Metallurgical design of high-strength austenitic Fe-C-Mn steels with excellent

formability (Metaldesign)

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Directorate-General for Research and InnovationInformation and Communication UnitEuropean Commission1049 Bruxelles/BrusselBELGIQUE/BELGIËFax +32 229-58220E-mail: research@ec.europa.euInternet: http://ec.europa.eu/research/rtdinfo.html

EUROPEAN COMMISSIONDirectorate-General for Research and InnovationResearch Fund for Coal and Steel Unit

Contact: RFCS publicationsAddress: European Commission, CDMA 0/178, 1049 Bruxelles/Brussel, BELGIQUE/BELGIË

Fax +32 229-65987; e-mail: rtd-steel-coal@ec.europa.eu

European Commission

Research Fund for Coal and SteelMetallurgical design of high-strength

austenitic Fe-C-Mn steels with excellent formability (Metaldesign)

A. Ferraiuolo, A. SmithCentro sviluppo Materiali

Via di Castel Romano 100, 00128 Rome RM, ITALY

J. G. Sevillano, F. de las CuevasCeit

Paseo de Manuel Lardizabal, 15, 20018 San Sebastián, SPAIN

P. KarjalainenUniversity of Oulu

Pentti Kaiteran Katu 1, FI-90014 Oulu, FINLAND

G. Pratolongoduferco

Rue des Rivaux 2, 7100 La Louvière, BELGIUM

H. Gouveia, M. Mendes RodriguesisQ

Taguspark-Oeiras, Av. Prof. Dr. Cavaco Silva 33, Porto Salvo¬, Portugal

Contract No RFSR-CT-2005-00030 1 July 2005 to 31 December 2009

Final report

Directorate-General for Research and innovation

2012 EUR 25063 EN

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Luxembourg: Publications Office of the European Union, 2012

ISBN 978-92-79-22205-4

doi:10.2777/14874

ISSN 1831-9424

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Table of Contents Final summary.............................................................................................................................. 5 WP1 Testing material supply and basic metallurgical characterisation of cast materials ......... 13 Task 1.1 Definition of the TWIP steel compositions matrix to be investigated based on the reference TWIP steel composition Fe-Mn-Al-Si and newer TWIP steel compositions (high C and N)......................................................................................................................................... 13 Task 1.1.1 Selection of TWIP steels to be investigated ............................................................. 13 Task 1.1.2 Definition of the TWIP steels compositions matrix ................................................. 14 Task 1.2 Laboratory VIM ingot casting..................................................................................... 16 Task 1.3 - Solidification structure characterisation.................................................................... 16 WP2: Fundamental investigations on the physical metallurgy of TWIP steels ......................... 23 Task 2.1 Measure of SFE by means of Transmission Electron Microscopy observations of extended nodes. Evaluation of Ms γ−>ε temperature. .................................................................. 23 Task 2.1.1: Measure of SFE by means of TEM observation of extended nodes ....................... 23 Task 2.1.2 Evaluation of the Md γ ε of TWIP steel variants................................................... 26 Task 2.2: Characterization of recrystallization behaviour ......................................................... 27 Task 2.2.1 Critical strain for DRX initiation, flow stress, peak stress, ...................................... 27 Task 2.2.2 Recrystallization kinetics under different cold rolling schedule .............................. 41 Task 2.2.3 Modelling of recrystallization behaviour of TWIP steels by modifying the available mathematical models for austenitic steels and the relevant constitutive equations ................... 45 Task 2.3 Study of the precipitation at equilibrium by means isothermal treatments. Precipitates analysis by means of electron microscopy techniques (SEM-EDX, extraction replica for TEM)..................................................................................................................................................... 48 Task 2.4 Industrial feasibility of hot and cold processing route of TWIP steels ....................... 58 WP3 Study of the deformation mechanisms and strain hardening behaviour ........................... 69 Task 3.1 Characterization of mechanical properties in relation with the microstructure, dominating deformation mechanism, strain rate and temperature. ............................................ 69 Task 3.1.1 Static uniaxial tensile tests to investigate the deformation mechanisms transition (deformation twinning --> dislocation glide) and strain hardening behaviour .......................... 69 Task 3.1.1.1 Quasi static tensile tests and strain hardening behaviour analysis ........................ 69 Task 3.1.1.2 Characterization of austenite phase stability and microstructural evolution in..... 75 Task 3.1.1.3 Investigation on TWIP steel embrittlement........................................................... 83 Task 3.1.2 Dynamic Tensile Properties .................................................................................... 91 Task 3.1.3. Torsion tests in the 20ºC-450ºC temperature range: stress-strain behaviour at large strains ......................................................................................................................................... 94 Task 3.1.4 Hot ductility curves in the temperature range 700 ÷ 1300 °C by means of a Gleeble simulator................................................................................................................................... 101 Task 3.1.5. Study of the influence of grain size in the tensile stress-strain behaviour of TWIP steels (Hall-Petch behaviour) ................................................................................................... 105 Task 3.2 Bending fatigue tests to determine the fatigue strength and cyclic softening/hardening behaviour and to analyse the crack initiation/propagation stages ............................................ 109 Task 3.3 –Evaluation of impact strength by means of Charpy tests at different temperature . 111 Task 3.4 Plain strain compression tests.................................................................................... 113 WP 4 – Basic characterisation of application properties: formability, weldability and coating ability........................................................................................................................................ 115 Task 4.1 Formability characterisation by means of Erichsen test and High-velocity forming tests........................................................................................................................................... 115 

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Task 4.1.1 Erichsen test .......................................................................................................... 115 Task 4.1.2 – High-velocity forming tests................................................................................. 116 Task 4.2 – Laboratory coating tests and coating layer characterisation................................... 117 Task 4.3 – Characterization of Weldability ............................................................................. 121 WP5: Industrial trial ................................................................................................................. 135 Task 5.1 Selection of the most interesting TWIP steel variant, for automotive applications, on the basis of the previous WPs results. ...................................................................................... 135 Task 5.2 Coil supplying and material processing (hot and cold rolling, annealing treatments, pickling and coating)................................................................................................................ 137 Objectives of the project .......................................................................................................... 143 Exploitation and impact of the research results ....................................................................... 145 List of figures and tables .......................................................................................................... 145 List of References..................................................................................................................... 151 

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Final summary Twinning induced plasticity (TWIP) steels are austenitic steels characterised by good combination between high tensile properties and high ductility. This behaviour is due to the characteristic high work hardening rate determined by the occurrence of deformation induced twinning during deformation. These properties are extremely attractive for automobile applications expecially for parts devoted to energy absorption (crashworthiness) and for structural reinforcement (body in white). TWIP steels are characterised by

• Austenite phase stability obtained with carbon and high manganese content (typically >16%).

• Low stacking fault energies (SFE) leading to the ability to be deformed by means of both dislocation slip and mechanical twinning. The relative intensity of deformation mechanisms depends upon the chemical composition (mainly the carbon and manganese contents) .

The main objectives of this project are:

1. To complete the understanding of the TWIP steel metallurgy particularly for what regard the following fields: • Solidification microstructure; • Recrystallization behaviour (dynamic and static) and texture formation. • Tensile properties and work hardening ability, precipitation behaviour as a function of

the steel chemical composition and its influence on the strip properties. • Influence of steel chemical composition on microstructure, mechanical properties,

strain hardening behaviour as well as on application properties such as formability, weldability and coatability.

2. To design a metallurgically based manufacturing route (hot/cold rolling process, annealing

treatment) taking into account the specific capability of the DUFERCO plants.

WP1: Testing material supply and basic metallurgical characterisation of cast materials This research started on defining the basic rules for the metallurgical design of TWIP steels. Five TWIP steel variants were selected on the basis of austenitic phase and low SFE. The variants were of two main type Fe-Mn-Si-C-N and Fe-Mn-Al-Si-C.

Table FS1: Chemical composition of the TWIP steels casted at vacuum induction melting (%wt).

Of these five TWIP variants the first variant (TWIP1) revealed a mixed primary solidification structure with austenite+ferrite. The ferrite phase is stable also at room temperature (Fv=6%). The remaining 4 TWIP steel variants (TWIP2-3-4-5) revealed a fully austenitic dendritic solidification. The austenitic phase is stable down to room temperature. Microsegregation associated to dendritic solidification determine zone with lower local SFE (low Mn and C). In these zones the presence of ε-martensite was detected.

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WP2: Fundamental investigation on the physical metallurgy of TWIP steels The investigations carried out on the physical metallurgy of TWIP steels revealed the following results: Task 2.1.1 The SFE of TWIP2 steel was measured by means of TEM of extended nodes size produced by interaction of extended dislocations. The average SFE value for TWIP2 steel resulted of 21.4 mJ/m2. The standard deviation of SFE values is consistent with other measurements reported in literature. The comparison with thermodynamical model reveal that the Dumay model gives the best correlation with the measured SFE for TWIP2. The model of Dumay predicts that the variant TWIP 1, 3, 4 and 5 should have SFE values below 18 mJ m-2. Task 2.1.2 The temperature for deformation induced transformation Md30

γ ε was evaluated for TWIP2, TWIP3 and TWIP5 and the values are respectively -170°C, -145°C and -72°C. These results again confirm that TWIP 2 and TWIP3 are the steel grades revealing the best microstructural stability and TWIP effect during deformation even at quite low temperature. Task 2.2.1 Hot deformation resistance of high-Mn TWIP steels is dependent on Mn content (strengthening about 2 MPa/wt%) and increases with increasing Al alloying up to 6% (strengthening about 12 MPa/wt%). Nb and N in the contents used here in steels have a minor influence. The austenitic high-Mn TWIP steels exhibit higher deformation resistance than those of low-C, C-Mn-Nb and austenitic stainless steels. Flow stress curves exhibit broad stress peaks at quite low strains. However, the completion of DRX occurs slowly. Very fine grain size is obtained as a result of DRX. SRX kinetics of TWIP steels is faster than that of Type 304 and C-Mn-Nb steels and slower than that of low-C steels. Mn is the main element retarding the rate of SRX and Al has only a minor contribution. The regression equation for the static recrystallization kinetics for TWIP steels can be used to predict the SRX rate under given conditions. Grain size is refined effectively by SRX. Task 2.2.2 • Recrystallized grain sizes of TWIP steel of 22% Mn - 0.6% C (in mass-%) cold rolled in the range

of 40%-70% reductions and isothermally annealed in the temperature range 600 °C ≤T ≤ 900 °C are very small, D ≤ 2 μm.

• At 450°C the effect of annealing is very weak. After some seconds there is some hardening (static ageing by solid solution segregation to dislocations) and a very weak softening thereafter, attributable to recovery, without any noticeable metallographic changes.

• Above 700 °C, recrystallization is complete in less than nine minutes and takes less than ten seconds above 800 °C. At 900 °C ≤ T ≤ 1100 °C, the kinetics observed only corresponds to grain growth.

• Recrystallization and grain growth textures are very weak. Consequently, the elastic and plastic anisotropies of annealed TWIP sheets will be negligible.

• Although there is no apparent texture change from recrystallization to grain growth, there is a strong change in the grain boundary composition. The fraction of Σ3 twin boundaries increases suddenly from 14% to 40% when grain growth starts and it remains constant thereafter independently of the grain size reached.

• An empirical equation of grain growth for this steel has been obtained with an apparent activation energy QGG = 363 ± 60 kJ/mol and an exponent nGG ≈ 3.9.

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Task 2.2.3 Physically based recrystallization models have been applied to the experimental data for cold rolled and annealed TWIP steel. Several models were considered with different assumptions for the nucleation behaviour and boundary velocity behaviour. All the models have assumed that grain boundary mobility was controlled by interaction with manganese solute atoms. The possible influence of twins on the recrystallization process was not considered. From the modelling results the following conclusions can be drawn:

• The model which best described the experimental data assumed that the nucleation was site saturated and the grain boundary velocity decreased with time due to static recovery.

• Comparison of this model with more experimental data revealed good agreement when the nucleation density was a free model parameter.

• Imposing an average nucleation density and recalculating recrystallization curves still gave reasonable agreement with experimental results.

Task 2.3 The precipitation behavior of TWIP steel depends on one side on the C, N, and other elements carbide-nitride formers (Ti, Al, Nb, V) content. On the other side the precipitation behavior is strongly affected by thermodynamic stability of austenite in the range 500-700°C. Infact the ternary system Fe-Mn-C under particular conditions of temperature and chemical composition forese that the austenite could partially transform in ferrite+carbides (pearlite like structure). The study carried out on TWIP steels by SEM-EDS and STEM-EDS allows to state the following conclusions:

• Both TWIP1 and TWIP4 show, at temperature below 700°C the tendency of austenite to destabilize and to forme ferrite+carbide structures. These observation is in agreement with other results achieved on TWIP1-4 and confirm that the austenite stability of these steel grades is lower with respect to reference grade TWIP2 (Fe-22Mn-0.6C);

• TWIP2 and TWIP3 show a similar behavior. Both steel variants revealed a massive carbide precipitation of cementite type (Fe, Mn)3C in the range 500-700°C. No ferrite phase was detected in the range 500-700°C.

• TWIP5 variant reveals a very stable structure: no massive carbide precipitation was detected in the range 500-700°C. The precipitated particles are constituted of cementite carbides and fine (Ti,Al) carbo-nitrides. No ferrite phase was detected even after long soaking time in the range 500-700°C.

Task 2.4 The task 2.4 was focused on demonstrating the feasibility of TWIP steel production at Duferco La Louviere steel works and to define the industrial conditions for TWIP steel processing. The results of the study allowed to conclude that TWIP steel hot rolling is feasible at Duferco La Louviere plant, provided that a dedicated operating practice, quite different from conventional low C-Mn steel, can be adopted. This is the result from the point of view of plant capability. On the other side the impact on the production schedule due to the insertion of a mini-program of TWIP steel has been evaluated in terms of costs, time, fuel consumption, etc. Summarising the operating practice for hot rolling process of TWIP steel production at DLL are:

• Slab dimensions: thickness 250mm, width 1000mm, length 5000mm; • Slab discharging temperature range: 1180-1220°C; • Entry finishing temperature range: >1040°C; • Finishing rolling temperature: 900-920°C;

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• Hot band thickness achievable > 4.0 mm; • Coiling temperature: 450-500°C. • Hot rolled strip pickling: conventional pickling process as adopted for C-Mn steel (HCl acid).

WP3: Study of the deformation mechanisms and strain hardening behaviour Task 3.1.1.1 All the TWIP steel variants investigated revealed excellent tensile properties. However significant differences are present in terms of deformation mechanisms and microstructural evolution during deformation. The tensile properties resulted to be quite sensitive to the specimen surface preparation but the investigation of this phenomena is reported in the task 3.1.1.3. Temperature effect The tensile properties of TWIP steel grades show a quite similar behaviour at increasing temperature with only some quantitative differences. The yield stress is slightly affected by temperature while the tensile strength is significantly influenced by temperature. This means that at higher temperature the Ys/UTS ratio tends to increase due to progressive change of the deformation mechanism (twinning dislocation glide) and a consequent reduction of the strain hardening ability and elongation is detected. Strain rate effect For equivalent plastic strain below 0.25, the behaviour of TWIP steels is very similar at room temperature and the strain rate influence on the strength is very weak. Only at very high strain rates ( 200>ε& s-1) there is a small strain rate induced increment of the flow stress. Strain hardening ability The main characteristic of a true TWIP steel (TWIP2 and TWIP3) is the presence of a intermediate stage in the strain hardening curve the instantaneous coefficient n increase with strain. This should be related to the effect of deformation induced twinning occurring after the first stage of the σ−ε curve predominated by dislocation glide. The new twins act as barriers to dislocation motion, and lead to an increase in strain hardening rate. Depending on the extent of twinning, this leads to the observed overall hardening rate. The other TWIP grades (TWIP1,4,5) are characterised by a lower austenite stability and this is demonstrated by the occurrence of martensite (ε or α) during deformation or presence of ferrite also at zero strain as for TWIP1. The strain hardening of TWIP 1,4,5 grades show a slight different behaviour as a function of strain, due to the complex combination of deformation mechanisms (twinning+dislocations) and deformation induced phases (ε-martensite and α-martensite).

Task 3.1.1.2 The results of this task can be summarized in the following topics:

• TWIP2 and TWIP3 are characterized by largest stability of austenite phase under deformation from 250°C down to -180°C. Within this temperature range the deformation induced twinning represent the main deformation mechanism. At temperature of 350°C both steels do not reveal deformation induced twinning.

• TWIP1, TWIP4 and TWIP5 XRD pattern analysis revealed the occurrence during deformation at room temperature in addition of deformation induced twinning (TWIP effect) even formation of second martensitic phases (α’ and ε). This tendency became stronger at low temperature due to decrease of SFE. Since the fraction of epsilon martensite increases continuously with strain, it is suggested that alpha martensite forms directly from austenite and not through the two-step transformation where epsilon martensite is the intermediate phase. The occurrence of deformation induced twinning disappears before 250°C.

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Task 3.1.1.3 The results achieved can be summarized in the following points:

• The results of the investigations carried out in the present task 3.1.1.3 allow to argue that the embrittlement problems of TWIP steels are mainly related to decarburization and Mn depletion producing a mixed γ+α’microstructure in the subsurface zone.

• Carbon and manganese concentration profiles on TWIP steels show that both C and Mn are depleted in the subsurface zone due to annealing treatment under decarburizing atmosphere. It is worthy to note that the depth of Mn depleted zone (typically <30μm) is significantly lower than C decarburized layer.

• All the TWIP samples annealed at 1000°C and 1200°C reveal in the subsurface zone a mixed microstructure γ+α' even without any deformation with a resultant magnetic behavior The presence of α'-martensite is due to the local lowering of SFE in the C and Mn depleted zone.

• TWIP2,3,4,5 steels reveal a similar behavior that is in good agreement with the theory of decarburization in austenitic phase (Birks-Jackson model) strictly related to bulk carbon diffusion. The decarburization depth is more sensitive to annealing temperature (exponential dependence) with respect to soaking time (t 1/2 dependence).

• The decarburization depth of TWIP steels is larger than carbon steel probably because for austenitic steels the carbon is in solid solution and is ready to react with oxygen. The kinetics of decarburization of a medium carbon steel (C45) on relatively short time is decreased by phase transformation α+pearlite γ. During this transient the formation of a ferrite layer on the strip surface could slow down the steel decarburization kinetics.

• In terms of decarburization the final annealing process (after cold rolling) could be critical. For this treatment a controlled annealing furnace atmosphere have to be considered in order to avoid the occurrence of decarburization and so the formation in the subsurface zone of a mixed γ+α’ microstructure.

• Tensile tests (ε’=10-2s-1) carried out on samples with different hydrogen content show similar tensile properties. The elongation to rupture of hydrogen charged specimen is markedly higher than the strip sample with a surface decarburized layer (impaired by presence of martensite in the microstructure).

• The high desorption temperatures indicate that the hydrogen is trapped by high energy bonding. This result would suggest that the diffusible hydrogen can be considered negligible. Therefore, under plastic deformation the hydrogen absorbed cannot move so easily to reach dislocation and/or crack tip, creating favourable condition to brittle fracture.

Task 3.1.2 The stress ratios (the yield stress to the tensile strength) of TWIP steels are much lower than those of the low C cold drawing steels under quasi-static conditions. This means TWIP steels promoted higher strain hardening potential. Under dynamic tensile conditions, stress ratios of TWIP steels increased as a result of increasing yield strength, but they are still lower than those of the low C cold drawing steels. Furthermore, the n-values of TWIP steels from dynamic tensile tests are much higher than those of conventional automotive steels. The enhancing of the ductility, strength and strain hardening of material is advantageous for crash energy absorbing characteristics. TWIP steels promoted higher ductility and higher strain hardening with high strength, and therefore TWIP steels have higher crash energy absorption and consequently higher crash safety than steels TRIP700, DP600 and H340LAD. Task 3.1.3 The hot torsion results show a strong temperature effect in the strength level and an important effect of strain rate in the ductility. The trend of the behaviour is common for the four steels investigated

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(TWIP2,3,4,5), although TWIP3 is clearly more ductile that the other three compositions. A change in the behaviour from cold-work to warm/hot-work occurs at about 250ºC. At the highest strain rate and at low temperature, the initial strain hardening rate is much higher than its static counterpart, a fact that could be linked to an enhancement of the twinning activity at high strain rates. However, only TWIP2 and TWIP3 below 250ºC clearly behave as expected from TWIP steels, with the typical hardening stage associated to profuse deformation twinning. In all other cases the work hardening is an approximately linearly decreasing function of the flow stress soon after the yield stress, a behaviour characteristic of dislocation-mediated plasticity controlled by dynamic recovery. Task 3.1.4 The hot ductility curves of the TWIP steel variants show quite good high temperature performances. The hot ductility in some cases, such as for TWIP 3 and TWIP2 (RA>60% between 700-1200°C), is higher than that of the austenitic AISI 304. Task 3.1.5 The yield stress for 0.2% plastic elongation clearly shows a grain size dependence. The Hall-Petch constant found for TWIP2 steel grade is =HPK 356.5 MPa μm1/2. The results of this task allowed to achieve a better description of the grain growth kinetics of the TWIP2 steel. It is worthy to note that for fine grain size (<10μm) the twin boundary fraction is significantly lower than for larger grain size. This means that the TWIP effect is dependent on grain size and the best performances are relevant to a grain size in the range 15-30μm. Task 3.2 The fatigue behavior of three high-Mn TWIP steels, with slightly different Mn contents (between 16 and 22 wt.%) and Nb or Al alloying were investigated using reversed bending loading and examining the cyclic damage features on surfaces. The main conclusions can be drawn as follows:

• Fatigue behavior of three TWIP steels is quite identical. Fatigue stress limit (the cyclic life beyond 2x106 cycles) is well above their yield strength values. The ratio of fatigue limit/tensile strength is 0.42-0.48 that is quite a similar value as commonly observed for various carbon steels and for Types 301LN and 316L austenitic stainless steels.

• During cyclic loading, planar slip bands are formed in an early stage of fatigue life consisting of extrusions and intrusions. With continuing cycling, the slip bands intersect with grain boundaries as well as annealing twin boundaries producing local strain concentrations that induce microcracks at these boundaries.

• Fatigue crack embryos nucleate at an early stage of fatigue life (≈25%) at sites of intersections of slip bands and grain boundaries as well as annealing twin boundaries.

• Crack propagation takes place along slip lines, grain and twin boundaries but the overall path is mainly transgranular in its character. In this stage, ductile striations are formed on fracture surfaces.

• Microcracks link and propagate readily along grain boundaries indicating some degree of inherent grain boundary weakness, as suggested in the literature for austenitic high-Mn steels.

• Neither mechanical twins nor ε-martensite are formed during cyclic loading in the investigated TWIP steels, so that the TRIP or TWIP effects seem to play no role in the course of high-cycle fatigue.

• The degree of cyclic hardening revealed by hardness is strongly dependent on the grain size decreasing with refined grain size.

• Refinement of the grain size improves significantly the fatigue strength of the 0.6C-22Mn TWIP steel.

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Task 3.3 Comparing the Charpy values of TWIP variants arise that TWIP1 reveals the best results at all tested temperature. TWIP2 and TWIP3 are slightly better than TWIP4 and TWIP5. Qualitatively the Charpy energy behavior of TWIP steels as a function of the temperature is quite similar to stainless steels such as AISI304. In fact the Charpy energy remains of the same order from +150°C down to -50°C, with a slight increase from RT down to -50°C. Task 3.4 The work hardening of TWIP2,3,4,5 steels is significantly higher than stainless steels; at a true strain of 0.5 the difference in terms of stress is about 38%. The above result implies difficulties in cold rolling in terms of loads, number of passes or reduction ratio at each stage. Depending on the hot strip thickness, an intermediate annealing treatment (two step cold rolling process) could be necessary for obtaining the aimed final cold strip thickness (<2.0 mm). WP 4: Basic characterisation of application properties: formability, weldability and coatability Task 4.1 It can be seen that variation among the steels is quite small. Generally, the Erichsen Index (IE) is higher if the elongation is higher, and the best values are obtained for TWIP2. However, better IE values were obtained at high-speed Erichsen testing. This can be attributed to the adiabatic heating of the sample during the tensile test increasing the SFE and resulting in decreasing density of mechanical twins. However, with increasing the strain rate up to 1000 s-1, the total elongation increases again reaching values above that in tensile testing at the strain rate of 0.1 s-1, It is possible to make comparison between the elongation in dynamic tensile tests ( 1000≈ε& s-1) using the Hopkinson split bar method and that in quasi-static tensile tests. Then, for example, for TWIP2 the total elongation is 83% under quasi-static tension and 80% under high-speed tensile testing, i.e. the total elongation in these two cases is almost equal. In the conventional Erichsen testing according to the DIN 50101, the typical strain rate can be estimated to be order of 0.01 s-1. (For instance, no 0.4 cylinder speed is 0.56 mm/s). In high-speed Erichsen testing, where the speed of impact front is about 200 m/s, the strain rate can be order of 5 s-1 (dome height 10 mm; at 200m/s it takes 0.05 s. Strain is 0.26, hence, the strain rate = 0.26/0.05s = 5/s), i.e. about 500 times higher. Task 4.2 The visual control of galvanized sample’s surfaces revealed a general uniformity of the coating layer. Coating adhesion test was made and as a results all the tests were acceptable according to the standard since no detachment of any small square was verified. No noticeable differences were detected between the two pre-treatments in the galvanized adhesion layer test. All the samples after the galvanization process revealed the typical galvanization structure. The batch hot-dip galvanized coating consists of a series of zinc-iron alloy layers with a surface layer of almost pure zinc. The alloy layer is as much as 50% of the total thickness and it consists of two or more distinct zinc/iron layers. Each layer has a specific amount of iron and zinc. Task 4.3 To assess the weldability as well as its main influencing factors, experimental laser welding analysis of four TWIP has been performed. The work allowed to conclude that the joining of TWIP steels is easily achievable by laser welding. The factors that determine the good welding behaviour and quality of TWIP steels are intrinsically related to the steel composition and austenitic structure.

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Concerning the dissimilar steel welding all the weld and heat affected zones present acceptable hardness values which indicates the mechanical properties of the base materials are matched or improved after laser welding. This study permitted to prove the good potentialities of TWIP materials on automotive applications where the welding ability is one the the main requirements.

WP 5: Industrial trial Task 5.1 TWIP3 is the selected TWIP grade to be casted at Duferco steelworks. TWIP3 steel variant (chemical composition reported in following table) is characterized by a markedly lower Mn content that is the most expensive element in TWIP steels and, as well known, the element that could have unwished impact on environment during steel making due to Mn evaporation. The properties revealed by this steel if, in some instance, are below the TWIP2 are really promising and could be further improved by means of a fine tuning of the alloy design and industrial processing conditions. Up to now no patent application has been submitted.

Fe Mn C Si Nb

min bal 16.00 0.60 0.20 0.015

max bal 17.50 0.65 0.30 0.030

Table FS2: Range of elements of the TWIP steel selected for industrial trial (TWIP3). Task 5.2 TWIP steel industrial heat was not performed due to the stronger than expected relevance of the following issues:

• Hydrogen embrittlement susceptibility of high Mn steel. To minimize this risk the heat was postponed in order to wait for the installation of a VD facility at Duferco steelworks aiming to achieve a significant reduction of hydrogen content in steel. Unfortunately several causes determined a larger delay of the VD facility completion inclusive, of course, the recent steel markets crisis that widened the time for VD installation and commissioning work.

