbV grade. Therefore, it can be expected, that the precipitation of Nb and V in air 
cooled plates are higher than that in DQ and ACC plates. 

Figure 84: Dissolution calculations for NbV grade, showing the mass fraction of free N, C, 
Nb and V as a function of temperature. 

Table 27: ICP-MS results on the precipitation of Nb (left) and V (right) after intercritical 
rolling and quenching. 

IC+DQ 

770oC 
750oC 

%Nb 
sol. 
0.033 
0.018 

%Nb 
pre. 
0.022 
0.034 

Tot. Nb 
iCAP 
0.055 
0.052 

Pre- 
cipitated 
40% 
65% 

IC+DQ 

770oC 
750oC 

%V 
sol. 
0.027 
0.028 

%V 
pre. 
0.004 
0.006 

Tot. V 
iCAP 
0.031 
0.034 

Pre- 
cipitated 
13% 
18% 

On a selection of plates detailed EBSD investigation at mid thickness were performed to determine 
the  fraction  of  ferrite,  martensite  and  bainite  as  well  as  the  average  grain  size  of  each  of  the 
constituent and of the complete microstructure. 

82 

 
 
 
 
 
 
 
 
For the CMn (0.12%C) grade three samples have been measured:  

  Direct quench after finish rolling at 820°C, i.e. from fully austenitic (820_DQ) 
  Direct  quench  after  air  cooling  to  intercritical  temperature  of  700°C  without  intercritical 

deformation (700_DQ) 

  Air cooling after intercritical rolling at 700°C (IC_700_Air) 

Based on Grain Average Image Quality and Grain Average Neighbor Misorientation, it was possible 
to distinguish the different phases. The microstructure of the 820_DQ sample contains 38% ferrite, 
13% martensite and 48% bainite. The fraction of ferrite increases to 85% for 700_DQ and further 
to 93% for IC_700_Air. The average grain size of martensite in all samples is rather constant with 
about 4µm. However, the ferrite grain size is increasing from ~6µm to ~15µm, as given in  Figure 
85a. 

Figure 85: Summary of grain size and phase fractions for CMn (0.12%C) grade (a) and for 
NbV grade (b). 

(a) 

(b) 

For the NbV grade only two samples have been measured: 

  Direct quench after finish rolling at 820°C, i.e. from fully austenitic (820_DQ) 
  Direct  quench  after  air  cooling  to  intercritical  temperature  of  750°C  without  intercritical 

deformation (750_DQ) 

Here, no further distinction between martensite and bainite were made. However, it is obvious that 
at  820_DQ  the  main  phase  is  “martensite  and  bainite”  while  at  750_DQ  it  is  about  50-50  in  the 
Appendix. A summary of grain size and phase fraction is given in Figure 85b. The average grain size 
after austenitic rolling is slightly higher for NbV grade than for the CMn (0.12%) grade (13.5µm vs. 
11.9µm). During slow air cooling to intercritical temperatures the average grain size for both grades 
is about the same  (9.1µm vs.  9.4µm). The  increase in grain size is for both  grades  linked to  the 
increased fraction of ferrite, which has a larger grain size. 

At  CEIT,  impact  of  chemical  composition  on  microstructures  of  intercritical  rolling  has  been 
investigated on deformed dilatometry samples of the CMn(0.06%C) and NbV steels. Details of this 
investigation are given in the Appendix. 

Microstructures  of  the  two  materials  at  different  intercritical  temperatures  without  intercritical 
deformation have been investigated on quenched samples. As can be seen in Table 28, the reduction 
of the intercritical temperature promotes the increment of the ferrite content gradually. The fractions 
of  the  intercritical  ferrite  prior  to  deformation  were  measured  and  given  in  Table  28.  For  both 
chemical compositions, three different austenite-ferrite balances (25%, 50% and 75% ferrite) are 
obtained  using  different  deformation  temperatures.  For  NbV  steel  (see  in  Appendix),  a 
microstructural  refinement  (finer  ferrite  grains)  is  noticed  when  the  transformation  occurs  from 
deformed  austenite.  The  effect  of  the  strain  accumulation  on  the  grain  refinement  is  due  to  the 
increase in the density of nucleation sites introduced by deformation. 

83 

 
 
 
 
 
Table 28: Optimum GOS values needed to obtain the desired ferrite fraction in each case 
(CMn(0.06%C),  NbV  Rec  and  NbV  def).  Measured  deformed  ferrite  fractions  were  also 
included, considering the proposed tolerance angle of 4º. 

Tdef 
(ºC) 

% Objective 
deformed ferrite 

Optimum GOS 
(º) 

CMn(0.06%C) 

NbV Deformed 
austenite 

NbV Rex. 
austenite 

750 

740 

730 
770 
760 
740 
750 
740 
720 

25 

50 

75 
25 
50 
75 
25 
50 
75 

4.5 

4.5 

3.5 
2 
4 
3 
4 
4 
4 

%Measured 
deformed ferrite 
(GOS criteria) 
37 

64 

63 
8 
53 
45 
24 
61 
71 

In addition to the characterization of the samples quenched before deformation, the microstructural 
characterization  of  the  air  cooled  microstructures  obtained  after  deformation  is  needed  with  the 
purpose  of  gaining  a  better  understanding  of  the  effect  of  intercritical  deformation  on  the  final 
microstructure. Therefore, the microstructures obtained after air cooling are characterized by means 
of  optical  microscopy  and  scanning  electron  microscopy,  for  both  chemical  compositions  and 
austenite  conditions  and  entire  range  of  deformation  temperatures  (25,  50  and  75%  ferrite).  In 
Figure 86 a- c optical micrographs related to CMn(0.06%C) steel are presented, while in Figure 86 
d-i, the microstructures obtained for both recrystallized and deformed austenite in the NbV steel can 
be  compared.  Combinations  between  non-deformed  ferrite  (NDF)  and  deformed  ferrite  (DF)  are 
clearly  distinguished  in  all  the  cases.  Moreover,  the  formation  of  pearlite  is  also  observed  in  the 
resulting microstructure,  mainly in  the  CMn(0.06%C) steel and the lowest ferrite content prior to 
deformation. The fraction of deformed ferrite increases as the deformation temperature is decreased, 
owing to the higher presence of ferrite in the microstructure prior to deformation. Figure 86 shows 
that the deformed ferrite is characterized by a significant presence of substructure. 

)
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(a) Tdef 750ºC 

(b) Tdef 740ºC 

(c) Tdef 730ºC 

(d) Tdef 750ºC 

(e) Tdef 730ºC 

(f) Tdef 720ºC 

(g) Tdef 770ºC 

(h) Tdef 750ºC 

(i) Tdef 740ºC 

84 

 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
Figure  86:  Optical  micrographs  obtained  after  air  cooling  at  different  deformation 
temperatures:  (a,d,g)  low  ferrite  content,  (b,e,h)  medium  ferrite  content  and  (c,f)  high 
ferrite  content.  The  microstructures  obtained  in  all  the  cases  are  shown:  (a,b,c) 
CMn(0.06%C), (d,e,f) NbV Recrystallized austenite and (g,h,i) NbV Deformed austenite. 

In the SEM micrographs shown in Figure 87 the differences between different ferrite morphologies 
can be appreciated more properly. Figure 87 illustrates the morphology of the phases formed after 
different intercritical deformation temperatures and different austenite conditions (recrystallized and 
deformed austenite) for NbV steel. The resulting microstructures are characterized by the presence 
of a fraction of deformed ferrite and a fraction of non-deformed ferrite, which is formed during the 
final air cooling step. The SEM images suggest that the reduction of deformation temperature leads 
to  the  presence  of  more  non-equiaxed  microstructure,  due  to  the  increment  of  deformed  ferrite 
content in the resulting microstructure. 

NbV (Recrystallized austenite) 

NbV (Deformed austenite) 

e
t
i
r
r
e
f

w
o
L

e
t
i
r
r
e
f

i

m
u
d
e
M

e
t
i
r
r
e
f

h
g
H

i

(a)  Tdef 750ºC 

(d) Tdef 770ºC 

(b) Tdef 730ºC 

(e) Tdef 750ºC 

(c) Tdef 720ºC 

(f) Tdef 740ºC 

Figure 87: SEM micrographs related to the samples obtained after different deformation 
temperatures  for  NbV  Recrystallized  and  Deformed  austenite:  (a,d)  Low,  (b,e)  medium 
and (c,f) high ferrite content. 

In  the  current  study,  the  microstructural  analysis  has  been  extended  using  the  Electron 
Backscattered  Diffraction  (EBSD)  technique,  in  order  to  quantify  the  effect  of  the  deformation 
temperature and the ferrite fraction before deformation. EBSD technique is required for  analysing 
85 

 
 
 
 
 
 
 
 
 
 
 
 
 
 
the  microstructural  features  of  a  sample  deformed  in  the  intercritical  region,  which  cannot  be 
correctly quantified by any other standard microstructural characterization technique such as optical 
microscopy  and/or  FEG-SEM.  In  Figure  88  Inverse  Pole  Figures  and  Image  Quality  maps  are 
superimposed, as well as high and low angle boundaries (coarse and fine black lines, respectively). 
Figure  88  shows  the  EBSD  maps  corresponding  to  the  NbV  steel  and  deformed  at  different 
deformation temperatures after different austenite conditioning strategies. The EBSD maps shown 
in Figure 88 suggest that depending on the deformation temperature and consequently, the fraction 
of ferrite prior  to intercritical  deformation, completely different microstructures are  formed. When 
the microstructure contains a low fraction of ferrite prior to deformation, the presence of relatively 
polygonal  ferrite  grains  is  clearly  observed.  As  the  deformation  temperature  decreases,  more 
elongated and coarser grains can be observed. The applied deformation promotes the modification 
of the formed ferrite, reflected in the presence of a higher substructure. 

Analyzing the effect of austenite condition on the resulting intercritically deformed microstructure, 
clear  differences  can  be  noticed  regarding  the  morphology  of  the  deformed  ferrite  population. 
Comparing the EBSD maps obtained for both recrystallized and deformed austenite (see Figure 89), 
completely different substructure is observed inside the deformed ferrite grains for each austenite 
condition (see the low angle boundaries drawn in the IPF maps). 

NbV (Recrystallized austenite) 

NbV (Deformed austenite) 

e
t
i
r
r
e
f

w
o
L

e
t
i
r
r
e
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m
u
d
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(a) 

(d) 

(b) 

(e) 

86 

 
 
 
 
 
 
 
 
 
 
e
t
i
r
r
e
f

h
g
H

i

(c) 

(f) 

Figure 88: IPF+IQ maps with high and low angle boundaries (coarse and fine black lines, 
respectively),  corresponding  to  (a,b,c)  NbV  Recrystallized  austenite  and  (d,e,f)  NbV 
Deformed austenite. Different ferrite levels prior to deformation have been included: (a,d) 
Low ferrite, (b,e) Medium ferrite and (c,f) High ferrite.  

(a) 

(b) 

Figure 89: IPF+IQ maps with high and low angle boundaries (coarse and fine black lines, 
respectively),  corresponding  to  (a)  NbV  Recrystallized  austenite  and  (b)  NbV  Deformed 
austenite (high ferrite fraction prior to intercritical deformation). 

Given  that  more  bainitic  phases,  such  as  quasi-polygonal  ferrite,  are  formed  when  microalloying 
elements  are  added,  the  differentiation  of  the  non-deformed  and  deformed  ferrite  grains  is  more 
complex. Therefore, when microalloyed  steels are  used,  the transformation occurs from  relatively 
coarse austenite grains. The methodology able to distinguish non-deformed ferrite from deformed 
ferrite needs to be re-assessed. For the distinction between deformed and non-deformed ferrite, the 
Grain Orientation Spread (GOS) parameter was taken into account, which is the average deviation 
between the orientation of each point in the grain and the average orientation of the grain. Based 
on this study, a threshold GOS value of 2º was defined for the differentiation between deformed and 
non-deformed polygonal ferrite grains. 

Nevertheless,  when  more  bainitic  phases  are  formed,  an  extension  of  the  proposed  procedure  is 
required. In the microstructures shown before, three different populations are identified: polygonal 
ferrite (PF), quasi-polygonal ferrite (QF) and deformed ferrite. For each GOS distribution, a threshold 
GOS value was defined for the differentiation between deformed and non-deformed ferrite grains. In 
Table 28, optimum GOS values needed to obtain the desired ferrite fraction (25, 50 and 75%) are 
shown. A common value of 4º was estimated as the average of all the threshold angles defined in 
each scan. This minimum GOS value is able to distinguish both non-deformed and deformed ferrite 
in the final microstructure for different compositions (CMn(0.06%C) and NbV), ferrite contents before 
deformation and austenite conditions (recrystallized and deformed ferrite). In this case, the ferrite 
population with GOS values lower than 4º is considered as non-deformed ferrite, where PF and QF 
are  included. GOS  values  higher than 4º  correspond to  deformed ferrite grains. Furthermore, the 
87 

Subgrains50 µmDeformationbands50 µm 
 
 
 
 
 
measured  deformed  ferrite  fractions,  considering  the  proposed  tolerance  angle  of  4º,  were  also 
included in Table 28. 

Based on the described GOS criterion, a quantification of high angle misorientation unit sizes for each 
ferrite type was carried out considering 15º misorientation criteria (D15º). In Figure 90a and b, unit 
size distributions (in terms of frequency) are presented for both non-deformed and deformed ferrite 
families,  respectively.  The  results  corresponding  to  50%  ferrite  content  prior  to  deformation  are 
included  in  Figure  90.  Comparing  the  grain  size  distributions  obtained  in  each  ferrite  population, 
significantly finer non-deformed distributions are measured for non-deformed ferrite (see Figure 90a) 
compared to deformed ferrite (Figure 90b). Additionally, in Figure 90c and d, mean grain size values 
are plotted for both ferrite types and different ferrite levels, as well as both chemistries and austenite 
conditions. Analyzing the results shown in Figure 90c and d, as mentioned previously, considerably 
finer microstructures are obtained in the non-deformed ferrite formed during final air cooling step, 
for entire ferrite content prior to intercritical deformation and all the studied cases. Concerning the 
mean  grain  size  trends  corresponding  to  non-deformed  ferrite  population,  no  significant  effect  of 
ferrite fraction before deformation and addition of microalloying elements on grain size is observed. 
Similar  mean  unit  size  values  are  measured  in  entire  ferrite  content,  both  steels  and  different 
austenite conditions. Nevertheless, a different behavior can be detected regarding the evolution of 
mean grain size regarding deformed ferrite family. The EBSD measurements plotted in  Figure 90d 
suggest  that  the  addition  of  Nb  and  V  promotes  microstructural  refinement,  mainly  when  the 
transformation  occurs  from  deformed  austenite.  Slightly  coarser  microstructures  are  obtained  for 
CMn(0.06%C) steel compared to NbV Recrystallized condition. However microstructural refinement 
caused by addition of microalloying elements is more severe when accumulation of deformation is 
ensured in the austenite prior to transformation. The benefit of the accumulation of deformation on 
the microstructural refinement  is associated with the  increase of the specific grain boundary area 
and  is  attributed  to  a  significant  increase  in  the  density  of  ferrite  nucleation  sites  introduced  by 
deformation. 

(a) 

(b) 

(c) 

(d) 

88 

0%10%20%30%40%50%60%Frequency (%)Unit size (μm)CMn(0.06%C)  50%NbV REC  50%NbV DEF  50%Nondeformed0%5%10%15%20%25%30%Frequency (%)Unit size (μm)CMn(0.06%C)  50%NbV REC  50%NbV DEF  50%Deformed ferrite46810121416182002040608010015º Mean Unit Size (μm)Non deformed ferrite fraction (%)6CMn Non Deformed ferriteNbV Rec Non Deformed ferriteNbV Def Non Deformed ferriteCMn(0.06%C) Non deformedferriteNbVRecNon deformedferriteNbVDefNon deformedferrite46810121416182002040608015º Mean Unit Size (μm)Deformed ferrite fraction (%)6CMn Deformed ferriteNbV Rec Deformed ferriteNbV Def Deformed ferriteCMn(0.06%C)  deformedferriteNbVRecdeformedferriteNbVDefdeformedferrite 
 
 
 
 
 
Figure 90: Grain size distributions corresponding to 50% ferrite prior to deformation and 
different  chemical  composition  (CMn(0.06%C),  NbV  Recrystallized  austenite  and  NbV 
Deformed  austenite):  (a)  Non-deformed  ferrite  and  (b)  Deformed  ferrite  population. 
Influence of austenite/ferrite balance, chemical composition and austenite condition on 
15º  mean  grain  size  of  both  ferrite  populations:  (c)  Non-deformed  ferrite  and  (d) 
Deformed ferrite population. 

TASK 6.3 Determination of the mechanical properties 

The mechanical properties of strength and Charpy  impact toughness  of the lab rolled  plates  have 
been tested in longitudinal and transverse direction.  As before, tensile samples were machined in 
A50 geometry and Charpy samples as full-size V-notch samples taken at the mid thickness. Due to 
a lack of material, only limited amount of test temperatures was tested by Charpy. 

In the following the results of CMn (0.12%C) and NbV grade are reported separately. 

CMn (0.12%C) grade 

Figure 91 shows the yield and tensile strength as well as the elongation for the different processing 
tested.  As  can  be  seen,  the  strength  levels  after  direct  quench  from  the  intercritical  temperature 
regime are about the same independent if intercritical rolling (IC) is applied or not. Furthermore, the 
strength is the highest with increasing cooling rate (arrow indicated by “high CR”). The elongation, 
however, is decreasing with increasing cooling rate. When fast cooling is applied, the strength level 
is lower for high ferrite fractions, due to the lower amount of martensite (arrow indicated by “less 
M”).  However, when  plates are  air  cooled after  intercritical rolling, the strength is higher for  high 
ferrite  fractions.  In  most  cases  a  low  intercritical  rolling  temperature  tends  to  decrease  the 
elongation.  

89 

 
 
 
 
 
Figure  91:  Tensile  strength  of  CMn  (0.12%C)  in  top:  longitudinal  (0)  and  bottom: 
transverse (90) direction for the different processing 

Figure 92 shows the transition curves of the CMn(0.12%C) grade with the different processing tested 
in longitudinal direction. As can be  seen,  the  overall toughness significantly drops with increasing 
cooling rate. For air cooled material and accelerated cooled from the low temperature intercritical 
regime (IC_700_ACC), the toughness is rather comparable down to -60°C. The few values below -
60°C indicate, that the best toughness is obtained for low fraction of intercritical deformed ferrite 
followed by air cooling.  

high CR 

90 

 
 
 
 
 
Figure 92: Charpy impact toughness of the CMn (0.12%C) grade with different processing 
tested in longitudinal direction, (top) with intercritical rolling; (bottom) with direct quench 

NbV grade 

Figure 93 shows the yield and tensile strength and elongation for the different processing tested. As 
can  be  seen,  the  strength  levels  are  rather  constant  for  direct  quenched,  independent  the  start 
cooling temperature, however, the strength level is slightly higher with intercritical rolling. As already 
seen for CMn (0.12%C) the strength level increases with increasing cooling rate. In contrast to the 
high C grade, for lower intercritical rolling temperatures the strength levels are in all cases (slightly) 
higher.  The  elongation  is  decreasing  with  increasing  cooling  rate  and  lower  intercritical  rolling 
temperatures tend to lower the elongation as well.  