• Definition of the steelmaking route of high Mn steel using a direct current electric arc furnace (DC-EAF). The feasibility of TWIP steel on a direct current arc furnace required a in depth study of interactions between high Mn steel and furnace soils electrodes to avoid furnace damagement. The conclusion of this study yielded to state the feasibility of a high Mn steel. However some countermeasures were introduced in the steelmaking operating practice to increase the EAF safety, among which, minimization of the Mn melting time and hence the interaction time between Mn and EAF soil electrodes and, before the heat, assessment of the consumption status of the soil electrodes.

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Scientific and technical description of the results

WP1 Testing material supply and basic metallurgical characterisation of cast materials Task 1.1 Definition of the TWIP steel compositions matrix to be investigated based on the reference TWIP steel composition Fe-Mn-Al-Si and newer TWIP steel compositions (high C and N). Task 1.1.1 Selection of TWIP steels to be investigated The first task of the project consisted in the definition of the TWIP steel compositions to be investigated. After a in depth discussion between the partners it was decided to select 2 reference TWIP steel chemical compositions on the basis of the literature survey. The reference TWIP steels were selected for two order of reasons:

1. To investigate the basic metallurgical properties of TWIP steel according to the technical annex and

2. To evaluate the effect of different alloying solutions on TWIP steel properties.

The reference TWIP steel selected were: 1. Fe-22Mn-3Al-3Si 2. Fe-22Mn-0.6C

In addition to the reference TWIP steel, further 3 TWIP variants were proposed. The scope of the introduction of additional TWIP variants to be investigated was to investigate the effect of different Mn, Al, Si, C, N content on:

• metallurgical properties such as SFE, tensile properties, hot and cold deformation properties, etc.

• application properties: coating ability, welding ability; • manufacturing route: from hot rolling up to final cold processing route.

On the basis of the above considerations it was decided to investigate on TWIP steels variants with the following alloying modifications: a) Lower Mn content (16-18%) could be exploited in order to reduce the heat costs and the enviromental problems related to high Mn evaporation during steelmaking process; b) Al addition in order to stabilize γ with respect to ε phase. c) Addition of C and N to achieve the optimal SFE value and austenite stability but avoiding precipitation of carbide/nitride during heat treatments or high temperature treatments. The addition of N, in partial substitution of C, will be tested also to verify the influence on the mechanical properties. d) Nb addition: Nb is reported in literature to have a positive effect on SFE but the effectiveness is still unclear. Summarising the steels to be investigated are reported in the following table. The reference steel are well defined in terms of composition while for the additional TWIP variant TWIP3-4-5 calculations of the SFE, by means of thermodynamical models, are necessary.

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TWIP variant 

Mn  C  N  Al  Si  Nb 

TWIP1 (ref.) 

22 0.01 - 3.0 3.0 -

TWIP2 (ref.) 

22 0.6 - - 0.2 -

TWIP3  18-22 0.6 - - 0.2 0.02 TWIP4  16-18 0.6 - 1.5 - - TWIP5  18-22 0.2-0.3 0.2-0.3 - 0.2 - Table 1.1.1.1 Selected TWIP steel compositions (Fe balance).

Task 1.1.2 Definition of the TWIP steels compositions matrix Thermodynamic model for SFE The SFE plays a strategic role on selection of the deformation mode and therefore on design of TWIP composition. A thermodynamical model for SFE calculation was developed mainly based on thermodynamic models for the chemical driving force of the εγ → martensitic transformation. The SFE itself depends upon two important metallurgical and thermodynamic parameters: the composition and the temperature. Therefore, SFE can be modeled as a thermodynamic function, if each stacking fault is considered as a double phase boundary between the γ-austenite (FCC) and the ε-martensite phase (HCP). SFE of γ in the Fe-Mn base alloys has been calculated using the following equation proposed by Olsen and Cohen [1]: εγεγ σργ 22 +Δ== →GSFE fcc (1) ρ is the planar packing density of a closed packed plane.

εγσ is the interfacial energy of a coherent boundary between the γ and ε phases. The established values of εγσ in transition metals are from 10 to 15 mJ/m2. In the present calculations, 15 mJ/m2 was used. Based on the regular solution model, the free energy change for the εγ → phase transformation

εγ →ΔG can be calculated as a function of temperature and composition, and then SFE can be estimated based on the Olsen and Cohen’s model (Eq. 1). Two ways were used to calculate εγ →ΔG .

• Using the equations found in literature giving εγ →ΔGi for each elements; • Using ThermoCalc to calculate the total free energy change for the εγ → phase transformation

εγ →ΔG . At moment of the present project start-up, two thermodynamical models for SFE calculations were reported in literature: Grässel model [2] and Allain model [3]. In the table 1.1.2.1 are shown the parameters for calculation of SFE for both models. It can be noticed that using these equations somewhat different results for SFE would be obtained. It can be seen that, while Allain’s equations give lower SFE for low C with low Mn, e.g. 0.1C with 15-20Mn, Grässel’s equations seems to give more satisfying calculations of SFE at the interesting ranges

15

of C and Mn. Furthermore Allain’s model has no equations for including Al and Si, so that they cannot be applied for Al or Si-bearing steels. For this reason it was decided to adopt the Grässel model in order to define the steel chemical composition to be investigated in the project. Parameter Grässel et al. [2] Allain et al. [3]

εγ →Δ FeG 200222.0685.185.821 TT ++− -2243.38+4.309T εγ →Δ MnG 200455.07.20.3925 TT +− -1000.00+1.123T εγ →Δ CG 12.24595− -22166 εγ →Δ AlG 5481.04-1.79912T εγ →Δ SiG T+−1800 εγ →ΔΩFeMn MnX1.152825.9135 +− 2180+532(XFe-XMn) εγ →ΔΩ FeC 42500 42500 εγ →ΔΩ FeAl 3323 εγ →ΔΩ FeSi 1780

σ 15 9 Table 1.1.2.1: Numerical values and functions used for the calculations of SFE in two models. Definition of the steel chemical compositions to be cast According with the Grässel equation, the calculated SFE values of reference steels (TWIP1) Fe-22Mn-3Al-3Si and (TWIP2) Fe-22Mn-0.6C are respectively 26 and 42 mJ/m2. On the basis of the Grässel model the effect of each element on SFE was evaluated and the three new variants were selected as reported in the following table.

TWIP variant Steel composition Grässel et al SFE (mJ m-2)

TWIP 1- ref. Fe-22Mn-3Al-3Si-0.01C

42

TWIP 2 – ref. Fe-22Mn-0.6C-0.2Si

26

TWIP 3 Fe-18Mn-0.6C-0.2Si

23

TWIP 4 Fe-16Mn-1.5Al-0.2Si-0.3C

28

TWIP 5 Fe-21Mn-0.2Si-0.2C-0.2N

17

Table 1.1.2.2: Calculated SFE for TWIP steels at 298 K. It is worthy to note that in Grässel model the nitrogen effect is not taken in consideration.

16

Task 1.2 Laboratory VIM ingot casting Five heats were performed at CSM VIM facility with the following chemical compositions. The ingots produced were about 80kg in weigth.

Table 1.2.1.1 Ingots chemical analysis.

The ingots were hot rolled and afterthat cold rolled in order to produce the strip samples needed for the start up of the investigations. Task 1.3 - Solidification structure characterisation TWIP1 ingot sample The ingot sample of TWIP1 shows a columnar macrostructure at the ingot surface, approx. 1.5 cm thick, and a core region of equiaxed grains, fig. 1.3.1a). Its microstructure is a two-phase structure: austenite matrix with an interdendritic ferrite phase that represents about 10% in volume, see fig. 1.3.1b) and fig.1.3.2.

a) b) Fig. 1.3.1: Images of TWIP1 ingot. a) Macrostructure of the TWIP1 ingot sample, polished and Nital 2% etched b) SEM image of structure 2% Nital etch. The sample is ferro-magnetic and this is in agreement with ferrite presence in the microstructure. The solidification structure consists of austenite γ-fcc + ferrite α-bcc (fig. 1.3.2).

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a) b) Fig. 1.3.2: a) Optical microstructures of TWIP1 in the cast state of lathy (black) and vermicular ferrite; b) XRD of TWIP1 in the cast state showing the presence of duplex structure. The ferrite phase appears both with lathy and vermicular morphology. XRD-analysis was performed to confirm the presence of the austenite and ferrite phases and the diffractogram is shown in Fig. 1.3.2b). Concentration profiles from the cast TWIP1 structure by SEM-EDS mapping indicate different contents of the alloying elements of Mn and Al between the ferrite and austenite phases, as shown in fig. 1.3.3. As seen in figs. 1.3.3.b-c, the ferrite phase is enriched in Al and depleted in Mn.

a) b) c) Fig. 1.3.3: The element distribution (X-ray map) in the austenite and ferrite phases in cast TWIP1, (a) SEM image, (b) Al distribution, and (c) Mn distribution.

Fig. 1.3.4: SEM-EDS analysis locations on the cast TWIP1.

The presence of ferrite should be related to the presence of both high Si content in connection with high Al, which are strong ferrite formers. The C content in the steel is not enough to stabilize the austenite in spite of the high Mn content. Figure 1.3.4 shows the dual-phase microstructure with five locations of

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analysis. As shown, there is a significant difference in Al and Mn contents between the austenite and ferrite phases. However, the difference in the Si content is low. The partition ratio PD of Al, Mn and Si, defined as the ratio of the concentration of an element in ferrite to that in the adjacent austenite, are given in table 1.3.1. The partition ratios of Al and Si are higher than 1 and that of Mn is lower than 1. It is well known Al and Si are ferrite formers and Mn favours the austenite in the Fe-Mn-Al system. The partition of Si is quite slight. Element Al Mn Si PD 1.62 0.77 1.18

Table 1.3.1: Partition ratios PD of alloying elements between austenite and ferrite in TWIP1. TWIP2 The microstructure of the cast TWIP2 (Fe-22Mn-0.6C) is fully austenitic with large elongated grains. The microstructure displayed the dendritic structure, as shown in Fig. 1.3.5. The chemical etching suggested clear composition variations in the structure which fact was confirmed by EPMA. A fine criss-cross pattern of dendrites can be detected within each grain; dendrite cores being lower in Mn and appear relatively bright. Between the dendrites (interdendritic regions) the Mn content is higher and these regions are more heavily etched appearing darker in colour.

Fig.1.3.5: OM photo of the microstructure of the cast TWIP-2.

Electron probe microanalysis (EPMA; JeolJXA 8200) with the pixel size of 0.2 μm was applied to investigate the microsegregation of the alloying elements. A Vickers hardness tester was used to mark the start and end points of the line scan. Concentration profiles recorded from the dendritic cast structure of TWIP2 show several maxima and minima. The minima of both Mn and C are located at the dendrite cores. The segregation ratio (S), i.e. the ratio of an element concentration in the interdendritic regions to that at the dendrite cores, was estimated after the quantitative calibration of EPMA. It was found that SMn and SC are 1.6 and 1.5 respectively. TWIP3 The microstructure of TWIP3 (Fe-18Mn-0.6C-0.019Nb) steel, shown in fig. 1.3.6, revealed a austenitic solidification with dendritic morphology.

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Fig. 1.3.6: OM photo of the microstructure of the TWIP3 in the cast state.

The microsegregation of the steel was examined using the EPMA similarly as for TWIP2. C and Mn have simultaneous segregation, peaks locating at the interdendritic regions, similarly as in the case of TWIP2. The segregation ratio SMn is equal to 1.5. The Nb content is so low that it could not be recorded. With careful focusing, the microstructure of TWIP3 revealed the presence of ε−martensite phase with the lath morphology, as shown in Fig. 1.3.7. Martensite plates extend across the whole austenite grains, i.e. crossing several dendrites. It can been suggested that the low stacking fault energy of this steel favours the formation of ε-martensite in the cast ingot, similarly as in TWIP 1.

Fig. 1.3.7: Presence of ε-martensite in the cast TWIP3.

The presence of ε−martensite in TWIP3, as seen in Fig. 1.3.7, may be attributed to the segregation of Mn. Martensite plates are always located in dendrite cores, where Mn has its minimum concentration. TWIP4 The microstructure of the cast TWIP4 (Fe-16Mn-1.5Al-0.3C) displayed dendritic structure, as shown in Fig. 1.3.8. The maximum distance between two dendrite cores is about 900 μm. However, at high magnifications, the presence of fine plate phase, i.e. ε-martensite, in the dendritic regions can be seen. However, this phase is present in the dendritic core regions, as shown in Fig. 1.3.8.

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Fig. 1.3.8: OM photo of TWIP-4 in the cast state. Fine plates of ε-martensite in dendrite cores (right). It can be concluded that segregation of alloying elements plays a strong role on the generation of preferential nucleation sites for the ε-martensite. From the elemental analysis, it is known that in the dendrite core regions have lower Mn and C contents, i.e. the SFE is locally lower there. According to Olson and Cohen’s thermodynamic model, the lower Mn content would decrease SFE in the dendrite core regions to 24 mJ/m2, while the higher Mn content in the interdendritic regions would mean a SFE of 45 mJ/m2. This explains the observation the ε-martensite can exist in the dendrite core. TWIP-5 The cast microstructure of TWIP5 (Fe-22Mn-0.2C-0.2N) displays a dendritic segregation, as seen in Fig. 1.3.9. At high magnifications, the fine plate-like phase of ε-martensite is present in the dendrite cores and arms, but such a phase is not found in the interdendritic regions, as seen in Fig. 1.3.9(b). This situation is exactly similar to that for TWIP4. However, the ε-martensite fraction in TWIP5 is much higher.

a) b) Fig. 1.3.9: OM photos of the microstructure of the cast TWIP5.

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EPMA analysis was conducted to record the micro-segregation using a line crossing two dendrite cores, with the length of 250 um. The C and Mn concentration profiles begin at the peak followed by a minimum, thereafter there is second maximum that corresponds to the interdendritic region (etched dark in the figure). However, no variation of N can be revealed. The segregation ratio of Mn is equal to 1.5, similarly as for TWIP2 and TWIP4. To study further the short-range segregation, a distance between two secondary dendrite arms was analysed using EPMA. As observed, the maxima of Mn and C correspond to the interdendritic region between the two secondary arms. As the dendrite grows into the melt, and as secondary arms spread from the main dendrite stem, while the solute Mn is rejected. In the N distribution, no segregation could be detected. Determination of the liquidus and solidus temperature of the TWIP steel variants The evaluation of liquidus and solidus temperature are of paramount importance in order to define the operating practice steelmaking and continuous casting process. On the basis of these critical temperatures is calculated the liquid steel overheating (in tundish and mould). At the same way the knowledge of solidification range is also important to define the optimal continuous casting process. Differential thermal analysis (DTA) was carried out in order to evaluate the liquidus and solidus temperature of TWIP2 steel. For other TWIP variants the critical temperature were qualitatively evaluated by Thermocalc model. The results of measures and calculations are summarized in the table 1.3.2. Comparing the measured and calculated critical temperatures the following considerations can be argued for TWIP 2 steel:

• The measured liquidus temperature is quite close to calculated value (the difference is only 0.4%);

• The calculated solidus temperature is underestimated of 30°C (about 2%).

Steel grade Tsolidus (°C) Tliquidus (°C) TWIP 2 1322 (DTA measure) 1406 (DTA measure) TWIP2 1292(calculated) 1400 (calculated) TWIP 3 1312(calculated) 1420 (calculated) TWIP4 1314(calculated) 1418(calculated)

C, Mn steel 1495(calculated) 1525 (calculated) Table 1.3.2: Tsolidus and Tliquidus of TWIP steel variants.

Conclusion of Task 1 Five TWIP steel variants were selected on the basis of stable austenitic phase and low SFE. The variants were of two main types Fe-Mn-C-N and Fe-Mn-Al-Si-C.

Table 1.3.3 Ingots chemical analysis.

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Of these five TWIP variants the first variant revealed a mixed primary solidification structure with austenite + ferrite. The ferrite phase is stable also at RT. The remaining 4 TWIP variants revealed a fully austenitic solidification. The austenitic phase is stable down to RT. Microsegregation associated to dendritic solidification determine zone with lower local SFE (low Mn and C). In these zones the presence of ε-martensite was detected.

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WP2: Fundamental investigations on the physical metallurgy of TWIP steels Task 2.1 Measure of SFE by means of Transmission Electron Microscopy observations of extended nodes. Evaluation of Ms γ−>ε temperature. Task 2.1.1: Measure of SFE by means of TEM observation of extended nodes Background Mechanical twinning and ε-martensite formation are competitive deformation mechanisms that are very similar from the morphology and origin point of view. Whelan first [1] and Brown later in 1964 [2] showed that the dissociation of dislocation a/2 <110> into partials of burger vector a/6<112> lying on {111} planes creates stacking fault on successive parallel {111} planes. The crystallographic structure in the faulted region is still FCC but in a twin orientation compared to the matrix. Epsilon martensite formation occurs when the same dislocations glide on every second {111} plane. In this case the deformed region takes the form of a thin lamella or platelets with a close packed hexagonal structure. Both mechanisms are strongly related to the SFE, which controls the energetic cost for creating such defects. The methodology adopted for the SFE measure is based on the work of Remy [3] and A. W. Ruff and L.K. Ives [4] and consists in the measure of SFE by means of TEM measurement of extended nodes size produced by interaction of extended dislocations. Sample preparation methodology The SFE measurement was carried out on a TWIP2 steel specimen. The selected strip sample was a cold rolled strip annealed, in N2 atmosphere, at 1000°C for 5 min. (Fig. 2.1.1.1).

Fig 2.1.1.1: TWIP2 microstructure after annealing treatments at 1000°Cx5 min (average grain size 32 μm). The samples were polished first mechanically up to a thickness of 200 μm, then further thinning was carried out electrolytically down to 80-100 μm. After this step the specimen was deformed by means of a double bending (to produce ‘few’ dislocations). The final step consisted in the jet spray electrolytic thinning to achieve the thin foil specimen suitable for TEM observation. Results Different thin foils were produced and observed by means of a STEM (JEOL -100KV). As a whole 10 nodes were selected during TEM examination and were judged suitable for SFE measurement. In fig

24

2.1.1.2 are shown some images of the extended nodes selected. In tab. 2.1.1.1 are reported the values of the SFE obtained taking into consideration only the suitable extended nodes.

Node W (m) γ mJ/m2

A1 2.5E-08 19.8A 2.0E-08 24.6B 2.0E-08 24.6C 2.7E-08 18.5D 3.1E-08 15.8E 2.4E-08 20.7F 1.7E-08 28.7H 2.8E-08 17.7I 1.9E-08 25.7L 2.7E-08 18.1

average 21.4dev. Std. 4.2std error 0.4

Fig. 2.1.1.2: Extended nodes detected on a TWIP2 sample. Table 2.1.1.1: SFE results achieved on TWIP2.

The average value of SFE for TWIP 2 resulted in 21.4 mJ/m2. The standard deviation of SFE values is consistent with other measurements reported in literature. Comparison of TEM result with empirical models The stacking fault energy (SFE) of high Mn TWIP steels was evaluated using thermodynamic models due to Grässel and Allain. In addition ThermoCalc code was used to assess the SFE. The analysis revealed that ThermoCalc and the model due to Grässel gave good agreement, whilst that due to Allain et al. gave significant differences also because this model did not consider the influence of Al and Si on stacking fault energy. However, recently the model of Allain et al has been extended to consider Al and Si by the work of Dumay et al [5]. Thus, to assess the SFE of TWIP grades 1 to 5 the new model will be used and compared to the results of the model due to Grässel. The new model due to Dumay et al is reported in the table below.

25

Table 2.1.1.2: Model parameters for calculation of SFE in Dumay et al model [5].

Steel composition Dumay et al SFE (mJ m-2)

Grassel et al SFE (mJ m-2)

Fe-21.3Mn-3Al-3Si-0.01C (TWIP 1)

11 42

Fe-22.3Mn-0.59C-0.22Si (TWIP 2)

22 26

Fe-17.8Mn-0.6C-0.2Si-0.02Nb (TWIP 3)

17 23

Fe-16.4Mn-1.54Al-0.21Si-0.29C (TWIP 4)

10 28

Fe-21Mn-0.23C-0.2Si-0.2N (TWIP 5)

7 17

Table 2.1.1.3: Calculated SFE for TWIP steels at 298 K. Table 2.1.1.3 shows the calculated SFE from the Dumay model and that due to Grassel for TWIP grades 1 to 5. Comparing the results in table 2.1.1.3 arises that the model of Dumay gives the best correlation with the measured SFE for TWIP2. Conclusions of task 2.1.1 The SFE of TWIP2 steel was measured by means of TEM adopting the extended nodes method. The average SFE value for TWIP2 steel resulted of 21.4 mJ/m2. The standard deviation of SFE values is consistent with other measurements reported in literature. The comparison with thermodynamical model reveals that the Dumay model gives the best correlation with the measured SFE for TWIP2. The model of Dumay predicts that the variants TWIP 1, 3, 4 and 5 should have SFE values below 18 mJ m-2. However the contribution of N and Nb is not taken into account.

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Task 2.1.2 Evaluation of the Md γ ε of TWIP steel variants In low SFE austenitic steels the temperature Msγ ε indicates the temperature at which the austenite phase is thermodinamically instable and hcp-martensite start to form. This temperature is of particular importance because it is strictly related to SFE value and gives an indication of the stability of the austenitic phase during deformation against the formation of hcp-martensite. The volume fraction formed of hcp-martensite depends on the strain and on temperature through the SFE. Generally for austenitic steels is more significant and easier to evaluate the temperature at which the strain induced hcp-martensite achieves a threshold value. For austenitic stainless steel the martensite transformation temperature Md30 is defined as the temperature at which 50% of martensite (α’+ε) is formed after a straining of 0.3 true strain. In order to define also for TWIP steel a similar parameter related to deformation induced martensite the following activities were carried out :

• Two series of TWIP samples were deformed at true strain 0.1 and 0.3 at temperatures in the range 20°C ÷ -180°C using liquid nitrogen to cool down and control the sample temperature;

• The presence of α’-martensite and ε-martensite as a function of the temperature was evaluated metallographically by means of special color etching and measuring the magnetic behavior (ferritoscope).

XRD pattern on deformed samples were also carried out but the low fractions of phases to be measured and also the small difference between the different samples suggested to evaluate metallographically the fraction of α’-martensite and ε-martensite as a function of the temperature by means of special color etching and measuring the magnetic behavior by ferritoscope. On the basis of the results achieved and reported in fig. 2.1.2.1 the martensite transformation temperature Md

γ ε for TWIP steels was defined in these terms:

Md30γ ε = the temperature at which the fraction of hcp-martensite achieve a fraction of 2% after a

deformation of 0.30. Adopting this definition the value for the Md30

γ ε of TWIP2, TWIP3 and TWIP5 was evaluated and reported in the following table 2.1.2.1.

Fig. 2.1.2.1: Hcp-martensite fraction on sample deformed at 0.30 at low temperature.

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TWIP variant Md30 γ ε (°C)

TWIP2 -170

TWIP3 -145

TWIP5 -72

Table 2.1.2.1: Md30γ ε evaluation from fig.2.1.2.1.

Conclusions of task 2.1.2 The temperature for deformation induced transformation Md30

γ ε was evaluated for TWIP2, TWIP3 and TWIP5 and the values are respectively -170°C, -145°C and -72°C. These results confirm, again, that TWIP 2 and TWIP3 are the steel grades revealing the best microstructural stability and TWIP effect during deformation even at quite low temperature. Task 2.2: Characterization of recrystallization behaviour Task 2.2.1 Critical strain for DRX initiation, flow stress, peak stress, maximum softening rate by using a thermo-mechanical simulator (Gleeble) and Hot torsion tests at different temperature (900-1100°C) and strain rates. To study the hot deformation characteristics of the steels, the flow resistance and softening behaviour, the specimens were compression tested using a Gleeble 1500 thermo-mechanical simulator. In hot compression tests, specimens were reheated at 1200 °C for 2 min and then cooled to the test temperature (between 900 °C and 1100 °C) at the cooling rate of 5 °C/s. After 15 s of soaking at the test temperature, the specimens were compressed in a single hit to the true strain of 0.8 at the constant true strain rate 0.1 s-1. The static recrystallization (SRX) kinetics of the steels was studied by employing the double-hit compression test technique at temperatures between 900 °C and 1100 °C and at the constant strain rate of 0.1 s-1. The applied strain was 0.2 and the holding times between the passes were 1 to 1000 s. The typical test schedule is shown in Fig. 2.2.1.1. The 5% total strain reloading method was adopted in determining the recrystallized fraction in order to exclude the effect of recovery from the softening data [1,2]. As conventionally, the softening fraction, X, is calculated from the equation:

( )( )13

23

σσσσ

−−

=X (1)

where σ1 and σ2 are the offset stresses due to the first and second hit, respectively, and σ3 is the flow stress of the work-hardened material.

28

Fig. 2.2.1.1: Schedule used in double-hit compression tests. Hot deformation behavior: TWIP-2 (Fe-22Mn-0.6C) The microstructure of TWIP2 (22Mn-0.6C), characterized using the SEM-EBSD technique, as heat-treated at 1200°C for 2 min on the Gleeble simulator and cooled very fast by water spray to room temperature, is shown in fig. 2.2.1.2. The microstructure exhibits coarse austenite grains with the size about 97 μm.

Fig. 2.2.1.2: SEM-EBSD photo of TWIP-2 steel as heated at 1200°C for 2 min and water quenched. Red lines reveal high-angle (>15°) grain boundaries. Flow stress curves at high temperatures Typical true stress-true strain curves of TWIP2, as compressed at 1100, 1000 and 900 °C at the strain rate of 0.1 s-1, are shown in fig. 2.2.1.3. For comparison, the flow strain curves of a low-carbon steel (0.10C-0.45Mn) are displayed in the same figure. It can be seen that the flow curves are featured by rapid work hardening at the initial strains and a broad stress peak at all test temperatures revealing the occurrence of dynamic recrystallization (DRX). The peak stress (σp) and peak strain (εp) depend on the temperature. For example, at 1100°C, the peak stress (112 MPa) is attained at the strain 0.26 and at 1000°C, the peak (σp = 157 MPa) is at 0.38. The high-Mn TWIP steels have much higher hot deformation resistance than that of low-C steel, as reported previously [3], and also displayed in fig. 2.2.1.3 at all testing temperatures.

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Fig. 2.2.1.3: Typical true stress-strain curves of TWIP-2 (Fe-22Mn-0.6C) steel at high temperatures and constant strain rate of 0.1 s-1. Curves for the low-carbon steel are included for comparison. The peak stresses and peak strains for the both steels are plotted as a function of the inverse temperature in fig. 2.2.1.4. It is seen that the peak strain is higher for the high-Mn steel compared to that of the low-C steel. Also the peak strain is lower for the low-C steel. For example, at 1100°C, the peak strains are 0.16 and 0.26 for the low-C steel and TWIP-2, respectively, indicating a significant delaying effect of high Mn content on DRX. The peak strain εp is an important variable. The critical strain (εc) at which dynamic recrystallization starts can be approximately determined from εp , for Pc εε 8.0≈ (1). The dependence of (εp) on the Zener-Holloman parameter (Z) and initial grain size (d0) has been found as the following: εp = A nd0 Zm (2) εp = A1 Zm (3)

Z = ε′ exp(Qdef/RT) (4) where the parameters A, n, m and A1 depend on the alloy composition. Where ε′ is the strain rate, Qdef is the activation energy of hot deformation, R the gas constant and T the absolute temperature. In the present work, the peak strain (εp) was determined for all the flow curves where the peak stress (σp) was clearly discernable. The results were then fitted with Eq. (3), and shown Fig. 2.2.1.5. The activation energy of deformation Qdef for low-C and high-Mn TWIP steels have been determined in other works [3-5] to be 315 kJ/mol and about 380 kJ/mol for low-C and 25Mn-Al type steels, respectively. It is seen the slopes vary to some extent, so that the Zener-Hollomon exponents (m) values are 0.197 and 0.135 for low-C and 22Mn-0.6C steels, respectively. These values are in good agreement with the values reported for C-steels in the range 0.12-0.22 [6]. Elwahabi et al. [7] measured the Zener-Hollomon exponents for Type 304 austenitic stainless steel to be in the range 0.125-0.156.