91 

 
 
 
 
Figure  93:  Tensile  strength  of  NbV  in  top:  longitudinal  (0)  and  bottom:  transverse  (90) 
direction for the different processing 

Figure  94  shows  the  transition  curves  of  the  NbV  grade  with  the  different  processing  tested  in 
longitudinal  direction.  As  can  be  seen,  the  overall  toughness  significantly  drops  with  increasing 
cooling rate but not as strong as for the CMn (0.12%C) (high C). The air cooled and the accelerated 
cooled  from  high  intercritical  rolling  temperature  (IC_770_ACC)  transition  curves  are  rather 
comparable, while the curve of low intercritical rolling with ACC (IC_750_ACC) is similar to the direct 
quench ones. The differences observed for the direct quench curves are not dependent on intercritical 
rolling applied or not, but probably on the realized cooling rates.  

high CR 

92 

 
 
 
 
Figure 94: Charpy impact toughness of the NbV grade with different processing tested in 
longitudinal direction, (top) with intercritical rolling; (bottom) with direct quench 

TASK 6.4 Quantification of the impact of C, Nb and V on the microstructural and 
property relationship 

To evaluate the effect of microalloying elements on precipitation dilatometer tests were run and TEM 
analyses on fine precipitates was carried out with carbon extraction replicas. The characterization of 
the  precipitation  was  carried  out  with  dilatometric  samples  obtained  after  the  thermomechanical 
schedule shown in the Appendix. Figure 95 shows differences in precipitation for both recrystallized 
and  deformed  austenite,  when  75%  ferrite  content  prior  to  intercritical  deformation  is  formed. 
Regarding the effect of austenite conditioning, TEM characterization clearly reveals that completely 
different  precipitation  is  formed  depending  on  the  applied  hot  rolling  strategy  (recrystallized  or 
deformed austenite). When the transformation occurs from recrystallized austenite, a high density 
of fine precipitates can be detected (see Figure 95a). These precipitates are considered to be formed 
in ferrite during or after the intercritical deformation. Nevertheless, in the sample corresponding to 
deformed  austenite  strain  induced  precipitates  formed  in  austenite  are  also  observed  (see  Figure 
95b). Therefore, the micrographs shown in Figure 95 suggest that the followed strategy for austenite 
conditioning strongly modifies the density and size of precipitates. Most of the particles are Nb rich 
precipitates, as well as Nb-V and Nb-Ti-V rich precipitates. 

(a) 

(b) 

93 

 
 
 
 
 
 
Figure 95 TEM images obtained in (a) NbV Recrystallized austenite and (b) NbV Deformed 
austenite. 

In  Figure  96a  and  b  the  precipitate  size  distributions  are  plotted,  in  terms  of  area  fraction  and 
accumulated  area  fraction,  respectively.  The  precipitate  diameter  distributions  corresponding  to 
recrystallized austenite and deformed austenite conditions (obtained for 75% fraction of ferrite prior 
to intercritical deformation) can be compared. The results indicate that noticeably finer precipitates 
have been measured when transformation occurs from recrystallized austenite (mean precipitate size 
is decreased from 15.9 to 10.2 nm). A lower intercritical deformation temperature is applied in Cycle 
A (recrystallized austenite), leading to the formation of finer precipitates in ferrite. In addition, as 
mentioned previously, a higher density of fine precipitates is observed. 

(a)  

(d) 

Figure 96 Grain size distributions measured for NbV Rec and NbV Def conditions in terms 
of  (a)  area  fraction  and  (b)  accumulated  frequency  (75%  of  ferrite  prior  to  intercritical 
deformation). 

Besides the evaluation of the influence of the accumulation of deformation in the austenite prior to 
transformation, the effect of ferrite fraction before intercritical deformation was also analyzed. As an 
example, Figure 97 shows the comparison between TEM images corresponding to low and high ferrite 
levels (25 and 75% prior to deformation). Furthermore, in Figure 98 the precipitate size distributions 
are  also presented. Quantified mean precipitate sizes are  also include in the graphs. A significant 
precipitate size refinement is clearly noticed as ferrite content prior to deformation increases and 
intercritical  deformation  temperature  decreases.  Furthermore,  as  shown  in  Figure  97,  a  higher 
fraction of fine precipitates is noticed for high ferrite content compared to low ferrite fraction. Mean 
precipitate  size  decreases  from  13.8  to  10.2  nm,  when  the  fraction  of  ferrite  before  intercritical 
deformation increases from 25 to 75%.  

(a)  

(b) 

94 

0%5%10%15%20%25%30%Area Fraction (%)Precipitate size (nm)DEF 75%REC 75%Dppt=10.2 nmDppt=15.9 nm00,20,40,60,81Accumulated FrequencyPrecipitate size (nm)Deformed austeniteRecrystallized austenite75% FERRITE 
 
 
 
 
 
Figure  97  TEM  images  obtained  in  NbV  Recrystallized  austenite  and  different  ferrite 
contents prior to deformation: (a) 25% ferrite and (b) 75% ferrite. 

(a)  

(d) 

Figure  98  Grain  size  distributions  measured  for  NbV  Rec  and  ferrite  fraction  of  25  and 
75%: (a) area fraction and (b) accumulated frequency. 

Table 29 summarize the key parameters on process, mechanical properties and microstructure after 
intercritical  rolling.  The  mechanical  properties  are  limited  to  tensile  strength and  ductile to brittle 
transition temperature (50% USE), both tested in rolling direction. The microstructure is represented 
by the amount of 2nd phase, which includes martensite, bainite and pearlite, and was obtained by 
optical image analysis. Hence, the amount of ferrite is equal to the reduction of 2nd phase from 100%.  
Figure 99 and Figure 100 are plotting the results. Composition independent, it can be seen, that with 
increasing  cooling  rate  from  air  to  direct  quench,  strength  is  increasing  and  the  ductile  to  brittle 
temperature is increasing, i.e. toughness is decreasing. This can be linked to the fraction and kind 
of second phase – small amount of pearlite and high fraction of martensite. With lower intercritical 
temperature, that means with higher amount of ferrite, the strength is slightly increasing due to the 
higher strength of the 2nd phase caused by its C enrichment.  

With respect to composition, it can be observed, that the microalloying of NbV is increasing strength 
by at least 100MPa by solid solution and precipitation hardening, independent of cooling rate. The 
addition  of  C  is  increasing  strength  significantly  for  direct  quench.  For  air  cooling,  the  strength 
increases by increasing pearlite fraction is lower than by adding NbV.  

NbV  significantly  improves  ductile  to  brittle  transition  temperature.  The  results  also  indicate  that 
toughness becomes less dependent on cooling rate.  

The addition of C is deteriorating the toughness, i.e. ductile to brittle temperature is increasing and 
upper shelf energy is decreasing (Figure 92). Toughness improves only by reducing the cooling rate 
to  room  temperature,  transforming  the  hard  martensite  phase  into  ferrite  and  pearlite.  Against 
expectations, the high C grade CMn (0.12%C) shows better ductile to brittle transition temperatures 
for air cooling than the reference grade CMn (0.06%C), although the double pearlite is present in 
banded structure. 

Table 29: Summary of process-property-microstructure for the three investigated grades 
after intercritical rolling. DBTT = 50%USE, 2nd phase identification by optical microscopy 
(OM) with P for pearlite, M for martensite and B for bainite. 

95 

0%5%10%15%20%25%30%Area Fraction (%)Precipitate size (nm)REC 25%REC 75%Dppt=10.2 nmDppt=13.8 nm00.20.40.60.81Accumulated FrequencyPrecipitate size (nm)REC 25%REC 75% 
 
 
 
 
CR 

Figure 99: Link Process-Tensile Strength-Fraction of Ferrite 

96 

CMn (0.06%C)Strength (Rm)Toughness (DBTT)2nd phaseStrength (Rm)Toughness (DBTT)2nd phaseStrength (Rm)Toughness (DBTT)2nd phase[MPa][°C][%][MPa][°C][%][MPa][°C][%]790°Clow α686-6060549-5260429-5613770°Clow α700-6650448-639750°Chigh α711-3725588-4020481-729730°Cfull α693-2625482-999avgerageRDAirACCDQCMn (0.12%C)Strength (Rm)Toughness (DBTT)2nd phaseStrength (Rm)Toughness (DBTT)2nd phaseStrength (Rm)Toughness (DBTT)2nd phase[MPa][°C][%][MPa][°C][%][MPa][°C][%]750°Clow α888-3058622-2051486-9022700°Chigh α882-3027563-7022539-8022averageRDAirACCDQNbVStrength (Rm)Toughness (DBTT)2nd phaseStrength (Rm)Toughness (DBTT)2nd phaseStrength (Rm)Toughness (DBTT)2nd phase[MPa][°C][%][MPa][°C][%][MPa][°C][%]770°Clow α771-7565627-9050568-8510750°Chigh α783-7534715-8044566-8510averageRDAirACCDQ 
 
 
 
CR 

Figure 100: Link Process-Ductile to Brittle Transition Temperature-Fraction of Ferrite 

Most of the property changes are explainable by the fraction of martensite or presence of pearlite 
instead. A high fraction of martensite is increasing strength but decreasing toughness, while ferrite 
and pearlite is tougher than ferrite and martensite. Furthermore, to increase strength by intercritical 
rolling a high fraction of ferrite is beneficial, while for toughness reason a low fraction of ferrite seems 
better. 

Table 30: Overview of the EBSD results on the investigated samples 

EBSD  was  performed  on  several  samples,  determining  the  fraction  of  2nd  phase  to  be  able  to 
determine the grain size of the single ferrite. Both, the grain size of the total microstructure measured 
as  well  as  the  ferrite  grain  size  is  given  in  Table  30.  Except  for  the  austenitic  rolled  and  direct 
quenched set, the ferrite grain size determines the total grain size. When having a look on the ferrite 
grain evolution along the investigated process, independent of the composition, the ferrite grain size 
is increasing with decreasing  temperature. The average ferrite grain size is also increasing, when 
intercritical rolling is applied (see CMn (0.06%C)). This indicates, that no recrystallization is taking 
place.  During  air  cooling,  see  CMn  (0.12%C),  the  ferrite  grain  size  is  even  further  increasing. 
Furthermore, the grain size of the NbV alloyed material is smaller than of CMn (0.12%C). 

97 

 
 
 
 
 
 
 
98 

 
 
 
 
WP7: Definition of the optimum processing conditions – Link between thin gauge 
homogeneous material behavior and real heavy gauge process 
The objectives of the WP7 are: 

- 
- 

- 

To model the microstructure evolution and relations with data from previous WPs. 
To provide guidelines for the optimum process window for the production of high strength 
steel plates with an optimum strength-toughness balance based on all the results gathered 
in the previous WPs. 
To validate the guidelines with pilot tests. 

TASK 7.1 Micromechanical modelling 

From previous studies, it has been known that the evolution of microstructure and texture during 
intercritical rolling is not a simple phenomenon. Intercritical rolling is not just a hot rolling process 
for materials of dual phase ferrite-austenite microstructures. More than frequent, steels are rolled at 
high  temperature,  in  the  fully  austenitic  condition.  Depending on  the  thickness  reduction  and  the 
rolling temperature, after the rolling, austenite could be presence as deformed or recrystallized state. 
Even if the rolling temperature for fully austenitic rolling is below the non-recrystallized temperature 
(Tnr),  the  holding  or  slow  cooling  from  high  temperatures  could  promote  recovery  and/or  static 
recrystallization of austenite. Apparently, for intercritical rolling, when the steel is cooled down to a 
specific intercritical temperature after fully  austenitic rolling, transformation of austenite to ferrite 
will occur. The rolling at intercritical temperatures, therefore, will lead to deformation of both ferrite 
and austenite. Finally, the cooling from intercritical temperature to room temperature will define the 
final microstructure and texture of intercritical rolled samples. Along the whole complicate process, 
there  are  three  important  physical  phenomena  occurring  in  the  materials:  plastic  deformation, 
recrystallization and phase transformation. To simulate all phenomena during this complete process 
is  beyond  the  scope  of  this  study.  Nevertheless,  with  the  main  focus  on  intercritical  rolling  step, 
crystal plasticity models to simulate the plastic deformation and strain partitioning between ferrite 
and austenite will be intensively used. Recrystallization and phase transformation models to simulate 
texture evolution of ferrite and austenite during intercritical rolling, although presented here, are just 
used for illustration purposes. 

FCT 

VPSC 

ALAMEL 

Figure 101: Schematic representations of different crystal plasticity models, which can be 
applied to predict evolution of deformation textures in rolled plates. 

To  account  for  the  development  of  deformation  textures,  various  crystal  plasticity  models  are 
available.  From  the  full  constraints  Taylor  (FCT)  model  to  the  visco-plastic  self-consistent  (VPSC) 
model  and  the  advanced  Lamel  (ALAMEL)  model,  crystal  deformation  and  rotation  are  argued  by 
dislocation  activities  responding  to  geometry  constraints  of  mechanical  forces  (Figure  101). 
Differences  among  these  models  are  the  consideration  of  crystal  aggregates.  In  the  FCT  model 
[Tayl38],  polycrystals  are  assumed  to  be  non-interactive  sets  of  single  crystals.  Deformation  of 
individual  crystals,  therefore,  is  equal  for  each  grain  and  exactly  equal  to  the  macroscopic 
deformation. This condition, obviously, is non-realistic, particularly in the case of rolling. Since the 
single crystals in the polycrystalline aggregate are interacting systems, stress equilibrium and strain 
compatibility must exist between neighboring grains. Conceptually, the micro-stress and micro-strain 
conditions  need  to  be  maintained  along  the  grain  boundaries.  Therefore,  the  idea  of  the  ALAMEL 
model [VanH05] is that the polycrystal can be divided into pairs of co-deformation regions. For pairs 
of crystals along boundary segments, stress equilibrium and strain compatibility are satisfied, and 
the deformation applied to pairs of crystals is equal to the macroscopic deformation, irrespective of 
99 

 
 
 
 
the  misorientation  and  the  inclination  of  the  boundary  between  the  two  crystals.  Moreover,  if 
deformation  of  a  crystal  is  affected  by  all  other  crystals  in  the  same  aggregate,  regardless  their 
physical  locations,  crystals  can  be  considered  as  (spherical/ellipsoid)  inclusions  in  the  averaged 
medium formed by all crystals of the aggregate. Therefore, deformations of individual crystals in the 
VPSC model [Lebe93] are self-consistent. Stress equilibrium and strain compatibility are maintained 
on entire grains, not just along the grain boundaries. This stress equilibrium and strain compatibility 
are  between  individual  crystals  and  the  averaged  medium,  which  hardly  ever  occurs  in  real 
microstructures. Nevertheless, in this study, the VPSC model is chosen to simulate the evolution of 
deformation textures after rolling. A particular advantage of this model is the ability to predict the 
formation of deformation textures in multi-phase materials. Evolution of grain morphology and strain 
partitioning among grains and between phases are output of the model as primary results of the self-
consistent problem. 

For the first prediction of austenite texture after rolling, the initial texture is assumed to be random. 
This initial texture is represented by a set of 10000 discrete orientations. Deformation and rotation 
of  austenite  grains  are  accounted  for  by  glide  of  the  {111}<110>  slip  systems.  The  material  is 
considered to be fully plastic and have no hardening behavior. The macroscopic deformation is plane 
strain  compression  (PSC)  with  the  total  equivalent  strain  (εvM)  of  1.61,  corresponding  to  80% 
thickness reduction of roughing and finish passes. The simulated deformation texture for austenite 
after the rolling of 80% thickness reduction is given in Figure 102. The simulated deformation texture 
of austenite includes components of the β-fiber. The maximum intensity peak in this texture is located 
at the Copper ({112}<111>) orientation, with 8.3 mrd. Both S-({123}<634>) orientation and Brass 
({110}<112>) orientation are observed in this simulated texture with the intensities of 7.0 and 6.8 
mrd respectively. The β-fiber components are typical features of the rolling texture in face centered 
cubic (FCC) structured materials. For low carbon steels, this texture might be observed in austenite, 
if this phase is retained at room temperature and the volume fraction is high enough for statistical 
measurements. 

Figure  102:  Simulated  deformation  texture  for  austenite  after  the  rolling  to  reduce  the 
plate  thickness  of  80%.  Reference  orientations:  ●  {112}<111>,  ▲  {110}<112>  and  ■  
{110}<001>. From left to right 2 sections at 45°, 65° and 90° (see upper right corner of 
the ODF). 

Since the total strain imposed by rolling is quite high, the stored energy in deformed austenite grains 
is  expected  to  be  large  enough  for  recrystallization.  The  recrystallization  of  austenite  does  not 
necessarily occur during rolling (dynamic recrystallization). It could happen during the slow cooling 
to room temperature. Therefore, the formation of recrystallization texture in austenite needs to be 
considered. To predict the development of recrystallization texture, a statistical model proposed by 
Kestens  &  Jonas  [Kest96]  is  employed.  Schematically,  this  recrystallization  model  includes  two 
stages;  nucleation  and  growth.  In  each  stage,  probabilities  for  individual  crystals  in  deformed 
samples to nucleate and then grow as recrystallized crystals are evaluated. In the first stage, the 
nucleation, chances to nucleate for all deformed crystals are calculated. Nucleation probabilities of 
all  deformed  crystals  are  assumed  to  be  dependent  on  the  stored  energy  of  plastic  deformation, 
which in turn is proportional to the Taylor factor of the crystal plasticity  [Hutc99],[Hutc07]. In the 
second stage, probabilities to grow into other deformation orientations are calculated for the nuclei. 
Depending on the growth rule, the growth probabilities of nuclei are different and thus the formation 

100 

 
 
 
of recrystallization texture depends on the selected growth rule. This growth rule could be simply the 
equal migration rate for all high angle boundaries, or the high migration rate for special boundaries. 

Figure  103:  Simulated  recrystallization  texture  of  austenite  according  to  the  low  stored 
energy nucleation and the high growth rate of the <111>40o grain boundary. Reference 
orientations:  ●  {001}<100>,  ▲  {110}<112>  and  ■   {110}<001>.  From  left  to  right  2 
sections at 45°, 65° and 90° (see upper right corner of the ODF). 

For FCC structured materials, the higher nucleation probability is commonly attributed to crystals of 
the lower Taylor factor, meaning the ones with less stored energy [Dohe97]. The higher the Taylor 
factor of a crystal the lower chance it has to act as potential nucleus. In addition, the high mobility 
rate is usually defined for the <111>40o boundary. These criteria, therefore, are used to predict the 
development  of  recrystallization  texture 
in  austenite.  Figure  103  shows  the  simulated 
recrystallization  texture  of  austenite,  according  to  the  low  stored  energy  nucleation  and  the  high 
mobility  of  the  <111>40o  boundary.  The  Cube  orientation  has  the  highest  intensity  in  the 
recrystallization texture, of 14 times of the random level. This recrystallization texture is well known 
for FCC materials after the conventional rolling and annealing. 

Independently  from  the  recrystallization  of  austenite,  the  transformation  to  ferrite,  bainite  or 
martensite always occurs during the cooling of the studied material to room temperature. Morphology 
and chemical composition of product crystals are functions of the transformation mechanisms, which 
depends on  the cooling rate for each specific material. Nevertheless, the requirement to keep energy 
of the new phase lower than that of the old phase guides the development of product crystals into 
preferential  directions.  Specifically,  during  nucleation  and  growth,  product  crystals  gain  energy 
advantages when they have an orientation relationship to the parent crystals. For both diffusional 
and  displacive  transformations  in  steels,  the  (1-10)bcc  plane  and  the  [111]bcc  direction  of  product 
crystals are frequently observed to coincide with the (111)fcc plane and the [01-1]fcc direction of the 
parent  crystals.  This  most  frequently  observed  orientation  relationship,  so  called  the  Kurdjumov-
Sachs orientation relationship (K-S OR), may vary slightly by steel composition and/or micro-strain 
state of crystals. However, for texture simulations, the tolerances from the K-S OR do not bring any 
significant effects, and thus in this study the K-S OR will be applied to predict transformation textures. 
The orientation transformation from FCC to BCC is simply described as a rotation of 90o around the 
<112> axes of parent crystals. Because of the crystal symmetry, this rotation is not unique. For each 
individual austenite crystal, there will be 24 possible product crystals of ferrite. Orientations of these 
products  on  the  parent  coordinate  system  and  the  K-S  orientation  relationship  rotation  are 
schematically given in Figure 104. 

101 

 
 
 
(a) 

Figure 104: Orientations of 24 ferrite products on the parent austenite coordinate system 
(a) and the K-S OR rotation between austenite and ferrite (b) 

(b) 

In  martensite  microstructures,  all  24  product  crystals  are  frequently  observed,  whereas  in 
polygonal ferrite microstructures, only few variants are present. The number of variants in ferrite 
microstructures is less than 24 due to variant selection during phase transformations. There are 
several  rules  for  variant  selection  during  phase  transformation.  However,  since  it  is  not  clear 
which rule is the most probable in this specific material,  variant selection is not considered in 
texture predictions. 
Figure 105 shows the simulated transformation textures of ferrite from deformed austenite and 
recrystallized  austenite.  The  transformation  ferrite  texture  from  deformed  austenite  (Figure 
105a)  includes  transformed  Copper  and  transformed  Brass  orientations.  The  {113}<110> 
orientation  has  the  highest  texture  intensity  of  2.1  mrd.  This  orientation  is  a  product  of  the 
Copper orientation, which has the highest intensity in the deformed austenite texture. Another 
product of the Copper orientation is the rotated Goss ({110}<110>) orientation is also present 
in the transformation ferrite texture, although the intensity of this orientation is rather weak, 1.2 
times  the  random  level.  Transformed  Brass  orientation  is  clearly  observed  at  the  Dillamore 
component  as  a  local  high  intensity  peak  of  1.8  mrd.  The  transformation  ferrite  texture  from 
recrystallized austenite (Figure 105b) contains the rotated Cube ({001}<110>) orientation, the 
Goss  ({110}<001>)  orientation  and  the  rotated  Goss  orientation.  All  three  orientations  are 
products of the Cube austenite orientation. Among them, the rotated Cube orientation has the 
highest intensity of 1.5 mrd. 

(a) 

(b) 

Figure 105: Simulated transformation textures of ferrite from deformed austenite (a) and 
from recrystallized austenite (b). Reference orientations: ● {112}<110>, ▲ {554}<225> 
and ■  {110}<110>. ODFs in  2= 45° section of the Euler space. 

When  intercritical  rolling  is  applied,  the  evolution  of  microstructure  and  texture  follows  the  most 
complicate  scenario  because  both  ferrite  and  austenite  will  be  deformed.  The  strain  partitioning 
between  these  two  phases  at  various  intercritical  temperatures  will  lead  to  different  evolution  of 
microstructure and texture. Deformations of prior-deformed and recrystallized austenite grains are 
not  similar.  Transformation  of  these  austenite  grains,  during  or  after  the  intercritical  rolling,  also 
adds a factor of complexity to the evolution of microstructure and texture. In addition, intercritically 
deformed ferrite may recrystallize during or after the rolling. All these phenomena make quantitative 
description  of  ferrite  microstructure  and  texture  formation  after  intercritical  rolling  extremely 

102 

 
 
 
 
 
 
 
 
complex problem. Simplification, therefore, is necessary to analyze ferrite textures measured at room 
temperature of the intercritically rolled samples. 

It is assumed that phase interaction and dynamic processes do not occur during intercritical rolling. 
This assumption implies independent evolutions of ferrite and austenite during intercritical rolling. It 
largely  reduces  the  complexity  of  microstructure  and  texture  evolution  after  intercritical  rolling. 
Nevertheless, it also ignores most of essential features for intercritical processes. Strain partitioning 
between  ferrite  and  austenite  is  suppressed,  and  thus  each  phase  undergoes  the  same  strain  as 
imposed by the rolling. Phase fractions and grain states (deformed or recrystallized) are constant 
during the intercritical rolling and the microstructures after intercritical rolling is described only  by 
the deformation state of both ferrite and austenite after the intercritical rolling. Regardless the initial 
state of deformation or recrystallization, the rolling pass of 25% thickness reduction always brings 
the  β-fiber  texture  to  the  austenite  after  the  intercritical  rolling.  The  transformation  to  ferrite  of 
intercritically deformed austenite is expected to be similar to the one of fully deformed austenite and 
thus results in the ferrite texture as given in Figure 105a. Developments of intercritical ferrite after 
the intercritical rolling, however, depend on the state of the prior austenite. Rolling simulations of 
25%  thickness  reduction  for  intercritical  ferrite  have  been  carried  on  by  the  VPSC  model.  Initial 
textures of these simulations are from recrystallized and from deformed austenite. Figure 106 shows 
simulated  textures  of  the  intercritical  ferrite  after  intercritical  rolling.  Both  of  the  two  simulated 
textures contain components of the α-(<110>//RD) and the γ-(<111>//ND) fibers, which are typical 
for rolling texture in BCC structured materials. However, because of differences in initial stages, high 
intensity  peaks  in  these  textures  are  located  at  various  orientations.  For  the  intercritical  ferrite 
texture  originated  from  recrystallized  and  transformed  austenite,  the  highest  intensity  peak  is 
observed  near  the  rotated  Cube  orientation,  with  3.1  mrd.  On  the  contrary,  the  highest  intensity 
peak in the intercritical ferrite texture from deformed and transformed austenite is located at the 
{112}<110> orientation, with 3.8 mrd. Along the γ-fiber, the highest intensity peak in the former 
texture  is  measured  at  the  {111}<110>  orientation,  while  in  the  later  at  the  {111}<112> 
orientation. 

(a) 

(b) 

Figure 106: Simulated deformation textures of intercritical ferrite after the rolling of 25% 
thickness reduction from (a) recrystallized and transformed austenite and (b) deformed 
and transformed austenite. Reference orientations: ● {112}<110>, ▲ {111}<110>. ODFs 
in 2= 45° section of the Euler space. 

TASK 7.2 Effect of temperature gradient 

The  previous  tests  (WP3-6)  have  been  performed  in  well  controlled  conditions  aiming  at  a  good 
understanding of each individual process parameter. In industrial practice however, there is not one 
γ/α balance during deformation, but the fractions of the phases will vary along the product thickness 
due to the temperature gradient present.  

The effect of this temperature gradient will be assessed by performing trials on the fully automated 
pilot mill producing samples with a higher end-thickness. In the end, rolling of 25mm was realized 
for the NbV grade as they have most industrial relevance. Also, rolling of even 36mm thickness was 

103 

 
 
 
 
 
realized  for  the  CMn  (0.06%C)  grades on  which  most  investigations  were  made.  The  plates  were 
instrumented with thermocouples at subsurface top and bottom as well as at mid thickness, to have 
a detailed mapping of the temperature over the thickness. Furthermore, temperature reading of the 
top surface temperature by infrared cameras are available. This information will be coupled to the 
information on the phase fractions in function of temperature from WP3 and WP6.  

In the previous work packages it was seen, that rolling at low intercritical temperature is good for 
strength  while  high  intercritical  temperatures  are  better  for  toughness.  However,  industrially  one 
single  pass  with  high  intercritical  reduction  is  not  possible  but  multiple  passes  are  given  instead. 
During the interpasses a natural drop in temperature is occurring. Hence, it was decided, to first split 
the intercritical rolling into 2 passes and 3 passes for 25mm and 36mm, respectively; secondly to 
perform the first intercritical rolling pass at high intercritical temperature. After rolling, the plates 
were air cooled to room temperature. The rolling schemes were designed in that way, that the total 
as well as the sequence rolling reduction ratios are comparable for the two plates. 

Table 31 gives the target slab dimension as well as the final plate dimension in mm and the different 
rolling reduction ratios. A width of >300mm was targeted to machine BDWTT samples directly from 
the plate. However, rolling such a high width limited the pass reduction due to mill force reasons. 
Figure 107 shows the successive rolled plate of 36mm thickness. Figure 108 illustrates the target 
core temperature and pass reduction, and Figure 109 shows the realized temperatures measured by 
thermocouples at core, top and bottom subsurface and surface temperature measured by infrared 
(IR) camera. It is worth noting, that the temperature gradient is not symmetrically since the bottom 
surface is slightly colder than the top surface. The difference in temperature is caused by the contact 
of the plate with the rolling table.  

Table 31: Slab and final plate target dimension for rolling to (left) 25mm and (right) 36mm 

Figure 107: Rolled plate of CMn (0.06%C) with final thickness of 36mm after air cooling 

104 

 
 
 
 
Figure 108: Overview of target core temperature (blue) and pass reduction (orange) for 
the plates rolled to 25mm (left) and 36mm (right) 

Figure  109:  Temperature  recording  across  thickness  of  each  pass  for  25mm  (left)  and 
36mm (right) 

The target was to roll multiple plates per thickness to have sufficient material to sample tensile and 
Charpy  samples  at  different  thickness  in  longitudinal  and  transverse  direction.  However,  due  to 
circumstances, only 1 good plate per thickness was finally realized. Hence, it was decided to sample 
transverse direction only and at top, core and bottom as well as full thickness. Battelle testing was 
already intentionally targeted in transverse direction and full thickness but needed to be limited to 
one test temperature. 

TASK 7.3 Microstructural characterization over thickness 

The 25mm plate of NbV was investigated through thickness at subsurface top, quarter, mid, three 
quarter  thickness  and  subsurface  bottom.  The  additional  36mm  plate  of  CMn  (0.06%C)  was 
investigated only at subsurface top, quarter and mid thickness. The light optical micrographs together 
with the EBSD maps of high and low angle grain boundaries and Kernel are shown in Figure 110 and 
Figure 111. For both grades and thickness about 12-13% pearlite is measured, due to the same C-
level.  For  both  grades  a  bimodal  grain  structure  is  seen  which  appears  more  pronounced  for  the 
25mm NbV plate. Furthermore, the overall average grain size is smaller and heavily substructured 
in bands. For the 36mm CMn (0.06%C) plate the grains are larger and less substructured, indicating 
a higher level of restoration. Grain size distributions measured at top, quarter, mid-thickness regions 
of  the  CMN-0.06%C  steel  are  given  in  Figure  112.  An  overview  of  the  phase  fraction  and  overall 
average grain size is given in Table 32. It must be stated, that for the determination of grain size 
the  pearlite  could  not  be  excluded.  Furthermore,  for  the  grain  determination  a  grain  boundary 
completion  down  to  2°  was  allowed  with  a  minimum  area  of  4  pixels  (1  µm²).  This  causes  the 
detection of a high amount of small grains, resulting in a too small average grain size in comparison 

105 

 
 
 
 
to  the  size  expected  based  on  optical  images.  However,  the  procedure  of  scanning  and  post 
processing was in line with the EBSD measurements performed in the previous work packages.  

Using the GOS criteria, the grains were differentiated into recrystallized ferrite for GOS≤2°, above 
2° the ferrite grains are assumed to be substructured and deformed. For both plates and all positions, 
a recrystallized ferrite fraction of ~10% was determined with an average size of about 3µm for the 
25mm  NbV  and  4-5µm  for  the  CMn  (0.06%C)  plate.  The  un-recrystallized  average  grain  size  is 
slightly  larger  for  NbV  with  ~6µm  but  significantly  larger  for  CMn  (0.06%C)  with  12-13µm.  The 
different grain sizes measured through thickness are rather constant.   

Both materials have in common the absence of pearlite at the top and bottom surfaces. This is caused 
by  superficial  decarburization.  This  phenomenon  is  often  observed,  reaching  a  depth  of  typically 
~100µm. Hence, in combination with the lower surface temperature, the microstructure appears as 
cold deformed.  

25mm 

NITAL – LOM 

EBSD 

top 

¼ 

½ 

¾ 

106 

 
 
 
 
 
 
 
 
 
 
bottom 

Figure 110: Through thickness microstructure  of the 25mm NbV plate  

36mm 
top 

NITAL - LOM 

EBSD 

1/4 

1/2 

¾ 

107 

 
 
 
 
 
 
 
 
 
 
 
 
bottom 

Figure 111: Through thickness microstructure  of the 36mm CMn (0.06%C)  

Figure 112: Grain size distribution of 36mm CMn (0.06%C) plate 

Table 32: Overview on phase fraction and grain size 

108 

 
 
 
 
 
 
 
 
25mm 

NITAL – LOM 

GB 5 red GB 15 black 

Kernel (0-5°) 

EBSD 

top 

¼ 

½ 

¾ 

bottom 

Figure 113: Through thickness microstructure of the 25mm NbV plate, ND-RD orientation 

109 

 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
110 

 
 
 
 
36mm 

NITAL - LOM 

GB 5 red GB 15 black 

Kernel (0-5°) 

EBSD 

top 

1/4 

1/2 

¾ 

bottom 

Figure  114:  Through  thickness  microstructure  of  the  36mm  CMn  (0.06%C),  ND-RD 
orientation 

111 

 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
The precipitation fraction of Nb and V was determined by ICP-MS and the results are given in Table 
33. From previous investigation (WP6) it is known that about 2/3 of Nb and 1/5 of V is precipitated 
at 750°C intercritical rolling temperature, followed by direct quench. At that moment of investigation, 
it was suggested, that precipitation level can further increase during slow air cooling. The results on 
heavy gauge proofs this thesis with precipitation levels for V with about 26% and for Nb with 91%. 
Due to technical circumstances it was impossible to measure the particle size through thickness by 
TEM. 

Table 33: Average precipitation fraction of 25mm NbV plate 

TASK  7.4  Mechanical  characterization  over  thickness  and  the  correlation  with 
microstructure 

The  mechanical  properties  were  determined  by  tensile  and  Charpy  testing  at  subsurface  top  and 
bottom and full thickness with samples of 10mm thickness. Full thickness was tested by tensile and 
Battelle. Test direction in all cases was transverse to the rolling direction due to the unforeseen lack 
of material. 

Figure  115  shows  the  tensile  properties  in  transverse  direction  measured  at  top  and  bottom 
subsurface,  mid  thickness  and  full  thickness.  For  both  plates  the  tensile  strength  is  less  position 
dependent, that means only small variation of ±10 MPa for 36mm and ±23MPa for 25mm; with core 
and full thickness slightly weaker than the subsurface positions. Yield strength behaves comparable 
but  with  bigger  differences,  i.e.  ±25MPa  for  36mm  and  ±32MPa  for  25mm.  The  rather  small 
differences through thickness are explainable by the small differences in grain size. The cause of the 
weaker  core  can  be  explained  by  the  slower  cooling,  causing  better  restauration  processes,  i.e. 
recovery and reduction of dislocation densities. The elongation is for both plates constant through 
thickness, while full thickness testing results in even better elongation. The 36mm CMn (0.06%C) 
grades shows better elongation than the 25mm NbV grade. Also, the Charpy toughness behavior for 
the 36mm CMn (0.06%) plate is better with ductile to brittle transition temperatures at or even below 
-80°C, Figure 116. However, the top and bottom position shows less good toughness than the core. 
In comparison, for 25mm the three positions tested are rather comparable with DBTT at around  -
60°C. The difference in toughness between the 36mm and 25mm plate is probably caused by the 
bimodal  microstructure,  which  is  severely  substructured  for  the  25mm  but  well  restored  for  the 
36mm. The larger average grain size of the 36mm plate is in this case less detrimental.  

112 

 
 
 
Figure 115: Average tensile strength in transverse direction  

Figure 116: Charpy transition curves in transverse direction 

Battelle  drop  weight  test  was  performed  at  -10°C.  None  of  the  samples  fulfilled  the  linepipe 
requirement of min 85% shear area fracture. Furthermore, the Battelle set-up (1.6t hammer falling 
from 2.47m height) was for one sample of 36mm not powerful enough to even break the sample. 
Pictures of the broken samples are given in Figure 117. The poor Battelle drop weight toughness can 
be  linked  to  the  strong  bimodal  microstructure,  in  particular  for  the  25mm  NbV  plate,  and  large 
average grain size, in particular for the 36mm CMn (0.06%C) plate. 

113 

 
 
 
 
 
 
 
25mm NbV at -10°C 

36mm CMn (0.06%C) at -10°C 

34% SA 

27%SA 

Figure 117: Fractured BDWTT 

22% SA 

TASK 7.5 Definition of the guidelines 

Heavy gauge structural and line pipe steel plates are usually rolled just above or in the two-phase 
region, particularly in the last rolling pass. Hot rolling below the Ar3 temperature is known to increase 
the strength of materials. However, the strength and toughness of steels usually develop in opposite 
directions,  meaning  the  gain  in  strength  always  goes  with  a  loss  in  toughness  of  the  materials. 
Compromised values of the two properties, on the other hand, are not well established, particularly 
for  steels  produced  in  the  intercritical  region.  Through  various  work  packages  from  WP2  to  WP6, 
different  aspects  of  intercritical  rolling  have  been  systematically  considered.  Impacts  of  the 
processing  parameters  on  the  strength  and  toughness  of  materials  are  also  evaluated.  With  the 
complete data set collected in this project, a guideline for steel production in two-phase region could 
be given as following: 

d, Ac1

Among the most important parameters for hot rolling, especially for intercritical rolling, are critical 
d). In WP2, critical temperatures of CMn-0.06%C, CMn-0.12%C 
temperatures (Tnr, Ac3, Ac1, Ac3
and NbV steels were carefully and systematically investigated. Equilibrium critical temperatures of 
A1 and A3 for the CMn-0.06%C and NbV steels are calculated as 682°C, 841oC, 681°C and 849oC. 
The corresponding temperatures for the CMn-0.12%C steel are 692oC and 826oC. However, these 
temperatures will never be reached in real thermomechanical experiments. The faster the cooling 
rate, the lower Ac1 and Ac3 temperatures of each materials. Additionally, the austenite states after 
deformation or recrystallization also affect the austenite to ferrite transformation and thus the critical 
temperatures.  Generally,  critical  temperatures  increase  with  the  amount  of  stored  energy  in 
austenite  grains  after  plastic  deformation.  When  plastic  strain  of  austenite  is  considered,  the 
transformation from austenite to ferrite becomes more complicated to consider. Not only the stored 
strain energy in austenite and the cooling rate are the defining factors for the phase transformation. 
Strain  rate  by  which  austenite  is  deformed  is  also  one  of  the  important  factors  to  define  critical 
temperatures  for  the  studied  materials.  Uniaxial  compressive  experiments  on  plasto-dilatometer, 
although allow highly accurate temperature determination, could not simulate the high strain rates 
of the real hot rolling trials. Interrupted hot torsion experiments have advantages on high strain rate 
and  large  strain  magnitude,  as  compared  to  dilatometry  experiments.  However,  the  shear  strain 
mode  and  the  discrete  nature  of  hot  torsion  experiments  do  not  account  for  the  true  physical 
phenomenon during hot rolling. Additionally, resolution for critical temperature determination by hot 
torsion is not as high as by dilatometry. Therefore, critical temperatures of the three materials were 

114 

 
 
 
 
 
 
 
systematically  measured  and  validated  by  all  three  partners  (CRM,  OCAS  and  CEIT).  Besides  the 
critical  temperatures  determined  for  each  individual  material,  an  important  feature  is  the 
temperature range for intercritical region. For all three materials after the fully austenitic deformation 
d  are  mostly  within  50°C.  For 
of  0.3,  the  difference  between  critical  temperatures  of  Ac3
d temperatures of the CMn-0.06%C after a fully austenitic deformation of 
d and Ac1
instance, the Ac3
0.3 are 790oC and 750°C. 

d  and  Ac1

Rolling temperature is a particularly important parameter for intercritical rolling. The lower rolling 
temperature, the larger volume fraction of intercritical ferrite and thus the higher rolling force need 
to  be  applied  to  deform  material.  However,  the  higher  strength  of  steels  by  this  way  could  be 
achieved.  In  WP3,  single  intercritical  rolling  pass  of  25%  thickness  reduction  at  different 
temperatures  were  applied  to  the  CMn-0.06%C  material.  The  tensile  and  yield  strength  of  the 
samples rolled at lower temperatures are indeed higher than the ones rolled at higher temperatures. 
Ductility and toughness of the material, however, decrease with the reduction of intercritical rolling 
temperature. Mechanical properties of intercritical rolled samples, additionally, are influenced by the 
cooling rate applied to the material after rolling. The strength of the material in general increases 
with the cooling rate from air cooling to accelerated cooling and direct quench. By applying the higher 
cooling rate, the yield strength of CMn-0.06%C steel produced by intercritical rolling gains 50-100 
MPa, while the tensile strength increases 200-250 MPa. The increase in cooling rates, however, brings 
detrimental effects on the toughness of material. The faster cooling rate after intercritical rolling at 
lower  temperature  reduces  the  upper  shelf  energy  while  increases  the  ductile-brittle  transition 
temperature.  Only  in  the  case  of  intercritical  rolling  at  730°C  and  then  air  cooling  to  room 
temperature, the USE is still similar to the one of higher intercritical rolling temperature, while the 
DBTT is the lowest among all values of samples cooled by air at different intercritical temperature. 