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Fig. 2.2.1.4: Dependence of peak stress (σp) and peak strain (εp) on the inverse temperature at the strain rate of 0.1 s-1 for TWIP-2 and low-C steels.

Fig. 2.2.1.5: Plot of ln(εp) vs ln Z for the low-C and TWIP-2 (22Mn-0.6C) steels.

Recrystallization kinetics of TWIP2 A selection of interrupted double-hit compression stress-strain curves determined at 1000°C, at the constant true strain rate of 0.1 s-1, an applied strain 0.2 and interpass times of 3 to 50 s. During the unloading time, static softening occurs by the static recovery at short times and mainly by static recrystallization (SRX) at longer times. The fractional SRX curves at three temperatures (900-1100°C) as a function of holding time after compression to the 0.2 strain at the strain rate of 0.1 s-1 are plotted in fig. 2.2.1.6.

Fig. 2.2.1.6: SRX rates of TWIP-2 at constant strain rate of 0.1 s-1. Double-hit compression data and fitted curves. Times for 50% recrystallization, t50, and the Avrami exponents are listed. As seen in fig. 2.2.1.6, the data can be fitted well with the sigmoidal curves, corresponding to the Avrami-type relationship [8], as follows:

( )[ ]nttX 50693.0exp1 −−= (5)

31

where X is the recrystallized fraction, t time, 50t time for the 50% recrystallization and n is the Avrami exponent. It was found that the values for n of the TWIP-2 are quite low (≈ 0.8-1.2) compared with the values 1.5-2 typical of the carbon steels [2]. TWIP-3 (Fe-18Mn-0.6C-0.02Nb) Initial microstructure The microstructure of TWIP3 sample, after soaking at 1200°C for 2 min on the Gleeble simulator and cooled very fast by water quenching to room temperature, was characterized using the SEM-EBSD technique. The microstructure exhibits coarse austenite grains of the size of about 80 μm. Flow stress curves at high temperatures Typical true stress-true strain curves of the present steel as compressed at 1100, 1000 and 900 °C at the strain rate of 0.1 s-1 are shown in fig. 2.2.1.7. For comparison, the flow stress curve of Fe-26Mn-0.14C at 1000°C/0.1 s-1, tested in another study [9], is displayed in the same figure. It is seen that TWIP3 steel exhibits hot deformation behaviour similar to that of TWIP2. For example, at 1100°C, the flow stress curve has a peak at 0.26, which peak strain is identical for that of TWIP2. Also, at 1000°C, both steels exhibited the peak at the same strain of 0.36. Comparison with the Fe-26Mn-0.14C steel shows that the both steels have the same behaviour and flow resistance up to the peak strain of TWIP3 ( Pε = 0.36), but then the steel with the higher Mn level delaying DRX so that the peak stress is at 0.46. Hence, as observed earlier by comparing 22Mn-0.6C steel with the low-C steel (with 0.45Mn), it is seen that a higher Mn content in steels retards DRX (assuming that C content has no influence). Comparing TWIP steels having equal peak strains, it can be concluded that Nb of 0.022% retards DRX to the same amount than the difference in the Mn contents of the steels (22 and 18%) enhances.

Fig. 2.2.1.7. Typical true stress-strain curves of TWIP-3 (Fe-18Mn-0.6C-0.02Nb) steel at high temperatures and the constant strain rate of 0.1 s-1. A curve for Fe-26Mn-0.14C steel is included for comparison.

32

Recrystallization kinetics of TWIP3 The softening kinetics of TWIP3 was investigated at three temperatures 900°C, 1000°C and 1100°C after compression to the 0.2 strain at a constant strain rate 0.1 s-1. The fractional SRX curves as a function of holding time are plotted in fig. 2.2.1.8. For comparison, the softening of Fe-26Mn-0.14C steel, investigated elsewhere [10], are included in the figure. Also the data are being fitted well with the sigmoidal curves, corresponding to the Avrami-type relationship [8].

Fig. 2.2.1.8: SRX of TWIP-3 (0.2 strain at constant strain rate 0.1 s-1). Double compression data and fitted curves.

Fig. 2.2.1.9: SRX of TWIP-3 determined in a relaxation test at 1000°C/0.2/0.1 s-1.

It can be observed that the SRX kinetics of TWIP3 is high at 1000°C and 1100°C with Avrami exponent n = 1.1 and 1.2 and t50 11 s and 4 s, respectively. However, at 900°C SRX is very slow. SRX kinetics of Fe-26Mn-0.14C is slightly slower than that of TWIP3. This can be attributed to the higher Mn content of TWIP3, Mn retarding SRX [3,9]. Another hot deformation testing method, stress relaxation, was also employed to measure the SRX kinetics of TWIP3. In the stress relaxation test, the displacement of the anvils is held constant after deformation and the drop in the stress recorded as a function of time. The SRX fraction from a stress relaxation test at 1000°C is shown in fig. 2.2.1.9 with the Avrami-type fitting. A good agreement between the double-hit compression data and the data from the stress relaxation test are observed. Both methods displayed the same SRX kinetics with the Avrami exponent of 1.2 and t50 = 11 s. TWIP4 (Fe-16Mn-1.5Al-0.3C) Flow stress curves The flow stress curves of TWIP4 (Fe-16Mn-1.5Al-0.3C) at three temperatures and at the constant strain rates of 0.1 s-1 are shown in fig. 2.2.1.10.

33

Fig. 2.2.1.10: Typical true stress-strain curves of TWIP4 (Fe-16Mn-1.5Al-0.3C) at the constant strain rate of 0.1 s-1. Flow stress curves for 25Mn1Al [9] steel are included for comparison. It can be seen that the flow stress curves of TWIP4 display distinct peaks at 1100°C and 1000°C, but it is difficult to discern any peak at 900°C. The peak strains are quite equal. For comparison, the flow stress curves for the 25Mn1Al steel are included in the same figure [9]. The hot deformation resistance of TWIP4 is lower than that of the 25Mn1Al steel. This can be attributed to the lower Mn content in TWIP4. From literature the strengthening effect of Mn is about 2 MPa/wt% and that of Al about 12 MPa/wt%. Hence, the lower Mn content of 9% in TWIP4 would mean the stress of 18 MPa lower and 0.5% higher Al, in turn, about 6 MPa higher stress level. As a result, the flow stress could be 12 MPa lower for TWIP4 compared to that of 25Mn1Al. In fig. 2.2.1.10, the difference seems to be about 20 MPa, i.e. somewhat more than predicted. Recrystallization kinetics of TWIP4 The softening kinetics of TWIP4 was investigated at three temperatures 900°C, 1000°C and 1100°C after compression to the 0.2 strain at the constant strain rate of 0.1 s-1. The fractional SRX curves as a function of holding time are plotted in fig. 2.2.1.11. Also, the data was fitted well with the sigmoidal curves, corresponding to the Avrami-type relationship [8]. It can be observed that SRX of TWIP4 is fast at 1000°C and 1100°C with t50 = 8 s and 3 s, respectively. However, at 900°C SRX kinetics is slower with t50 = 55 s. 25Mn1Al has identical recrystallization kinetics, with t50 = 8 s at 1000°C and the Avrami exponent of 1.

34

Fig. 2.2.1.11: SRX rates of TWIP-4 at the constant strain rate of 0.1 s-1. Double-hit compression data and fitted curves. Data for 25Mn1Al at 1000°C from [9]. TWIP-5 (Fe-22Mn-0.2C-0.2N) Fig.2.2.1.12 shows the flow stress curves of TWIP-5 at high temperatures (900-1100°C) at the constant strain rate of 0.1 s-1, compared with the flow stress of TWIP-2 (Fe-22Mn-0.6C). Both steels show the same hot deformation behaviour, displaying the peak at the identical strains 0.27 and 0.35 at 1100°C and 1000°C, respectively. However, TWIP-5 has lower deformation resistance (about 10 MPa) at highest temperatures, but at 900°C, TWIP-5 shows a higher flow stress.

Fig. 2.2.1.12: True stress-strain curves of TWIP-5 (Fe-22Mn-0.2C-0.2N) steel at high temperatures and at the constant strain rate of 0.1 s-1. Curve for TWIP-2 (Fe-22Mn-0.6C) is included for comparison. The results from the double-hit compression tests for TWIP-5 in fig. 2.2.1.13 show that TWIP-5 has the slowest recrystallization kinetics among the tested TWIP steels, because t50 at 1100°C (7 s) 1000°C (22

35

s) and 900°C (110 s) are about two times longer than those of the other steels (about 3-4, 8-11 and 55-75 s, respectively). Hence, N alloying seems to retard the softening kinetics. Also the Avrami exponents are smaller than typical of the other steels.

Fig. 2.2.1.13: SRX of TWIP5 at the constant strain rate of 0.1 s-1. Double-hit compression data and fitted curves. Effect of C on high temperature deformation behavior Typical true stress-true strain curves of TWIP steel, as compressed at 1100, 1000 and 900 °C at the strain rate of 0.1 s-1 have been presented above. The flow stress curves of TWIP steels with different alloying contents and a low-carbon steel (0.10C-0.45Mn) have been compared. Complexively the steels considered are: TWIP steels of Metaldesign project Fe-22Mn-0.6C Fe-18Mn-0.6C-0.02Nb Fe-16Mn-1.5Al-0.3C Fe-22Mn-0.2C-0.2N TWIP steels from literature Fe-25Mn-0.16C-1Al Fe-26Mn-0.14C

Low C steel from literature (0.10C-0.45Mn) As can be realized, the C content ranges between 0.10% up to 0.6% and the Mn content range is between 0.45% and 26%.

36

Qualitatively the value of peak stress at different temperatures can be considered as an indicator of the hot deformation resistance behavior of the steel considered. In the following table, only for example, the peak stress at 1000°C of TWIP steels are listed together with the other literature data. On the basis of these scarce data, a multiple correlation analysis between alloying contents (C, Mn, Al and N) and peak stress at different temperatures (at 900, 1000 and 1100°C) was carried out. Some conclusions can be drawn and summarized in these terms:

• The high-Mn TWIP steels have a much higher hot deformation resistance than that of low-C steel. This difference seems to be mainly related to the higher Mn content in TWIP steels.

• Considering only TWIP steels, the effect of alloying element contents (C, Mn, Al and N) on the peak stress is not so distinct and this is confirmed by quite poor correlation factors found (<0.7).

Steel Peak stress (MPa) at 1000°C Fe-22Mn-0.6C

148

Fe-18Mn-0.6C-0.02Nb

145

Fe-16Mn-1.5Al-0.3C

148

Fe-22Mn-0.2C-0.2N

140

Fe-25Mn-0.16C-1Al

162

Fe-26Mn-0.14C

148

C-Mn steel (0.10C-0.45Mn)

90

Table 2.2.1.1: Peak stress at 1000°C of different TWIP steels. Effect of C on recrystallization behaviour On the basis of the results of double-hit compression tests carried out on the TWIP steel variants, a qualitative study of the effect of alloying elements (C, Mn) was evaluated by means of multiple correlation analysis between C and Mn content with t50 values (time for 50% of recrystallization) at different temperatures. In the following table, for example, the values of t50 at 1000°C and 1100°C relevant to TWIP steels variants are reported.

Steel t50 at 1000°C (s) t50 at 1100°C (s) Fe-22Mn-0.6C

10 4

Fe-18Mn-0.6C-0.02Nb

11 4

Fe-16Mn-1.5Al-0.3C

8 3

Fe-22Mn-0.2C-0.2N

22 7

Table 2.2.1.2: Values of t50 at 1000°C and 1100°C relevant to TWIP steels variants. Some conclusions can be inferred and summarized in these terms:

37

• No clear trend arises from the analysis of the effect of C content on rex kinetics. This is confirmed by quite a poor correlation coefficient. This is consistent with C-steels, where, according to literature, C has no influence on static recrystallization kinetics.

• TWIP-5 revealed the slowest recrystallization kinetics among the tested TWIP steels. N alloying seems to retard the softening kinetics.

Static recrystallization kinetics The results shown above indicated that all the investigated high-Mn TWIP steels have almost equal recrystallization kinetics that is identical to that of 25Mn1Al. Only the TWIP5 has a distinctly slower softening behaviour. The steel 25Mn1Al has been used to investigate more extensively the power of various variables on the SRX kinetics by applying double-hit compression technique. These results have been published in [9] but briefly described here to complete the characterization of SRX behaviour of high-Mn TWIP steels. Temperature and alloying SRX fractional softening curves after the deformation to the 0.2 strain (i.e. smaller than the peak strain) at 0.1 s-1 at five temperatures are plotted in fig. 2.2.1.14. As shown, sigmoidal Avrami-type curves can be fitted with the experimental data. It is well established that the time for the 50% recrystallized fraction t50 can be described by the following empirical relation [16]:

t50 = A εp ε•

q ds exp(Qapp/RT) (6) where A is a constant, ε the strain, ε

•

strain rate, d the grain size, Qapp the apparent activation energy of recrystallisation, R the universal gas constant and T the absolute temperature. Material constants p, q and s are the strain, strain rate and grain size exponents, respectively.

Fig. 2.2.1.14: SRX rates (strain 0.2, strain rate 0.1 s-1) for the 25Mn1Al steel. Double-hit compression data and fitted curves.

38

The recrystallization kinetics of low-C (0.09C-0.45Mn) and Nb-microalloyed (C-Mn-0.03Nb) steels as well as one austenitic stainless steel (Type 304) has been investigated previously, e.g. [11,12,17]. Some new tests were performed for low-C and austenitic stainless steels, and these data at 1050°C are plotted in fig. 2.2.1.15 for comparison. For the Nb-steel, SRX data obtained by the double-hit and stress relaxation methods at 1050°C/0.1s-1 was taken from a paper of Perttula and Karjalainen [2], suggesting the t50 time of 7 s. However, the grain size of this steel was 60 μm. The grain size difference can be accounted by using the power of grain size s = 2.13*d-0.105 [18,19]. The predicted data and the curve are plotted in fig. 2.2.1.15. It is seen that the 25Mn1Al steel exhibits faster SRX kinetics than the austenitic stainless or Nb-bearing steels but slower than that of low-C steel. The t50 times are 1, 4, 16 and 15 seconds for low-C, 25Mn1Al, Nb and Type 304 steels, respectively. For high Mn-Al steels, the Avrami exponent seems to decrease with decreasing temperature. For the Avrami curves the exponents are 1.2, 1.1, 0.8, 0.7 and 0.7 at 1100, 1050, 1000, 950 and 900°C, respectively. These values are quite low compared with the commonly reported values for C-Mn steels (between 1-2) [13]. It is known that recovery tends to decrease the value. Hence, it can be concluded that there might still be some contribution of recovery, even though the 5% total strain method was adopted. It was realized that in high-Mn steels, dynamic recovery is very intense; particularly in steels with high Al levels that have higher stacking fault energy [5].

Fig. 2.2.1.15: Comparison of SRX kinetics of 25Mn1Al to low-carbon, Nb and Type 304 steels. Data for the Nb-steel from [2]. Effect of strain and strain rate on SRX The effects of strain and strain rate on t50 times for the 25Mn1Al steel are shown in figs. 2.2.16 and 2.2.17. From the slopes of the lines, values of -2.7 and -0.3 are obtained for the strain and strain rate exponents, respectively. The strain exponent for C-Mn steels has been reported to be in the range -2.5 to -4 at the strain range of 0.1-0.2 [18,20-22]. Values between -2 and -3 have been measured by the stress relaxation technique for some microalloyed steels [23]. Somani et al. [18,19] used values of -2.8 for C/C-Mn/Nb/Ti/Nb-Ti steels and -2.5 for Mn-V/other steels in developing a regression model based on chemical composition for predicting SRX kinetics of a large variety of steels. Koskiniemi et al. [12] determined values of -3 and -1.5 for Type 304 and a 12Cr stainless steels, respectively. Barraclough and Sellars [24] found a strain exponent of -4 at small strains, but the value decreased as the strain

39

approached the critical strain for the onset of DRX. Most of these strain exponent values are of the same order, and hence, a strain exponent of -2.7 can be considered reasonable for the 25Mn1Al steel as well as for other TWIP steels, owing to the low effect of Al on the SRX characteristics and kinetics. As compared to strain, the effect of strain rate is relatively small at small strains, and the computed exponent from the experimental data for 25Mn1Al and some limited data for the 25Mn, 25Mn3Al and 25Mn6Al steels is about -0.3. In comparison, Koskiniemi et al. [12] obtained the strain rate exponents for Type 304 and 12Cr stainless steels as -0.30 and -0.33, respectively. Strain rate exponents equal to -0.38 are commonly reported for Type 304 and Type 316 steels in the literature, too (e.g. [25-27]). All these values are relatively high compared to the values obtained for C and C-Mn steels (-0.11), medium carbon steels (-0.13) and also Ti-steels (-0.12) [28]. In between the two, the strain rate exponents for Nb and Nb-Ti and Mo-steels are measured to be about -0.23 [13,19], which value has also been used in the development of regression models for kinetics and activation energy of SRX, essentially based on chemical composition [18,19]. All these values indicate the dependence of strain rate effect on the extent and type of alloying. Thus, the exponent of -0.3, experimentally measured for the 25Mn1Al steel, can be considered reasonable for the all TWIP steels, being closely equal to that of Type 304 stainless steel, as mentioned earlier [12].

Fig. 2.2.1.16: Plot of t50 vs strain to calculate Fig. 2.2.1.17: Plot of t50 vs strain to calculate the power of strain for 25Mn1Al. the power of strain rate.

Activation energy of SRX The temperature dependence of SRX kinetics for the 25Mn1Al steel is shown in fig. 2.2.18, where t50 times are plotted against the inverse absolute temperature. The apparent activation energy of SRX (excluding data above 1075°C) deduced from the data is about 257 kJ/mol, which is lower than that reported for Type 304 (285 kJ/mol) and close to that of the 12Cr stainless steel (265 kJ/mol) [12].

40

Fig. 2.2.1.18: Time t50 of SRX for 25Mn1Al vs the inverse absolute temperature. The activation energy of hot deformation Qdef for the 25Mn and 25Mn3Al steels has earlier been determined to be 380-405 kJ/mol [3], so that Qdef of 385 kJ/mol can be assumed for 25Mn1Al. Then, Qrex for this steel can be computed to be about 373 kJ/mol. In comparison, Koskiniemi et al. [12] reported equal Qrex values of 400 kJ/mol for both Type 304 and 12Cr stainless steels. In the literature, largely varying values for Qrex have been reported for Type 304 steel, e.g. in the lists of Barbosa [29] (360-460 kJ/mol) and Ryan and McQueen [27,30] (334-550 kJ/mol). Anyhow, all these values are of the same order, so that Qrex of 373 kJ/mol computed for the 25Mn1Al steel can be considered to be reasonable also for the other TWIP steels, even though a small effect of Al on Qdef cannot be ignored. This level of Qrex is comparable to those obtained for Ti-steels (255-275 kJ/mol) and Nb-steels (345-400 kJ/mol) [18,19] but it is substantially higher than the ordinary C/C-Mn and medium-C steels (200-240 kJ/mol). With the powers of strain (-2.7) and strain rate (-0.3) and the apparent activation energy of SRX (257 kJ/mol), the following fractional softening equation for SRX has been determined: t50 = 3.2*10-15 ε-2.7 ε'-0.3 ds exp(257000/RT) (7) where the power of grain size, taken from [18,19], is assumed to be s = 2.13 d-0.105 (8) Grain size evolution in SRX Several equations given in the literature for predicting the statically recrystallized grain size in carbon and austenitic steels were tried to apply. It was found that some common equations developed for C-Mn steels give quite reasonable predictions, particularly those in refs. [21,32]. One of these equations is as follows: drex =0.35 ε -1 67.0

0d (9) where do is the austenitic grain size before hot deformation.

41

Conclusions of Task 2.2.1 The main results and conclusions from the work carried out regarding flow stress and SRX behaviour are as follows: Hot deformation resistance of high-Mn TWIP steels is dependent on Mn content (strengthening about 2 MPa/wt%) and increases with increasing Al alloying up to 6% (strengthening about 12 MPa/wt%). Nb and N in the contents used here in steels have a minor influence. The austenitic high-Mn TWIP steels exhibit higher deformation resistance than those of low-C, C-Mn-Nb and austenitic stainless steels.Flow stress curves exhibit broad stress peaks at quite low strains. However, the completion of DRX occurs slowly. Very fine grain size is obtained as a result of DRX. SRX kinetics of TWIP steels is faster than that of Type 304 and C-Mn-Nb steels and slower than that of low-C steels. Mn is the main element retarding the rate of SRX and Al has only a minor contribution. The regression equation for the static recrystallization kinetics for TWIP steels can be used to predict the SRX rate under given conditions. Grain size will be refined effectively by SRX, and certain equations published for C-Mn steels can be used to predict the grain size. Task 2.2.2 Recrystallization kinetics under different cold rolling schedule Introduction The investigation on recrystallization behavior of TWIP steels has been focused on a thorough characterization of the isothermal annealing behaviour of cold rolled TWIP2 steel (22% Mn, 0.6% C, hot rolled band, 40% to 70% cold rolling reductions) in the temperature range 450-1100 ºC . The assessment of the annealing effects has been made by following the Vickers macro-hardness, structure (recrystallized fraction, grain size, proportion of twin boundaries) and texture evolution (microtexture). Material and annealing treatments Hot rolled TWIP2 samples (5.4 mm thick) were reduced by cold rolling at a CSM laboratory mill to different thicknesses: 3.24 mm (R = 40 %), 2.7 mm (R = 50 %), 2.16 mm (R = 60%) and 1.6 mm (R = 70%). An OIM of the hot rolled structure of the starting material is shown in fig. 2.2.2.1. A non-equilibrium grain structure typical of high temperature deformation with some residual deformed intra-grain structure (subgrains) is evidenced by this OIM image. The grain size (mean linear intercept) is 23.3 μm (an upper bound, as twin boundaries were not counted). The structure of the cold rolled samples is too distorted for being completely resolved by OIM. The micro-texture of the resolvable regions is a typical FCC rolling texture. Practically it contains only the Brass orientation, {110}<112>, plus very weak Goss {110}<001> and S {123}<634> components. The cold rolled samples subjected to the annealing treatments summarized in Table 2.2.2.1. The 450 ºC ≤ T ≤ 900 ºC treatments were performed in a salt bath furnace and were interrupted by water quenching after different annealing times. The heating time was controlled by treating dummy samples with inserted thermocouples. Heating time was always a small fraction of the time spent at each treating temperature but for the nominal t ≤ 12 s ones.

42

Fig. 2.2.2.1: OIM image of hot rolled TWIP2 sample (inverse pole figure and grain maps, rolling plane normal. Treatments at 1000 ºC and 1100 ºC were performed in a resistance furnace under Ar protective atmosphere and were followed by water quenching. The fraction recrystallized for each annealing time, temperature and prior cold rolling reduction has been deduced from hardness measurements complemented with metallographic and micro-texture observations. T (ºC) Time of annealing (minutes) at target T before water quench

450 0.02 (1 s) 0.2 (12 s) 0.5 (30 s) 1 3 9 27 81 243 729 600 0.02 (1 s) 0.2 (12 s) 0.5 (30 s) 1 3 9 27 81 243 729 700 0.02 (1 s) 0.2 (12 s) 0.5 (30 s) 1 3 9 27 81 243 729 800 0.02 (1 s) 0.2 (12 s) 0.5 (30 s) 1 3 9 27 81 243 850 0.02 (1 s) 0.2 (12 s) 0.5 (30 s) 1 3 9 27 81

In a salt bath furnace

900 0.2 (12 s) 0.5 (30 s) 1 3 9 27 1000 3 9 27 In a mufla

furnace (Only for R = 60 %) 1100 3 9 27

Table 2.2.2.1. Annealing treatments Recristallization kinetics The evolution of static recrystallization fraction at two temperatures, 600 ºC, 650°C and 700 ºC for 60% cold rolling reduction is plotted in fig. 2.2.2.2 in order to check their agreement with an Avrami behaviour. For 600 ºC, recrystallization has not ended for the maximum time tested and the fraction recrystallized being known from the metallographic observations, ( )[ ] %7.8615.12 600 ≅Crex ohX . In fig. 2.2.2.3 is shown the recrystallization kinetics at 700°C for different cold rolling reductions. As can be noted at 700 ºC, recrystallization ends after 3 minutes, data for longer times are for grain growth. Final grain size is very small: many nucleation sites are available because of the tremendous number of twin intersections after 60% cold rolling deformation. Same observations have been made for recrystallized grain sizes of austenitic stainless steel (metastable austenite) and TWIP steels [16].

43

0

0.2

0.4

0.6

0.8

1

1.2

0.00001 0.01 10 10000 1000000

t (s)

X sof

60 % REDUCTION ■ 700 ºC● 650 ºC▲ 600 ºC

0

0.2

0.4

0.6

0.8

1

1.2

0.001 0.1 10 1000 100000

t (s)

X sof

t

■ 40%♦ 50%▲ 60%● 70%

700 ºC

Figure 2.2.2.2: Recrystallization kinetics for 60% reduction derived from the softening by annealing temperature of 600ºC, 650ºC and 700ºC as function of time. TWIP 2 steel with composition 22% Mn- 0.6% C

Figure 2.2.2.3: Recrystallization kinetics for all reductions derived from the softening by annealing temperature of 700ºC as function of time. TWIP 2 steel with composition 22% Mn- 0.6% C

Table 2.2.2.2 shows the individual values of ksoft and B (constants of the Avrami equation). The values of the exponent are of the typical order of magnitude reported for other alloys. No intent of a refined characterization of the kinetic equation of recrystallization has been pursued because the recrystallised grain size is so small (less than 2 μ in all cases) and the time for complete recrystallization so small, above 650°C, that a finer analysis does not seem relevant for practical purposes.

Table 2.2.2.2. Values of ksoft and B (constants of the Avrami equation) obtained from Avrami fittings to the recrystallization results. With all the observations carried out, it is possible to draw an annealing map (table 2.2.2.3) in the explored T-t sub-space that summarizes the different processes at work. A small region of static ageing (hardening) is observed at low annealing temperatures and times. Other authors have observed dynamic strain ageing from room temperature to 300°C [11, 12]. No structural changes are detected at 450°C; a weak softening (weak recovery) takes place for long annealing times beyond the ageing region.

44

11001000900850800700650600450

T [ºC] vs. t [s] 1 12 30 60 180 540 1620 4860 14580 43740

Recrystallization. Negligible grain growthStatic ageing by solid solution segregation to dislocationsRecovery. No recrystallization

Grain growth. Recrystallization is completed

Table 2.2.2.3. Annealing processes map of TWIP steel in T-t space. Not explored regions in blank. Grain growth kinetics According to the measurements of grain sizes performed by EBSD-OIM, grain growth rate at 700 °C is very low. Recrystallized grain size is 1.5 μm and recrystallization is reached after nine minutes of annealing but grain size is only 2.27 μm after twelve hours. It is still rather low, 4.2 μm, after 27 minutes at 900 °C. Only above 1000 °C grain growth rate increases enough for reaching grain sizes larger than 20 μm in 27 minutes of annealing. From isochronal plots, it is possible to estimate the activation energy for the grain growth processes QGG. Combining the data from annealing at 900 °C, 1000 °C, 1100 °C, and assuming as activation energy the value QGG = 363 ± 60 kJ/mol calculated from the isochronal plots, a good fitting of all the grain growth data is obtained with an exponent nGG ≈ 3.5, fig. 2.2.2.4.

y = 3.5407x - 38.315R2 = 0.8867

-40

-35

-30

-25

-20

-15

0 1 2 3 4 5Ln (D*)

ln[e

xp(-Q

GG/R

T)t]

QGG = 363 ± 60 kJ/mol

Fig. 2.2.2.4: Grain growth equation. D* is calculated from mean linear intercept method, twin boundaries not counted as grain boundaries. TWIP2 steel with composition 22% Mn- 0.6% C.