Strain  partitioning  between  ferrite  and  austenite  during  intercritical  deformation,  although  an 
important phenomenon during intercritical rolling, is a complex micromechanical behaviour. As the 
first order estimation, plasticity of CMn-0.06%C steel at intercritical temperature is investigated with 
“single  phase  materials”.  These  single-phase  materials  have  the  same  chemical  compositions  as 
equilibrium  ferrite  and  austenite  of  corresponding  composite  materials  at  desired  intercritical 
temperatures.  Mechanical  behaviour  of  single-phase  materials  are  separately  measured.  The 
mechanical  behaviours  of  composite  materials  afterward  are  calculated  by  weighted  average 
behaviours of the corresponding single-phase materials. Although its principle is true, this approach 
does not work in reality. Austenite and ferrite at any intercritical temperatures are results of phase 
decomposition either during cooling or during reheating. Phase interaction and chemical composition 
redistribution between austenite and ferrite are continuous phenomena. Therefore, the equilibrium 
state  at  a  certain  intercritical  temperature  of  the  composite  material  is  far  to  be  reached  within 
practical time frame. On the other hand, the equilibrium state of  single-phase materials could not 
represent  for  the  state  of  corresponding  ferrite  and  austenite  phases  at  desired  intercritical 
temperature. As a result, mechanical behaviours of ferrite and austenite within the intercritical region 
are  better  estimated  by  extrapolation  of  mechanical  behaviours  of  each  phase  from  outside  the 
intercritical  region.  This  extrapolation  method,  however,  neglects  the  presence  of  second  phases 
(cementite, carbide …) below the intercritical region and only accounts for the strain partitioning at 
the beginning of intercritical deformation. Evolution of the strain partitioning between austenite and 
ferrite during deformation is still not able to capture. 

Strain  distribution  and  strain  accumulation  during  intercritical  rolling  although  expected  to  have 
significant  effect  on  microstructures  and  mechanical  properties  of  the  CMn-0.06%C  steel  did  not 
bring  any  special  result.  Not  many  different  features  in  microstructures  of  the  intercritical  rolled 
samples have been observed. Ferrite grains in these microstructures are all in the range of 12-17 
μm. Intercritical rolling of multi deformation passes, however, produces slightly finer microstructures 
than the single pass intercritical rolling. In-grain orientation gradients and low angle grain boundaries 
are visible in ferrite grains of the samples deformed by multi intercritical rolling passes and at lower 
rolling temperature. The increase in number of intercritical rolling passes from 1 to 3 passes brings 
50 MPa to the tensile strength of the CMn-0.06%C material. However, the total elongation of this 
material reduces about 2%, and the upper shelf energy (USE) measured the Charpy V-notched tests 
decrease 70 J. For the effect of strain magnitude, microstructures and mechanical properties of the 
two  intercritical  rolled  samples  with  25%  and  40%  thickness  reduction  in  single  rolling  passes  at 
745oC were compared. As the thickness reduction increases from 25% to 40%, ferrite grains in the 
115 

 
 
CMn-0.06%C steel become more elongated. Sub-grain structures in deformed ferrite grains are also 
visible, particularly in the sample of 40% thickness reduction. The yield strength and tensile strength 
of  the  sample  with  40%  thickness  reduction  are  both  higher  than  those  of  the  sample  with  25% 
thickness reduction. On the contrary, upper shelf energy measured by Charpy V-notch test of the 
later sample is higher than the corresponding value of the former sample. 

Chemical  composition  (C  and  micro  alloying)  content  effects  microstructure  and  thus  mechanical 
properties of intercritical rolled samples. Irrespective to chemical composition, with the increase of 
cooling rate from air to direct quench, strength is increasing and the ductile to brittle temperature is 
increasing, i.e. toughness is decreasing. This can be linked to the fraction and kind of second phase 
– small amount of pearlite and high fraction of martensite. With lower intercritical temperature, that 
means with higher amount of ferrite, the strength is slightly increasing due to the higher strength of 
the 2nd phase caused by its C enrichment. With respect to composition, it can be observed, that the 
microalloying of NbV is increasing strength by at least 100MPa by solid solution and precipitation 
hardening, independent of cooling rate. The addition of C is increasing strength significantly for direct 
quench. For air cooling, the strength increases by increasing pearlite fraction is lower than by adding 
NbV. NbV significantly  improves ductile to  brittle transition temperature. The results also indicate 
that  toughness  becomes  less  dependent  on  cooling  rate.  The  addition  of  C  is  deteriorating  the 
toughness,  i.e.  ductile  to  brittle  temperature  is  increasing  and  upper  shelf  energy  is  decreasing. 
Toughness improves only by reducing the cooling rate to room temperature, transforming the hard 
martensite phase into ferrite and pearlite. Against expectations, the high C grade shows better ductile 
to brittle transition temperatures for air cooling than the reference grade. 

116 

 
 
 
 
 
117 

 
 
 
 
Conclusions 

CONCLUSION WP1 

In general, the whole project has been carried out according to proposed time and cost schedules. 
Periodical meetings and conference calls were regularly organized to keep all partners well informed 
on activities and progress of the project. The annual, mid-term and final reports are delivered on-
time. Results of the project have been appropriately disseminated. 

CONCLUSION WP2 

Although material production was unexpectedly prolonged until the mid-term of the project, no delay 
has  been  observed  on  any  other  activities.  Massive  laboratory  casting  activities  were  shared  and 
carried on by CRM and OCAS to provide enough materials for all partners. 

Critical temperatures of all three studied materials (CMn-0.06%C, CMn-0.12%C and NbV steels) are 
intensively  investigated  by  CEIT,  CRM  and  OCAS.  These  temperatures  are  determined  by  both 
dilatometry  and  hot  torsion  experiments.  Impacts  of  different  austenite  microstructures,  i.e. 
deformed or recrystallized, are also considered in this study. 

General,  the  determination  of  the  transformation  temperatures  during  cooling  remains  a  difficult 
issue because transformation usually occurs in dynamic conditions and far from equilibrium. It has 
been shown that different techniques and various thermo-mechanical cycles lead to different values 
of critical temperatures. This is due to volume fractions of intercritical ferrite and austenite not only 
depending  on  the  intercritical  temperature,  but  also  on  the  state  of  initial  austenite  before  the 
intercritical deformation. For instance, due to the high Mn content, transformation can take up to 1 
hour  before  achieving  a  steady  state,  which  is  still  far  from  the  equilibrium  one.  It  follows  that 
thermo-mechanical treatments to reach desired intercritical microstructures of 25%, 50% and 75% 
ferrite  may  not  be  reached  precisely  and  homogeneously.  Additionally,  the  austenite  to  ferrite 
transformation  in  all  three  materials  occur  in  a  rather  narrow  range  of  temperatures  (~30°C). 
Therefore,  the  uncertainty  in  temperature  measurement,  as  well  as  thermal  gradient  across  the 
sample thickness, were considered carefully in the subsequent tasks. In the WP3 and WP4, different 
schedules were applied to achieve the desired intercritical ferrite volume fraction. 

CONCLUSION WP3 

The impact of intercritical rolling on mechanical properties, specifically the impact of the γ/α phase 
fraction on strength and toughness of the CMn-0.06%C steel is the objective the WP3. Two cooling 
strategies (i.e. air cooling and accelerated cooling) were applied for intercritical rolled plates. 

The intercritical rolling trials performed confirm the trends observed in the industrial practice. The 
YS increase is of the  same order  of  magnitude as observed industrially. The increase in YS being 
larger in TS, the TS/YS ratio reduces as the FRT decreases. The work-hardening ability is therefore 
somewhat reduced after intercritical of ferritic rolling. 

The  plate  strengthening  is  however  associated  to  a  decrease  in  toughness,  especially  when 
accelerated cooling or direct quench is applied. 

Texture  does  not  seem  to  be  the  main  parameter  dictating  toughness:  very  different  toughness 
properties  are  measured  for  very  similar  textures.  On  the  contrary,  a  clear  relationship  between 
DBTT and the grain size is established. The orientation unit θ>15° is found to be the effective grain 
size  governing  ductility.  Accelerated  cooling  does  not  allow  ferrite  recrystallization  and  the  ferrite 
grain size after IC rolling is larger than that obtained by accelerated cooling of deformed austenite.  

Applying the last pass in the fully ferritic range (which is usually not done in the industrial practice) 
seems  to  combine  plate  strengthening  and  toughness.  One  explanation  is  the  possible  ferrite 
recrystallization taking place during the slow cooling in air. It can be speculated that to refine the 

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ferritic grains (the only metallurgical strategy to improve at the same time strength and toughness), 
two processing routes are possible: 

1.  Rolling in the austenite range followed by accelerated cooling. 
2.  Rolling in fully ferritic range followed by fast cooling to suppress ferrite recrystallization and 

keep the refined ferrite microstructure. 

CONCLUSION WP4 

Monitoring  and  mapping  the  strain  partitioning  between  austenite  and  ferrite  during  intercritical 
deformation is the main goal of WP4. 

The study in this WP confirms that the mechanical behavior of dual-phase microstructures during hot 
deformation is more complex than a simple weight average of austenite and ferrite alone. However, 
the approach of studying each component of the intercritical deformation as “single phase materials” 
was  not  entirely  successful.  Because  the  γ  to  α  phase  transformation  takes  place  well  below  the 
equilibrium temperature, such “single phase materials” are not thermodynamically stable and even 
could transform during deformation. The flow behavior of pure austenite and pure ferrite afterward 
was investigated at several temperatures and extrapolated to the intercritical region. This approach 
would  finally  allow  the  establishment  of  a  strain  partitioning  law  between  γ  and  α  phases  during 
intercritical deformation. 

Although  a  more  complex  relationship  between  γ  and  α  phases  during  intercritical  deformation  is 
always expected, the restrictions of intercritical rolling to strains relatively low, a composite model 
under  iso-strain  assumption  was  shown  to  give  very  reasonable  predictions  of  the  flow  behavior 
during intercritical deformation. 

Furthermore,  regarding  the  microstructural  evolution  during  intercritical  deformation,  a  basic 
phenomenological study was developed understanding the impact that process parameters such as 
phase fractions, holding times, effect of deformation have on the intercritical deformation and the 
final microstructure. A quantitative procedure of microstructural characterization was developed in 
the frame of this project using EBSD technique. The quantification was made upon GOS parameter 
and  grain  boundary  maps,  and  a  successful  differentiation  between  deformed  and  non-deformed 
microstructures was achieved. 

CONCLUSION WP5 

The objective of WP5 is to investigate the effect of strain magnitude and strain path in intercritical 
region on microstructure and mechanical properties of the CMn-0.06%C steel. To complete the whole 
range  of  single  pass  intercritical  rolling  with  different  strain  magnitudes,  lab  rolling  trials  with 
reduction  of  10%  and  40%  were  implemented.  Rolling  trials  of  40%  thickness  reduction  with  2 
intercritical  rolling  passes  were  initially  proposed.  However,  intercritical  rolling  of  high  thickness 
reduction, for instance 20% and 40% per pass, is not technically feasible, particularly at low rolling 
temperatures. Therefore, intercritical rolling trials of 1, 2 and 3 rolling passes with 10% thickness 
reduction  per  pass  were  performed.  Microstructures  and  mechanical  properties  of  these  samples, 
afterward, were measured and analyzed. 

For the study on effect of strain magnitude, microstructures and mechanical properties of the two 
intercritical rolled samples with 25% and 40% thickness reduction in single rolling passes at 745oC 
were compared. As the thickness reduction increases from 25% to 40%, ferrite grains in the CMn-
0.06%C  steel  become  more  elongated.  Sub-grain  structures  in  deformed  ferrite  grains  are  also 
visible, particularly in the sample of 40% thickness reduction. The yield strength and tensile strength 
of  the  sample  with  40%  thickness  reduction  are  both  higher  than  those  of  the  sample  with  25% 
thickness reduction. On the contrary, upper shelf energy measured by Charpy V-notch test of the 
later sample is higher than the corresponding value of the former sample. 

Strain  distribution  and  strain  accumulation  during  intercritical  rolling  although  expected  to  have 
significant  effect  on  microstructures  and  mechanical  properties  of  the  CMn-0.06%C  steel  did  not 
bring  any  special  result.  Not  many  different  features  in  microstructures  of  the  intercritical  rolled 

119 

 
 
 
samples have been observed. Ferrite grains in these microstructures are all in the range of 12-17 
μm. Intercritical rolling of multi deformation passes, however, produces slightly finer microstructures 
than the single pass intercritical rolling. In-grain orientation gradients and low angle grain boundaries 
are visible in ferrite grains of the samples deformed by multi intercritical rolling passes and at lower 
rolling temperature. The increase in number of intercritical rolling passes from 1 to 3 passes brings 
50 MPa to the tensile strength of the CMn-0.06%C material. However, the total elongation of this 
material reduces about 2%, and the upper shelf energy (USE) measured the Charpy V-notched tests 
decrease 70 J. 

CONCLUSION WP6 

Understanding and quantifying the effects of chemical composition (C and micro-alloying) content 
on microstructures and mechanical properties of  intercritical rolled samples is the aim of WP6. To 
achieve this goal, the systematic approach as applied in WP3 an WP5 was used. The number of rolling 
trials in this WP, however, were reduced as compared to the previous ones. Only two intercritical 
temperatures, corresponding to low and high ferrite volume fractions, have been selected for further 
studies.  These  temperatures  of  the  CMn-0.12%C  and  NbV  steels  were  selected  based  on  the 
“relaxation  like”  torsion  experiment,  carried  on  in  WP2.  Besides  the  two  cooling  strategies  of  air 
cooling and accelerated cooling, direct quench was also applied in this WP to investigate the state of 
materials before intercritical rolling. 

Microstructure  investigation  on  samples  of  hot  rolling  and  then  cooling  to  room  temperature  by 
various cooling rates shows different behaviors between CMn-0.12%C and NbV steels. For NbV steel, 
ferrite volume fractions at 770°C and 750°C are mostly similar, irrespective to cooling conditions. 
Only after intercritical rolling and then direct quench, different ferrite volume fractions were obtained 
between the two intercritical temperatures. For CMn-0.12%C steel, equal volume fraction of ferrite 
was  obtained  in  samples  of  intercritical  rolling  at  750°C  and  700°C  and  then  air  cooling  to  room 
temperature. To evaluate the effect of microalloying elements on precipitation, TEM analyses on fine 
precipitates was carried out with carbon extraction replicas. When the transformation occurs from 
recrystallized austenite, a high density of fine precipitates has been detected. These precipitates are 
considered  to  be  formed  in  ferrite  during  or  after  the  intercritical  deformation.  In  the  sample 
corresponding  to  deformed  austenite  strain  induced  precipitates  formed  in  austenite  are  also 
observed. 

Irrespective  to  chemical  composition,  with  the  increase  of  cooling  rate  from  air  to  direct  quench, 
strength  is  increasing  and  the  ductile  to  brittle  temperature  is  increasing,  i.e.  toughness  is 
decreasing. This can be linked to the fraction and kind of second phase – small amount of pearlite 
and high fraction of martensite. With lower intercritical temperature, that means with higher amount 
of ferrite, the strength is slightly increasing due to the higher strength of the 2nd phase caused by its 
C  enrichment.  With  respect  to  composition,  it  can  be  observed,  that  the  microalloying  of  NbV  is 
increasing strength by at least 100MPa by solid solution and precipitation hardening, independent of 
cooling rate. The addition of C is increasing strength significantly for direct quench. For air cooling, 
the strength increases by increasing pearlite fraction is lower than by adding NbV. NbV significantly 
improves ductile to brittle transition temperature. The results also indicate that toughness becomes 
less dependent on cooling rate. The addition of C is deteriorating the toughness, i.e. ductile to brittle 
temperature  is  increasing  and  upper  shelf  energy  is  decreasing.  Toughness  improves  only  by 
reducing the cooling rate to room temperature, transforming the hard martensite phase into ferrite 
and  pearlite.  Against  expectations,  the  high  C  grade  shows  better  ductile  to  brittle  transition 
temperatures for air cooling than the reference grade. 

CONCLUSION WP7 

Objective of the WP7 is to link all research data collected in previous WPs, build a micromechanical 
model for microstructure and texture evolution during intercritical rolling and finally define 
guidelines for future steel production. From previous studies, it has been known that the evolution 
of microstructure and texture during intercritical rolling is not a simple phenomenon. Intercritical 
rolling is not just a hot rolling process for materials of dual phase ferrite-austenite microstructures. 
More than frequent, steels are rolled at high temperature, in the fully austenitic condition. 
Depending on the thickness reduction and the rolling temperature, after the rolling, austenite could 

120 

 
 
be presence as deformed or recrystallized state. Even if the rolling temperature for fully austenitic 
rolling is below the non-recrystallized temperature (Tnr), the holding or slow cooling from high 
temperatures could promote recovery and/or static recrystallization of austenite. Apparently, for 
intercritical rolling, when the steel is cooled down to a specific intercritical temperature after fully 
austenitic rolling, transformation of austenite to ferrite will occur. The rolling at intercritical 
temperatures, therefore, will lead to deformation of both ferrite and austenite. Finally, the cooling 
from intercritical temperature to room temperature will define the final microstructure and texture 
of intercritical rolled samples. Along the whole complicate process, there are three important 
physical phenomena occurring in the materials: plastic deformation, recrystallization and phase 
transformation. Simulating all phenomena during this complete process is beyond the scope of this 
study. Therefore, the visco-plastic self-consistence (VPSC) crystal plasticity model has been chosen 
to simulate the plastic deformation and strain partitioning between ferrite and austenite during 
intercritical rolling. 

121 

 
 
 
 
Exploitation and impact of the research results 

Publications 

  U.  Mayo,  N.  Isasti,  D.  Jorge-Badiola,  J.M.  Rodriguez-Ibabe,  P.  Uranga,  An  EBSD-based 
methodology for the characterization of intercritically deformed low carbon steel, 
Materials 
p31-42 
(https://doi.org/10.1016/j.matchar.2018.10.014) 

Characterization, 

(2019) 

147 

Conference presentations 

  M. Caruso, U. Lorenz, T. Nguyen-Minh, P. Uranga, U. Mayo, N. Isasti, R. Petrov, Intercritical 
rolling of heavy-gauge structural plates: which compromise between strength and 
toughness?, MSE Congress 2018, September 26-28, 2018 Darmstadt, Germany. 

  U. Lorenz, M. Gauvin, Advanced EBSD study of deformation mechanisms during single 
pass  intercritical  rolling,  MSE  Congress  2018,  September  26-28,  2018  Darmstadt, 
Germany. 

  U.  Mayo  Ijurra,  N.  Isasti  Gordobil,  P.  Uranga  Zuaznabar,  J.M.  Rodriguez  Ibabe, 
Microstructural characterization procedure using EBSD technique for intercritically 
deformed  low  carbon  steels,  MSE  Congress  2018,  September  26-28,  2018  Darmstadt, 
Germany. 

  N. Isasti Gordobil, U. Mayo Ijurra, P. Uranga Zuaznabar, J.M. Rodriguez Ibabe, Analysis of 
the microstructural evolution of an intercritically deformed NbV microalloyed steel, 
MSE Congress 2018, September 26-28, 2018 Darmstadt, Germany. 

 

T. Nguyen-Minh, M. Caruso, U. Lorenz, J. Sietsma, R. Petrov, Texture evolutions in heavy 
gauge  steel  plates  after  intercritical  rolling  at  high  temperatures,  MSE  Congress 
2018, September 26-28, 2018 Darmstadt, Germany. 

  M. Caruso, U. Lorenz, T. Nguyen-Minh, P. Uranga, R. Petrov, Impact of intercritical rolling 
on the strength/toughness balance in heavy gauge structural steel plates, THERMEC 
2018, July 9-13, 2018, Paris, France. 

  U. Mayo, N. Isasti, J.M. Rodriguez-Ibabe and P. Uranga, Effect of intercritical deformation 
on  final  microstructure  in  low  carbon  grades,  MS&T'17  Conference,  October  08-12, 
2017, Pittsburgh, USA. 

122 

 
 
 
 
 
 
 
 
 
 
 
 
 
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List of figures 
Figure 1: Technical program  on the INCROHSS  seminar at  the  MSE Congress  2018 on 26th  -  28th 
September 2018 in Darmstadt, Germany. .............................................................................. 9 
Figure  2:  Schematics  of  the  thermomechanical  schedules  performed  at  the  dilatometer  (non-
deformed austenite in Cycle A and deformed austenite in Cycle B). .......................................... 15 
Figure  3:  Transformation  start  and  finish  temperature  as  a  function  of  cooling  rate  for  (a)  Non 
deformed and (b) Deformed austenite. ................................................................................. 16 
Figure 4: Schematics of the thermomechanical schedules for the MFS torsion trials .................... 17 
Figure 5: Mean flow stress as a function of temperature from the multipass torsion test on the NbV 
material. The strain of each torsion pass is 0.3, while the inter-pass times are 20 s, 10 s and 5 s 
respectively. ...................................................................................................................... 17 
Figure 6: Influence of composition on critical temperatures (Tnr, Ar3 and Ar1): Mean flow stress curves 
corresponding  to  the  steels  CMn-0.06%C  and  NbV  steel,  using  an  interpass  time  of  20 s  and  a 
deformation per pass of 0.2. ................................................................................................ 19 
Figure  7:  (a)  Influence  of  chemical  composition  and  interpass  time  in  the  non-recrystallization 
temperature. (b) Effect of deformation per pass in the Tnr. ...................................................... 19 
Figure 8: Schedule “relaxation-like” torsion trials ................................................................... 20 
Figure 9: Extrapolation of the stress temperature dependence for of the pure ferrite and austenite 
and determination of the ferrite transformation kinetics .......................................................... 21 
Figure 10: Determination of the temperatures corresponding to 25%, 50% and 75% ferrite formed
 ........................................................................................................................................ 21 
Figure 11: Experimental plan of the first rolling trials of WP3. .................................................. 23 
Figure 12: Summary of the ferrite fraction formed as function of temperature as estimated by torsion 
and dilatometry, and comparison with the fraction determined by metallography. ...................... 24 
Figure 13: Microstructures at quarter thickness as obtained in the set C after etching with Nital 2%, 
i.e. quenched just before (a-c) and after (d-f) the last intercritical pass (OCAS trials) ................. 26 
Figure 14: Comparison of the martensite fraction after direct quench from intercritical temperature 
range  before  (full  symbols)  and  after  intercritical  rolling  (cross  symbol),  determined  from  optical 
micrographs ...................................................................................................................... 26 
Figure 15: Microstructures as obtained in the set C after etching with Nital 2%, i.e. quenched just 
before the last intercritical pass (CRM trials) .......................................................................... 27 
Figure 16: Microstructures obtained after intercritical rolling at various temperatures and accelerated 
cooling (simulated by quench in hot water at 60 °C) (set A and set B-ACC) ............................... 29 
Figure 17: Microstructures obtained after intercritical rolling at various temperatures followed by air 
cooling .............................................................................................................................. 29 
Figure 18: Strength as a function of start cooling temperature for plates rolled intercritical followed 
by air cooling. .................................................................................................................... 30 
Figure 19: Charpy impact toughness of the plates air cooled after intercritical rolling, as a function 
of finish rolling / start cooling temperature ............................................................................ 30 
Figure  20:  Tensile  properties  as  function  of  the  last  pass  temperature  (FRT=795  °C/  austenitic 
rolling), (FRT= 755, 745 and 735 °C – intercritical rolling). Accelerated cooling conditions. ......... 31 
Figure 21: Charpy impact toughness as function of the last pass temperature (FRT=795 °C/ austenitic 
rolling), (FRT= 755, 745 and 735 °C – intercritical rolling) ...................................................... 31 
Figure 22: Comparison of tensile results before (full) and after (empty) intercritical rolling reduction 
followed by direct quench in both, transverse and longitudinal direction .................................... 32 
Figure 23: Comparison of the Charpy impact toughness with direct quench before or after intercritical 
rolling as a function of start cooling temperature .................................................................... 32 
Figure 24: Direct comparison of Charpy impact toughness for a selection of start cooling temperature 
- 770°C and 750°C ............................................................................................................. 33 
Figure 25: Comparison of the tensile properties obtained after intercritical rolling for different cooling 
strategies .......................................................................................................................... 33 
Figure  26:  Comparison  of  the  impact  toughness  properties  obtained  after  intercritical  rolling  for 
different cooling strategies .................................................................................................. 34 
Figure 27: Strength vs.  toughness  maps: Upper  Shelf Energy (USE) and Ductile Brittle Transition 
Temperature (DBTT) as function of the Yield Strength (YS) level .............................................. 34 

124 

 
 
 
Figure 28: The change of average ferrite grain size and area fraction of deformed ferrite grains which 
have GOS > 2.0° with finish rolling temperatures in ACC samples. ........................................... 35 
Figure  29:  The  link  between  microstructure  and  mechanical  properties  of  intercritically  rolled  and 
ACC samples. .................................................................................................................... 36 
Figure 30: Thermomechanical cycle applied in torsion. ............................................................ 37 
Figure  31:  Microstructures  obtained  after  a  quenching  at  different  deformation  temperatures  and 
both compositions (torsion tests): (a,d) low ferrite content, (b,e) medium ferrite content and (c,f) 
high ferrite content. ............................................................................................................ 38 
Figure  32:  Mean  Flow  Stress  curves  (single  hit  tests  until  rupture)  obtained  at  the  different 
deformation temperatures for the (a) CMn(0.06%C) and (b) CMn(0.12%C) steels. .................... 39 
Figure 33: Optical micrographs obtained after quenching at different deformation temperatures and 
compositions: (a) quenching at 740ºC in J577 steel (100% austenite), (b) quenching at 780ºC in 
J572 steel (austenite + ferrite) and (c) quenching at 760ºC in J571 steel (100% ferrite). ............ 39 
Figure 34: Mean Flow Stress curves (single hit tests until rupture) obtained for the CMn steels and 
pure  phase  materials  (100%  austenite  and  100%  ferrite):  (a,b,c)  CMn(0.06%C)  and  (d,e,f) 
CMn(0.12%C). Different deformation temperatures are applied, obtaining different ferrite content: 
(a,d) low ferrite, (b,e) medium ferrite and (c,f) high ferrite. .................................................... 40 
Figure  35:  Microstructures  observed  after  rupture  in  the  austenitic  (J577),  ferritic  (J571)  and 
composite material (CMn(0.12%C)). Deformation temperature of 740ºC is applied. ................... 41 
Figure  36:  Optical  micrographs  obtained  after  rupture  and  quenched  at  different  deformation 
temperatures  and  CMn(0.12%C)  steel  (torsion  tests):  (a)  low  ferrite  content,  (b)  medium  ferrite 
content and (c) high ferrite content. ..................................................................................... 41 
Figure 37: Experimental (N refers to normal and C refers to compression mode of dilatometer) and 
theoretical fractions of ferrite formed during isothermal holding. .............................................. 42 
Figure 38: Dilatometer plots of specimens 770-10-q and 770-3000-q. ...................................... 43 
Figure 39: OpticalOptical micrograph of specimens (a) 770-10-q and (b) 770-3000-q. ................ 44 
Figure 40: Compression stress-strain curve of specimens 770-10-q and 770-3000-q .................. 44 
Figure 41: Optical micrograph of specimens (a) 770-10-0.4d-q, (b) 770-3000-0.4d-q, (c) 770-10-
0.4d-10-q, (d) 770-3000-0.4d-10-q, (e) 770-10-0.4d-600-q and (f) 770-3000-0.4d-600-q. ........ 45 
Figure  42:  Isothermal  formation  of  ferrite  before  and  after  deformation  at  different  intercritical 
temperatures. .................................................................................................................... 45 
Figure 43: Thermomechanical cycle applied in dilatometer. ..................................................... 46 
Figure  44:  Optical  micrographs  obtained  after  a  quenching  at  different  deformation  temperatures 
and both compositions (dilatometry tests): (a,d) low ferrite content, (b,e) medium ferrite content 
and (c,f) high ferrite content. ............................................................................................... 47 
Figure 45: Optical micrographs obtained after air cooling at different deformation temperatures and 
both compositions (dilatometry tests): (a,d) low ferrite content, (b,e) medium ferrite content and 
(c,f) high ferrite content. ..................................................................................................... 48 
Figure  46:  FEG-SEM  micrographs  related  to  the  samples  obtained  after  different  deformation 
temperatures for CMn(0.12%C) steel: (a,b) 760ºC (low ferrite content) and (c,d) 740ºC (high ferrite 
content). Different cooling strategies have been used: (a,c) air cooling and (b,d) quenching. ...... 49 
Figure 47: Influence of deformation. Optical micrographs corresponding to 0.12%C-CMn and high 
ferrite content (75%): (a) ε=0.1, (b) ε=0.25 and (c) ε=0.4. ................................................... 50 
Figure  48:  Influence  of  deformation.  FEG-SEM  micrographs  corresponding  to  0.06%C-CMn  and 
medium ferrite content (50%): (a) ε=0.25 and (b) ε=0.4. ...................................................... 50 
Figure  49:  Optical  micrographs  obtained  after  a  quenching  at  different  deformation  temperatures 
and both austenite conditions (Recrystallized and Deformed): (a,d) low ferrite content, (b,e) medium 
ferrite content and (c,f) high ferrite content. .......................................................................... 51 
Figure 50: IQ maps corresponding to the samples quenched before deformation for CMn(0.12%C) 
and  different  ferrite  content:  (a)  low  ferrite,  (b)  medium  ferrite,  (c)  high  ferrite.  Image  Quality 
distributions corresponding to: (d) low ferrite, (e) medium ferrite, (f) high ferrite contents. ........ 51 
Figure  51:  (a)  Image  Quality,  (b)  IPF  and  (c)  grain  boundary    maps  (black  lines  for  high  angle 
misorientation criteria, bigger than 15º,and red lines for low angle misorientation criteria, between 
4º and 15º) corresponding to CMn(0.06%C) steel and high ferrite content (deformation of 0.4 and 
deformation temperature of 740ºC). ..................................................................................... 52 
Figure  52:  Image  Quality  maps  corresponding  to  CMn(0.06%C)  Steel  and  low  ferrite  content 
(deformation of 0.25 and cooling rate of 1ºC/s): (a) with pearlite and (b) without pearlite. ......... 53 

125 

 
 
Figure 53: Kernel maps (pearlite is deleted) corresponding to CMn(0.06%C) steel and deformation 
of 0.25 and different ferrite contents: (a) low ferrite (deformation temperature of 760ºC) and (b) 
high ferrite (deformation temperature of 740ºC). ................................................................... 53 
Figure 54: GOS  parameter  distributions for  different ferrite contents (low, medium  and high) and 
both chemical compositions (deformation of 0.4): (a) CMn(0.06%C) and (b) CMn(0.12%C). ....... 54 
Figure 55: EBSD maps (grain boundary and Kernel maps) corresponding to CMn(0.06%C) steel and 
high ferrite content (deformation of 0.4): (a,b) Non-deformed ferrite population (GOS<2º) and (c,d) 
deformed ferrite population (GOS>2º). ................................................................................. 55 
Figure  56:  Influence  of  chemical  composition  and  deformation  on  mean  unit  size  for  both  ferrite 
populations (deformed and non-deformed ferrite) and different deformations  (0.1, 0.25 and 0.4). 
Different misorientation criterion are considered (4º and 15º). ................................................ 56 
Figure 57: Influence of deformation per pass on mean unit size for both ferrite populations: (a, c) 
deformed and (b,d) non-deformed ferrite. Different misorientation criterion are considered: (a,b) 4º 
and (c,d)15º. ..................................................................................................................... 57 
Figure 58: Influence of chemical composition and deformation on Kernel Average Misorientation for 
both ferrite populations (deformed and non-deformed ferrite) and different deformations (0.1, 0.25 
and 0.4). ........................................................................................................................... 58 
Figure  59:  Thermomechanical  cycles  applied  in  torsion  for  (a)  austenitic,  (b)  ferritic  and  (c) 
intercritical region. ............................................................................................................. 59 
Figure  60:  Mean  Flow  Stress  curves  (single  hit  tests  until  rupture)  obtained  at  the  different 
deformation temperatures for the (a) CMn(0.06%C) and (b) CMn(0.12%C) steels. .................... 59 
Figure 61: The evolution of saturation stress values as a function of deformation temperature strain 
rate for both steels: (a) CMn(0.06%C) and (b) CMn(0.12%C ................................................... 60 
Figure 62: Quenched microstructures for CMn(0.12%C) obtained before and after single hit rupture 
tests and different deformation temperatures (760, 750 and 740ºC): (a,d) low, (b,e) medium and 
(c,f) high ferrite content. ..................................................................................................... 61 
Figure 63: Quenched microstructures for (a,b) CMn(0.06%C) and (c,d) CMn(0.12%C) obtained after 
a strain of 0.65 and 0.75. Different deformation temperatures have been included (760 and 750ºC): 
(a,c) low and (b,d) medium ferrite content. ........................................................................... 61 
Figure  64:  Comparison  between  experimental  curves  and  predicted  curves  for  (a,b)  ferritic,  (c,d) 
austenitic and (e,f) intercritical regions: (a,c) CMn (0.06%C) and (b,d) CMn (0.12%C). .............. 63 
Figure 65 Comparison between flow intercritical curves obtained by the law of mixtures (assumption 
of iso-strain) and the stress-strain curves considering constitutive equations for both steels  (a,b,c 
CMn(0.06%C) and d,e,f CMn(0.12%C)) and different deformation temperatures. ....................... 65 
Figure  66:  Thermomechanical  schedules  applied  in  the  deformation  dilatometer:  (a)  Deformation 
pass of 0.4 and (b) two deformation passes of 0.2 (interpass time of 0, 1 and 5s). .................... 69 
Figure 67: Optical micrographs obtained in intercritically deformed and air cooled microstructures: 
(a) Deformation pass of 0.4, (b,c,d) two deformation passes of 0.2 and interpass times of 0, 1 and 
5s, respectively. ................................................................................................................. 70 
Figure  68:  Optical  micrographs  obtained  in  intercritically  deformed  and  quenched  microstructures 
obtained after an interpass time of (a) 1 and (b) 5s. ............................................................... 70 
Figure  69:    Kernel    maps  obtained  in  intercritically  deformed  and  air  cooled  microstructures:  (a) 
Deformation pass of 0.4, (b,c,d) two deformation passes of 0.2 and interpass times of 0, 1 and 5s, 
respectively. ...................................................................................................................... 71 
Figure 70: Effect of reduction rate (25% vs 40%) on microstructure of intercritical rolled plates .. 71 
Figure 71: Microstructures of intercritically rolled samples, after deformation by 1 (a, b), 2 (c, d) and 
3 (e, f) rolling passes of 10% thickness reduction per pass at 770oC (a, c, e) and 745oC (b, d, f) and 
air cooling to room temperature. .......................................................................................... 73 
Figure 72: Inverse Pole Figure (IPF) maps measured at the mid-thickness regions of the intercritically 
rolled samples at 745oC after 1, 2 and 3 rolling passes of 10% thickness reduction. ................... 73 
Figure  73:  Crystallographic  textures  measured  at  the  mid-thickness  regions  of  the  intercritically 
rolled samples at 745°C after 1, 2 and 3 rolling passes of 10% thickness reduction. ................... 73 
Figure  74:  Effect  of  reduction  rate  (25%  vs.  40%)  on  the  tensile  properties  of  intercritical  rolled 
plates ............................................................................................................................... 74 
Figure 75: Effect of reduction rate (25% vs 40%) on the toughness properties of intercritical rolled 
plates. .............................................................................................................................. 74 

126 

 
 
Figure 76: Tensile properties of the 3  samples which  were intercritical rolled by (a) 1 pass, (b)  2 
passes and (c) 3 passes of 10% thickness reduction at different temperatures between 710oC and 
800oC. .............................................................................................................................. 75 
Figure 77: Toughness measured in longitudinal and transversal directions of the samples which were 
intercritically rolled at different temperatures between 710oC and 800oC in 1, 2 and 3 rolling passes 
of 10% thickness reduction. ................................................................................................ 76 
Figure 78: Tensile properties of the samples which were intercritically rolled to reduce the thickness 
10% at different temperatures between 710oC and 800oC and cooled in (a) air and (b) warm water 
(ACC) to room temperature. ................................................................................................ 76 
Figure  79:  Toughness  properties  of  the  samples  which  were  intercritically  rolled  to  reduce  the 
thickness 10% at different temperatures between 710oC and 800oC and cooled in (a, b) air and (c, 
d) warm water (ACC) to room temperature. .......................................................................... 77 
Figure  80:  The  strength-toughness  map  versus  the  number  of  intercritical  rolling  passes  on  all 
samples which were intercritically rolled and air cooled to room temperature. ............................ 78 
Figure 81: Schematic of the experimental plan of the rolling trials of WP6. ................................ 80 
Figure 82: Thermal profiles during rolling and cooling: (a-b) for grade NbV and (c-d) for grade CMn 
(0.12%C). Graph (a & c) with direct quench before and after IC-rolling; Graph (b & d) with ACC. 81 
Figure  83:  Fraction  of  Ferrite  as  a  function  of  realized  start  cooling  temperature  for  (left)  CMn 
(0.12%C) grade and (right) NbV grade at ¼ thickness ............................................................ 82 
Figure 84: Dissolution calculations for NbV grade, showing the mass fraction of free N, C, Nb and V 
as a function of temperature. ............................................................................................... 82 
Figure 85: Summary of grain size and phase fractions for CMn (0.12%C) grade (a) and for NbV grade 
(b). .................................................................................................................................. 83 
Figure  86:  Optical  micrographs  obtained  after  air  cooling  at  different  deformation  temperatures: 
(a,d,g)  low  ferrite  content,  (b,e,h)  medium  ferrite  content  and  (c,f)  high  ferrite  content.  The 
microstructures obtained in all the cases are shown: (a,b,c) CMn(0.06%C), (d,e,f) NbV Recrystallized 
austenite and (g,h,i) NbV Deformed austenite. ....................................................................... 85 
Figure  87:  SEM  micrographs  related  to  the  samples  obtained  after  different  deformation 
temperatures for NbV Recrystallized and Deformed austenite: (a,d) Low, (b,e) medium and (c,f) high 
ferrite content. ................................................................................................................... 85 
Figure  88:  IPF+IQ  maps  with  high  and  low  angle  boundaries  (coarse  and  fine  black  lines, 
respectively),  corresponding  to  (a,b,c)  NbV  Recrystallized  austenite  and  (d,e,f)  NbV  Deformed 
austenite. Different ferrite levels prior to deformation have been included: (a,d) Low ferrite, (b,e) 
Medium ferrite and (c,f) High ferrite. .................................................................................... 87 
Figure  89:  IPF+IQ  maps  with  high  and  low  angle  boundaries  (coarse  and  fine  black  lines, 
respectively),  corresponding  to  (a)  NbV  Recrystallized  austenite  and  (b)  NbV  Deformed  austenite 
(high ferrite fraction prior to intercritical deformation). ........................................................... 87 
Figure 90: Grain size distributions corresponding to 50% ferrite prior to deformation and different 
chemical composition (CMn(0.06%C), NbV Recrystallized austenite and NbV Deformed austenite): 
(a) Non-deformed ferrite and (b) Deformed ferrite population. Influence of austenite/ferrite balance, 
chemical composition and austenite condition on 15º mean grain size of both ferrite populations: (c) 
Non-deformed ferrite and (d) Deformed ferrite population. ...................................................... 89 
Figure 91: Tensile strength of CMn (0.12%C) in top: longitudinal (0) and bottom: transverse (90) 
direction for the different processing ..................................................................................... 90 
Figure 92: Charpy impact toughness of the CMn (0.12%C) grade with different processing tested in 
longitudinal direction, (top) with intercritical rolling; (bottom) with direct quench ....................... 91 
Figure 93: Tensile strength of NbV in top: longitudinal (0) and bottom: transverse (90) direction for 
the different processing....................................................................................................... 92 
Figure 94: Charpy impact toughness of the NbV grade with different processing tested in longitudinal 
direction, (top) with intercritical rolling; (bottom) with direct quench ........................................ 93 
Figure 95 TEM images obtained in (a) NbV Recrystallized austenite and (b) NbV Deformed austenite.
 ........................................................................................................................................ 94 
Figure 96 Grain size distributions measured for NbV Rec and NbV Def conditions in terms of (a) area 
fraction and (b) accumulated frequency (75% of ferrite prior to intercritical deformation). .......... 94 
Figure 97 TEM images obtained in NbV Recrystallized austenite and different ferrite contents prior to 
deformation: (a) 25% ferrite and (b) 75% ferrite. .................................................................. 95 
Figure 98 Grain size distributions measured for NbV Rec and ferrite fraction of 25 and 75%: (a) area 
fraction and (b) accumulated frequency. ............................................................................... 95 