Texture changes The texture after cold rolling is the typical brass texture {011}<211>. The recrystallization texture (without perceptible grain growth, grain size D < 2 μm) is rather similar to the rolling one, with a strong brass component, plus a weaker {011}<100> component. After grain growth these texture components remain but the texture intensity is even weaker, with intensity maxima smaller than 1.5. Fig. 2.2.2.5 shows the pole figures corresponding to the 60% cold rolled sample.

45

Fig. 2.2.2.5 Macro-textures, pole figures: a) TWIP steel 60% cold rolled with the typical brass-type texture {011}<211>; b) Recrystallization texture, D ≤ 2 μm; c) Grain growth texture with D ≈ 9 μm. Conclusions task 2.2.2 • Recrystallized grain sizes of TWIP steel of 22% Mn - 0.6% C (in mass-%) cold rolled in the range

of 40%-70% reductions and isothermally annealed in the temperature range 600 °C ≤T ≤ 900 °C are very small, D ≤ 2 μm.

• At 450°C the effect of annealing is very weak. After some seconds there is some hardening (static ageing by solid solution segregation to dislocations) and a very weak softening thereafter, attributable to recovery, without any noticeable metallographic changes.

• Above 700 °C, recrystallization is complete in less than nine minutes and takes less than ten seconds above 800 °C. At 900 °C ≤ T ≤ 1100 °C, the kinetics observed only corresponds to grain growth.

• Recrystallization and grain growth textures are very weak. Consequently, the elastic and plastic anisotropies of annealed TWIP sheets will be negligible.

• Although there is no apparent texture change from recrystallization to grain growth, there is a strong change in the grain boundary composition. The fraction of Σ3 twin boundaries increases suddenly from 14% to 40% when grain growth starts and it remains constant thereafter independently of the grain size reached.

• An empirical equation of grain growth for this steel has been obtained with an apparent activation energy QGG = 363 ± 60 kJ/mol and an exponent nGG ≈ 3.9.

Task 2.2.3 Modelling of recrystallization behaviour of TWIP steels by modifying the available mathematical models for austenitic steels and the relevant constitutive equations Introduction In modelling static recrystallization of cold deformed austenite three different approaches due to Zurob and co-workers has been used [1-3]. Assuming that the grain boundaries interact with Mn atoms in this

46

TWIP steel, the velocity of the grain boundaries vg during recrystallization, can be related to the driving force F, using Cahns theory of solute drag, quoted in [3]. In the following section, models for the fraction recrystallized evolution are described which all consider the interaction of manganese solute atoms with grain boundaries. The three models have been initially tuned using the experimental fraction softened curve of previous section. Next the best model is then applied to a range of experimental fraction softened curves for various cold rolling reductions and the results discussed. The models considered were:

1. Model n.1: Constant grain boundary velocity with site saturated nucleation. The equation for the fraction recrystallized Xv as a function of time t is the following

( )33exp1 tfNvX gv −−= , where f is a shape factor (= 4π/3 for spheres), vg grain boundary velocity, N is the recrystallized nuclei density.

2. Model n.2: Constant grain boundary velocity with a decreasing nucleation rate. The equation for the fraction recrystallized Xv as a function of time t is the following

( )3exp1 gv dgfvX −−= ,

Where g is a constant of the nucleation rate equation Bt

gNr +=

1 ,

f is a shape factor and d is given by:

( )( ) ( )46

1ln6236B

BtBtBtBtd +−+−+=

3. Model n.3: Decreasing boundary velocity with site saturated nucleation.

The static recovery model used, is that due to Verdier et al [12] later expanded by Zurob. In the model the stress due to dislocations σd is related to time through [2]:

tTkGbMk

RTQ

GM B

d

tdd⎟⎟⎠

⎞⎜⎜⎝

⎛⎟⎠⎞

⎜⎝⎛−+=

=

310

230,

sinhexp3

411 αα

υσσ

,

where σd,t=0 is the initial stress, υd is the Debye frequency (= 1 X 1013 s-1), Q0 is the activation energy for bulk self diffusion in austenite (= 270 kJ/mol [13] ), kB is Boltzmanns constant and k1 is a constant less than unity. Finally from [7] the fraction recrystallized is given by:

⎟⎟

⎠

⎞

⎜⎜

⎝

⎛⎟⎟⎠

⎞⎜⎜⎝

⎛−−= ∫

3

0

exp1t

gv dtvfNX

Comparison of models with experimental results As can be seen from figures 2.2.3.1-2-3, the model n.3, i.e. that in which the boundary velocity decreases with time due to recovery, gives the best agreement with experimental data. Thus the modelling of experimental data only model 3 will be used.

47

Model n.3 is applied to experimental data for different cold reductions with the annealing temperature kept constant at 600°C. The results are shown in fig. 2.2.3.4.

00,10,20,30,40,50,60,70,80,9

1

1 10 100 1000 10000 100000

time (s)

Frac

tion

rexe

d

00,10,20,30,40,50,60,70,80,9

1

1 10 100 1000 10000 100000

time (s)

Frac

tion

rexe

d

Figure 2.2.3.1: Comparison of model 1 (black line) with experimental data for a cold reduction of 60% with an annealing temperature of 600°C.

Figure 2.2.3.2: Comparison of model 2 (red line) with (corrected) experimental data for a cold reduction of 60% with an annealing temperature of 600°C.

00,10,20,30,40,50,60,70,80,9

1

1 10 100 1000 10000 100000

time (s)

Frac

tion

rexe

d

Figure 2.2.3.3: Comparison of model 3 (blue line) with (corrected) experimental data for a cold reduction of 60% with an annealing temperature of 600°C.

As can be seen from fig. 2.2.3.4 generally good agreement between model and experiment is obtained for different cold reductions with N as the only free model parameter. The agreement for the cold reduction of 40% is less good than for the other reductions however. Comparison with the average N (N as 7.23 X 1017 m-3) shows good agreement for reductions of 60% and 70% but not for the intermediate reductions. Applying the model n.3 with an average N (with A as 0.657 and k1 as 0.0685) still gives acceptable overall agreement with the experimental data. The fact that the model gives acceptable agreement for a single value of N seems to suggest that the recrystallized grain size does not change significantly in the range of reductions investigated.

48

00,10,20,30,40,50,60,70,80,9

1

1 10 100 1000 10000 100000

time (s)

Frac

tion

rexe

d

00,10,20,30,40,50,60,70,80,9

1

1 10 100 1000 10000 100000

time (s)

Frac

tion

rexe

d

A) B)

00,10,20,30,40,50,60,70,80,9

1

1 10 100 1000 10000 100000

time (s)

Frac

tion

rexe

d

00,10,20,30,40,50,60,70,80,9

1

1 10 100 1000 10000 100000

time (s)

Frac

tion

rexe

d

C) D) Fig. 2.2.3.4: Comparison of model 3 with a range of experimental data, for a constant annealing temperature of 600°C. A) 40% reduction B) 50% reduction C) 60% reduction and D) 70% reduction. Conclusions task 2.2.3 Physically based recrystallization models have been applied to the experimental data for cold rolled and annealed TWIP steel. Several models were considered with different assumptions for the nucleation behaviour and boundary velocity behaviour. All models assumed that grain boundary mobility was controlled by interaction with manganese solute atoms. The possible influence of twins on the recrystallization process was not considered. From the modelling results the following conclusions can be drawn:

• The model which best described the experimental data was the model n.3 in which nucleation was site saturated and the grain boundary velocity decreased with time due to static recovery.

• Comparison of this model with more experimental data revealed good agreement when the nucleation density was a free model parameter.

• Imposing an average nucleation density and recalculating recrystallization curves still gave reasonable agreement with experimental results.

Task 2.3 Study of the precipitation at equilibrium by means isothermal treatments. Precipitates analysis by means of electron microscopy techniques (SEM-EDX, extraction replica for TEM). The aim of this task is to study the precipitation behaviour of TWIP steel at different temperatures (from 500°C to 1000°C) at long soaking time (approaching equilibrium conditions). The results will be used to design the hot rolling processing conditions (coiling, annealing). The hot rolled strip samples used for this activity are in the solution treated condition (1200°C x 1h). In table 2.3.1 are reported the annealing treatments carried out for all the TWIP variants.

49

Soaking Temperature

(°C)

Soaking time (h)

500 10h 600 10h 700 10h 800 10h 900 1h

Table 2.3.1: thermal treatments for precipitation study.

The isothermal section of Fe-Mn-C [1] system in the range of 800°C - 900°C does not reveal the precipitation of carbides. At temperature below 700°C the γ-phase became more restricted and the reduction of C solubility coincides with the occurrence of a complex series of equilibria between γ-phase and Fe-Mn carbides. At temperature below 700°C and for Mn content in the range 16-22%wt and C = 0.4-0.6%wt the most stable carbide is M3C cementite type (with C content about 7%wt). The isothermal section of Fe-Mn-C and Fe-Mn-Al systems [1] highlights that at temperature lower than 600°C, and in correspondence of segregated zone with lower Mn content (dendrite centre), the austenite can be thermodynamically unstable and the formation of pearlitic phases containing cementite and ferrite can occur [2]. This phenomena should occur on very long time (>10h) and is more pronounced during massive carbide precipitation due to local lowering of the C in solid solution. Precipitation analysis at scanning electron microscope A first screening of the precipitation behaviour of the TWIP variants was carried out by means of LOM and SEM aiming to evaluate the occurrence of carbide precipitation at temperature below 900°C and the differences in relation with the TWIP steel chemical composition. The aim of this qualitative analysis is first of all to define the range in which the carbide precipitation has a massive aspect. The occurrence of so huge precipitation must be absolutely avoided, in any step of the production route, because the tensile properties of TWIP steel could be significantly affected. The results of the SEM-EDS analysis carried out are summarised as follows. Precipitation after 10h of soaking at T=800-900°C In fig. 2.3.1 are shown the typical microstructures found for each TWIP variant after soaking at 800°C for 10h. As can be noted no precipitation was observed by OM and SEM analysis. Only oxide particles and MnS were detected. Tha same situation was found in the samples heated at 900°C for 1h.

50

TWIP 1 TWIP2

TWIP3 TWIP4

TWIP5

Fig.2.3.1: Microstructures of the TWIP variants after soaking at 800°C for 10h. Temperature range 500-700°C Carbides precipitation is detected in the range 500-700°C but with remarkable differences between the TWIP variants. In the following sections are summarised the results for each variant. TWIP 1 steel TWIP 1 grade is characterised by the lowest C content (0.01wt%) and by a significant second phase ferrite content (about 6%). This characteristic suggests a lower austenitic stability with respect the other TWIP variants. Very long permanence (10h) at 700°C resulted in additional ferrite formation. The ferrite particles reveal a higher Al and Si content and lower Mn content with respect to austenitic matrix (fig. 2.3.2). The formation of ferrite was also confirmed by ferritoscope measurement.

51

Fig. 2.3.2: TWIP1 - 700°Cx10h.

TWIP2 - TWIP3 steels TWIP2 and TWIP3 are characterised by a Mn content in the range 18-22wt% and C content 0.6wt%. Carbides precipitation in these steel variants is quite abundant in the temperature range 500-600°C. In the fig. 2.3.3 is shown a SEM image of carbides precipitated on the austenitic grain boundaries of TWIP2 after a soaking at 700°C for 10h. The semi-quantitative analysis gave a C content of about 7%wt consistent with carbide of cementite type. In fig.2.3.4 is shown a SEM image of carbides precipitated on the austenitic grain boundaries of TWIP3 after a soaking at 700°C for 10h. In fig.2.3.5-6 is shown carbide precipitated on austenitic grain boundaries at 600°Cx10h on both TWIP2 and TWIP3.

Fig. 2.3.3: SEM image of the TWIP2 sample after soaking at 700°C for 10h.

Fig. 2.3.4: SEM image of the TWIP3 sample after soaking at 700°C for 10h.

52

Fig. 2.3.5: SEM image of the TWIP2 sample after soaking at 600°C for 10h.

Fig. 2.3.6: TWIP3: SEM + EDS image of the cementite precipitated on the austenitic grain boundaries after a soaking at 600°C for 10h.

TWIP4 TWIP4 grade is characterised by a Mn, Al and C content respectively of 16%, 1.5% and 0.4%. In the temperature range of 500-700°C TWIP4 grade revealed a formation of ferrite and precipitation of carbide. In fig 2.3.7 is shown the microstructure of TWIP4 after 10h soaking at 700°C in which can be noted the ferrite formed. The carbides were not resolved probably due to fine dimensions but the presence of a higher carbon content with respect to matrix suggests that the phase formed are ferrite+carbides. In the fig. 2.3.8-9 is shown the microstructure of TWIP4 after 10h soaking at 500-600°C in which can be noted the ferrite formed.

Fig. 2.3.7: TWIP4 at 700°C: within some grains is detected the precipitation of ferrite + carbides.

53

   matrix  particle  Wt %  Wt % 

C -K 0.5  2.9 Al-K 1.4  1.7 Si-K 0.3  0.8 Mn-K 18.2  15.4 Fe-K 77.2  78.4 

Fig. 2.3.8: SEM + EDS image of the TWIP 4 samples after a soaking at 600°C for 10h.

Fig. 2.3.9: SEM + EDS image of the TWIP 4 samples after a soaking at 500°C for 10h. TWIP 5 steel TWIP5 is characterised by a Mn content of 22wt% and a C content of 0.2wt% and Nitrogen of 0.2wt%. The precipitation behaviour of TWIP5 is quite different with respect to previous variants. SEM analysis carried out on the sample annealed at 700°C and 600°C for 10h did not reveal any presence of carbides even at quite high magnification. A fine precipitation of carbide particles (fig.2.3.10) was detected after soaking at 500°C for 10h. Also in this case the small carbides were recognized as cementite because the (Fe, Mn)/C ratio is consistent with cementite carbides (C content about 7%wt).

54

Fig. 2.3.10: TWIP 5 specimen after soaking for 10h at 500°C. Extraction replica analysis by transmission electron microscopy Single stage carbon extraction replicas were prepared using the conventional method starting from a surfaces prepared as for optical microscopy and using a selective electrolytic dissolution method to remove the matrix around the carbides. The extracted carbon replicas were examined using a Jeol JEM 200CX (STEM) electron microscope operated at 200kV. The replica extraction was carried out on TWIP samples soaked for 10h at 600°C and 700°C. The results of the replica analysis are detailed as follows. TWIP 1 TEM extraction replica on TWIP1 steel confirmed the presence of ferrite+carbides. In fig.2.3.11 can be noted rows of particles extracted from larger island in which can be recognized particles of ferrite with a typical higher content of Al and Si and lower content of Mn with respect to matrix (Mn is about 22%). Carbides detected are of cementite type (FeMn)3C (fig. 2.3.12) and are characterized by finer size with respect to ferrite particles. The STEM-EDS analysis of ferrite and carbide particles are reported in fig.2.3.11.

Fig.2.3.11: TWIP1-700°C: Low magnification image.

55

Fig.2.3.12: TWIP1-700°C: Carbide of cementite type (FeMn)3C, the analysis is reported in fig. 2.3.11.

TWIP2 The precipitation features of TWIP2 at 600°C and 700°C, in terms of type and chemical composition, are the same. The remarkable difference is in terms of amount of precipitates: at 700°C the number of precipitates is significantly lower than at 600°C. Two type of precipitates were found:

1. (Fe, Mn)3C – cementite 2. Vanadium carbonitrides characterized by fine size ( <50nm).

The presence of Vanadium carbonitrides is due to unwished presence of V in ferroalloy during ingot casting. In fig.2.3.13-14 are reported the picture of the two carbide types with the relevant diffraction pattern.

Fig. 2.3.13: (Fe,Mn)3C carbide image with relevant diffraction

In the fig. 2.3.14 is reported the image of the Vanadium carbo-nitride precipitate type.

56

Fig.2.3.14: Vanadium carbo-nitride image with relevant diffraction.

TWIP3 TWIP3 steel is characterized in terms of chemical composition by Nb addition. As expected the precipitation of carbo-nitride is strongly affected by this micro addition. In fig. 2.3.15 is shown the difference in terms of number of Nb carbo-nitrides between 600°C and 700°C (the total number of precipitates analysed is 50). As can be noted the amount of NbCN at 700°C is significantly higher and even the particle size are bigger. This means that at relatively low temperature (below 600°C) the precipitation of Nb carbo-nitrides is strongly reduced. In addition to NbCN, carbides of cementite type (Fe, Mn)3C were detected. The latter are predominant at 600°C as can be noted in the fig.2.3.16.

TWIP3-600°C TWIP3-700°C Fig.2.3.15: Nb content of the precipitates selected on extraction replica.

57

Fig.2.3.16: TWIP3 – 600°C: (Fe,Mn)3C carbide image.

TWIP5 STEM-EDS analysis was carried out on TWIP 5 sample soaked at 700°C for 10h. A quite low fraction of fine (Ti, Al) carbo-nitrides (size less than 150nm) was detected. The presence of Ti and Al is due to unwished residual presence in ferroalloy during ingot casting. In fig. 2.3.17 is shown the image of a typical (Ti, Al) carbo-nitride.

Fig. 2.3.17: TWIP5 at 700°Cx10h - (Ti, Al) carbo-nitride.

58

Conclusions of task 2.3 The precipitation behavior of TWIP steel depends on one side on the C, N, and elements carbide-nitride formers (Ti, Al, Nb, V) content, but on the other side is strongly affected by thermodynamical stability of austenite in the range 500-700°C; Infact the ternary system Fe-Mn-C under defined conditions of temperature and chemical composition shows that the austenite could partially transform in ferrite+carbides (pearlite like structure). The study carried out on TWIP steels by SEM-EDS and STEM-EDS allows to state the following conclusions:

• Both TWIP1 and TWIP4 revealed, at temperature below 700°C the tendency of austenite to destabilize and to forme ferrite+carbide structures. These observation is in agreement with other results achieved on TWIP1-4 and confirm that the austenite stability of these steel grades is lower with respect to reference grade TWIP2 (Fe-22Mn-0.6C);

• TWIP2 and TWIP3 show a similar behavior. Both steel grades revealed in the range 500-700°C a massive carbide precipitation of cementite type (Fe, Mn)3C. No ferrite phase was detected in the range 500-700°C.

• TWIP5 grade revealed a very stable structure: no massive carbide precipitation was detected in the range 500-700°C. The precipitated particle are constituted of cementite carbides and fine (Ti,Al) carbo-nitrides. No ferrite phase was detected even after long soaking time in the range 500-700°C.

Task 2.4 Industrial feasibility of hot and cold processing route of TWIP steels Introduction The design of the hot and cold strip processing route of TWIP steel consisted in the definition of the operating practice of the following steel manufacturing steps:

1. Hot rolling process • Slab reheating process: the goals of the study are the definition of the slab

discharging temperatures, evaluation of impact on usual C-steel production, and minimization of fuel consumption.

• Thermo-mechanical treatment:

-design of hot rolling schedule; -cooling pattern on ROT; -coiling temperature;

• Hot rolled strip annealing conditions (if necessary);

2. Pickling process conditions 4. Cold rolling process (plain strain compression tests – task 3.4)

5. Cold rolled strip annealing conditions.

The definition of the industrial operating practice have adopted, on one side, the results arising from the basic investigations and, on the other side, additional activities were carried out as for basic creep strength evaluation, hot and cold deformation strength evaluation, characterization of pickling behavior. In the following are detailed the results find out.

59

Slab reheating process Qualitative creep strength evaluation One important aspect in the design of the manufacturing cycle consists in the definition of the optimal slab reheating temperature. The need of high finishing rolling temperature (900°C) impose as a consequence to choose a high entry temperature for roughing mill (1180-1200°C) and entry finishing mill (1100-1150°C). This means that the TWIP slabs must be reheated at a temperature very close to usual carbon steel slab reheating temperature, that is 1200-1250 °C. From CSM laboratory results and from thermo-dynamical calculations TWIP steels are expected to be characterized by a lower Tsolidus (ranging 1320-1350°C) with respect to a conventional C-Mn steel (Tsol= 1490-1500°C) selected as a reference. This means that the feasibility of slab reheating in the range 1220-1250°C has to be carefully verified in order to avoid problems in the reheating furnace (walking beam) and in following hot rolling process. To evaluate qualitatively the creep resistance behavior at high temperature (1200°C) a test based on the sag test specification was carried out. This test consisted in measuring the bending occurred after a 3h soaking at 1200°C of TWIP steel specimens with same geometry (200mm length, 20mm width, 2.0 mm thickness). For reference the tests was carried out also for a conventional carbon steel (C,Mn steel grade). The tests was carried out adopting a laboratory resistance furnace with a controlled oxygen atmosphere (O2 about 5%). The bending measured after treatment is strongly dependent on the sample geometry and of course the thinner is the specimen the higher is the bending. In the following are presented the results achieved adopting specimens with 2mm thickness. The results in terms of max bending in the centre of specimen are reported in table 2.4.1. As can be noted the creep resistance of TWIP2 at 1200°C is significantly lower than the reference C, Mn steel samples.

Sample size

NIM length (mm) thickness[mm]Average bending

[mm] TWIP 2 200 2.00 22 TWIP 3 200 2.00 15 TWIP4 200 2.00 16

C, Mn steel 200 2.00 12 Table 2.4.1: Bending after a sag test with 3h soaking at 1200°C.

TWIP3 as well as TWIP4 steel variants show an intermediate creep strength with respect to conventional C, Mn steel and TWIP2. The difference between TWIP2 and TWIP3/TWIP4 could be ascribable to the different Tsolidus as reported in the table 5.1.2. The results of this qualitative tests indicate that the TWIP slabs could show, during the reheating process in walking beam, a slightly larger bending at high temperature and this suggests, for safety of the plants, to adopt slabs with maximum feasible thickness (250mm) and reduced slab length (5m instead of usual 10m) in order to control the maximum bending of slabs in furnace. Due to their austenitic structure the expected thermal conductivity of TWIP steels, below about 800°C, should be significantly lower than ferritic steels. In order to avoid too long reheating time, loss of productivity and to minimize the fuel consumption, the slab hot charging at a temperature higher than 600°C is recommended.

60

For this scope hot charging tests were carried out in order to verify the industrial feasibility at the Duferco-La Louviere plant and hence to define the best practice. In parallel calculations were carried out in order to evaluate the effectiveness of hot charging in terms of productivity and to evaluate the impact for insertion of mini-program inside usual rolling schedule taking into consideration the expected total reheating time. In fig 2.4.1 is shown the typical temperature evolution of the conventional route. The slab from oxy-cutting, down the casting machine, until the piling up in the storage area. As can be noted the slab temperature decreases quite fast in the first part but after pile up the cooling rate is slower (about 20°C/h). The average time to achieve 500°C is, adopting the conventional route, about 14 h. An insulated box suitable for hot charging was realized setting-up a device already existing in steel works. To increase the thermal insulation a special attention was devoted to control the air circulation within the box. The slab temperature evolution adopting a hot charging route gave the following results. In fig.2.4.2 is shown the slab temperature evolution from the oxy-cutting until the storage within the insulated box. The slab monitored was the second as stacking order in the box. It can be noted that, a part the initial drop due to air cooling during transfer, the slab temperature drop (in the range 600-400°C) is strongly reduced (7°C/h) in the case of slab stored in box to be compared with 18°C/h for conventional slab piling up.

.

Fig.2.4.1 Typical slab temperature in conventional route.

Piling up

61

Fig.2.4.2 Typical slab cooling evolution adopting the hot charge route (piling up in a box).

Hot rolling schedule design The hot rolling schedule design consisted in the following main activities:

1. Evaluation of the high temperature mean flow stress of TWIP steels from hot compression tests; 2. Calculation of hot rolling schedule by means of CSM model taking into account the technical

specification of the DLL hot rolling mill; 3. Definition of the strip cooling pattern on run out table and coiling temperature.

Evaluation of mean flow stress from hot compression test results Calculation of the TWIP steel flow stress on the basis of flow stress measurements carried out by means of compression tests with a strain rate of 0.1s-1. Assuming that the maximum flow stress curve is an estimation of the mean flow stress it was applied the Ford-Alexander equation to calculate the rolling forces. This assumption is considered reasonable even if the measurements were made at strain rate significantly lower than the strain rate experimented in industrial hot rolling conditions.

Piling up in box

62

Fig. 2.4.3: Comparison of the MFS of TWIP2 and a conventional low-C steel (strain rate 0.1s-1). In fig 2.4.3 is shown the mean flow stress of TWIP2 steel. In the same plot is reported also the mean flow stress of a reference low C-Mn steel. As can be noted from the comparison results that TWIP2 steel has a significantly higher (70%) flow stress compared with a conventional low-C steel. Calculation of HRM schedules In the following the results of rolling schedules calculations of TWIP steel (Fe-22Mn-0.6C) on both roughing and finishing mill at the DLL Hot Strip Mill are reported. The following conditions have been taken into account:

Discharging slab temperature: 1200 °C Tranfer time on the run-out table from reheating furnace to roughing mill stand: 35 sec. Fe-22Mn-0.6C TWIP steel. Initial slab thickness: 250 mm Initial slab length: 5000 mm Rolling schedules are calculated for the bar head threading in all the passes

The dimensions considered in the calculations are reported in the next tables. ************************************************************************************** ** Reversing Roughing Hot Strip Mill Preset - Calculated rolling schedule ** ************************************************************************************** Strip width [mm]: 1000 Pass reduction [%]: 13.52 19.06 23.43 25.15 29.21 34.51 39.78 Bar thickness [mm]: 250.00 216.20 175.00 134.00 100.30 71.00 46.50 28.00Surface bar rolling temperature [°C]: 1137 1097 1087 1068 1069 1063 1068 1066Bar strain [-]: 0.145 0.211 0.267 0.290 0.345 0.423 0.507 Bar strain rate [1/sec]: 1.4 1.9 2.5 3.0 3.7 4.8 6.1 Bar yield stress [N/mm2]: 130.0 136.1 141.4 145.1 151.1 159.9 169.4 Rolling Force [kN]: 18064 21745 23528 22779 23567 24992 26056 Bar threading speed [m/sec]: 1.25 1.35 1.40 1.40 1.40 1.40 1.35 High rolling speed [m/sec]: 1.25 1.35 1.80 2.80 3.30 4.20 4.50 Forward slip [-]: 0.020 0.025 0.033 0.042 0.057 0.081 0.116 Rolling torque [kNm]: 1544 2052 2216 1947 1881 1828 1663 Motor power [kW]: 4934 7042 7827 6818 6495 6174 5244 **************************************************************************************

Table 2.4.2: HRM Rolling schedule - TWIP 22Mn 0.6C steel- slab format 1000x250x28 mm.

63

Table 2.4.3: Calculated TWIP steel finishing rolling schedule for strip format: 1000x28x4. mm. Discussion of results The calculated MFS (Fig. 2.4.3) reaveals that TWIP2 steel is characterised by significantly larger flow stress than the MFS of conventional low C-Mn steel. The conclusions of the hot rolling schedule calculations can be summarised as follows:

• Final rolling temperature should be greater than 900 °C to avoid excessive rolling loads and torques;

• Incoming bar temperature at the finishing mill entry should be as higher as possible(>1040°C) to decrease the motor current;

• Hot rolling of TWIP steel strips seems to be feasible, in terms of rolling loads and torques, till to 1000 mm width and final thickness down to 4 mm.

ROT cooling pattern and coiling temperature TWIP steels are characterised by a stable austenitic structure in the whole temperature range. This means that the strip cooling pattern after hot rolling does not play any metallurgical role. Nevertheless after hot rolling, even if, no phase transformation occurs the following topics have to be considered:

• Permanence in the temperature range 600-900°C could provoke strip surface oxidation (scale growth, decarburization, intergranular oxidation);

• Huge carbide precipitation occurs during slow coil cooling in the range 500-600°C; Coiling at high temperature (700-750°C) could favour a partial recrystallization without the need of additional continuous or batch annealing cycle. Nevertheless the risk of carbide precipitation during coil cooling has to be considered. Infact under the typical coil cooling rate conditions the time to cross the range 500-600°C can reach few hours. So the best strip cooling conditions on the run-out-table are:

• ROT cooling pattern: continuous cooling pattern from finishing rolling temperature (900°C) down to coiling temperature (CT) using the water curtains cooling system on the run out table.

• Coiling temperature: coiling in the range 450- 500°C.

64

Hot rolled strip annealing conditions The hot rolled strip microstructure can be strongly affected by:

• Hot rolling schedule: a proper choice of the hot rolling schedule allows to activate the dynamic recrystallization.

• Run-out table strip cooling pattern and coiling temperature: High coiling temperature could activate static recrystallization.

This means that if dynamic recrystallization is activated or high coiling temperature (>700°C) is adopted, to promote static recrystallization, a recrystallized (even partially) structure can be achieved and the annealing treatment of hot rolled strip can be avoided. Conversely if during hot rolling the dynamic recrystallization is not enough and/or the coiling temperature is too low to promote static recrystallization, in this case for the follow up of the industrial manufacturing route (cold rolling) a annealing cycle is necessary. In the latter case on the basis of the results of task 2.2.2, the annealing conditions could be defined as follows: soaking temperature at 1000°C for a soaking time of 30 s (continuous annealing line). Pickling ability characterization The pickling behavior of TWIP steels was evaluated on strip samples having a scale that could be considered as close as possible to the industrial hot band. The strip samples used for the experimentation were obtained through the following steps:

1. Hot rolled strips (TWIP and reference low carbon steels) were mechanically polished in order to remove completely the scale and simulate the steel condition at the entry finishing mill (after descaler);

2. The samples were heated in inert atmosphere (N2); After about 2 min for temperature homogenization was introduced air in the furnace in order to form the scale layer on the samples. To simulate the coiling, the samples were cooled down slowly in a furnace from a temperature of 700°C down to RT.

TWIP steel scale characterization by SEM In the following sections are reported the results of scale analysis by means of SEM-EDS. TWIP1 and TWIP4 (Fe-Mn-Al-Si-C) Both TWIP1 and TWIP4, in terms of chemical composition, are characterized by a significant Al content. This common aspect determines a similar scale structure and composition. The results can be summarised as follows:

• Outer scale layer: the outer layer is constituted of a matrix of Fe-Mn oxide (Fe is the main constituent). Dispersed in the matrix were detected a second phase of spinel particles (Al2O3)(Fe,Mn)O.

• Internal oxidation layer (interface with steel): • For TWIP1 sample the base steel, at interface scale-steel, is Fe-enriched and a fine

precipitation of small Fe-Mn oxide particles is detected. • For TWIP4 sample the base steel, at interface scale-steel, is Fe-enriched and a fine

precipitation of small Fe-Mn oxide particles together with rod-like spinel phase particles is detected.

In fig. 2.4.4-5 are reported the scale structure and main phases compositions detected on TWIP1 and TWIP4 steel variant.

65

Fig 2.4.4: TWIP1 scale analysis. Fig 2.4.5: TWIP4 scale analysis.

TWIP2-TWIP3-TWIP5 (Fe-Mn-C-N) The scale characteristics of TWIP2, TWIP3 and TWIP5 are quite similar. The results of SEM-EDS analysis can be summarised as follows:

• Visual aspect : brittle and porous; • Outer scale layer: the outer scale layer is constituted of a matrix of Fe-Mn oxide. In terms of

second phase some difference are found: • For TWIP2 sample the second phase is constituted of Fe-Mn-Si oxide particles (Mn is

the main constituent). • For TWIP3 and TWIP5 the second phase is constituted of metallic Fe-rich particles;

• Internal oxidation layer (interface with steel): at the interface scale-steel the base steel is Fe-

enriched and a fine precipitation of small Fe-Mn oxide particles is present.

In fig. 2.4.6 is reported the scale microstructure and main phases detected on TWIP2 and TWIP3 steel samples.

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a) TWIP 2 scale. b)TWIP3 scale. Fig.2.4.6: Scale microstructure and main phases detected on: a) TWIP2 and; b) TWIP3 steel samples.

TWIP 5 scale analysis.

Fig. 2.4.7: Scale microstructure and main phases detected on TWIP5.

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Summary of results • All TWIP variants are characterized by a outer scale layer constituted of a Fe-Mn oxide matrix

in which are embedded second phases particles. The chemical composition of these particles depends on specific steel chemical composition and in particular on the presence of Al. Therefore in presence of Al (TWIP1 &TWIP4) the second phase particles are spinel oxide particles (Al2O3)(Fe,Mn)O. In TWIP2-3-5 variants (no Al) the second phase are constituted of metallic particles (Fe-rich).

• The internal oxidation layer, at the interface scale-steel, is constituted of Fe-enriched metallic matrix in which are precipitated Fe-Mn oxide particles. This layer plays an important role in pickling ability and for this reason the pickling ability of TWIP steels is expected to be generally good due to presence of metallic Fe-enriched layer together with an overall brittle aspect.

Pickling tests The aim of this activity is to evaluate the pickling ability of TWIP steels using as reference a commercial low C-Mn steel. The procedure adopted for pickling tests can be described, for each TWIP variant, in the following steps:

• Preparation of a temperature controlled tank (66°C) with a water solution of 140 g/l HCl and 5g/l Fe2+ and 1ml/l inhibitor Leuzolit extra 284M.

• Measure of weight and dimensions of samples; • Immersion in the pickling bath for a fixed time ti; • Washing, cleaning and weighting of samples; • If after ti the sample did not reveal a good pickled surface the procedure was repeated for a

longer time ti + Δti. The results of the tests carried out in terms of average scale loss per second are reported in the table 2.4.5.

Table 2.4.5: Results of the pickling tests.

Further tests would be required for scaling up to industrial size as the mechanical conditioning of scale was not carried out before pickling (to get a cracked scale that favours the acid percolation and reduce the pickling time). Summarising the results yield to the following conclusions:

• The steel TWIP1, TWIP2 and TWIP3 reveal a good pickling ability (in HCl) in terms of soundness of pickled strip surface, even if in comparison with the reference low carbon steel the pickling kinetics is slightly slower.

• The steel TWIP4 reveals zone with residual scale even at longer immersion time in the pickling solution.

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• The steel TWIP5 requires longer pickling time with respect to reference and also with respect to the other TWIP1,2,3 steels.

Conclusion of task 2.4 The task 2.4 was incentrated on demonstrating the feasibility of TWIP steel production at Duferco La Louviere steel works and to define the industrial conditions for TWIP steel processing. The attention, of course, was focused mainly on the hot rolling process because it is well recognized that TWIP steel are characterized by a quite high flow stress at high temperature. The results of the study allows to conclude that TWIP steel hot rolling is feasible at Duferco La Louviere plant, provided that a dedicated operating practice, quite different from conventional low C-Mn steel, is adopted. This is the result from the point of view of plant capability. Slab hot charging at a temperature higher than 600°C is recommended to avoid too long reheating time, loss of productivity and to minimize the walking beam fuel consumption. Summarising the operating practice for hot rolling process of TWIP steel production are:

• Slab dimensions: thickness 250mm, width 1000mm, length 5000mm; • Slab hot charging; • Slab discharging temperature range: 1180-1220°C; • Entry finishing temperature range: >1040°C; • Finishing rolling temperature: 900-920°C; • Hot band thickness achievable > 4.0 mm; • Coiling temperature: 450-500°C. • Hot rolled strip pickling: conventional pickling process as adopted for C-Mn steel (HCl acid).

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WP3 Study of the deformation mechanisms and strain hardening behaviour Task 3.1 Characterization of mechanical properties in relation with the microstructure, dominating deformation mechanism, strain rate and temperature. The activities of task 3.1 are articled in 5 subtasks. The results are reported in the following sections. Task 3.1.1 Static uniaxial tensile tests to investigate the deformation mechanisms transition (deformation twinning --> dislocation glide) and strain hardening behaviour The activities of the subtask 3.1.1. revealed a more than expected complexity mainly due to necessity of a in-depth investigation on the effect of hydrogen and decarburization on embrittlement of TWIP steel. For this reason the activities were organized in further three subtasks as reported below.

Task 3.1.1.1 Quasi static tensile tests and strain hardening behaviour analysis • Quasi static tensile tests for a qualitative study of the effect of strain rate and temperature

on tensile behaviour and strain hardening analysis.

Task 3.1.1.2 Characterization of austenite phase stability and microstructural evolution in TWIP grades during deformation at different temperature • Investigation on deformation mechanism transition at high temperature (deformation

twinning dislocation glide); • Characterization of austenite phase stability and microstructural evolution in TWIP grades

during deformation in the temperature range -180°C 350°C (LOM, XRD);

Task 3.1.1.3 Investigation on TWIP steel embrittlement Due to anomalies detected during test performance and unexpected results on tensile tests (poor elongation) it was decided to carry out a in depth investigation on the causes of the large scattering of tensile properties revealed by TWIP steel. This subtask was not included in the technical annex of the project. The activities carried out were: • Metallographic examination of selected tensile specimens in order to characterize the strip

microstructure, presence of defects and analysis of surface oxidation and decarburization; • Low strain rate tensile tests on H-charged samples and internal hydrogen behavior (TDA).

Task 3.1.1.1 Quasi static tensile tests and strain hardening behaviour analysis Tensile tests, at room temperature, were carried out on all the TWIP variants and the results are reported in the table 3.1.1.1. The tensile specimens are in the state of cold rolled and annealed at 1000°Cx5min under N2 atmosphere. The following features can be highlighted. TWIP2 steel is characterized by an excellent elongation 82-88%. However also the other steel are characterized by quite high elongation to rupture (A80%) around 60-65%. These quite high elongations are due to large strain hardening coefficient n ranging between 0.55-0.58. TWIP5 steel is the exception because the coefficient n is significanly lower (0.4). In terms of yield stress and tensile strength some considerations can be underlined: TWIP steels are characterized by relatively low yield stress (240-280Mpa) and quite high tensile strength (820-930Mpa). TWIP5 steel has a different behaviour because yield stress is higher and a tensile strength is lower than TWIP1-2-3-4 grades.

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Table 3.1.1.1: Tensile test at 20°C results (cold rolled and annealed at 1000°Cx5min).

Tensile properties: effect of temperature The tensile properties of TWIP variants (TWIP2-3-4-5) were characterized at different temperature 20-350°C and strain rate. The occurrence of anomalies of tensile tests performance, mainly in terms of scattering of elongation, was highlighted. The tensile tests were repeated with a different preparation of the specimens, i.e. the samples after annealing were mechanically polished in order to remove a surface layer of about 0.5 mm (table 3.1.1.2). The comparison of the tensile tests results with and without surface polishing is shown in fig. 3.1.1.1.

Table 3.1.1.2: Tensile test results at different temperature with mechanically polished surface.

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As can be noted the elongations of polished tensile specimens are less scattered and typically are located at the upper bound of scattering band. As can be noted the temperature influence on tensile properties of the five TWIP grades is very similar with only some quantitative differences. The yield stress is slightly affected by temperature while the tensile strength is significantly influenced by temperature. This means that at higher temperature the Rp0.2/UTS ratio tend to increase due to progressive change of the deformation mechanism (twinning dislocation glide) and a consequent reduction of the strain hardening ability and elongation is detected (figure 3.1.1.1a-b).

a) b)

Fig. 3.1.1.1: a) tensile tests results carried out at room temperature and higher temperature 150-350°C; b) tensile tests results carried out at room temperature and higher temperatures 150-350°C.

Tensile properties: Strain rate sensitivity The strain rate sensitivity of the five TWIP steel grades was evaluated collecting all the data available, torsion tests, tensile tests and Hopkinson bar tests as well. All the mechanical tests were carried out at room temperature and the grain size of the specimens were in the range of 10μm<D<19 μm taking into account twin boundaries as a grain boundaries. Merging all the data obtained in this project in a single figure (tension and torsion, Von Mises equivalencies, no distinction between steels, 8 orders of magnitude of equivalent strain rate), figure 3.1.1.2 is obtained. In fig 3.1.1.2a is shown the yield stress behaviour as a function of strain rate. As can be noted, the behaviour of the five TWIP grades is very similar but with only some quantitative differences. The yield stress is substantially unaffected by strain rate up to 102s-1. This means YS is almost constant, i.e. sensitivity is about zero, over the strain rate range from 0.001 up to 100 s-1. Only at very high strain rates ( 200>ε& s-1) there is a strain rate induced increment of the flow stress. This results is confirmed even at higher deformation up to 0.25 equivalent plastic strain. At higher equivalent plastic strain the behaviour of TWIP steels appears different (fig 3.1.1.2b). In this case a negative strain rate sensitivity can be detected in the strain rate range from 10-3 up to 10-1 s-1. Taking into account the negligible strain rate sensitivity of Ys appear reasonable to ascribe this behavior to adiabatic heating of specimen. Infact for yield strain (almost) no work has been done yet, so no adiabatic heating can be expected. As it is well known the adiabatic heating is related to heat production during the work done during the deformation process of a specimen. However, at low strain rate, part of heat will disappear during slow straining, so that temperature increment remains lower (zero at strain rate below 0.0001 /s). At strain rates, high enough, the heat loss will remain insignificant, so that temperature increase is about independent of strain rate. This limit seems to be about 0.01 /s. Hence, at higher strain rates than this, temperature rise is constant (50°C after 0.4 strain and about 100° C after 0.9 strain).

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In fig 3.1.1.2b the flow stress at 40% as a function of strain rate is shown. It can be noted that at strain rates 0.01 /s and above, the stress slightly decrease (from 1300 to 1100 MPa). This shows that something is causing negative strain rate dependence/sensitivity. This could be the evidence of an intrinsic characteristic as reported in literature in TWIP steels but, in this case, it appears as a result of adiabatic heating of the specimens.

a) b) Fig. 3.1.1.2 The strain rate sensitivity of the five TWIP steel grades was evaluated collecting all the data, torsion, tensile tests and Hopkinson bar tests as well. Tensile properties: Strain hardening behavior Case 1: room temperature The work hardening rates for each TWIP steel is reproduced in figure 3.1.1.3. This analysis of the work

hardening behaviour was obtained by plotting the work hardening rate (εσ

dd

) against strain. To obtain

meaningful plots the experimental data was first smoothed by calculating a moving average of stress and strain values. A 6 order polynomial law was then fitted to each curve (plastic region only) and the work hardening rate obtained as a function of strain from the fitted polynomial.

Fig. 3.1.1.3: Calculated work hardening curves for each TWIP grade.

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From figure 3.1.1.3 it can be seen that the work hardening curves can be divided qualitatively into 3 stages.The first stage (decreasing hardening rate) is attributed to dynamic recovery. The change in hardening rate at the onset of stage 2 has been found to be due to the start of twinning. The new twins act as barriers to dislocation motion, and lead to an increase in strain hardening rate. Depending on the extent of twinning, these leads to the observed overall hardening rate to decrease at a lower rate, remain constant or for extensive twin formation to actually increase [4,5]. The third stage (decreasing hardening rate), has been interpreted as a region where further twin formation becomes difficult i.e. a region where the rate of twin formation decreases [4]. What it seems really peculiar of TWIP2 and TWIP3 steel is the second stage in which the strain hardening increase with strain. This should be related to the effect of deformation induced twinning occurring after the first stage predominated by dislocation glide. The comparison of the stress-strain curves for TWIP 2 and TWIP3 led to the conclusion that the larger work hardening rate in TWIP 2 could be a consequence of a greater degree of twinning compared to TWIP 3. From table 3.1.1.1-2 it can be seen that the UTS and elongation are significantly larger for TWIP 2 compared to TWIP 3, again consistent with a greater degree of twinning in TWIP 2. For the TWIP4 grade on the basis of the previous results it is possible to argue that the increase of work hardening in the second stage is a mixed contribution of strain induced alpha and/or epsilon martensite together with twinning. The balance of these mechanisms is responsible for the different behavior. Modelling the work hardening This analysis was carried out on TWIP grades 2 and 3 in which can be assumed that during deformation no strain induced phases are produced. Thus for these steels the stress-strain curves can be analysed further by applying a physical model for deformation of TWIP steels due to Bouaziz et al. [7]. The model fitting parameters are σ0 (yield stress), m (parameter related to SFE), k (dislocation forest hardening coefficient) and a (coefficient related to dynamic recovery). The values of the other model parameters used are shown in table 3.1.1.3 below:

Parameter Value Reference M 3 [7] α1 0.35 [9] G 72 GPa [7] b 0.26 nm [10] e 1 μm [7]

Table 3.1.1.3: Values of physical parameters used in modelling In addition to the above parameters the initial grain size of each TWIP grade is needed by the model. From optical microscopy results for TWIP 2 and 3, the average austenite grain sizes have been determined to be 32 and 25 μm respectively. The initial dislocation density, for both TWIP grades, was taken on reference [10] equal to 1012 m-2. A good agreement is obtained between model and experiment for both steel grades. The model fitting parameters obtained for each grade are shown in table 3.1.1.4. Steel grade σ0 (MPa) k a M TWIP 2 213 0.017 0.041 2.27 TWIP 3 235 0.009 0.040 2.06

Table 3.1.1.4: Obtained model fitting parameters As can be seen from table 3.1.1.4, the fitted yield stress is similar for both TWIP grades. Table 3.1.1.4 also shows the contributions of dislocation-forest hardening, dynamic recovery and twinning, through

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the parameters k, a and m respectively. As can be seen the dislocation forest hardening is significantly larger for the TWIP 2 grade. Furthermore TWIP 2 displays a slightly larger m value. In general a higher m (parameter related to SFE) value indicates a greater twin fraction for a particular strain. This analysis allows to conclude that TWIP2 steel developes a slightly larger amount of deformation induced twinning with respect to TWIP3.

Case 2: Strain hardening at elevated temperature

a) True stress-true strain curves for TWIP

3. b) True stress-true strain curves for TWIP

4. Fig. 3.1.1.4: True stress-true strain curves for TWIP3-4.

The interesting features of high temperature strain hardening behavior are:

• The appearance of a sharp yield point and the presence of serrations. Figure 3.1.1.4 shows the stress-strain behaviour of TWIP 3 and TWIP4 at 3 different temperatures. The occurrence of serrations is related to dynamic strain ageing via carbon locking of dislocations. This phenomenon is observed at room temperature but it is more pronounced at 250°C and almost absent at 350°C. This result is consistent with reference [7] discussed above, where in Fe-30Mn-1C alloy serrations were observed even at at room temperature.

• The work hardening is more extensive at room temperature. This is a consequence of the fact that the stacking fault energy increases with temperature. This leads in turn to a change in deformation mechanism depending on the SFE value. As temperature increases the tendency for twinning reduces and dislocation glide is preferred and became the main deformation mechanism.

Conclusions task 3.1.1.1 All the TWIP steel grades investigated revealed excellent tensile properties. However significant differences are present in terms of deformation mechanisms and microstructural evolution during deformation. The tensile properties resulted to be quite sensitive to the specimen surface preparation but the investigation of this phenomena is remainded in the task 3.1.1.3. Temperature effect The tensile properties of TWIP steel grades show a quite similar behaviour at increasing temperature with only some quantitative differences. The yield stress is slightly affected by temperature while the tensile strength is significantly influenced by temperature. This means that at higher temperature the Rp0.2/UTS ratio tends to increase due to progressive change of the deformation mechanism (twinning dislocation glide) and a consequent reduction of the strain hardening ability and elongation is detected. Strain rate effect

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For equivalent plastic strain below 0.25, the behaviour of TWIP steels is very similar at room temperature and the strain rate influence on the strength is very weak. Only at very high strain rates ( 200>ε& s-1) there is a small strain rate induced increment of the flow stress. Strain hardening ability The main characteristic of a true TWIP steel (TWIP2 and TWIP3) is the presence of a intermediate stage in the strain hardening curve the instantaneous coefficient n increase with strain. This should be related to the effect of deformation induced twinning occurring after the first stage of the σ−ε curve predominated by dislocation glide. The new twins act as barriers to dislocation motion, and lead to an increase in strain hardening rate. Depending on the extent of twinning, these leads to the observed overall hardening rate. The other TWIP grades (TWIP1,4,5) are characterised by a lower austenite stability and this is demonstrated by the occurrence of martensite (ε or α) during deformation or presence of ferrite also at zero strain as for TWIP1. The strain hardening of TWIP 1,4,5 grades shows a sligth different behaviour as a function of strain, due to the complex combination of deformation mechanisms (twinning+dislocations) and deformation induced phases (ε-martensite and α-martensite). Task 3.1.1.2 Characterization of austenite phase stability and microstructural evolution in TWIP grades during deformation at different temperature It is well understood that TWIP steel tensile properties are strictly related to SFE and therefore to temperature. The present task was focused on the study of the deformation mechanisms and austenite stability of TWIP variants in a broad range of temperature from -180°C up to 350°C.

Case 1: Room temperature microstructure evolution during straining For each TWIP grade extra tensile tests were performed for true strains of 0.1, 0.2 and 0.4. Samples for optical microscopy were then taken from each tensile specimen in both the “head” region (non-deformed) and the centre region (deformed). The samples were cut and sectioned to view the microstructure in the longitudinal direction i.e. parallel to the deformation axis. The microstructures were revealed with two etching techniques. The first (etchant I) involved a pre-etch with 2 % nital solution followed by etching with a solution of 20g sodium metabisulphite dissolved in 100ml of water. This etchant revealed the grain contrast. The second (etchant II), involves etching in a solution of 50ml saturated aqueous sodium thiosulphate and 5g of potassium metabisulphite. This reagent (also known as Klemm’s II reagent), gives a colour etch, with austenite grains yellow to brown or light to dark blue, whilst α-martensite is shown as dark brown and finally ε-martensite is revealed as white [1]. Figures 3.1.1.5-3.1.1.8 shows a summary of the microstructures obtained. Note that for clarity only the micrographs for 2 levels of strain are shown i.e. 0 (head region) and 0.4. TWIP1 microstructure in undeformed state is constituted of austenite + ferrite (6%). The deformed samples reveal flattened grains in which can be detected intersecting shear bands. In some regions dark plate like structures are visible. From their colour (brown) these plate structures are suggested to be α-martensite. From figure 3.1.1.5 a) TWIP 2 displays a equiaxed austenitic structure with a larger grain size compared to TWIP 1. The microstructure appears to be completely austenitic. The deformed structure (figure 3.1.1.5b) reveals intersecting shear bands in some grains.

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a) b) Fig. 3.1.1.5: TWIP 2 x500 a) head region etchant I, b) strain = 0.4, etchant I. The microstructure of TWIP 3 in the undeformed state is shown by figures 3.1.1.6 a). The structure is completely austenitic with equiaxed grains showing some shear banding and annealing twins as for TWIP 2 in the same condition. The grain size appears to be smaller than for TWIP 2. In the deformed state figures 3.1.1.6b) shows intersecting shear bands, as for TWIP 2.

a) b) Fig. 3.1.1.6: TWIP 3 x500 a) head region etchant I, b) strain = 0.4, etchant I.

Figure 3.1.1.7 a) shows the TWIP4 structure constituted of equiaxed grain structure with some annealing twins, as for the previous grades in the undeformed state. In addition the grain size appears to be similar to that for TWIP 2. After a strain of 0.4 it can be seen from figure 3.1.1.7b) that most grains contain shear bands which are mostly obscured by a darker constituent which seems to be present as thin parallel plates in some grains. The brown plate like features are observed. From their colour this constituent seems to be α-martensite. This result is confirmed by ferritoscope measurements.

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a) b) Fig. 3.1.1.7: TWIP 4 x500 a) head region (strain=0) etchant I, b) strain = 0.4 etchant II.

The undeformed microstructure of TWIP 5 is shown in figure 3.1.1.8 a). In the deformed state the microstructure as revealed by figures 3.1.1.8 b) shows some intersecting shear bands. In addition as for the deformed TWIP 1 and 4 grades, some brown plate like features are visible i.e. α-martensite.

a) b) Fig. 3.1.1.8: TWIP 5 x500 a) head region etchant I, b) strain = 0.4 etchant II. XRD analysis X-ray Diffraction Analysis (XRD) was carried out on TWIP2, TWIP3, TWIP4, TWIP5, strip samples before and after interrupted tensile test at 10%, 20% and 40% of elongation (always at 1/4 strip thickness). X-Ray diffraction experiments were performed with Mo-kα radiation, the volume fraction of phases and constituents were quantified by the Rietveld method (spectrum fitting in the range 18-35° 2θ). From XRD diffraction spectra 4 constituents were identified with varying intensities: austenite, ferrite (only for TWIP1), ε-martensite and α’-martensite.

10 μm

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Although the accuracy of the calculated volume fractions is low (high background noise level), the results obtained can be used as an indication of the microstructural evolution during tensile straining, and the differences in behaviour observed between the TWIP grades are expected to be still significant. Figure 3.1.1.9 shows the results obtained.

A) B)

C) D)

E)

Fig.3.1.1.9: Calculated volume fractions of constituents in TWIP steels during tensile straining. A) TWIP 1, B) TWIP 2, C) TWIP 3, D) TWIP 4 and E) TWIP 5. From figure 3.1.1.9A it can be seen that for the TWIP 1 steel, a significant amount of epsilon martensite forms during straining (compared to TWIP 2, 3 and 4). In addition the amount of alpha increases significantly during straining i.e. strain-induced alpha martensite has formed in addition to the ferrite content present at at zero strain. Figures 3.1.1.9B and 3.1.1.9C (TWIP 2 and TWIP3) show similar behaviour i.e. very low levels of epsilon martensite and alpha which remain approximately constant with increasing strain. For TWIP 4, figure 3.1.1.9D shows that a negligible amount of epsilon martensite is formed during straining, whilst the fraction of alpha i.e. alpha-martensite, increases with strain.