127 

 
 
Figure 99: Link Process-Tensile Strength-Fraction of Ferrite ..................................................... 96 
Figure 100: Link Process-Ductile to Brittle Transition Temperature-Fraction of Ferrite ................. 97 
Figure 101: Schematic representations of different crystal plasticity models, which can be applied to 
predict evolution of deformation textures in rolled plates. ........................................................ 99 
Figure 102: Simulated deformation texture for austenite after the rolling to reduce the plate thickness 
of 80%. Reference orientations: ● {112}<111>, ▲ {110}<112> and ■ {110}<001>. From left to 
right 2 sections at 45°, 65° and 90° (see upper right corner of the ODF). .............................. 100 
Figure  103:  Simulated  recrystallization  texture  of  austenite  according  to  the  low  stored  energy 
nucleation  and  the  high  growth  rate  of  the  <111>40o  grain  boundary.  Reference  orientations:  ● 
{001}<100>, ▲ {110}<112> and ■ {110}<001>. From left to right 2 sections at 45°, 65° and 
90° (see upper right corner of the ODF). ............................................................................. 101 
Figure 104: Orientations of 24 ferrite products on the parent austenite coordinate system (a) and 
the K-S OR rotation between austenite and ferrite (b)........................................................... 102 
Figure  105:  Simulated  transformation  textures  of  ferrite  from  deformed  austenite  (a)  and  from 
recrystallized  austenite  (b).  Reference  orientations:  ●  {112}<110>,  ▲  {554}<225>  and  ■ 
{110}<110>. ODFs in  2= 45° section of the Euler space. .................................................. 102 
Figure 106: Simulated deformation textures of intercritical ferrite after the rolling of 25% thickness 
reduction  from  (a)  recrystallized  and  transformed  austenite  and  (b)  deformed  and  transformed 
austenite. Reference orientations: ● {112}<110>, ▲ {111}<110>. ODFs in 2= 45° section of the 
Euler space. ..................................................................................................................... 103 
Figure 107: Rolled plate of CMn (0.06%C) with final thickness of 36mm after air cooling .......... 104 
Figure 108: Overview of target core temperature (blue) and pass reduction (orange) for the plates 
rolled to 25mm (left) and 36mm (right) .............................................................................. 105 
Figure 109: Temperature recording across thickness of each pass for 25mm (left) and 36mm (right)
 ...................................................................................................................................... 105 
Figure 110: Through thickness microstructure  of the 25mm NbV plate ................................... 107 
Figure 111: Through thickness microstructure  of the 36mm CMn (0.06%C) ............................ 108 
Figure 112: Grain size distribution of 36mm CMn (0.06%C) plate ........................................... 108 
Figure 113: Through thickness microstructure of the 25mm NbV plate, ND-RD orientation ........ 109 
Figure 114: Through thickness microstructure of the 36mm CMn (0.06%C), ND-RD orientation . 111 
Figure 115: Average tensile strength in transverse direction .................................................. 113 
Figure 116: Charpy transition curves in transverse direction .................................................. 113 
Figure 117: Fractured BDWTT ............................................................................................ 114 

128 

 
 
 
 
 
129 

 
 
 
 
List of tables 
Table 1: Nominal chemical compositions measured in wt.% of the investigated materials. ........... 12 
Table  2:  Chemical  composition  (OES)  of  the  CMn-0.06%C  grade,  measured  in  wt.%  on  each 
individual casting blocks. ..................................................................................................... 13 
Table  3:  Chemical  composition  (OES)  of  the  CMn-0.12%C  grade,  measured  in  wt.%  on  each 
individual casting blocks. ..................................................................................................... 13 
Table 4 : Chemical composition (OES) of the NbV grade, measured in wt.% on each individual casting 
blocks. .............................................................................................................................. 13 
Table 5: Chemical compositions measured in wt.% of the single phase materials, corresponding to 
25% - 50% - 75% ferrite and austenite for the CMn-0.06%C grade ......................................... 13 
Table 6: Chemical compositions measured in wt.% of the single phase materials, corresponding to 
25% - 50% - 75% ferrite and austenite for the CMn-0.12%C grade ......................................... 14 
Table 7: Transformation start (Ar3) and finish (Ar1) temperatures for different steels (CMn-0.06%C, 
CMn-0.12%C  and  NbV),  selected  cooling  rates  (0.1,  1  ,10  and  100  ºC/s)  and  different  austenite 
conditions prior to transformation (recrystallized and deformed austenite) ................................ 15 
Table 8: Deformation conditions defined in the multipass torsion tests. ..................................... 17 
Table  9:  Critical temperatures  (Tnr,  Ar3  and  Ar1)  obtained  from  multipass torsion  test  for  all the 
studied materials with various inter-pass times and deformation per torsion passes. .................. 18 
Table 10 : Critical temperatures determined by torsion testing (ε=0.3) by three partners (CEIT, CRM 
and  OCAS),  in  a  comparison  with  results  of  dilatometry  measurements,  calculated  equilibrium 
temperatures and empirical formulae. ................................................................................... 20 
Table 11: Detailed rolling schedule of the trials foreseen in WP3. ............................................. 24 
Table 12: Comparison of the rolling process temperatures targeted by OCAS and CRM ............... 25 
Table 13: Measured ferrite fraction in each deformation temperature (torsion tests) ................... 38 
Table 14: Ferrite fraction formed and carbon concentration of the phases at isothermal temperature.
 ........................................................................................................................................ 42 
Table 15: Measured ferrite fraction in each deformation temperature  (dilatometry tests) ........... 46 
Table  16:  Pearlite  fractions  measured  in  each  deformation  and  chemical  composition  after 
deformation and air cooled. ................................................................................................. 46 
Table 17: Measured fractions of deformed ferrite by means of EBSD (microstructures obtained after 
Cycle 4A). GOS values higher than 2º has taken into account. ................................................. 54 
Table 18 Constitutive equations obtained for describing hot deformation behavior. ..................... 62 
Table  19  Material  constants  obtained  in  each  region  (ferritic,  austenitic  and  intercritical)  for  both 
chemistries. ....................................................................................................................... 62 
Table 20. Rolling schedule for single intercritical pass of 10% thickness reduction. ..................... 67 
Table 21. Rolling schedule for single intercritical pass of 40% thickness reduction. ..................... 67 
Table 22. Rolling schedule for 3 intercritical passes of 10% thickness reduction per pass............. 68 
Table 23: Chemical composition of the CMn (0.12%C) grade heats cast .................................... 79 
Table 24: Chemical composition of the NbV grade heats cast ................................................... 79 
Table 25: Key rolling and cooling parameters of the rolling trials in WP6. .................................. 80 
Table 26: Realized rolling and cooling parameters, temperatures measured by thermocouples at plate 
mid thickness .................................................................................................................... 80 
Table 27: ICP-MS results on the precipitation of Nb (left) and V (right) after intercritical rolling and 
quenching. ........................................................................................................................ 82 
Table  28:  Optimum  GOS  values  needed  to  obtain  the  desired  ferrite  fraction  in  each  case 
(CMn(0.06%C),  NbV  Rec  and  NbV  def).  Measured  deformed  ferrite  fractions  were  also  included, 
considering the proposed tolerance angle of 4º. ..................................................................... 84 
Table  29:  Summary  of  process-property-microstructure  for  the  three  investigated  grades  after 
intercritical rolling. DBTT = 50%USE, 2nd phase identification by optical microscopy (OM) with P for 
pearlite, M for martensite and B for bainite. ........................................................................... 95 
Table 30: Overview of the EBSD results on the investigated samples ........................................ 97 
Table 31: Slab and final plate target dimension for rolling to (left) 25mm and (right) 36mm ...... 104 
Table 32: Overview on phase fraction and grain size ............................................................. 108 
Table 33: Average precipitation fraction of 25mm NbV plate .................................................. 112 

130 

 
 
 
 
131 

 
 
 
 
List of acronyms and abbreviations 

ACC 
ALAMEL 
Ar1 
Ar3 
B 
BCC 
BCT 
CMn 
DQ 
EBSD 
FCC 
FCT 
FEG-SEM 
FRT 
GF 
GOS 
HAGB 
HSLA 
IPF 
IQ 
KAM 
K-S 
LAGB 
LE 
LENP 
LOM 
M 
MA 
MFS 
mrd 
ND 
N-W 
ODF 
OES 
P 
PF 
PSC  
RD 
TD 
TM 
TMCP 
Tnr 
TS 
UFC 
VPSC 
WT3 
YS 
CRM 
CEIT 
OCAS 
TUDT 

Accelerated Cooling 
Advanced Lamel 
Finish ferrite transformation temperature 
Start ferrite transformation temperature 
Bainite 
Body-Centered Cubic 
Body-Centered Tetragonal 
Carbon-Manganese grade 
Direct Quench 
Electron BackScatter Electron Diffraction 
Face-Centered Cubic 
Full constraints Taylor 
Field Emissions Gun Scanning Electron Microscopy 
Finish Rolling Temperature 
Granular Ferrite 
Grain Orientation Spread 
High Angle Grain Boundary 
High Strength Low Alloy 
Inverse Pole Figure 
Image Quality 
Kernel Average Misorientation 
 Kurdjumov-Sachs orientation relationship 
Low Angle Grain Boundary 
Local Equilibrium 
Local Equilibrium Negligible Partitioning 
Light Optical Microscopy 
Martensite 
Martensite/Austenite constituent 
Mean Flow Stress 
multiple of random density 
Normal Direction 
Nishiyama-Wasserman orientation relationship 
Orientation distribution function 
Optical Emission Spectroscopy 
Pearlite 
Polygonal Ferrite 
Plain Strain Compression 
Rolling Direction 
Transvers Direction 
Thermo Mechanical 
Thermo-Mechanical Control Processing 
No recrystallization temperature 
Tensile Strength 
Ultra Fast Cooling 
Visco-Plastic Self-Consistent 
WalsTuig3 (Rolling Mill 3) 
Yield Strength 
Centre de Recherches Metallurgiques  
Asociacion Centro de Estudios e Investigaciones Tecnicas  
Onderzoekscentrum voor Aanwending van Staal  
Technische Universiteit Delft  

132 

 
 
 
 
 
133 

 
 
 
 
List of references 

[Chen14]  H. Chen, X. Xu, W. Xu, S. Van Der Zwag, Metall and Mat Trans A, 45 (2014) 1675 

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Coates DE. Diffusion-controlled precipitate growth in ternary systems I, Metall. Trans. 
1972;3:1203 

[Dohe97] 

Doherty RD, Hughes DA, Humphreys FJ, Jonas JJ, Jensen DJ, Kassner ME, King WE, 
McNelley TR, McQueen HJ, Rollett AD. Current issues in recrystallization: a review. Materials 
Science and Engineering: A. 1997;238:219. 

[Hill53] 

M. Hillert: Introduction to Paraequilibrium, Internal Report, Swedish Institute of Metals 
Research, Stockholm, 1953 

[Hult47] 

A. Hultgren: Trans. ASM, 39 (1947) 915 

[Hutc07] 

Hutchinson B. Deformation substructures and recrystallisation. Mater Sci Forum. 2007;558-
559:13. 

[Hutc99] 

Hutchinson B. Deformation microstructures and textures in steels. Philos. Trans. R. Soc. A 
Math. Phys. Eng. Sci. 1999;357:1471 

[Kest96] 

Kestens L, Jonas JJ. Modeling texture change during the static recrystallization of interstitial 
free steels. Metallurgical and Materials Transactions A. 1996;27:155. 

[Kirk58] 

J.S. Kirkakdy: Can. J. Phys., 1958, vol. 36, pp. 907. 

[Lebe93]  

Lebensohn RA, Tomé CN. A self-consistent anisotropic approach for the simulation of plastic 
deformation and texture development of polycrystals: Application to zirconium alloys. Acta 
Metallurgica Et Materialia. 1993;41:2611 

[Tayl38] 

Taylor GI. Plastic Strain in Metals. Journal of the Institute of Metals. 1938;62:307 

[VanH05] 

Van Houtte P, Li S, Seefeldt M, Delannay L. Deformation texture prediction: From the Taylor 
model to the advanced Lamel model. Int J Plast. 2005;21:589. 

[Zuro08] 

H. S. Zurob, C. R. Hutchinson, A. Béché, G. R. Purdy, Y. J. M. Bréchet, Acta Mater, 56 
(2008) 2203 

134 

 
 
 
 
 
 
 
135 

 
 
 
 
Appendices 
A. - Dilatometry tests 

Uniaxial compression tests were performed both at CEIT and OCAS in a Bähr DIL805A/D quenching 
and deformation dilatometer, where solid cylinders (5 mm diameter, 10 mm long) were tested. The 
dilatometry specimens were subjected to the thermomechanical schedules shown in Figure A 1. Two 
different thermomechanical schedules were defined, with the purpose of obtaining different austenite 
conditions  (non-deformed  and  deformed  austenite,  in  Cycle  A  and  Cycle  B,  respectively).  After  a 
soaking at 1250 ºC for 10 min, a deformation step at 1150º C which includes a deformation of 0.3 
with the deformation rate of 1/s and a holding time of 10 s was performed. Afterwards, the specimens 
were  cooled  down  (5 ºC/s)  to  900 ºC.  Cycle  A  is  followed  by  selected  cooling  rates  varying  from 
0.1 ºC/s  to  the  maximum  cooling  rate  of  ~100  °C/s.  In  Cycle  B,  the  specimens  are  deformed  at 
900 ºC using the strain of 0.3 with the strain rate of 1/s. Finally, the same cooling rates (0.1, 1, 10 
and ~100  ºC/s) as applied in Cycle A are used.  

Figure A 1: Schematics of the thermomechanical schedules performed at the dilatometer 
(non-deformed austenite in Cycle A and deformed austenite in Cycle B). 

A  systematic  analysis  of  dilatation  curves  is  necessary  to  determine  critical  transformation 
temperatures.  The  interpretation  of  the  dilatometry  need  to  be  done  in  the  conjunction  with 
microstructural characterization performed by optical and scanning electron microscopes.  Figure A 
2a shows measured dilatation curves of the CMn(0.12%C) steel with different cooling rates (0.1 ºC/s, 
1 ºC/s, 10 ºC/s and ~100 ºC/s). Figure A 2b depicts the evolution of transformed fraction obtained 
from dilatometry curves shown in Figure A 2a. For that purpose, the level rule is used, which is based 
on  extrapolating  the  linear  expansion  behavior  from  the  temperature  regions  where  no 
transformation  occurs,  and  subsequently  assuming  proportionality  between  the  fraction  of 
decomposed austenite and the observed length change. Figure A 2b shows that the transformation 
temperatures vary depending on the cooling rate. As cooling rate increases the transformation start 
and  finish  temperatures  decrease.    Moreover,  when  the  cooling  rate  is  increased  the  shifting  of 
transformed  fraction  to  lower  transformation  temperatures  is  observed  (see  Figure  A  2b).  For 
0.1ºC/s, 1 ºC/s and 10 ºC/s cooling rates, transformed fraction curves present slope changes, which 
indicates  that  different  transformation  products  are  formed  in  the  resultant  microstructures.  The 
slope  changes  observed  in  Figure  A  2b  are  related  to  transitions  from  polygonal  ferrite  to 
combinations between bainitic and pearlitic phases. However, for the highest cooling rate (maximum 
cooling), the evolution of transformed fraction doesn’t exhibit any slope change; this is due to the 
formation of a martensitic microstructure.  

136 

1250 ºC, 10 minTEMPERATURE TIMECoolingrates: 0.1 ºC/s, 1 ºC/s, 10 ºC/s, maxcoolingCycleACycleB=0.3; =1 s-10 s=0.3; =1 s-110 sRefinement5ºC/s 
 
 
 
(a) 

(b) 

Figure  A  2:  (a)  Dilatometry  curves  and  (b)  evolution  of  transformed  fraction  in 
CMn(0.12%C)  steel  and  selected  cooling  temperatures  (transformation  from  non-
deformed austenite): 0.1 ºC/s, 1 ºC/s, 10 ºC/s and maximum cooling. 

The final microstructures obtained after dilatometry tests were analyzed after etching in 2% Nital by 
optical microscopy. Dilatometry samples were sectioned along their longitudinal axis, selecting the 
region corresponding to a maximum area fraction of nominal strain and reduced strain gradient. In 
Figure A 3a, b, c and d optical micrographs corresponding to the CMn(0.12%C) steel, after different 
cooling  strategies  (0.1  ºC/s,  1 ºC/s,  10 ºC/s  and  maximum  cooling,  respectively)  are  presented. 
When the lowest cooling rate of 0.1ºC/s is applied, PF (polygonal ferrite) and pearlite is formed. For 
the cooling rate of 1 ºC/s, in addition to the combination between PF and P (pearlite), GF (granular 
ferrite) is also observed (see Figure A 3b). It is evident that the presence of GF is more notorious 
when a higher cooling rate of 10 ºC/s is applied. Furthermore, the microstructural characterization 
suggests  that  the  increasing  of  cooling  rate  promotes  the  refinement  of  the  final  microstructure. 
Finally, after a maximum cooling, a microstructure constituted by martensite is obtained (see Figure 
A 3d). 

(a) 

(b) 

(c) 

 (d) 

Figure  A  3:  LOM  micrographs  corresponding  to  the  microstructures  transformed 
CMn(0.12%C)  steel:  (a)  0.1 ºC/s,  (b)  1 ºC/s,  (c)  10ºC/s  and  (d)  max  Cooling  (non-
deformed γ). 

137 

-0.009-0.007-0.005-0.003-0.0010.00102004006008001000ΔL/LoTemperature (ºC)0.1ºC/s1ºC/s10ºC/sMax CoolingNon-deformedɣCMn (0.12%C)02040608010002004006008001000Transformed fraction (%)Temperature (ºC)0.1ºC/s1ºC/s10ºC/sMax CoolingNon-deformedɣCMn(0.12%C)PPFPFGFPGFPFBM 
 
 
 
 
 
 
 
 
When  non-polygonal  complex  microstructures  are  formed,  a  more  detailed  microstructural 
characterization is required by means of electron microscopy. FEG-SEM micrographs corresponding 
to the CMn(0.12%C) and different cooling rates (10 ºC/s and max cooling) are presented in Figure 
A 4. For the 10 ºC/s cooling rate (see Figure A 4a) the microstructure is composed by granular ferrite 
accompanied  by  martensite-austenite  islands  (MA).  After  maximum  cooling  a  martensitic 
microstructure is achieved, as shown in Figure A 4b. 