0 5

10 15 20 25 30 35

0 0.1 0.2 0.3 0.4

True strain

Fraction (%)

alphaepsilon martensite

0

5

10

15

20

0 0.1 0.2 0.3 0.4

True strain

Fraction (%)

alpha epsilon martensite

0 5

10 15 20

0 0.1 0.2 0.3 0.4 True strain

Fraction (%)

alpha epsilon martensite

0

5

10

15

20

0 0.1 0.2 0.3 0.4

True strain

Fraction (%)

alpha epsilon martensite

0

5 10

15

20

0 0.1 0.2 0.3 0.4True strain

Fraction (%)

alphaepsilon martensite

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Finally for TWIP 5, it can be seen from figure 3.1.1.9E that there is a low level of alpha that remains approximately constant with increasing strain. On the other hand the fraction of epsilon martensite appears to increase, before reaching a constant level at a strain of 0.2. Discussion of results TWIP 1 shows presence of both strain induced epsilon and alpha martensite. Since the fraction of epsilon martensite increases continuously with strain, it is suggested that alpha martensite forms directly from austenite i.e. different from the two-step transformation where epsilon martensite is the intermediate phase. This direct transformation from austenite to alpha martensite has previously been observed in another high manganese steel [1]. For TWIP grades 2 and 3, the X-Ray diffraction results suggest that the microstructure for both steels is almost entirely austenite during straining. This is consistent with conclusion that both steels deform via twinning. For TWIP4 grade, the X-Ray diffraction results indicate that strain induced alpha martensite forms. Since the fraction of epsilon martensite remains very low, the transformation proceeds directly from austenite as for TWIP 1. Finally for TWIP 5, a small amount of alpha phase is present during straining as indicated via X-Ray diffraction. From the previous optical microscopy results this phase is suggested to be alpha-martensite. Since the epsilon martensite fraction is shown to increase as strain proceeds (figure 3.1.1.9E), this would indicate that the alpha martensite forms directly from austenite as for TWIP 1 and 4. From the above results the microstructural evolution during room temperature deformation of these steels can be summarised in the following table:

Steel composition Undeformed microstructure

Deformed microstructure (strain = 0.4)

Fe-21Mn-3Al-3Si (TWIP 1)

.ferrBCCFCC αγ + Ms

HCPMsBCC

ferrBCCFCC εααγ +++ .

Fe-22Mn-0.6C-0.2Si (TWIP 2)

FCCγ FCCγ

Fe-18Mn-0.6C-0.2Si-0.02Nb (TWIP 3)

FCCγ FCCγ

Fe-16Mn-1.5Al-0.2Si-0.3C (TWIP 4)

FCCγ MsBCCFCC αγ +

Fe-21Mn-0.2C-0.2Si-0.2N (TWIP 5)

FCCγ MsHCP

MsBCCFCC εαγ ++

Note that in above table Ms = martensite and ferr. = ferrite. Table 3.1.1.5: Microstructural constituents/phases before and after testing of TWIP steels at room temperature.

Case 2: Low temperature range (0° -180°C) The characterization of the microstructure evolution during deformation at low temperature was carried out on specimens deformed (strain = 10% and 30%) at different temperature ranging between 0°C down to -180°C. The microstructure was examined by means of selective color etching typically used for high Mn steel (known as Klemm’s II). This reagent gives a color etch, with austenite grains yellow to brown or light to dark blue, while α’-martensite is shown as dark brown and finally ε-martensite is revealed as white.

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The fraction of the ferro-magnetic α’bcc(bcc-martensite), as a function of the temperature, was measured also by means of a ferritoscope. The results are shown in the following figure 3.1.1.10-11. On the basis of the selective metallographic etching the fraction of hcp-martensite was measured by means of LOM image analysis. The results of the measurements as a function of temperature are plotted in fig. 3.1.1.10. As can be noted the second phases (ε and α’) produced during deformation are below 3% for all TWIP variants. These circumstances together with low resolution related with deformed specimens does not allow the use of XRD or EBSD techniques to evaluate quantitatively the phases formed after straining.

Fig. 3.1.1.10: Ferritoscope measurements of α’bcc martensite on strained TWIP steels samples (deformation 10%).

Fig. 3.1.1.11: Ferritoscope measurements of α’bcc martensite on strained TWIP steels samples (deformation 30%).

The results can be summarized as follows: TWIP2: This variant revealed a quite stable austenitic structure in the range 20°C ÷ -180°C. In the whole temperature range tested only very low magnetic phase (α'-martensite) was detected. For what concern the hcp-martensite low amount (≤2%) was detected only after high deformation (30%) in the temperature range -150°C÷-180°C. TWIP3: The behavior of this variant is quite similar to previous steel (TWIP2). TWIP3 steel revealed a quite stable austenitic structure even at very low temperature. It is worthy to note that at a temperature below -120°C, increasing deformation from 10% to 30% the fraction of both α'-martensite and hcp-martensite phase increase even if the volume fraction remains quite low (<3%). TWIP4: The behavior of the steel TWIP4 is quite different from the other TWIP steels. The austenite phase of TWIP4 steel is markedly less stable even at room temperature due to higher formation rate of α’- martensite during deformation. As can be noted from comparison of fig. 3.1.1.10-11, on decreasing the test temperature the fraction of α’-martensite increases quickly while the fraction of ε-martensite decrease. This can be ascribed to temperature dependence of SFE and hence at lower temperature (lower SFE) α’-martensite is favored. TWIP5: in the temperature range 20°C ÷ 0°C the austenite phase is stable and only a low amount of hcp-martensite can be detected (<1%) even after straining at 30% . Some presence of α’-martensite is detected but the percentage remains very low even at the lowest temperature tested. At temperatures below -50°C the occurrence of hcp-martensite increases markedly with respect to the other TWIP variants. Nevertheless even at the lowest temperature (-180°C) the amount of second phase is below 5% (after deformation at 30%).

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Case 3: High temperature microstructure evolution during straining (RT 350°C) The study of deformation mechanisms operating during deformation at temperature in the range 250°C-350°C was carried out by means of tensile specimens interrupted after a true strain = 0.3 (strain rate 0.01 s-1). In the fig. 3.1.1.12-15 are shown the microstructure of TWIP1,3,4,5 deformed at high temperature.

Fig. 3.1.1.12: TWIP1: deformed at 250°C (strain 0.3 - strain rate = 0.01 s-1).

a) 250°C b) 350°C

Fig.3.1.1.13: TWIP3: microstructure deformed at: a) 250°C; and b) 350°C (strain 0.3 - strain rate = 0.01 s-1).

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a) 250°C b) 350°C

Fig.3.1.1.14: TWIP4: deformed structure at: a) 250°C; b) 350°C (strain 0.3 - strain rate = 0.01s-1).

Fig. 3.1.1.15: TWIP5: deformed at 250°C (strain 0.3 - strain rate = 0.01 s-1). The effect of temperature on deformation induced twins is quite evident. The results can be summarized in the following table in terms of occurrence of deformation induced twins.

Table 3.1.1.6: Mechanical twins occurrence during deformation at high temperature (strain rate 0.01s-1).

Conclusions of task 3.1.1.2 The results of this task can be summarized in the following topics:

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• TWIP2 and TWIP3 are characterized by largest stability of austenite phase under deformation from 250°C down to -180°C. Within this temperature range the deformation induced twinning represent the main deformation mechanism. At temperature of 350°C both steels do not reveal deformation induced twinning.

• TWIP1, TWIP4 and TWIP5 XRD pattern analysis revealed the occurrence during deformation at room temperature in addition of deformation induced twinning (TWIP effect) even formation of second martensitic phases (α’ and ε). This tendency became stronger at low temperature due to decrease of SFE. Since the fraction of epsilon martensite increases continuously with strain, it is suggested that alpha martensite forms directly from austenite and not through the two-step transformation where epsilon martensite is the intermediate phase. The occurrence of deformation induced twinning disappears before 250°C.

Task 3.1.1.3 Investigation on TWIP steel embrittlement The activities carried out in this subtask are:

• Metallographic examination of selected tensile specimens in order to characterize the strip microstructure, presence of defects and analysis of surface oxidation and decarburization;

• Low strain rate tensile tests on H-charged samples and internal hydrogen behavior and consequences on TWIP tensile properties.

3.1.1.3.1 Metallographic examination of selected tensile specimens in order to characterize the strip microstructure, surface decarburization and oxidation The microstructure of tensile specimens revealing low tensile properties was fully characterized. In fig 3.1.1.3.1 a) is shown the cross section of a tensile specimen of TWIP3 steel. The annealing treatments were carried out under N2 fluxing (O2= 4-5% and dew point 0°C). As can be noted the strip surface shows a lot of cracks with a penetration up to a depth of about 200 μm. Fig 3.1.3.1.1b) shows a close up of the microstructure in correspondence of the strip surface. As can be noted a huge presence of martensite can be detected in the microstructure. This is also confirmed by magnetic behavior revealed by means of a ferritoscope.

Fig. 3.1.1.3.1 a) Longitudinal cross section of a TWIP3 tensile specimen close the fracture surface.

Fig.3.1.1.3.1b): Subsurface microstructure after tensile test (TWIP3) – Klem II etching.

3.1.1.3.2 Study of TWIP steels decarburization during high temperature annealing The effect of surface decarburization on TWIP steel performances was studied through a series of annealing tests reported in the table 3.1.1.3.1.

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Samples Temperature Soaking time TWIP 2,3,4,5 1000°C, 1200°C 300s, 1100s C45 (reference) 1000°C, 1200°C 300s, 1100s

Table 3.1.1.3.1: Annealing tests carried out to study the strip surface decarburization. The annealing treatments were carried out in a muffle furnace with air (dew point 10°C) in order to study the effect of decarburizing conditions on TWIP steel. A conventional C45 steel sample was used, as a reference, to compare the decarburization behavior of TWIP (austenitic steel) at two high temperatures (1000°C and 1200°C). All the strip samples (hot rolled state) were mechanically polished (about 1mm was removed for each surface) in order to remove completely the scale layer and also the previous decarburized layer. On selected samples (TWIP2 and TWIP3) were carried out GDOES analysis in order to study quantitatively the carbon and manganese concentration profiles. Micro-hardness profiles were carried out on all samples in order to evaluate the evolution of the local tensile strength as a function of the depth from the surface. A ferritoscope was used in order to detect the presence of martensite α’ in the microstructure close the strip surface. All the TWIP samples annealed at 1000°C and 1200°C revealed in the subsurface zone a ferro-magnetic behavior. This confirms that the austenitic phase in the decarburized zone is destabilized and the microstructure is a mixture of γ+α' even without any deformation. The presence of α'-martensite is due to the local lowering of SFE in the C and Mn depleted zone. In the following fig. 3.1.1.3.2-3 are shown the microstructure of the TWIP2 samples after annealing treatment at 1000 and 1200°C. The effect of decarburization is quite clear and typically can be noted in all TWIP samples two different zones moving from surface towards the bulk (different etching behavior): the outer zone is characterized by higher amount of martensite (low C and Mn content). Below, towards the bulk, the presence of martensite is lower (Mn approach bulk content) but the local SFE value could be lower than bulk material and additional martensite can be formed during deformation.

Fig. 3.1.1.3.2: TWIP2 microstructure after annealing at 1000°C x 300s.

Fig. 3.1.1.3.3: TWIP2 microstructure after annealing at 1200°C x 1100s.

As can be noted from the values reported in the table 3.1.1.3.2, for TWIP steels, the zone with mixed microstructure reveals a quite high hardness with respect to bulk values. The huge difference in terms of hardness between bulk and surface highlights also that the ductility of the subsurface zone is quite low. On contrary it is worthy to note that, for conventional C45 steel, the decarburized zone is softer as revealed by lower hardness with respect to bulk.

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   1000°C  1200°C    Bulk (γ)  subsurface zone (γ+α')  Bulk (γ)  subsurface zone (γ+α') 

TWIP2  178  265  189  320 TWIP3  181  300  187  330 TWIP4  190  240  210  260 TWIP5  182  274  180  292 C45  270  190  275  180 

Table 3.1.1.3.2: Comparison of Vickers hardness (HV200g) average values on bulk and decarburized zone.

In the following fig. 3.1.1.3.4 the depth of decarburized layer (measured as depth of subsurface zone with microstructure (γ+α’) are plotted together with the Birks-Jackson model. As can be noted the agreement is quite good considering the half of depth arising from model calculations. In the same figure the decarburization behavior of TWIP steels is compared with the C45 medium carbon steel. As can be noted, in the range of soaking time explored 300-1100s, the decarburization depth of TWIP steels is larger than carbon steel. This difference could be originated by a different behavior of these steels in the temperature range below 910°C. In fact for the austenitic TWIP steels the carbon is in solid solution and is ready to react with oxygen even at relatively lower temperatures. For a medium carbon steel (C45) the amount of decarburization at lower temperatures (during heating) is low due to phase transformation α+pearlite γ. During this transient the formation of a ferrite layer (low C solubility) on the strip surface could slow down the steel decarburization kinetics.

Fig. 3.1.1.3.4: Comparison of the decarburization depth of different TWIP steel variants (TWIP2,3,4,5) and C45 steels at 1000°C and 1200°C. 3.1.1.3.3 Decarburized layer analysis by means of GDOES In fig. 3.1.1.3.5 is shown the carbon and manganese concentration profiles on TWIP2,3,4 samples in the annealed state. As can be noted the both C and Mn are depleted in the subsurface zone due to annealing

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treatments. The depth of Mn depleted zone (<30μm) is significantly lower than C decarburized layer. The decarburized layer is quite independent from TWIP variant and is typically in the range 120-150μm for annealing at 1000°C and 300-350μm for annealing at 1200°C.

TWIP 2: decarburization after annealing at 1200°C x

300s. TWIP 2: Mn concentration profile after annealing at

1200°C x 300s.

TWIP3: decarburization after annealing at 1000°Cx

300s. TWIP3: Mn concentration profile after annealing at

1000°C x 300s

TWIP4 - decarburization in the final state: (final annealing at 1000°Cx300s).

TWIP4 - Mn concentration profile in the final state (annealing at 1000°Cx300s).

Fig. 3.1.1.3.5: GDOES Carbon and manganese concentration profiles on TWIP2,3,4.

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3.1.1.3.4 Study of the effect of hydrogen pre-charging on low strain rate tensile tests The characterization of TWIP steels susceptibility to hydrogen embrittlement represent a huge work and require a dedicated project of investigations and tests. Nevertheless, even if in the present project this topic was not foreseen in the technical annex, selected tests were planned and carried out. In the following the activities carried out are detailed:

1. Hydrogen charging of tensile specimens; 2. Low strain rate tensile tests on hydrogen charged TWIP specimens (charged at 900°C) in

which the decarburized layer was mechanically removed. 3. Hydrogen thermal desorption analysis (TDA) was carried out in order to investigate on the

hydrogen trapping sites in TWIP steel in the temperature range 400-1100°C. Hydrogen charging of tensile specimens Tensile specimens of TWIP3 steel were hydrogen charged at 900°C using a suitable furnace working under a pure hydrogen flux (10 l/h). After soaking the samples were quenched and the final hydrogen content was evaluated and the results are reported in the following table 3.1.1.3.3.

Sample condition Hydrogen content (ppm)

Mechanically polished (reference) 1.3

Mechanically polished and Hydrogen charged at 900°C 4.5

Table 3.1.1.3.3: Hydrogen content of the reference and charged TWIP3 samples.

1. Tensile tests on hydrogen charged TWIP3 specimens Tensile tests (strain rate=10-2 s-1) on reference sample and hydrogen charged TWIP3 specimen, with the decarburized layer mechanically removed, were carried out in order to evaluate the effect of high hydrogen content on TWIP3 tensile properties (fig. 3.1.1.3.6).

Fig. 3.1.1.3.6: Tensile test curve of hydrogen charged TWIP3 sample.

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Table 3.1.1.3.4: Tensile tests reference and hydrogen charged TWIP3 sample.

As can be noted (fig. 3.1.1.3.7) the tensile properties of hydrogen charged sample and reference blank sample show similar tensile properties. The strip sample with the decarburized layer shows a significantly lower elongation to rupture.

Fig. 3.1.1.3.7: Results of tensile tests carried out on sample with different hydrogen content. 3.1.1.3.5 Thermal desorption analysis (TDA) on TWIP steel Hydrogen determination by LECO RC-412 Hydrogen content of metallic materials is measured at CSM by thermal desorption technique, using the LECO RC-412 instrument. Such instrument is basically made up by a resistance heated, temperature-controlled furnace, a thermal conductivity detector and several filters to separate hydrogen from other emitted gases. Instrument operation can be summarized as follows:

- Specimen heating in the furnace, under constant nitrogen flow; - Hydrogen and other adsorbed gas (H2O, CO2, etc.) evolution; - Emitted gas transport by nitrogen flow (carrier gas); - Gas filtration to separate hydrogen from other emitted gases; - Determination of metal hydrogen content by means of the thermal conductivity detector.

LECO RC-412 main features are presented in table:

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Furnace temperature range 80 – 1200 °C

Powder 0,25 grams Sample size Strip or tubular samples up to 25,4 mm x 101,6 mm Analysis time 2400 seconds maximum Heating rate 15 – 200 °C/minute

Number of selectable phase steps 10 Maximum time for each phase 600 seconds

Accuracy 0,1 ppm Instrument range 0,1 -1000 ppm

Table 3.1.1.3.5: LECO RC-412 main features. The hydrogen determination is made setting the starting temperature to 400°C (hold for 100 s), the ending temperature to 1000°C (hold for 300 s) using a heating rate of 120°C/minute. This procedure provides the operator with a graphic plot, hydrogen signal vs. analysis time. In this plot one or more peaks will appear and their area can be related to metal hydrogen content. For example, analysing austenitic steels or nickel super alloys three different peaks can be found, one at about 400 °C, a second between 500 - 600 °C and a third above 800 °C. The position and the height of each peak indicates that the relative amount of hydrogen in each different trapping site is different. According instrument features, a different procedure can be proposed in order to identify dominant types of trapping sites and their relative hydrogen content. A proper sample preparation could, also, permit an estimation of the amount of diffusible hydrogen. A slow heating rate should be used (verified minimum rate 15°C/minute) and the start temperature should be set to about 100°C (verified minimum temperature 80°C), this type of analysis will show the operator the differing types of hydrogen effusing the sample. The point is to design a temperature profile based on slow rate, broad temperature range and on instrument time constraints. For example, using a constant heating rate of 15 °C/min, can be investigated just a 700°C range, i.e. if the starting temperature is set to 80°C is not possible to overcome 700°C in one sweep. Optimization of these testing parameters would allow to understand different trapping site behaviour. To implement hydrogen differentiation, the temperature profile is programmable into 10 different phase. The operator enters start temperature, end temperature, heating rate, and amount of time to hold at the ending temperature for each phase selected. The results of thermal desorption analysis (TDA) of blank sample and hydrogen charged sample are shown in the following fig. 3.1.1.3.8. The characteristic main peaks, partially overlapping, of the reference sample are located at 800°C and 1000°C. The hydrogen charged sample is characterized by a main broad peak ranging between 800-1100°C. These high de-sorption temperatures indicate that the hydrogen is trapped into the metal structure by high energy bonding. The temperature profile indicates, also, that the diffusible hydrogen is practically absent. Therefore, under plastic deformation the hydrogen absorbed cannot move so easily to reach dislocation and/or crack tip, creating favourable condition to fragile fracture. These observations are confirmed by the mechanical testing results, where the specimen with about 5 ppm of hydrogen behaved in the same way of that having 1 ppm. Therefore, the conclusion is that the hydrogen absorbed into the TWIP alloy did not determined any embrittlement and/or loss in ductility.

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Fig. 3.1.1.3.8: DTA plot on blank and hydrogen pre-charged sample.

Conclusions Task 3.1.1.3 The results achieved can be summarized in the following points: 1. The results of the investigations carried out in the present task 3.1.1.3 allow to argue that the

embrittlement problems of TWIP steels are mainly related to decarburization and Mn depletion producing a mixed γ+α’microstructure in the subsurface zone.

2. Carbon and manganese concentration profiles on TWIP steels show that both C and Mn are depleted in the subsurface zone due to annealing treatment under decarburizing atmosphere. It is worthy to note that the depth of Mn depleted zone (typically <30μm) is significantly lower than C decarburized layer.

3. All the TWIP samples annealed at 1000°C and 1200°C reveal in the subsurface zone a mixed microstructure γ+α' even without any deformation with a resultant ferro-magnetic behavior. The presence of α'-martensite is due to the local lowering of SFE in the C and Mn depleted zone.

4. TWIP2,3,4,5 steels reveal a similar behavior that is in good agreement with the theory of decarburization in austenitic phase (Birks-Jackson model) strictly related to bulk carbon diffusion. The decarburization depth is more sensitive to annealing temperature (exponential dependence) with respect to soaking time (t 1/2 dependence).

5. The decarburization depth of TWIP steels is larger than carbon steel probably because for austenitic steels the carbon is in solid solution and is ready to react with oxygen. The kinetics of decarburization of a medium carbon steel (C45) on relatively short time is decreased by phase transformation α+pearlite γ. During this transient the formation of a ferrite layer on the strip surface could slow down the steel decarburization kinetics.

6. In terms of decarburization the final annealing process (after cold rolling) could be critical. For this treatment a controlled annealing furnace atmosphere has to be considered in order to avoid the occurrence of decarburization and so the formation in the subsurface zone of a mixed γ+α’ microstructure.

7. Tensile tests (ε’=10-2s-1) carried out on sample with different hydrogen contents shown similar tensile properties. The elongation to rupture of hydrogen charged specimen is markedly higher than the strip sample with a surface decarburized layer.

8. The high desorption temperatures indicate that the hydrogen is trapped into the metal structure by high energy bonding. This result would suggest that the diffusible hydrogen can be considered negligible. Therefore, under plastic deformation the hydrogen absorbed cannot move so easily to reach dislocation and/or crack tip, creating favourable condition to fragile fracture.

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Task 3.1.2 Dynamic Tensile Properties High strain rate deformation of TWIP steels were carried out on cold rolled sheets of high-Mn TWIP steels with 1.5-2 mm thickness using a tensile split-Hopkinson bar method. Dynamic tensile testing of sheet steels is becoming more important due to the need for more optimized vehicle crashworthiness analysis in the automotive industry. The deformation characteristics of sheet steels can be measured precisely to an order of 1000 s-1. Positive strain rate sensitivity, i.e. the strength increases with strain rate, offers a potential for improved energy absorption during a crash event. Specimens of varying thickness 1.5 < t < 2 mm were cut longitudinally to the rolling direction of the annealed strips. All tests were carried out at room temperature with two different strain rates 855 and 1250 s-1. Results and Discussion In the fig. 3.1.2.1a) the quasi-static tensile properties of TWIP1-5 grades are compared with the typical tensile properties of TRIP700, DP600, H340LAD. Dynamic tensile mechanical properties, shown in Fig. 3.1.2.1b), showed raising yield strength for all steels, for instance for TWIP2, Rp0.2 increased from 250 MPa for quasi-static to 560 MPa at high strain rates. This positive strain rate sensitivity revealed by TWIP is also detected in other steels nevertheless must be stressed that the elongation of TWIP steels is still much higher than those of the other three automotive steels (TRIP, DP600 and H340LAD). Fig. 3.1.2.2 shows the ratio of the yield stress to the tensile strength in quasi-static and dynamic tensile tests for TWIP steels and the automotive steels. In addition, the plot includes strain hardening coefficient n-value for full curve in dynamic tensile tests.

Fig. 3.1.2.1 a): Comparison between the quasi-static mechanical properties, yield strength Rp0.2, tensile strength Rm, and fracture elongation A of the present TWIP steels [15] and those of TRIP700, DP600 and H340LAD [17] at RT (the strain rate ≈ 10-3 s-1).

Fig. 3.1.2.1 b): Comparison between the dynamic mechanical properties, yield strength Rp0.2, tensile strength Rm, and fracture elongation A of the present TWIP steels and those of TRIP700, DP600 and H340LAD [17] at RT (the strain rate of order of 103 s-1).

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Fig. 3.1.2.2: Comparison between the stress ratio in quasi-static and dynamic tensile tests of the present TWIP steels and those of TRIP700, DP600 and H340LAD [17]. In the following sections the details of the behaviour of each TWIP variant are reported. TWIP2 (reference steel) The strain rate dependence of the stress-strain behaviour at 870 s-1 and 1210 s-1 appears similar. The yield strengths at dynamic strain rates are almost identical 550 and 562 MPa, respectively. However, the yield strength of that steel under quasi-static tensile tests at RT is much lower (273-280 MPa).TWIP2 steel exhibited higher uniform elongation (68%) than the other steels in SHB tests. These results are consistent with the results obtained under quasi-static tensile tests, while TWIP2 showed highest elongation to rupture (A = 80%). Fig. 3.1.2.3 shows the work hardening coefficient (n-values) of TWIP2 as tested at 1210 s-1. It can be seen that the n-value increases continuously with increasing strain and reaches maximum values about 0.6 at the uniform elongation.

Fig. 3.1.2.3: Strain-hardening coefficient (n-value) from the Hollomon equation vs. True plastic strain for TWIP2 tested at 1210 s-1 and RT.

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The fracture surface of TWIP2 after SHB test at 1210 s-1 was examined using SEM. The fracture surface shows a clear ductile morphology. TWIP1 The dynamic tensile true stress-true strain curves at 855 s-1 and 1250 s-1 reveal that the strain rate has small effect on the dynamic flow curves. The yield strengths are 504 MPa and 575 MPa at 855 s-1 and 1250 s-1, respectively. However, the yield strength of that steel under quasi-static tensile tests at RT is much lower (245-260 MPa). For both dynamic and quasi-static tensile tests, it is observed that after yielding, the flow stress increases gradually with further straining and it reaches almost the same ultimate (engineering) tensile strength of 780 MPa and uniform elongation (48-50%). Thus the tensile strength and failure strains are independent of strain rate. The n-values (strain hardening coefficients) of TWIP1 while tested at the highest strain rate of 1250 s-1 increases continuously with increasing strain and reaches the maximum value higher than 0.45 at true strain of 0.4. TWIP3 Dynamic tensile flow curves of TWIP3 at two strain rates show similar behaviour with small effect of the strain rate. For example, yield strengths are 480 and 510 MPa and ultimate tensile strengths are 835 and 900 MPa at 830 and 1280 s-1, respectively. However, the yield strength and tensile strength of that steel are 280 and 850 MPa, respectively, under quasi-static tensile tests at RT. The work hardening coefficient, n-values, of TWIP3 tested at 1280 s-1 shows that the steel exhibits a similar behaviour for the other investigated steels. While strain hardening increases with strain and reaches maximum n-value 0.55 at high uniform elongation. TWIP4 TWIP4 exhibited the lower yield strength (420 MPa) than TWIP2 and TWIP3 steels. This behaviour was similar also on quasi-static tensile tests with a yield strength of 240 MPa. It can be attributed to its lowest Mn content. Microstructure examination of TWIP4 after dynamic tensile tests at 1260 s-1 displays highly distorted grains with very high density of mechanical twins. The work hardening coefficient of TWIP4 tested at 1260 s-1 shows that the steel exhibits a similar behaviour as for the other investigated steels. The strain hardening increases with strain and reaches the same maximum n-value of 0.55 as in the case of TWIP3.

TWIP5 The steel exhibited high dynamic tensile mechanical properties with yield strength 510 MPa and high uniform elongation 52%.

Conclusion of task 3.1.2 The stress ratios (the yield stress to the tensile strength) of TWIP steels are much lower than those of the automotive steels under quasi-static conditions. This means TWIP steels promoted higher strain hardening potential. Under dynamic tensile conditions, stress ratios of TWIP steels increased as a result of increasing yield strength, but they are still lower than those of the three automotive steels. Furthermore, the n-values of TWIP steels from dynamic tensile tests are much higher than those of the automotive steels. The enhancing of the ductility, strength and strain hardening of material is advantageous for crash energy absorbing characteristics.