Figure  A  4:  FEG-SEM  micrographs  corresponding  to  the  microstructures  transformed 
CMn(0.12%C) steel: (a) 10 ºC/s and (b) max cooling (non-deformed γ). 

(a) 

(b) 

In Table A 1, transformation start (Ar3) and finish (Ar1) temperatures measured from dilatometry, 
for different steels, austenite conditions (non-deformed and deformed austenite) and selected cooling 
rates are listed. The transformation start and finish temperatures were determined as the 5 and 95% 
transformed fractions, defined from curves such as the ones shown in Figure A 2b. Additionally, the 
equilibrium austenite-ferrite transformation temperatures calculated by means of ThermoCalc have 
been included. The results obtained by CEIT and OCAS are in quite good agreement.  

In Figure A 5a and b transformation start temperatures have been plotted as a function of cooling 
rate  for  both  austenite  conditioning  (non-deformed  and  deformed  austenite,  respectively).  In 
reference to the effect of chemical composition, it is observed that the addition of C (from 0.06% to 
0.12%C) and the presence of microalloying elements (Nb and V) retard phase transformations, and 
consequently decrease the transformation temperatures (mainly transformation start temperature). 
For a similar carbon content (0.06%C) the addition of Nb and V causes the decrease of transformation 
start  and  finish  temperatures,  leading  to  the  formation  of  more  bainitic  phases  and  reducing  the 
presence of polygonal phases. As expected and in terms of the effect of cooling rate, the increase of 
cooling  rate  leads  to  the  formation  of  more  bainitic  microstructures,  transformed  at  lower 
transformation temperatures.  

Regarding  the  effect  of  austenite  conditioning  prior  to  transformation,  the  accumulation  of 
deformation  of  austenite  accelerates  phase  transformation,  increasing  transformation  start 
temperatures (see Figure A 5). The effect of the strain accumulation on the transformation is related 
to the increase in the density of nucleation sites introduced by deformation. 

Table  A  1:  Transformation  start  (Ar3)  and  finish  (Ar1)  temperatures  for  different  steels 
(CMn-0.06%C, CMn(0.12%C) and NbV), selected cooling rates (0.1, 1 ,10 and 100 ºC/s) 
and  different  austenite  conditions  prior  to  transformation  (non-deformed  and  deformed 
austenite) 

138 

 
 
 
 
Figure A 5: Transformation start and finish temperature as a function of cooling rate for 
(a) Non deformed and (b) Deformed austenite. 

Based on the transformation start and finish temperatures estimated from dilatometry, continuous 
cooling transformation diagrams (CCT diagrams) can also been built (see Figure A 6). In addition to 
phase transformation regions, Vickers hardness values measured have been included. As mentioned 
before, similar transformation products are observed in each CCT diagram. For the slow cooling rates 
(0.1 and 1ºC/s) combinations between PF, P and DP are noticed. As the cooling rate is increased, 
the  presence  of  more  bainitic  phases,  such  as  QF  and  GF,  increases  significantly.  Finally,  the 
formation of BF and M is limited for cooling rates higher than 10ºC/s. 

(a) 

(d) 

139 

Ts (ºC)Tf (ºC)Ts (ºC)Tf (ºC)Ts (ºC)Tf (ºC)841682826692849681CEIT799636771606801689OCAS821646768625778623CEIT766614708582709556OCAS761628720583CEIT700559669513651546OCAS733579673495630538OCAS100 ºC/s549417489338539408CEIT824720783614804714OCAS818716796616818718CEIT769624756571774554OCAS779632751581768541CEIT753571725528704540OCAS753565744493721537OCAS100 ºC/s569375417330475325Equilibrium Temperatures for ferrite (Thermocalc)CycleLabCooling RateNon Deformed ɣCycle A0.1 ºC/s1 ºC/s10 ºC/sDeformed ɣCycle B 0.1 ºC/s1 ºC/s10 ºC/sCMn (0.06%C)CMn (0.12%C)NbV020040060080010000.1110100100010000Temperature (ºC)Time (s)ºC/s         100       30       10    5               1     0.5           0.1        HV      ºC/s             Q                  10                     1                    0.1        185                  175                 171 P+DPPFBFºC/s         100       30       10    5               1     0.5           0.1        HV      ºC/s             Q                  10                     1                    0.1        185                  175                 171 P+DPPFGF+DPMsCMn (0.06%C)Recrystallized austenite020040060080010000.1110100100010000Temperature (ºC)Time (s)ºC/s         100       30       10    5               1     0.5           0.1        HV      ºC/s               Q                10                     1                    0.1        HV                                   209                 163                  166P+DPPFQF+GFBFMsCMn (0.06%C)Deformed austenite 
 
 
 
 
 
 
(b) 

(e) 

(c) 

(f) 

Figure  A  6:  Continuous  Cooling  Transformation  diagrams  (CCT  diagrams)  for  different 
compositions  (CMn-0.06%C,  CMn(0.12%C)  and  NbV)  and  austenite  conditioning  (non-
deformed and deformed austenite). 

B.- Torsion tests 

In order to determine the critical temperatures Tnr, Ar3, Ar1, each partner (CEIT, OCAS and  CRM) 
carried out multipass torsion tests following the thermomechanical schedule illustrated in Figure A 7. 
After a soaking at 1250 ºC for 10 min, a deformation step was performed at 1150º C (deformation 
of 0.3, deformation rate of 1 s-1 and a holding time of 10 s), in order to obtain a fine recrystallized 
austenite. Afterwards, 26 deformation passes were applied at decreasing temperature in the 1150 ºC 
and  650 ºC  range.  The  decrease  between  passes  is  20 ºC.  Different  strains  per  pass  (ɛ )  and 
interpass times (tip) has been used, as shown in Figure A 7. A strain rate of 1s -1 is defined in all the 
cases.  

140 

020040060080010000.1110100100010000Temperature (ºC)Time (s)ºC/s         100       30       10    5               1     0.5           0.1        HV      ºC/s              Q                 10                     1                    0.1        HV                                   202                  182                 159P+DPPFQF+GFBFMsCMn (0.12%C)Recrystallized austenite020040060080010000.1110100100010000Temperature (ºC)Time (s)ºC/s         100       30       10    5               1     0.5           0.1        HV      ºC/s             Q                 10                     1                    0.1        HV                                  235                  180                 169P+DPPFQF+GFBFMsCMn (0.12%C)Deformed austenite020040060080010000.1110100100010000Temperature (ºC)Time (s)ºC/s         100       30       10    5               1     0.5           0.1        HV      ºC/s             Q                  10                     1                    0.1        HV                                   219                  214                 187P+DPPFQF+GFBFMsNbVRecrystallized austenite020040060080010000.1110100100010000Temperature (ºC)Time (s)ºC/s         100       30       10    5               1     0.5           0.1        HV      ºC/s             Q                  10                     1                    0.1        HV                                   233                  214                 169P+DP+QF+GFPF+QFBFMsNbVDeformed austenite 
 
 
 
 
 
Figure A 7: Schematics of the thermomechanical schedules for the MFS torsion trials 

Table A 2: Deformation conditions defined in the multipass torsion tests. 

Soaking temperature (ºC) 
Deformation per pass 
Interpass time (s) 
Cooling rate (s-1) 

1250 
0.2, 0.3, 0.4 
5, 10, 20 
1, 2, 4 

TORSION TESTS / CEIT 

The  analysis  of  the  austenitic  structure  prior  to  transformation  is  required  in  order  to  a  suitable 
definition  of  torsion  schedule.  Average  austenite  grain  sizes  were  measured  in  all  the  samples 
quenched  from  1150 ºC.  Samples  were  etched  using  saturated  pricric  acid  and  HCl.  Austenite 
distributions and average austenite grain sizes before transformation were measured using the mean 
equivalent  diameter  method.  Optical  micrographs  of  the  specimens  quenched  from  1150 ºC  and 
different  steels  (CMn-0.06%C,  CMn(0.12%C)  and  NbV)  are  presented  inFigure  A  8.  A  relatively 
coarse austenite structure is obtained in all the cases, measuring a mean austenite size of 97 µm, 
135 µm and 79 µm, in the steel CMn-0.06%C, CMn(0.12%C) and NbV, respectively. In Figure A 9 
grain size distributions are  plotted  for the different  steels in terms of frequency and accumulated 
area fraction, in order to better identify of the presence of coarse grains. Figure A 9 confirms the 
formation of a coarser and more heterogeneous austenite structure in CMn(0.12%C). Furthermore, 
a slight refinement is observed when microalloying elements (Nb and V) are added.  

(a) 

(b) 

141 

1250 ºC, 10 minTEMPERATURE TIME=0.3; =1 s-110 s1; =1 s-1; =1 s-1RefinementMultipassDeformation () = 0.2, 0.3, 0.4Interpasstimes (tip) = 5 s, 10 s, 20 stiptipDɣ= 97 µmDɣ= 135 µm 
 
 
 
 
(c) 
Figure A 8: Optical micrographs corresponding to austenite formed in different steels: (a) 
CMn-0.06%C, (b) CMn(0.12%C) and (c) NbV. 

Figure  A  9:  Austenite  grain  size  distributions  for  the  steels  CMn-0.06%C,  CMn(0.12%C) 
and NbV (specimens quenched from 1150 ºC). 

In  Figure  A  10,  an  example  of  stress-strain  curves  obtained  after  multipass  torsion  testing  is 
presented for the NbV grade. A strain per pass of 0.3 (ε) and different interpass times (20 s, 10 s 
and 5 s) have been applied. In all the cases, at high temperature it is clearly evident that the stress 
increases as the temperature decreases. Furthermore, three regions can be distinguished in Figure 
A 10, called as Region I, Region II and Region III. In Region I, complete recrystallization between 
passes is achieved and increasing of stress from pass to pass is only attributed to the temperature 
drop. Nevertheless, in Region II, recrystallization between passes is inhibited and strain accumulation 
on  austenite  is  obtained.  This  trend  is  more  pronounced  in  stress-strain  curves  related  to  NbV 
microalloyed  steel  shown  in  Figure  A  10.  This  could  be  associated  to  solute  drag  effect  owing  to 
solute  atoms,  combined  with  pinning  effect  due  to  strain  induced  precipitates.  In  Region  III 
transformation takes place and consequently, a softening is observed.  

From the stress-strain curves obtained shown previously, mean flow stress (MFS) corresponding to 
each pass has been calculated for all the chemical composition, strain per pass and interpass time. 
The mean flow stress (MFS), defined as the area under each stress-strain curve divided by the pass 
strain, has been estimated by numerical integration. In Figure A 11 MFS is plotted as a function of 

142 

Dɣ= 79 µm00.20.40.60.810%5%10%15%20%25%30%0-2525-5050-7575-100100-125125-150150-175175-200200-225225-250250-275275-300300-325325-350>350Accumulated Area FractionFrecuencyGrain size (µm)CMn-0.06%CCMn-0.12%CNbVDɣmean (CMn-0.06%C)= 97 µmDɣmean (CMn-0.12%C)= 135 µmDɣmean (NbV)= 79 µm 
 
 
 
 
temperature, for NbV steel (strain per pass of 0.3). Results corresponding to different interpass times 
(20 s, 10 s and 5 s) are included. 

(a) 

(b) 

(c) 
Figure A 10: Stress strain curves obtained in a multipass torsion test for steel NbV (ε=0.3) 
and different interpass time: (a) 20 s, (b) 10 s and (c) 5 s. 

143 

0501001502002503003500246810Stress (MPa)StrainNbVtip=20sɛ=0.31150ºC1150ºC1110ºC1060ºC1020ºC980ºC820ºC900ºC940ºC860ºC700ºC740ºC780ºC660ºCRegionIIRegionIIIRegionI0501001502002503003500246810Stress (MPa)StrainNbVtip=10sɛ=0.31150ºC1150ºC1110ºC1060ºC1020ºC980ºC820ºC900ºC940ºC860ºC700ºC740ºC780ºC660ºCRegionIIRegionIRegionIII0501001502002503003500246810Stress (MPa)StrainNbVtip=5sɛ=0.31150ºC1150ºC1110ºC1060ºC1020ºC980ºC820ºC900ºC940ºC860ºC700ºC740ºC780ºC660ºCRegionIIRegionIRegionIII 
 
 
 
 
 
 
(a) 

(b) 

(c) 
Figure A 11: Mean flow stress as a function of temperature from the multipass torsion test. 

The  value  of  Tnr  was  determined  from  the  intersection  between  the  regression  lines  of  the  points 
corresponding to regions I and II (see Figure A 11). Apart from Tnr, transformation start (Ar3) and 
finish (Ar1) temperatures are also indicated in each MFS curve. The critical temperatures related to 
different steels and interpass times are listed in Table A 3. For the highest deformation of 0.4, in 
some cases the thermocouple was broken during the tests, due to too many deformation passes and 
high  deformation  per  pass.  Even  though  the  determination  of  Tnr  has  been  possible,  Ar3  and  Ar1 

144 

050100150200250300350600700800900100011001200MEAN FLOW STRESS (MPa)T (ºC)Tnr=1027ºCNbVtip=20sɛ=0.3Ar3=810ºCAr1=750ºC050100150200250300350600700800900100011001200MEAN FLOW STRESS (MPa)T (ºC)Tnr=1032ºCNbVtip=10sɛ=0.3Ar3=810ºCAr1=750ºC050100150200250300350600700800900100011001200MEAN FLOW STRESS (MPa)T (ºC)Tnr=1039ºCNbVtip=5sɛ=0.3Ar3=810ºCAr1=750ºC 
 
 
 
 
weren’t obtained in these cases. Regarding the non-recrystallization temperature, it is evident that 
NbV microalloyed steel presents higher Tnr values. The effect of microalloying elements on the non-
recrystallization  temperature  may be  due  to two mechanisms: the solute drag effect regarding to 
solute atoms and the pinning effect due to strain induced precipitates, the latter usually exerting the 
strongest effect.  

Table A 3: Critical temperatures (Tnr, Ar3 and Ar1) obtained from multipass torsion test for 
all the steels and interpass times (deformation per pass of 0.2, 0.3 and 0.4). 

In  terms  of  transformation  temperatures,  the  results  shown  in  Table  A  4  suggest  that  chemical 
composition modifies transformation start and finish temperature significantly. On the one hand, the 
addition of carbon (from 0.06% to 0.12%) promotes the reduction of transformation start (Ar3) and 
finish (Ar1) temperatures. For example, when a deformation of 0.3 is applied, transformation takes 
place  in the temperature range  of  790 ºC-750 ºC for CMn-0.06%C,  whereas in  CMn(0.12%C)  the 
transformation occurs in the temperature interval of 770 ºC-730 ºC (see Table A 3).  

In Figure A 12 MFS curves corresponding to CMn-0.06%C and NbV steels are compared (deformation 
of 0.2 and interpass time of 20 s), in order to a better understanding of the influence of microalloying 
elements  on  the  critical  temperatures.  It  is  observed  that  although  the  presence  of  microalloying 
elements affect notoriously in the non-recrystallization temperature, the transformation kinetics are 
not affected. 

145 

Tnr(ºC)Ar3(ºC)Ar1(ºC)Interpasstimetip (s)5s10s20s5s10s20s5s10s20sɛ=0.2CMn(0.06%C)927922913810790790770750750CMn(0.12%C)915913914770770770730730730NbV103010371036790810790750770750ɛ=0.3CMn(0.06%C)910908903790790790750750750CMn(0.12%C)914902902770770770730730730NbV104210261019810810810750750750ɛ=0.4CMn(0.06%C)899895898-810810-750770CMn(0.12%C)895877859------NbV102110191004810810810750750750050100150200250300350600700800900100011001200MEAN FLOW STRESS (MPa)T (ºC)CMn-0.06%CNbVTnr=913ºCtip=20sɛ=0.2Ar3=790ºCAr1=750ºCTnr=1036ºC 
 
 
 
 
 
Figure A 12: Influence of composition on critical temperatures (Tnr, Ar3 and Ar1): Mean flow 
stress curves corresponding to the steels CMn-0.06%C and NbV steel, using an interpass 
time of 20 s and a deformation per pass of 0.2. 

In  order  to  evaluate  the  effect  of  interpass  time  and  deformation  per  pass,  non-recrystallization 
temperature have been plotted as a function tip and deformation in Figure A 13 a and b, respectively. 
When  a  deformation  of  0.3  is  used  (see  Figure  A  13a),  even  though  its  influence  on  Tnr  is  not 
significant in CMn steels, in NbV microalloyed steel slightly higher non-recrystallization temperatures 
are measured as the interpass time is decreased. Similar trends are observed for the deformation of 
0.2 and 0.4. Concerning the influence of deformation per pass, the results shown in  Figure A 13b 
suggest  that  the  application  of  higher  deformation  leads  to  a  slight  reduction  of  the  non-
recrystallization temperature.  

Figure  A  13:  (a)  Influence  of  chemical  composition  and  interpass  time  in  the  non-
recrystallization temperature. (b) Effect of deformation per pass in the Tnr. 

(a) 

(b) 

TORSION TESTS / CRM 

At  CRM, the tests were carried-out on the torsion machine  located at Liège.  Conventional torsion 
samples having a 6-mm diameter were used. A thermocouple was spot-welded on one of the shoulder 
for the regulation of the temperature. At the same time, the temperature at the centre of the sample 
was  continuously  monitored  by  means  of  a  pyrometer.  The  latter  value  is  considered  as 
representative of the actual sample temperature. Indeed, especially by applying deformation at low 
temperatures and/or short interpass time, a thermal gradient up to 30 °C can appear on the sample. 

An example of the stress-strain curves realized is presented in Figure A 14. Those tests are carried-
out on the reference CMn(0,06%C), CMn(0,12%C) and NbV grades by applying a deformation  ɛ = 
0,3  and  several  interpass  times.  A  very  good  reproducibility  is  obtained:  the  flow  curves  overlap 
quite  good  if  full  recrystallization  occurs,  i.e.  during  deformation  at  high  temperature.  The  main 
features of the flow stress evolution during multipass deformation can be found in Figure A 15: 

1.  Recrystallization  occurs  easily  in  both  CMn(0,06%C)  and  CMn(0,12%C)  grades,  down  to 

temperatures close to Ar3.  Here, the interpass time has any or little influence. 

2.  The on-set of the ferrite transformation is recognizable by a drop in the stress level despite 
the  decrease  in  temperature.  Furthermore,  the  shape  of  the  flow  curve  of  ferrite  is 
significantly different because of the lower work-hardening comparing to austenite. 

3.  The Nb/V addition retards significantly recrystallization. Therefore, the interpass time has a 
significant influence. However, NbC precipitation and the formation of a ferrite layer on the 
γ grain boundaries can reduce the hot ductility and lead to premature fracture. 

146 

80085090095010001050110011500510152025Non-Recrystallization Temperature, Tnr (ºC)Interpass Time, tip (s)DEFORMATION 0.3CMn-0.06%CCMn-0.12%CNbV80085090095010001050110011500.10.20.30.40.5Non-Recrystallization Temperature, Tnr (ºC)Deformation per pass INTERPASS TIME 10sCMn-0.06%CCMn-0.12%CNbV 
 
 
 
 
 
Figure A 14: Stress strain curves in multipass torsion tests (ε=0.3, interpass time=10 s) 
for different grades 

147 

 
 
 
Figure A 15: Stress strain curves in multipass torsion tests (ε=0.3) with different interpass 
times. 

148 

 
 
 
 
The Tnr, Ar3, Ar1 values determined from the MFS vs 1/T plots are summarized in Table A 4 for each 
condition investigated. 