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TWIP steels promoted higher ductility and higher strain hardening with high strength, and therefore TWIP steels have higher crash energy absorption and consequently higher crash safety than steels TRIP700, DP600 and H340LAD. Task 3.1.3. Torsion tests in the 20ºC-450ºC temperature range: stress-strain behaviour at large strains The behaviour of four TWIP steels (TWIP steels 2, 3, 4 and 5) up to large plastic strains has been studied by performing torsion tests at five temperatures in the 20ºC-450ºC range and at three tensile equivalent strain rates ranging from 1.4·10-3 s-1 to 3.73 s-1. The stress-strain curves were calculated from the torque-twist curves using the method of Fields and Backofen. The initial structure was the structure after hot rolling. Figures 3.1.3.1 to 5 present the torsion shear stress vs. plastic shear strain curves for the four steels. As no extensometer was used in the torsion tests, the small strain range of the results (e.g., the yield stress) is not reliable. The results show a strong temperature effect in the strength level and an important effect of strain rate in the ductility. Also, the apparent strain rate effect has a superposed adiabatic heating effect at the two highest strain rates used. The trend of the behaviour is common for the four steels, although TWIP3 is clearly more ductile than the other three compositions. A change in the behaviour from cold-work to warm/hot-work occurs at about 250ºC. At the highest strain rate and at low temperature, the initial strain hardening rate is much higher than its static counterpart, a fact that could be linked to an enhancement of the twinning activity at high strain rates. However, only TWIP2 and TWIP3 below 250ºC clearly behave as expected from TWIP steels, with the typical hardening stage associated to profuse deformation twinning. In all other cases the work hardening is an approximately linearly decreasing function of the flow stress soon after the yield stress, a behaviour characteristic of dislocation-mediated plasticity controlled by dynamic recovery. At the higher temperatures and smaller strain rates tested, the stress-strain curves show a maximum (and even oscillations) that could be due to dynamic recrystallization, but the confirmation or rejection of this question by microstructural study is still pending. At room temperature, the strain rate dependence observed in torsion agrees very well with the equivalent values obtained in tension up to high strain rates. In fig. 3.1.3.6 the maximum shear stress in torsion as a function of temperature for different strain rates are shown. In fig. 3.1.3.7 the torsion shear failure strain of the four TWIP steels are shown as a function of temperature and strain rate.

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0

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Fig. 3.1.3.1: Torsion shear stress-plastic shear strain curves at =ε 1.4·10-3 s-1 and corresponding Kocks-Mecking plots (strain rate vs. flow stress).

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Fig. 3.1.3.2: Torsion shear stress-plastic shear strain curves at =ε 0.113 s-1 and corresponding Kocks-Mecking plots (strain rate vs. flow stress).

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(d)

TWIP 5

Fig. 3.1.3.3: Torsion shear stress-plastic shear strain curves at =ε 3.73 s-1 and corresponding Kocks-Mecking plots (strain rate vs. flow stress).

98

0

200

400

600

800

1000

0 100 200 300 400 500T (ºC)

τ γ (M

Pa)

6.5 s-11.2 s-10.0024 s-1

τ0.520

τ0.173

TWIP2

τ1.040

(a)

0

200

400

600

800

1000

0 100 200 300 400 500

T (ºC)

τ γ (M

Pa)

(b)

6.5 s-11.2 s-10.0024 s-1

τ0.520

τ0.173

TWIP3

τ1.040

Fig. 3.1.3.4: Torsion shear stress for three shear strain levels, 0.173, 0.520 and 1.040, as a function of temperature, for the three strain rates tested (TWIP2, TWIP3).

99

0

200

400

600

800

1000

0 100 200 300 400 500T (ºC)

τ γ (M

Pa)

(c)

6.5 s-11.2 s-10.0024 s-1

τ0.520

τ0.173

TWIP4

τ1.040

0

200

400

600

800

1000

0 100 200 300 400 500T (ºC)

τ γ (M

Pa)

(d)

6.5 s-11.2 s-10.0024 s-1

τ0.520

τ0.173

TWIP5

τ1.040

Fig. 3.1.3.5. Torsion shear stress for three shear strain levels, 0.173, 0.520 and 1.040, as a function of temperature, for the three strain rates tested (TWIP4, TWIP5).

100

TWIP 2

0200400600800

10001200

0 100 200 300 400 500T (ºC)

τ M (M

Pa)

0.0024 s-10.196 s-16.46 s-1

(a)

TWIP 3

0200400600800

10001200

0 100 200 300 400 500T (ºC)

τ M (M

Pa)

0.0024 s-10.196 s-16.46 s-1

(b)

TWIP 4

0200400600800

10001200

0 100 200 300 400 500T (ºC)

τ M (M

Pa)

0.0024 s-10.196 s-16.46 s-1

(c)

TWIP 5

0200400600800

10001200

0 100 200 300 400 500T (ºC)

τ M (M

Pa)

0.0024 s-10.196 s-16.46 s-1

(d)

Fig. 3.1.3.6: TWIP steels, maximum shear stress in torsion as a function of temperature for different strain rates.

TWIP 2

012345678

0 100 200 300 400 500T (ºC)

γ F

0.0024 s-10.196 s-16.46 s-1

(a)

TWIP 3

012345678

0 100 200 300 400 500T (ºC)

γ F

0.0024 s-10.196 s-16.46 s-1

(b)

TWIP 4

012345678

0 100 200 300 400 500T (ºC)

γ F

0.0024 s-10.196 s-16.46 s-1

(c)

TWIP 5

012345678

0 100 200 300 400 500T (ºC)

γ F

0.0024 s-10.196 s-16.46 s-1

(d) Fig. 3.1.3.7: Torsion shear failure strain of the four TWIP steels as a function of temperature and strain rate.

101

Conclusion Task 3.1.3 The hot torsion results show a strong temperature effect in the strength level and an important effect of strain rate in the ductility. The trend of the behaviour is common for the four steels, although TWIP3 is clearly more ductile than the other three compositions. A change in the behaviour from cold-work to warm/hot-work occurs at about 250ºC. At the highest strain rate and at low temperature, the initial strain hardening rate is much higher than its static counterpart, a fact that could be linked to an enhancement of the twinning activity at high strain rates. However, only TWIP2 and TWIP3 below 250ºC clearly behave as expected from TWIP steels, with the typical hardening stage associated to profuse deformation twinning. In all other cases the work hardening is an approximately linearly decreasing function of the flow stress soon after the yield stress, a behaviour characteristic of dislocation-mediated plasticity controlled by dynamic recovery.

Task 3.1.4 Hot ductility curves in the temperature range 700 ÷ 1300 °C by means of a Gleeble simulator Experimental Hot ductility of TWIP steels was carried out at a constant strain rate 1 s-1 on a Gleeble 1500 thermo-mechanical simulator. The diameter of tensile rod was 7 mm and the length 120 mm. In the middle there is a thinned gauge length zone with 8 mm in length and 6 mm in diameter. In hot ductility tests, specimens were reheated in vacuum at the rate of 20°C/s to 1150°C-1250°C for 2 min, to dissolve any precipitates present, and then cooled to the test temperature (between 700 and 1100°C) at the cooling rate of 5°C/s. After 30 s of soaking at the test temperature, long enough to stabilize the temperature field, the specimens were tensile strained until fracture at the constant true strain rate of 1 s−1. At the testing temperatures higher than the reheating temperatures, the specimens were heated directly to the test temperature and held for 60 s before tensile straining. Results and Discussion TWIP1 (Fe-22Mn-3Al-3Si) The dependence of the reduction of area (RA) on the deformation temperature, i.e. a hot ductility curve for TWIP_1 is shown in Fig. 3.1.4.1.

Fig. 3.1.4.1: Hot ductility curve for TWIP_1 (Fe-22Mn-3Al-3Si).

102

In the temperatures range 700-900°C, RA increases with decreasing temperature and the ductility has the lowest value of 28% at 900°C. In the temperature range 1000-1100°C, the RA increases with increasing temperature and it shows the highest RA of 80% at 1100°C. In the temperature range 1200-1300°C, the ductility decreases with increasing temperature and reaches zero at 1300°C. TWIP2 (Fe-22Mn-0.6C) The effect of deformation temperature on tensile behaviour of TWIP_2 is displayed that the strength decreases as temperature increases. However, tensile elongation to fracture decreases as temperature increases up to 900°C. Further increase in temperature results in increased elongation up to 1200°C. Hot ductility curve of TWIP_2 is shown in Fig. 3.1.4.2. In addition to the influence of deformation temperature, the effect of reheating temperature on the hot ductility was investigated. In most tests, the reheating temperature of 1250°C/2 min was used, but in addition two tests were carried out with the reheating temperature of 1150°C/2 min. Hot ductility curve for Type 304 austenitic stainless steel has been taken from Ref. [5] and included in Fig. 3.1.4.2 for comparison.

Fig. 3.1.4.2: Hot ductility curve of TWIP-2. Curve of AISI 304 is included for comparison.

Hot ductility curve of TWIP2 shows slight decrease of RA from 68% to 60% with increasing temperature from 700 to 900°C. This is similarly as the case for TWIP1 and is typical to the existence of ductility trough for C-Mn and microalloyed steels [6,7]. However, the minimum RA is quite high, about double compared to that of TWIP1, indicating excellent ductility even at 900°C. Further increase in temperature results in improvement of RA to 87% at 1200°C. At 1300°C, ductility decreased to the lowest value 44% in hot ductility curve. The differential thermal analysis of TWIP2 exhibits that the solidus temperature lies between 1320°C and 1340°C and the melting point is 1378°C. Thus the occurrence of local melting leads to decrease in hot ductility at 1300°C. As seen from the hot ductility curves in Fig. 3.1.4.2, the TWIP2 steel has even better ductility than that of the austenitic AISI 304. It can be attributed to more delayed dynamic recrystallization of AISI 304, occurring at higher temperatures than that of high-Mn TWIP steels [9]. TWIP3 (Fe-18Mn-0.6C-0.02Nb)

103

Hot ductility curve of TWIP_3 is shown in fig. 3.1.4.3. All hot ductility tests were carried out with the reheating temperature of 1150°C for 2 min. For comparison hot ductility curve of TWIP2 is included in the same plot. However, hot ductility tests of TWIP2 have been carried out with reheating temperature 1250°C for 2 min. The former RH provided austenitic structure with about 80 μm grain size and the latter 120 μm. It can be observed that both steels show similar behaviour with shifting curve of TWIP3 to higher RA values. TWIP3 displayed its highest RA value (98%) at 1100°C, while TWIP2 showed its highest RA value (87%) at 1200°C. It seems that finer grain size of TWIP3 promoted excellent hot ductility by enhancing dynamic recrystallization isolating micro-voids and consequently hindering the micro-voids coalescence and propagation.

Fig. 3.1.4.3: Hot ductility curve of TWIP3. Curve of TWIP2 is included for comparison.

TWIP4 (Fe-16Mn-1.5Al-0.3C) The hot ductility curve for TWIP4 is shown in fig. 3.1.4.4. For comparison, the hot ductility curve of TWIP2 is included. It is seen that both steels have almost same RA values up to 1000°C, regardless slightly lower ductility of TWIP4 at 800°C. Above 1000°C, TWIP4 exhibits excellent hot ductility up to 1300°C, with better values than those of TWIP2.

104

Fig. 3.1.4.4: Hot ductility curve of TWIP4. Curve of TWIP2 is included for comparison.

TWIP5 (Fe-22Mn-0.2C-0.2N) The hot ductility curves of TWIP5 and TWIP2 are shown in fig. 3.1.4.5. It can be seen that hot ductility behaviour of both steels is almost identical with minor differences in RA values. In TWIP5, RA decreases from 64% to 60% with increasing temperature from 700°C to 900°C, respectively. However, with further increasing temperature, hot ductility increases to the maximum value (95%) at 1200°C. At 1300°C, hot ductility has dropped to the minimum value 52%. In comparison with TWIP2, the hot ductility value of TWIP5 is slightly higher (RA = 95%) than that of TWIP2 (RA = 87%) at 1200°C. This may be attributed to the strengthening effect of nitrogen alloying. The influence of reheating temperature in TWIP5 is identical to that in TWIP2. With decreasing reheating temperature from 1250°C to 1150°C, the hot ductility increased from 82% to 98% at 1100°C.

Fig. 3.1.4.5: Hot ductility curve of TWIP5. Curve of TWIP2 is included for comparison.

105

Conclusions Task 3.1.4 The hot ductility curves of the TWIP steel variants show quite good high temperature performances. The hot ductility in some cases, such as for TWIP 3 and TWIP2 (RA>60% between 700-1200°C), is higher than that of the austenitic AISI 304. Task 3.1.5. Study of the influence of grain size in the tensile stress-strain behaviour of TWIP steels (Hall-Petch behaviour) The influence of grain size in the room-temperature tensile stress-strain behaviour of TWIP steels, TWIP1 and TWIP2 has been studied for two of the chemical compositions of the project at a strain rate of 10-3 s-1. TWIP1 is a 22% Mn, 3% Al, 3% Si, 0.01% C steel with a Duplex structure (less than 10% of ferritic phase). Its austenite partially transforms to ε-martensite upon plastic deformation (instead of deforming by twinning), because of its very low SFE. TWIP2 is a fully austenitic steel of composition 22% Mn, 0.06% C, that twins upon deformation. A Hall-Petch relationship has been found for the yield and flow stresses of both steels, with almost the same Hall-Petch slopes. The influence of strain rate on the Hall-Petch behaviour has been studied for steel TWIP2 by performing room temperature tensile tests at three strain rates, 10-3 s-1, 9.4 s-1 and 265 s-1. The yield stress or the flow stresses for a fixed tensile plastic strain obey Hall-Petch relationships. Surprisingly, the slopes are almost the same for both steels despite the strain-induced transformation occurring in TWIP1 steel or the twinning occurring in TWIP2 steel, as figure 3.1.5.1 shows for the yield stress. For true plastic strains ε < 0.005 the lamellae are practically absent from both steels, i.e., the yield stress should correspond to slip by dislocation glide.

y = 356.53x + 157.16R2 = 0.992

y = 320.16x + 176.82R2 = 0.9967

0100200300400500600700800900

1000

0 0.2 0.4 0.6 0.8 1 1.2 1.4

[D (μm)]-0.5

σ y (M

Pa)

TWIP 1TWIP 2

(a)

10-3 s-1

y = 356.53x + 157.16R2 = 0.992

y = 320.16x + 176.82R2 = 0.9967

0100200300400500600700800900

1000

0 0.2 0.4 0.6 0.8 1 1.2 1.4

[D (μm)]-0.5

σ y (M

Pa)

TWIP 1TWIP 2

(a)

10-3 s-1

Fig. 3.1.5.1: Hall-Petch plot of the tensile yield stress at 10-3 s-1 (flow stress for 0.002 tensile plastic strain) of steel TWIP1 and TWIP2.

106

0

100

200

300

400

500

600

0 0.03 0.06 0.09 0.12 0.15 0.1εp

KH

P (M

Pa μ

m1/

2 )

TWIP 1TWIP 2

10-3 s-1

KHP

(b)

0

100

200

300

400

500

600

0 0.03 0.06 0.09 0.12 0.15 0.1εp

KH

P (M

Pa μ

m1/

2 )

TWIP 1TWIP 2

10-3 s-1

KHP

(b)

0

100

200

300

400

500

600

700

0 0.1 0.2 0.3 0.4 0.5 0εp

KH

P (M

Pa μ

m1/

2 )

0.001 s-19.4 s-1265 s-1

Fig. 3.1.5.2: Hall-Petch slopes for the flow stress at fixed true plastic tensile strain values, steels TWIP1 and TWIP2 as a function of plastic true tensile strain, up to ε = 0.15.

Fig. 3.1.5.3: Hall-Petch slopes of steel TWIP2 at room temperature and three tensile strain rates, as a function of tensile plastic strain.

The influence of strain and strain rate on the Hall-Petch slopes is presented in figure 3.1.5.2-3. The great similitude of the Hall-Petch slopes of the two steels is evident. The Hall-Petch slope increases as strain increases, although the strain hardening of TWIP2 is larger than the hardening of TWIP1. The change observed is compatible with an effect of increasing heating of the samples during the tests at high strain rates. Such change is better seen in figure 3.1.5.3, showing the Hall-Petch slopes as a function of strain and strain rate. Numerical data corresponding to the tensile tests of TWIP2 are gathered in table 3.1.5.

dε/dt [s-1] D (μm) σy (ΜPa) σu (ΜPa) εu Θ0.1 (MPa) A%1.5 445 1624 0.42 3327 501.5 456 1586 0.37 3324 50.86.2 303 1481 0.44 2727 56.666.2 311 1377 0.38 2745 46.8312.2 260 1502 0.52 2400 73.312.2 254 1409 0.46 2394 66.719 250 1431 0.5 2340 69.219 225 1379 0.49 2368 66

40.7 213 1234 0.42 2732 56.640.7 209 1170 0.4 2651 52.661.5 407 1189 0.29 2757 40.31.5 407 1145 0.26 2760 35.3312.2 325 1032 0.31 2467 47.540.7 270 1021 0.39 2198 52.540.7 260 979 0.36 2337 51.661.5 530 1234 0.27 2413 45.212.2 365 1148 0.32 2555 53.3312.2 356 1120 0.32 2703 56.6640.7 300 1076 0.37 2078 62.5

9.4 s-1

265 s-1

10-3 s-1

Table 3.1.5: TWIP2, numerical results of tensile tests. The columns are: strain rate, grain size, yield stress, flow stress for maximum uniform strain, maximum uniform plastic strain, strain hardening rate at a strain of 0.1 and area total elongation. Influence of grain size on work hardening rate and ductility. For the TWIP steels the transition from the elastic domain to the near-linear plastic behaviour occurs very abruptly. The initial branch of the curves where the hardening rate decreases rapidly and makes a

107

knee is not a stage III (as it has been often quoted in the literature for twinning-deforming materials) but, merely the elasto-plastic transition in the polycrystal. The linear stage occupies almost all the uniform strain of the room temperature tensile tests and there is some influence of the grain size on its slope (fig. 3.1.5.4), being larger for smaller grain sizes. The slope ranges from 2250 MPa to 3200 MPa. After about 30% plastic elongation the hardening rate decreases, an indication that dynamic recovery is taking increasing importance and, consequently, an indication that the predominance of the twinning deformation that induces the “dynamic Hall-Petch effect” responsible of the TWIP initial linear stage is becoming less important.

1000

1500

2000

2500

3000

3500

4000

200 400 600 800 1000 1200 1400 1600σ (MPa)

d σ/d

ε (M

Pa)

12.2 μm40.7 μm

1.5 μm

19 μm

6.2 μm

Fig. 3.1.5.4: Work hardening rate vs. true flow stress of the TWIP 2 steel with composition 22% Mn-0.6% C. Tensile tests at room temperature and 10-3 s-1 for a equiaxed grain size in the range 1.5 μm < D < 50 μm. Annealing treatments. Grain growth kinetics Grain growth kinetics equation For studying the grain size dependence, the range of thermal treatments has been expanded with respect to the range covered for studying the annealing kinetics. The new results have allowed for a better description of the grain growth kinetics of the steel. The annealing treatments have covered a wide range from 10 min at 700ºC to 12.5 h at 1100ºC, all of them implying some grain growth after complete recrystallyzation of cold rolled structures (recrystallization is complete after 9 min at 700ºC). All the data are plotted in figs. 3.1.5.5a and 3.1.5.5b assuming an apparent activation energy of 363 kJ mol-1 . A reasonable fitting valid for the range covered by the experimental measurements (1.5 μm < D < 50 μm) is obtained:

( ) ( ) ( )⎟⎠⎞

⎜⎝⎛ −

⋅⋅=RT

molkJstD /363exp102516.2m 168894.3μ

108

y = 3.5407x - 38.315R2 = 0.8867

-40

-35

-30

-25

-20

0 0,5 1 1,5 2 2,5 3 3,5 4 4,5 5

ln (D*)

ln[e

xp(-Q

gg/R

T)t]

y = 3,8894x - 37,653R2 = 0,9023

-40

-35

-30

-25

-20

0 0,5 1 1,5 2 2,5 3 3,5 4 4,5 5

ln D

ln[e

xp(-Q

GG/R

T)·t]

Fig. 3.1.5.5a: D* grain size (mean linear intercept)

without considering twin boundaries, after annealing

time t (s) .

Fig. 3.1.5.5b: D id., with twin boundaries counted as grain boundaries.

Grain boundary analysis of annealed samples (EBSP-OIM) The new data available the twin boundary fraction of high angle boundaries can be added to the graph of twin boundary fraction vs. grain size, fig. 3.1.5.6. After an initial very rapid growth, the twin boundary fraction increases slowly as the grain size increases above 10μm.

Twin Boundary Fraction vs Grain Size D* (twin boundaries not counted)

0

0,05

0,1

0,15

0,2

0,25

0,3

0,35

0,4

0,45

0,00 10,00 20,00 30,00 40,00 50,00 60,00 70,00 80,00

Grain size D* (μm)

Twin

bou

ndar

y fr

actio

n

Fig. 3.1.5.6: Twin Boundary Fraction vs. Grain size D* (twin boundaries not counted).

Conclusions of task 3.1.5 The yield stress for 0.2% plastic elongation clearly shows a grain size dependence. The Hall-Petch constants found for TWIP2 steel grade is =HPK 356.5 MPa μm1/2. The results of this task allowed to achieve a better description of the grain growth kinetics of the TWIP2 steel. It is worthy to note that for fine grain size (<10μm) the twin boundary fraction is

109

significantly lower than for larger grain size. This means that the TWIP effect is dependent on grain size and the best performances are relevant to a grain size in the range 15-30μm. Task 3.2 Bending fatigue tests to determine the fatigue strength and cyclic softening/hardening behaviour and to analyse the crack initiation/propagation stages Experimental Cold rolled samples were annealed at different conditions to produce specimens with different grain sizes respectively of 4.5 μm, 13 μm, 32 μm and 55 μm. Fatigue tests were carried out at room temperature under normal atmospheric condition using a Schenk flexural bending fatigue machine driven at a frequency of 23 Hz with a zero mean stress. The standard hourglass-shape of flat fatigue specimens with a thickness of 2 mm were used in all tests. In order to verify the properties of the surface features of the fatigued steels, cyclic damage and crack formation were examined using three techniques: an optical microscope and a field emission gun scanning electron microscope FEG-SEM (Carl Zeiss Ultra plus) applying either a electron channeling contrast (SEM-ECC) imaging or SEM-EBSD. Results Fig. 3.2.1 shows the experimental fatigue data in the form of stress amplitude - number of cycles to failure plot (S-N curve). From the data, the fatigue limit of all investigated TWIP steels, corresponding no failure in 2*106 cycles, is about 400 MPa. For comparison, the S-N curves of two commercial austenitic stainless steels, Type 301LN and 316L, taken from Ref. [1] and a high-strength (Rm = 821 MPa) TRIP steel, taken from Ref. [2], are also inserted. The fatigue limits of 301LN, 316L and TRIP steels, determined in reversed plane bending, are about 350, 300 MPa and 400 MPa, respectively. Hence, in absolute values the fatigue limit of the TWIP steels are considerably higher than those of annealed austenitic stainless steels and equal to that of the 780-grade TRIP steel. It is known that the fatigue limit of steel is roughly related to its tensile strength, so that the ratio fatigue limit/tensile strength (FL/TS) is between 0.4-0.6, see e.g. [3]. It was seen that the ratio FL/TS of the TWIP steels varies between 0.42 and 0.48 (for AISI 304 is typically 0.42).

Fig. 3.2.1: S-N curves of the investigated TWIP steels and those of 301LN and 316L and high-strength TRIP steel are included for comparison [1,2].

110

Crack nucleation and propagation A few specimens were examined after the 25, 50 and 100 percents of the expected fatigue life to follow the fatigue crack initiation and propagation. It was seen that many cracks emerge from the ordinary grain boundaries (intergranular cracking) as well as from annealing twin boundaries (twin cracking). To investigate the type of crack propagation in relation to the grain structure, a fatigued sample was cycled at the stress level of 500 MPa until failure. A high density of intense slip bands in the region near the main fatigue fracture was observed, as shown in Fig. 3.2.2. In addition, high densities of intergranular cracks and annealing twin boundary cracks could be observed.

Fig. 3.2.2: The final fatigued structure of TWIP (Fe-22Mn-0.6C) steel after failure.

Based on the microscopy observations on tensile strained TWIP steels, mechanical twinning is an active deformation mechanism under monotonic loading. By SEM-EBSD technique, mechanical twins can be identified based on their orientation difference in relation to the surrounding matrix. The SEM-EBSD samples analysis indicates that no mechanical twinning takes place in fatigue of TWIP steels, but only intense slip bands are formed as a result of cyclic stressing. This means that the TWIP effect has no essential role in cyclic deformation of these steels. Influence of grain size The results of flexural bending fatigue tests at the stress amplitude level of 500 MPa are shown in Fig. 3.2.3. The figure indicates that a significant beneficial effect on the fatigue life is obtained by refining the grain size. Moreover, it can be seen that the fatigue limit of the structure with the smallest grain size of 4.5 μm equals or is higher than 500 MPa, because no failure occurred within 2.2 million cycles.

111

Fig. 3.2.3: Effect of the grain size on the fatigue life of the TWIP2 steel (Fe-22Mn-0.6C). All tests were carried out at the amplitude of 500 MPa. Conclusions task 3.2 The fatigue behavior of three high-Mn TWIP steels, with slightly different Mn contents (between 16 and 22 wt.%) and Nb or Al alloying were investigated using reversed bending loading and examining the cyclic damage features on surfaces. The main conclusions can be drawn as follows: (1) Fatigue behavior of three TWIP steels is quite identical. Fatigue stress limit (the cyclic life beyond 2x106 cycles) is well above their yield strength values. The ratio of fatigue limit/tensile strength is 0.42-0.48 that is quite a similar value as commonly observed for various carbon steels and for Types 301LN and 316L austenitic stainless steels. (2) During cyclic loading, planar slip bands are formed in an early stage of fatigue life consisting of extrusions and intrusions. With continuing cycling, the slip bands intersect with grain boundaries as well as annealing twin boundaries producing local strain concentrations that induce microcracks at these boundaries. (3) Fatigue crack embryos nucleate at an early stage of fatigue life (≈25%) at sites of intersections of slip bands and grain boundaries as well as annealing twin boundaries. (4) Crack propagation takes place along slip lines, grain and twin boundaries but the overall path is mainly transgranular in its character. In this stage, ductile striations are formed on fracture surfaces. (5) Microcracks link and propagate readily along grain boundaries indicating some degree of inherent grain boundary weakness, as suggested in the literature for austenitic high-Mn steels. (6) Neither mechanical twins nor ε-martensite are formed during cyclic loading in the investigated TWIP steels, so that the TRIP or TWIP effects seem to play no role in the course of high-cycle fatigue. (7) The degree of cyclic hardening revealed by hardness is strongly dependent on the grain size decreasing with refined grain size. (8) Refinement of the grain size improves significantly the fatigue strength of the 0.6C-22Mn TWIP steel.

Task 3.3 –Evaluation of impact strength by means of Charpy tests at different temperature In this task 3 the absorbed energy of the five compositions being studied was determined by charpy tests in four different temperatures. The impact tests were performed according to the test standard NP EN 10045-1, three specimens were taken from each material. It is possible to see the results in table 3.3.1 and the graphic on figure 3.3.1.