Table A 4: Tnr, Ar3, Ar1 temperature values (°C) as determined by MFS vs 1/T plots 

ε=0,3 
5 s 
10 s 
5 s 
tip 
840 
825 
Tnr  830 
770 
775 
Ar3  790 
750 
745 
Ar1  750 
845 
850 
Tnr  850 
730 
760 
Ar3  770 
Ar1  740 
715 
730 
Tnr  1025  1020  1020  990 
795 
Ar3  795 
745 
Ar1 

20 s 
825 
775 
750 
840 
760 
735 

X 
X 

X 
X 

X 

ε=0,4 
10 s 
825 
760 
735 
X 
X 
X 
995 
770 
735 

20 s 
790 
755 
725 
805 
740 
715 
970 
780 
730 

CMn(0,06%C) 

CMn(0,12%C) 

NbV 

More insight on the transformation kinetics can be obtained from the plots σ vs T presented in Figure 
A 16. This approach, complementary to the more classical analysis from MFS vs 1/T plots, further 
confirms that: 

1.  The  additional  increase  strengthening  in  the  microalloyed  grade  due  to  incomplete 

recrystallization. 

2.  A drop of the stress level is observed as ferrite is formed. This drop is less pronounced in the 
CMn(0.12%C) grade as austenite as almost the same strength level as the ferrite-pearlite 
mixture. 

3.  Similar transformation temperatures for the CMn(0.06%C) and NbV grades and about  20 

°C lower for the CMn(0.12%C) grade. 

149 

 
 
 
  
  
  
 
 
Figure A 16: Stress vs temperature plots obtained in multipass torsion testing (ε=0.3) and 
determination of the critical temperatures 

ROUND ROBIN: COMPARISON MFS TORSION RESULTS CEIT/OCAS/CRM 

The values of the critical temperature determined by the partners are compared in Table A 5.  
The Ar3 and Ar1  values found by CEIT and CRM are in quite good agreement for all the conditions 
investigated  and  similar  to  those  predicted  by  the  empirical  formulae.  Larger  differences  are 
encountered for Tnr, which are probably due to the not unique definition of this temperature. In any 
case,  the  temperatures  determined  by  CRM  are  closer  to  those  reported  in  literature  for  similar 
compositions by Boratto et al.  
As experimental issues occurred during the torsion testing at OCAS, some instabilities were detected, 
which could not be solved within the needed timeframe. As a result the accuracy of the results might 
be impacted. 

Table  A  5:  Critical  temperatures  determined  by  torsion  testing  (ε=0.3)  by  each  partner 
and comparison with dilatometry, equilibrium temperature and literature values. 

150 

 
 
 
 
 
 
“RELAXATION-LIKE” TORSION TRIALS 

The γ-to-α kinetics is strongly influenced by the preparation of austenite (grain size, recrystallization 
state, carbonitrides precipitation) prior to transformation. As it is  impossible to simulate an entire 
rolling schedule by plasto-dilatometry, torsion can be used to simulate a multipass rolling schedule 
and  monitor  the  subsequent  transformation  kinetics  by  means  of  a  “relaxation-like”  test.  As 
illustrated in Figure A 17, after having applied the same thermomechanical schedule investigated in 
WP3, the sample is cooled down at 0,7 °C/s, i.e. the cooling rate measured at mid-thickness in a 12 
mm-thick plate in WP3. 

Figure A 17: Schedule “relaxation-like” torsion trials. 

During cooling, a continuous deformation is applied at an extremely low strain rate (0.001 s-1 ) in 
order to monitor continuously the stress level.  The formation of softer ferrite is detected by a drop 
in the stress level, as it is shown in Figure A 18. After having determined the stress-temperature 
dependence  of  pure  austenite  and  pure  ferrite  (nicely  described  of  an  Arrhenius-type  law)  by 
extrapolation, it is also  possible  to determine the entire transformation kinetics by a simple lever 
rule, as illustrated in Figure A 19. 

151 

 
 
 
 
 
 
 
 
 
Figure  A  18:  Relaxation-like  torsion  trials  after  simulation  of  the  thermomechanical 
schedule of the rolling trials foreseen in WP3. 

Figure  A  19:  Extrapolation  of  the  stress  temperature  dependence  for  of  the  pure  ferrite 
and austenite and determination of the ferrite transformation kinetics. 

Despite the large assumptions made (no influence of continuous straining, linear stress partitioning 
law  between  ferrite  and  austenite),    the  results  obtained  by  the  “relaxation-like”  method  are  in 
excellent agreement with the dilatometry trials carried-out with a similar  prior austenite grain size 
and  with  the  metallography  determined  by  CEIT  in  WP4.  On  the  contrary,  MFS  trials  seem  to 
overestimate the transformation temperatures by about 30 °C. This is probably a consequence of 
the large deformation performed during the MFS trial (ε=8), leading to intense austenite refinement 
and work-hardening. This may also explain the fact the transformation temperatures determined by 
MFS trials are almost the same for both CMn(0.06%C) and NbV grades. 

152 

 
 
 
 
Figure A 20: Validation of the kinetics determined by the “relaxation-like” torsion trials 

The  transformation  kinetics  determined  for  the  same  austenite  conditioning  as  during  the  rolling 
trials in WP3 allow to determine  the temperature for 25%, 50% and 75% ferrite formation which 
will  be  used  in  WP3.  The  NbV  grade  has  the  highest  transformation  temperatures  because  of 
austenite pancaking occurring during the last deformation passes (Figure A 21). Strain accumulation 
does not occur in the reference CMn(0.06%C) grade and consequently the transformation takes place 
about  20  °C  lower.    Increasing  the  carbon  content  (CMn(0.12%C)  grade)  decreases  further  the 
transformation temperatures.  

Figure  A  21:  Determination  of  the  temperatures  corresponding  to  25%,  50%  and  75% 
ferrite formed. 

153 

 
 
 
 
Microstructure characterization on samples of WP6 

Figure  A  22:  Overview  of  the  optical  micrographs  of  CMn  (0.12%C)  grade  in  ND-RD 
direction at ¼ thickness, indicating the process (black) and the fraction of ferrite (red). 
Scale bar of 50µm. 

Figure A 23: Overview of the optical micrographs of NbV grade in ND-RD direction at ¼ 
thickness,  indicating  the  process  (black)  and  the  fraction  of  ferrite  (red).  Scale  bar  of 
50µm. 

154 

 
 
 
 
 
 
Figure A 24: EBSD IQ-map of L591A: CMn (0.12%C) grade quenched after rolling at 820°C, 
with grain size distribution of the complete data set. 

Martensite: Light blue 

Bainite: Yellow; Ferrite: Dark blue 

Figure A 25: Phase determination and grain size distribution of the separate phase. Colour 
code: Light blue: Martensite; Yellow: Bainite and Dark blue: Ferrite. 

DQ_820 

750_DQ 

155 

 
 
 
 
 
 
 
 
 
 
Figure A 26: Phase determination of NbV grade. 

Uniaxial  compression  tests  were  performed  using  a  Bähr  DIL805D  deformation  dilatometer.  Solid 
cylinders of 5 mm in diameter and 10 mm in length were machined for the experiments. The samples 
were subjected to the thermomechanical cycles depicted schematically in Figure A 27. Both schedules 
include  a  solubilisation  treatment  at  1250ºC  for  15  minutes.  Afterwards,  two  different 
thermomechanical  schedules  were  defined,  with  the  purpose  of  obtaining  different  austenite 
conditions  (recrystallized  and  deformed  austenite,  in  Cycle  A  and  Cycle  B,  respectively)  prior  to 
transformation. For  CMn(0.06%C) steel, Cycle  A was applied, whereas for NbV microalloyed steel 
both  schedules  were  applied.  The  samples  were  cooled  down  slowly  (1ºC/s)  to  three  different 
deformation  temperatures,  in  order  to  obtain  different  ferrite  fractions  before  the  intercritical 
deformation.  Different  deformation  temperatures  are  defined  depending  on  the  chemical 
composition.  A  careful  analysis  of  the  dilation  curves  is  required  for  the  definition  of  deformation 
temperatures,  which  will  be  shown  in  detail  later.  In  the  current  study,  in  order  to  measure  the 
evolution of transformed fraction, the lever rule is applied [32] to the dilatometry curves. This rule 
is  based  on  extrapolating  the  linear  expansion  behavior  from  the  temperature  regions  where  no 
transformation  occurs,  and  subsequently  assuming  proportionality  between  the  fraction  of 
decomposed austenite and the observed length change. Finally, a deformation of 0.4 was applied in 
the intercritical region. After that, the specimens were cooled down to room temperature (1ºC/s) in 
both steels.  

156 

 
 
 
 
 
 
 
 
(a) 

(b) 

Figure A 27: Thermomechanical schedule performed at the dilatometer: (a) CMn(0.06%C) 
steel and (b) NbV microalloyed steel. 

For the aim of defining deformation temperatures that ensure the formation of different ferrite levels 
prior  to  intercritical  deformation,  an  interpretation  of  dilation  curves  obtained  by  dilatometry  is 
required. For that purpose, the samples are cooled down slowly (1ºC/s) from 1050ºC and 900ºC to 
room  temperature,  depending  on  the  applied  thermomechanical  cycle  (Cycle  A  and  Cycle  B, 
respectively). The dilation during the cooling step was registered. In Figure A 28, a dilation curves 
obtained during air cooling are plotted for both steels and austenite conditions , represented as ΔL/L0 
as a function of temperature. Figure A 28b illustrates the evolution of the transformed fraction for 
both  CMn(0.06%C),  NbV  recrystallized  austenite  and  NbV  deformed  austenite  (these  curves  are 
obtained from the dilatometry curves shown in Figure A 28a). As mentioned previously, the lever 
rule was considered and proportionality between transformed fraction and measured sample length 
change is assumed in Figure A 28a. In Figure A 28 b the desired fractions of ferrite (25, 50 and 75%) 
are also drawn. As shown in Figure A 28, depending on the chemical composition, the transformation 
temperature varies. Based on the evolution of the transformed fraction presented in Figure A 28b, 
three temperature levels can be defined in each case, in order to meet the ferrite fraction objectives 
of 25, 50 and 75%. Using the predicted deformation temperatures coming from dilatometry study, 
several  interrupted  quenching  were  performed  and  the  deformation  temperatures  that  meet  the 
desired  ferrite  contents  were  defined.  With  the  aim  of  generating  different  ferrite  contents  of 
approximately 25, 50 and 75%, deformation temperatures of 750, 740 and 730ºC were selected for 
the CMn(0.06%C) steel. Meanwhile, for the NbV steel, the formation of different fractions of ferrite 
were achieved  at  deformation  temperatures of 750, 730 and 720ºC in Cycle  A and  770, 750 and 
740ºC in Cycle B (low, medium and high ferrite content, respectively). 

Figure A 28: a) Dilation curves and (b) evolution of transformed fraction in both steels and 
both austenite conditions. 

(a) 

(b) 

157 

1ºC/s1050ºC3minTemperature(ºC)Time (s)1250ºC, 15minMax Coolingɛ=0.4, 1s-11ºC/sRecrystallisedausteniteCycleAɛ=0.4, 1s-1Tdef25%Tdef50%Tdef75%NbVsteelCMn(0.06%C) steelDeformedausteniteCycleB1050ºC3minTemperature(ºC)Time (s)1250ºC, 15minMax Coolingɛ=0.4, 1s-1900ºC 1ºC/sɛ=0.4, 1s-1ɛ=0.4, 1s-1Tdef25%Tdef50%Tdef75%NbVsteel-0.008-0.006-0.004-0.0020.0000.0020.0040.0060.00802004006008001000ΔL/LoTemperature (ºC)6CMnNbV Recrytallized AusteniteNbV Deformed AusteniteCMn(0.06%C)0255075100300400500600700800900Transformed fraction (%)Temperature (ºC)6CMnNbV Recrystallized AusteniteNbV Deformed Austenite75%obj50%obj25%objCMn(0.06%C) 
 
 
 
 
 
 
 
With the purpose of analysing the microstructures prior to deformation, samples were quenched from 
the deformation temperature. In Figure A 29, the microstructures corresponding to NbV quenched 
specimens before intercritical deformation are shown for both austenite conditions and entire range 
of deformation temperature. Figure A 30 shows the quenched microstructures obtained after different 
intercritical deformation temperatures for CMn(0.06%C) steel. The microstructures shown in Figure 
A  29  and  Figure  A  30  illustrate  the  variation  of  the  ferrite-martensite  balance  caused  by  the 
modification of deformation temperature. 

NbV (Recrystallized austenite) 

NbV (Deformed austenite) 

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(b)  Tdef 750ºC 

(d) Tdef 770ºC 

(b) Tdef 730ºC 

(e) Tdef 750ºC 

(c) Tdef 720ºC 

(f) Tdef 740ºC 

Figure  A  29:  Optical  micrographs  obtained  after  a  quenching  at  different  deformation 
temperatures  and  both  austenite  conditions  (Recrystallized  and  Deformed):  (a,d)  low 
ferrite content, (b,e) medium ferrite content and (c,f) high ferrite content. 

158 

Real ferritefraction26% Real ferritefraction31% Real ferritefraction49% Real ferritefraction61% Real ferritefraction76% Real ferritefraction79%  
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
(a) Tdef 750ºC 

(b) Tdef 740ºC 

Figure  A  30:  Optical  micrographs  obtained  after  a  quenching  at  different  deformation 
temperatures and CMn(0.06%C) steel: (a) low ferrite content, (b) medium ferrite content 
and (c) high ferrite content. 

(c) Tdef 730ºC 

In Figure A 31, FEG-SEM micrographs corresponding to the NbV steel and the lowest deformation 
temperatures are presented at higher magnifications, equivalent to a ~75% ferrite formation prior 
to  intercritical  deformation.  Figure  A  31a  corresponds  to  the  microstructure  obtained  after 
transformation from recrystallized austenite (Cycle A) and intercritically deformed at 720ºC, whereas 
in Figure A 31b the micrograph is relative to the sample transformed from deformed austenite (Cycle 
B) and a intercritical deformation temperature of 740ºC.  Differences  between the microstructures 
formed from recrystallized austenite and deformed austenite are clear. The addition of microalloying 
elements  promotes  the  formation  of  more  bainitic  phases  and  reduces  the  presence  of  polygonal 
phases. Quasipolygonal ferrite is observed in NbV microalloyed steel and both austenite conditions. 
Furthermore,  and  when  the  transformation  occurs  from  recrystallized  austenite  (Figure  A  31a), 
deformation  bands  are  identified  inside  the  deformed  ferrite  grains,  reflecting  a  lack  of  ferrite 
restoration. The addition of Nb delays or suppresses the restoration of ferrite when transformation 
occurs from Cycle A (recrystallized austenite). Both the drag effect and/or strain induced precipitation 
at low temperatures (range between 740ºC and 720ºC), delay and suppress the restoration of ferrite 
during intercritical deformation. However, presence of substructure can be appreciated in the sample 
transformed  from  deformed  austenite  (Figure  A  31b)  and  some  subgrain  formation  can  be  seen, 
associated  with  the  activation  of  restoration  during  the  deformation  pass.  The  deformation  of 
austenite  below  the  non-recrystallization  temperature  (Tnr)  promotes  strain  induced  precipitates, 
which  are  effective  for  austenite  pancaking  but  they  reduce  the  Nb  available  during  and  after 
transformation to interact with ferrite restoration or recrystallization phenomena.  

159 

Real ferritefraction24% Real ferritefraction53% Real ferritefraction73%  
 
 
 
 
 
Figure A 31: FEG-SEM micrographs from the NbV and the lowest deformation temperature: 
(a)  Cycle  A  (transformation  from  recrystallized  austenite)  and  intercritical  deformation 
temperature  of  720ºC  and  (b)  Cycle  B  (transformation  from  deformed  austenite)  and 
deformation temperature of 740ºC. 

Concerning  the  effect  of  addition  of  microalloying  elements  and  austenite-ferrite  balances  before 
intercritical deformation, in Figure A 32 the grain boundary maps related to CMn(0.06%C) and NbV 
(recrystallized austenite) are shown for the lowest and the highest ferrite content. It is evident that 
the content of deformed ferrite increases as the deformation temperature is decreased, due to the 
higher  presence  of  ferrite  in  the  microstructure  prior  to  deformation.  Figure  A  33  shows  that  the 
deformed  ferrite  is  characterized  by  a  significant  presence  of  substructure,  reflected  in  a  higher 
fraction  of  low  angle  boundaries  drawn  in  red.  This  could  be  clearly  observed  for  both  chemical 
compositions. In terms of the influence of adding microalloying elements, more non-polygonal phases 
can be distinguished in the final microstructures corresponding to NbV microalloyed steel. This could 
be  mainly  associated  with  the  potential  of  Nb  retarding  phase  transformation  and  promoting  the 
formation of more bainitic phases.  

NbV (Recrystallized ɣ) 

CMn(0.06%C) (Recrystallized ɣ) 

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(b) 

(d) 

Figure A 32: Grain Boundary maps (low and high angle boundaries, in red and black colour, 
respectively) corresponding to (a,b) NbV Recrystallized austenite and (c,d) CMn(0.06%C). 
Different ferrite levels prior to deformation have been included: (a,c) Low ferrite and (b,d) 
High ferrite. 

As an example, several grain boundary maps (coarse and fine black lines, respectively) corresponding 
to medium ferrite content (~50%) prior to intercritical deformation are shown in Figure A 33. Both 
non-deformed  (polygonal  and  quasipolygonal  ferrite  in  Figure  A  33a,c,d)  and  deformed  ferrite 
families (in Figure A 33b,d,f) were split for each chemical composition and austenite condition. 

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Non-deformed ferrite 

Deformed ferrite 

(a) 

(b) 

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(d) 

161 

 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
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Figure A 33: Grain Boundary maps (low and high angle boundaries, in red and black colour, 
respectively)  corresponding  to  50%  ferrite  prior  to  deformation  and  different  chemical 
composition:  (a,b)  CMn(0.06%C),  (c,d)  NbV  Recrystallized  austenite  and  (e,f)  NbV 
Deformed austenite. 

162 

 
 
 
 
 
 
 
 
 
 
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Heavy  gauge  plates  of  structural  and  line-pipe 
steels  are  intentionally  and/or  partially  rolled 
in  the  two-phase  region  for  two  reasons:  1)  to 
improve of YS and TS 2) to avoid the deformation 
gradient  across  the  plate  thickness.  However, 
industrially  intercritical  rolling  has  been  known  for 
the detrimental effects on toughness of steels. The 
aim of this project, therefore, is to understand the 
effect of temperature gradient, α-γ balance, strain 
partitioning  between  phases  during  deformation 
and  phase  transformation  during  (accelerated) 
cooling on the strength-toughness balance. Results 
of this research will provide a wider and more stable 
processing window for steels near/at the two phase 
intercritical region.

For studies in this project, three C-Mn grades with 
different C and micro-alloying (Nb, V) contents have 
been  chosen.  Critical  temperatures  (Tnr,  Ar1,  Ar3, 
Ar1d, Ar3d) of these materials have been determined 
by  dilatometry  and  torsion  experiments.  Impacts 
of  various  γ/α  phase  fractions  on  microstructures 
and mechanical properties of the CMn-0.06 wt.%C 
steel are first studied with a single-pass intercritical 
rolling  at  different 
temperatures.  Afterward, 
thermomechanical  experiments  and  material 
characterizations  are  implemented  to  study  the 
effects  of  strain  partitioning  between  phases  and 
different  strain  paths  in  the  intercritical  region 
on  microstructures  and  mechanical  properties  of 
the  same  material.  Finally,  influences  of  chemical 
compositions  in  combinations  with  intercritical 
rolling on microstructures and mechanical properties 
are investigated with the CMn-0.12 wt.%C and the 
NbV  steels.  To  predict  the  behaviors  of  materials 
during intercritical rolling, a micromechanical model 
is set-up.