112

Material Temperature [ºC] Absorbed energy[Joules/cm2] Average Absorbed energy[Joules/cm2] -50 150.00 -50 175.00 -50 150.00

158.33

0 166.67 0 133.33 0 141.67

147.22

20 108.33 20 100.00 20 100.00

102.78

150 125.00 150 125.00

TWIP 1

150 141.67

130.56

-50 115.38 -50 115.38 -50 141.03

123.93

0 102.56 0 89.74 0 89.74

94.02

20 102.56 20 76.92 20 89.74

89.74

150 115.38 150 115.38

TWIP 2

150 96.15

108.97

-50 116.67 -50 138.89 -50 111.11

122.22

0 106.38 0 127.66 0 111.70

115.25

20 72.22 20 72.22 20 83.33

75.93

150 111.11 150 127.78

TWIP 3

150 100.00

112.96

-50 86.21 -50 68.97 -50 68.97

74.71

0 77.59 0 77.59 0 77.59

77.59

20 51.72 20 51.72 20 51.72

51.72

150 51.72 150 60.34

TWIP 4

150 60.34

57.47

-50 75.76 -50 83.33 -50 90.91

83.33

0 75.76 0 83.33 0 83.33

80.81

20 60.61 20 60.61 20 68.18

63.13

150 75.76 150 75.76

TWIP5

150 75.76

75.76

Table 3.3.1 – Charpy

113

0

20

40

60

80

100

120

140

160

180

-50 0 20 150Temperature (ºC)

Abs

orbe

d En

ergy

(J/c

m2)

TWIP 1TWIP 2TWIP 3TWIP 4TWIP 5

Figure 3.3.1 –Charpy tests results for the five TWIP compositions.

Conclusions task 3.3 Comparing the Charpy values of TWIP variants arise that TWIP1 reveals the best results at all tested temperature. TWIP2 and TWIP3 are slightly better than TWIP4 and TWIP5. Qualitatively the Charpy energy behavior of TWIP steels as a function of the temperature is quite similar to stainless steels such as AISI304. In fact the Charpy energy remains of the same order from +150°C down to -50°C, with a slight increase from RT down to -50°C.

Task 3.4 Plain strain compression tests Plain strain compression tests were completed and the stress-strain curve relevant to TWIP2 is shown in the fig. 3.4.1. The work hardening of TWIP2,3,4,5 steels is significantly higher than stainless steels (AISI304). At a true strain of 0.5 the difference in terms of stress is about 38%. TWIP1 reveals a softer behavior quite different to other TWIP2-3-4-5 steel grades and is more similar to AISI304 probably due to presence of softer ferrite in microstructure. In table 3.4.1 are reported the stress corresponding at a strain of 0.5.

Fig. 3.4.1: Plain strain compression tests results on TWIP2 variant.

Table 3.4.1: Stress after a true strain of 0.5.

114

Conclusions Task 3.4 The work hardening of TWIP2,3,4,5 steels is significantly higher than stainless steels; at a true strain of 0.5 the difference in terms of stress is about 38%. The above result implies difficulties in cold rolling in terms of loads, number of passes or reduction ratio at each stage. Depending on the hot strip thickness, an intermediate annealing treatment (two step cold rolling process) could be necessary for obtaining the aimed final cold strip thickness (<2.0 mm).

115

WP 4 – Basic characterisation of application properties: formability, weldability and coating ability Task 4.1 Formability characterisation by means of Erichsen test and High-velocity forming tests. Task 4.1.1 Erichsen test The Erichsen cupping test is a simple stretch forming test of a sheet clamped firmly between blank holders to prevent in-flow of sheet material into the deformation zone. The punch is forced onto the clamped sheet with tool contact (lubricated, but with some friction) until cracks occur. The depth (mm) of the punch is measured and gives the Erichsen index 'IE', standardised under DIN 50101. In the table 4.1.1.1 and figure 4.1.1.1 it can be seen the results of the Erichsen tests on the TWIP steels. These results demonstrate that TWIP steels are characterized by quite good forming ability.

Material Nº of tests

Thickness [mm] IE [mm] Average 1.28 12.4 1.28 12.3 TWIP 1 3 1.28 12.5

12.40

2.12 11.7 1.76 11.9 TWIP 2 3 2.11 12.4

12.00

2.18 10.4 2.18 10.3 TWIP 3 3 2.18 10.4

10.37

1.23 10.9 1.21 11 1.21 11.5 TWIP 4 4

1.18 10.8

11.05

1.37 10.9 1.34 11.3 TWIP 5 3 1.35 11.3

11.17

Table 4.1.1.1: Erichsen tests results.

9.00

9.50

10.00

10.50

11.00

11.50

12.00

12.50

13.00

TWIP 1

TWIP 2

TWIP 3

TWIP 4

TWIP 5

Erichsen test results

Fig. 4.1.1.1: Graphic representation of the Erichesen tests results for the five TWIP compositions.

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Task 4.1.2 – High-velocity forming tests During the project the formability of TWIP-steels at a high deformation rate [1] has been investigated and compared to the formability results using conventional Erichsen tests. Cold rolled sheets of high-Mn TWIP steels with 1.5 mm thickness were prepared and supplied by CSM for high-velocity forming testing. In the present work, a special high-speed Erichsen testing was employed using an electro-hydraulic impulse forming unit (at the Stainless Steel Studio, Tornio, Finland). In this instance, the Erichsen tests were carried out using a hemispherical punch with a diameter of 20 mm at room temperature to investigate the stretch formability in accordance with ASTM 643-84 [2]. Some grease was set as a lubricant between the punch and the sheet. The Erichsen index (IE) is the value of the punch penetration before cracking of the sheet. However, in a high-velocity Erichsen test, the punch cannot be stopped in the instance of cracking, but the displacement of the stroke must be adjusted beforehand. In the used equipment, the displacement can be adjusted at the intervals of 0.25 mm, so that the actual IE value is something between the two dome heights (stroke displacements) of highest crack-free and the first fractured Erichsen cups. However, in present tests the displacements varied much more and problems appeared to adjust it precisely. Therefore, intervals in the displacements obtained in the tests were often much longer. To analyze the strains on the sheet samples occurred in forming, a circle grid was etched by applying an electrochemical marking method on the samples to create arrays of overlapping circles (the diameter 2 mm). Results Figure 4.1.2.1 compares IE values (solid dome - first cracked dome heights) obtained for the investigated TWIP steels in high-speed Erichsen testing in relation to the total elongation obtained in quasi-static tensile tests. The figure also contains the corresponding data for AISI 304 austenitic stainless steel [5]. It can be seen that variation among the steels is quite small. Generally, IE is higher if the elongation is higher, and the best values are obtained for TWIP_2. However, better IE values were obtained at high-speed Erichsen testing.

Fig. 4.1.2.1: Comparison between the Erichsen index (IE) in high-speed (and conventional) testing and the total elongation (A) from quasi-static tensile tests [3] and that of AISI 304 [5].

In literature is reported that [7] the formability of a TWIP (Fe-15Mn-3Al-3Si) steel having a very pronounced TRIP effect at different strain rates was enhanced with increasing strain rate from 2x10-3 to 10-1 s-1. The explanation is that due to the adiabatic heating of the deformed sheets (heated up to 70°C)

117

at higher strain rates, the martensitic transformation was retarded, and consequently the ductility was improved by more intense twinning. In contrast the ductility of TWIP (Fe-25Mn-3Al-3Si) steel was found to decrease with increasing strain rate from 10-3 to 10-1 s-1 in tensile tests [8]. This can be attributed to the adiabatic heating of the sample during the tensile test increasing the SFE and resulting in decreasing density of mechanical twins. However, with increasing the strain rate up to 1000 s-1, the total elongation increases again reaching values above that in tensile testing at the strain rate of 0.1 s-1. It is possible to make comparison between the elongation in dynamic tensile tests ( 1000≈ε& s-1) using the Hopkinson split bar method and that in quasi-static tensile tests [3]. Then, for example, for TWIP_2 the total elongation is 83% under quasi-static tension and 80% under high-speed tensile testing, i.e. the total elongation in these two cases is almost equal (note: the specimen shapes are very different). In the conventional Erichsen testing according to the DIN 50101, the typical strain rate can be estimated to be order of 0.01 s-1. (For instance, no 0.4 cylinder speed is 0.56 mm/s). In high-speed Erichsen testing, where the speed of impact front is about 200 m/s, the strain rate can be order of 5 s-1 (dome height 10 mm; at 200m/s it takes 0.05 s. Strain is 0.26, hence, the strain rate = 0.26/0.05s = 5/s), i.e. about 500 times higher. Conclusions of task 4.1.2 Comparing quasi-static IE values with high-speed Erichsen testing arise that the total elongation variation is quite small. Generally, IE is higher if the elongation is higher, and the best values are obtained for TWIP_2. However, better IE values were obtained at high-speed Erichsen testing for all TWIP grades except for TWIP1 grade. Task 4.2 – Laboratory coating tests and coating layer characterisation Eight samples (four TWIP steel studied materials) were zinc coated using the batch hot-dip galvanizing process. The batch hot-dip galvanizing process, also known as general galvanizing, produces a zinc coating on steel products by immersion of the material in a batch of liquid zinc, which then forms a durable bond to the iron at the atomic level. Before the coating is applied, the steel is cleaned to remove all oils, greases, soils, mill scale, and rust. The cleaning cycle usually consists of a degreasing step, followed by acid pickling (heated sulphuric acid) to remove scale and rust, and fluxing, which inhibits oxidation of the steel before dipping in the molten zinc. In this project two different pre-treatment preparations were used: The conventional one (chemical picking - marked DQ) and the replacement of the two first chemical cleaning steps by mechanical pickling (marked DM). The reason for this variation is related to the observation of possible hydrogen contamination of the samples in the first situation. After galvanization, i.e., after the deposition of the Zn based coating layer, the samples were characterized. The visual control of sample’s surfaces revealed a general uniformity of the coating layer (eventually, in two cases, with an over-thickness of coating that was confirmed in coating layer thickness measurement) with no bulges. The samples were then subjected to a test to evaluate the adhesion of their galvanized layer. The galvanization adhesion test was made through the cross test method according to standards NF A91-121 and NP EN ISO 1461. As shown in figures 4.2.1-2 for TWIP2 (the results for other TWIP grades were similar) all the tests were acceptable according to the standard since no detachment of any small square was verified. No noticeable differences were detected between the two pre-treatments in the galvanized adhesion layer test.

118

Fig. 4.2.1: sample TWIP 2 DM. Fig. 4.2.2: sample TWIP 2 DQ.

The thickness of the coating layers was determined through its measurement in a transversal section of the samples. The results obtained are presented in the table below.

Samples ref. Thickness (μm) Averag

e

Standard

deviation

TWIP 2 DM

119,2 95,7 140,

5 126,7

140,4

105,4

166,0

140,4

138,3

106,4 127,9 21,4

TWIP 2 DQ

134,8

138,3

130,9

107,5

110,7

120,2

109,6

106,4

112,8

112,8 118,4 12,0

TWIP 3 DM

113,8

134,1

117,1

138,3

126,7

109,6

114,9

110,6

121,3

101,1 118,7 11,5

TWIP 3 DQ

103.4 97.7 104.

6 117.1

114.8

122.8

117.1

111.4 95,5 101.

2 108.6 9.3 TWIP 4 DM

140,4

128,7

121,3

131,9

124,5

130,9

128,7

156,4

125,5

140,6 132,9 10,4

TWIP 4 DQ 95,0 100,

1 93,8 100,0 98,9 104,

3 119,2

117,0

121,9

126,8 107,7 12,2

TWIP 5 DM

136,2

125,5

124,5

117,0

124,5

123,4 94,7 104,

3 112,9 98,0 116,1 13,4

TWIP 5 DQ

180,9

211,7

243,7

267,2

239,4

210,7

264,6

186,6

191,5

204,5 220,1 31,6

Table 4.2.1: Thickness of coating layers.

The coating’s layer thickness measurement shows that sample TWIP 5 DQ presents the highest value but, with this exception no noticeable difference exists between the two sets of samples obtained with two different pre-treatments. The most common reason for the different zinc coating thickness, when using the same galvanization conditions, is the chemistry of the steel pieces. There are two elements of steel chemistry which most strongly influences the final thickness and appearance; silicon and phosphorous. Both silicon and phosphorous promote coating growth.

119

The recommended silicon composition is either less then 0.04% or between 0.15% and 0.25%. Any steels outside these ranges are considered reactive steels and are expected to form zinc coatings that tend to be thicker resulting of the rapid zinc-iron intermetallic growth. In this project the steels under analysis have silicon contents near the maximum limit (between 0.2 and 0.22%); this could be a reason for the different thickness values in the galvanized samples. With the objective of determining the galvanisation composition EDS analysis in the pure zinc layer of the zinc coating was performed. As reported in the following table the used Zinc bath presented small additions of Phosphorus and Aluminium.

Weight % Element Weight % Sigma

Zn K 97.68 0.14 Al K 0.91 0.09 P K 0.75 0.07

TABLE 4.2.2: EDS ANALYSIS.

All the samples were metallurgically characterized after the galvanization process. In general, the typical galvanization structure was observed. The batch hot-dip galvanized coating consists of a series of zinc-iron alloy layers with a surface layer of almost pure zinc. The alloy layer is as much as 50% of the total thickness and it consists of two or more distinct zinc/iron layers. Each layer has a specific amount of iron and zinc. A representative photomicrograph of the alloy layer that forms while the steel is immersed in the bath is shown in Figure 4.2.3.

Fig. 4.2.3: Sample TWIP 5 DM (200X).

However some quality defects were observed. In all the chemical picking samples large black spots in the Zn coating defined as gas pockets can be observed. In the fig. 4.2.4-4.2.7 are shown some example of such defects relevant to TWIP2 and TWIP5. This could be an effect related to hydrogen absorption by steel overpickled before galvanizing. Hydrogen absorbed by overpickled steel, if not completely expelled at the temperature of the molten zinc bath, may be trapped in the coating, causing gas pockets and blisters trapped in the alloy layers. This defect was not present in the mechanical pickling samples since in this process sulphuric acid was not used.

120

Fig. 4.2.4: Sample TWIP 2 DM (200X). Fig. 4.2.5: Sample TWIP 2 DQ (500X).

Fig. 4.2.6: Sample TWIP 5 DM (200X). Fig. 4.2.7: Sample TWIP 5 DQ (500X).

Other defect was the lack of uniformity and reproducibility in the quality of the interface zinc/steel. Even after mechanical or chemical surface preparation it was impossible to eliminate all the superficial defects resulting from the samples production. For this reason in the same sample it was possible to find areas with a very good quality of the interface between the steel and the coating and others with adhesion defects caused by strip surface irregularities (cracks). In the TWIP steel material it was impossible to etch the area near the zinc coating due to the zinc coating cathodic effect. Nevertheless the galvanising temperature used (451ºC) do not seem to be enough to promote any phase transformation or zinc inter-grain diffusion on the TWIP material. Conclusions of task 4.2

• The visual control of sample’s surfaces revealed a general uniformity of the coating layer. • The galvanization adhesion test was made all the tests were acceptable according to the

standard since no detachment of any small square was verified. No noticeable differences were detected between the two pre-treatments in the galvanized adhesion layer test.

• All the samples after the galvanization process revealed the typical galvanization structure. The batch hot-dip galvanized coating consists of a series of zinc-iron alloy layers with a surface layer of almost pure zinc. The alloy layer is as much as 50% of the total thickness and it consists of two or more distinct zinc/iron layers. Each layer has a specific amount of iron and zinc.

• Some quality defects were observed. In all the chemical picking samples large black spots in the Zn coating defined as gas pockets can be observed. This is the effect of hydrogen absorption by steel overpickled before galvanizing. This defect was not present in the mechanical pickling samples since in this process sulphuric acid was not used.

121

Task 4.3 – Characterization of Weldability TWIP laser welding During the project it was decided to concentrate the welding tests on laser welding. This decision was taken having in mind that this is in fact a very common method in the automotive industry for welding thin sheet steel (<2.0 mm). Some preliminary tests with resistance spot welding were performed and were sufficient to conclude that the use of these steels in spot welding is adequate. The laser welding was the process chosen to be fully tested because the application of laser in the automotive industry has increased in the last years replacing several times resistance spot welding. The following study was focused on the understanding of TWIP steel behaviour in joining, particularly in laser welding. Four high-Mn based Fe-alloys with austenitic structure developed by alloying with carbon, aluminium and silicon elements aiming to obtain different grades of TWIP steels were studied. TWIP 1 was not characterized because it cannot be considered a TWIP steel due to its mixed microstructure (ferrite+austenite). Laboratory cast ingots were homogenized at 1200 ºC for 1 hour to remove the segregation of the alloying elements, especially those Mn based. Subsequently, they were hot rolled to 2 mm thick bands in a laboratory rolling mill. Finally, the bands have suffered a heat treatment at 1000 ºC for 5 min. The used laser system was a diode laser unit, with a maximum power of 3KW working at a wavelength of 940 nm and 808 nm. This laser is connected to a KUKA robot (model R125F/2) An in-depth detailed observation of the diode laser welded joint to explain how weldability assessment was done and confirm the absence of the heat affected zone as well as welding defects. Some conditions of the metallurgical structure were emphasized as important for the welding final structure. Four of the five TWIP materials welded with the following laser welding conditions were fully characterized Diode laser welded joints were made using 2 mm-thick plates in a butt joint configuration. The optimisation of the studied welds using the different steel compositions has been carried out following a sequence of trials. The criterion for optimisation was the full penetration and absence of defects. The final selection of the welding procedure has been based on the technical experience, visual analysis and X-ray tests. A full metallurgical characterization of each welded sample was made after the process optimisation. Table 4.3.1 summarizes the optimal laser welding parameters for each steel composition.

Gas protection Thickness With Power Speed

Argon [l/min] [mm] [mm] [W] [mm/s] 2 15 2.0 180 1500 10 3 15 1.5 140 1500 12 4 15 2.0 90 1600 15 5 15 2.1 140 1500 10

Table 4.3.1: Optimal laser welding parameters for each “TWIP” steel composition.

Some observations made during the welding tests are considered relevant: - high quantity of fumes (probably Mn evaporation) was observed especially in TWIP4; - oxide contamination of the welding surface existed even with inert gas atmosphere (argon). Vickers hardness tests were performed to evaluate the mechanical properties in the welded joints. Measurements were extended to the entire thickness along the centreline of the WZ (welded zone) with a load of 2.5 Kg.

122

Visual and Macrographic Analysis The samples submitted to visual and macrographic inspections confirmed the possibility of having sound welds with a good repeatability for all the studied compositions. In all cases, the penetration was complete along all the plate thickness, and the welds were free from cracks or other defects. Some of the detected defects were a consequence of the non-homogeneity of the sheets thickness and difficulty in joint positioning. Figure 4.3.1 and 4.3.2 illustrates the diode laser weld performed on the TWIP3 steel where we can emphasize the presence of surface oxide contamination even with gas protection and the full penetration of the welding. This oxide contamination proves not to have any consequences on the welding quality as we can see later on this article.

Fig. 4.3.1: Visual observation of the front surface of the TWIP3 diode laser weld.

Fig. 4.3.2: Visual observation of the root surface of the TWIP3 diode laser weld.

Three examples of typical cross sections of the welded joints are shown in figures 4.3.3 to 4.3.5. These allow to observe different regions of the welded joint, namely the base and welded materials and to conclude that no visible heat affected zone exists.

Fig. 4.3.3: Low-magnification light optical micrograph from TWIP2 steel showing a cross

section of the welded joint.

Fig. 4.3.4: Low-magnification light optical micrograph from TWIP3 steel showing a cross

section of the welded joint.

0.5 mm

0.5 mm

123

Fig. 4.3.5: Low-magnification light optical micrograph from TWIP4 steel showing a cross

section of the welded joint.

Fig. 4.3.6: Micrography of TWIP2 steel base material (500X).

Optical Observation All the four TWIP base materials present a microstructure representative of this steel grade. It was a fully austenitic structure with globular and fully recrystallised grains. Figure 4.3.6 presents TWIP2 base material microstructure. Figure 4.3.7-8 show the microstrutural variation in TWIP3 sample from the base material to the welded zone, making clear the transition zone. The absence of an “actual” heat affected zone derives from the low heat input of the laser process. Figure 4.3.9 presents, as an example, the welded zones microstructures as observed in TWIP2 samples. The typical microstructure observed in most of the welded samples is constituted of an austenitic structure superimposed in a cellular-dendritic solidification structure. However, figure 4.3.9 pertaining to the welded material of TWIP2 sample shows the presence of ε-martensite. The zone where the ε-martensite appeared corresponds nearly to the top face of weld axis, which is a zone of higher cooling rate.

Fig. 4.3.7: Micrography of TWIP3 welded

sample showing the transition base material- welded material (200X).

Fig. 4.3.8: Micrography of TWIP3 welded joint (200X).

The appearance of ε-martensite is not desirable due to the significant reduction that it implies in ductility. As the welding conditions were quite similar in all experiments then some reason should exist to justify the observed behaviour in TWIP2 sample.

0.5 mm

124

Fig. 4.3.9: Micrography of TWIP2 welded material showing the presence of ε-martensite (200X).

As referred before, the quite high ductility behaviour of TWIP steels is associated to the stacking fault energy (SFE) which is the ignitor of the twinning mechanism. A possible explanation of the occurrence of martensite in the welded zone could be related to a change in chemical composition in particular the content of Mn and C during welding. Although the difference in terms of composition could be small, the risk of reaching the austenite instability (deformation induced) could be the cause of ε-martensite formation.

ASTM Grain

Size 2 6.994 3 9.895 4 7.754 5 7.379

Table 4.3.2: ASTM grain size measurements using the circular interception technique according with the ASTM E 112 standard.

Vickers Hardness Measurements Figure 4.3.10 shows the Vickers hardness distribution through the weld transverse cross section in the different welded TWIP steel samples. The measurements were performed in three longitudinal rows (top, middle and bottom weld metal levels). The base materials hardness of TWIP3, TWIP4 and TWIP5 steels present similar values, and all the weld zones present acceptable values which indicate that the mechanical properties will be maintained. By contrast, TWIP2 shows lower hardness values in the base materials which seem to agree with the measured higher average grain size and, in addition, an abrupt increase in the weld zone hardness is observed. The later is in agreement with the observed presence of ε-martensite.

125

150160

170180

190200

210

BM1 WM BM2Welding Zones

Vick

ers

Har

dnes

s TWIP 1

TWIP 2

TWIP 3

TWIP 4

Fig. 4.3.10: Vickers hardness measurements.

Based on the experimental laser welding analysis of four TWIP steel compositions, weldability of these materials was established as well as the main influencing factors. The work allowed to conclude that the joining of TWIP steels is easily achievable by laser welding. The factors that could influence the steel welding behaviour and its quality are intrinsically related to the steel composition and crystallographic structure. All the factors influencing stacking fault energy of base materials are the key for a sound joining of TWIP steels. Dissimilar Welds Joining dissimilar materials became inevitable in engineering industries for both technical and economic reasons. The adoption of dissimilar-metal combinations provides possibilities for the flexible design of the product by using each material efficiently, i.e., benefiting from the specific properties of each material in a functional way. For this reason a study of the weldability of TWIP steels in dissimilar welds was carried out. After receiving from the industrial partners more TWIP material all the samples for welding were prepared to eliminate the bad superficial conditions of the as received material. High strength steels to perform the dissimilar welds were acquired (DOCOL 600 DP and DOCOL 800 DP). Docol DP are cold reduced dual phase steels subjected to special heat treatment in the continuous annealing line, which produces a two-phase structure in which the ferrite that imparts unique forming properties is one of the phases, and martensite that accounts for the strength is the other. The microstructure of DOCOL 800DP steel can be seen in figure 4.3.11 .

126

Figure 4.3.11 – Docol 800DP steel microstructure taken in a scanning microscope (X500).

The following tables describe the high strength steels characteristics (chemical composition and mechanical properties). In the DP steels, the difference between the yield strength and the tensile strength is wide, which means that the steel has good ability to distribute the strain during working. Typical applications of this material are safety components in cars like door beams, bumper reinforcement and seat tracks.

Steel grade C [%] Si [%] Mn [%] P [%] S [%] Alt [%] Nb [%]

Docol 600DP 0.10 0.20 0.80 0.010 0.002 0.040 0.015 Docol 800DP 0.13 0.20 1.50 0.010 0.002 0.040 0.015

Table 4.3.3 – Docol DP chemical composition. Due to serious malfunctions in ISQ diode laser it was necessary to perform the welds outside ISQ in a CO2 laser. The steels combinations were the ones described in table 4.3.4.

TWIP 2 TWIP 3 TWIP 4 TWIP 5 DOCOL 600 TWIP 2/600 TWIP 3/600 TWIP 4/600 TWIP 5/600 DOCOL 800 TWIP 2/800 TWIP 3/800 TWIP 4/800 TWIP 5/800

Table 4.3.4 – Steels combinations welded with CO2 laser.

CO2 laser welded joints were made using a tailored blank joint configuration. This option was taken to give an idea of the potentialities of this material in this kind of application. A TWIP steel plate of 3 mm was welded to a Docol DP steel plate of 1.5mm as illustrated in the figure 4.3.12.

Martensite

Ferrite

127

Fig. 4.3.12: Tailored blank weld configuration used in the welding tests.

The welds were processed using a 1.7 kW CW-CO2 laser consisting of the laser source, a beam delivery system, and a CNC worktable. The laser parameters used were: laser power 1500 W, speed 25 mm/s and frequency 50000Hz. Figure 4.3.13 illustrates the CO2 laser weld performed on the TWIP4-600 steels combination where we can emphasize the absence of superficial defects and the full penetration of the weld.

Fig. 4.3.13: Visual observation of the front surface of the TWIP4 – Docol 600DP CO2 laser weld. The samples submitted to macrographic inspections confirmed the possibility of having sound welds with a good repeatability for most of the studied compositions. In TWIP 2 and TWIP 3 welds, the penetration was complete along all the plate thickness, and the welds were free from cracks or other defects. However, in TWIP 4 and TWIP 5 welds with some porosity were detected. The porosity was present both in the welds with DOCOL 600 DP steel and the ones with Docol 800DP steel. Four examples of typical cross sections of the welded joints are shown in figures 4.3.14 to 4.3.17. It allows to observe the different regions of the welded joint, namely the base and welded materials and to conclude that heat affected zones were only present in Docol steel sides.

Docol DP steel TWIP steel

Welding position

Dissimilar weld

128

Fig. 4.3.14: Low-magnification light optical micrograph showing a cross section of the welded joint - TWIP2 and Docol 600DP steels.

Fig. 4.3.15: Low-magnification light optical micrograph showing a cross section of the welded joint - TWIP2 and Docol 800DP steels

0.3 mm 0.3 mm

129

.

Fig. 4.3.16: Low-magnification light optical micrograph showing a cross section of the welded joint - TWIP3 and Docol 600DP steels.

Fig. 4.3.17: Low-magnification light optical micrograph showing a cross section of the welded joint - TWIP3 and Docol 600DP steels.

Vickers hardness tests were performed to evaluate the mechanical properties in the welded joints. Measurements were extended to the entire thickness along the centreline of the WZ (welded zone) with a load of 0.2 Kg. As a reference figure 4.3.18 represents the hardness distribution curve of a Docol 800 DP laser weld.

0.3 mm 0.3 mm

130

150

200

250

300

350

400

450

500

-8 -6 -4 -2 0 2 4 6 8Distance (mm)

Har

dnes

s (H

V5)

Docol 800DP

Fig. 4.3.18: Hardness distribution curve of a Docol 800 DP laser weld.

Figures 4.3.19 to 4.3.24 show the Vickers hardness distribution through the weld transverse cross section in the different welded TWIP - Docol steels combinations samples. The measurements were performed in one longitudinal row.

Fig. 4.3.19: Graphic representation of the hardness measurements performed on the TWIP 2 – DOCOL

600DP weld.

150

200

250

300

350

400

0 2 4 6 8 10 12 14 16 18 20 22 24 26

Har

dnes

s (H

V 0,

2)

TWIP 2 - 600

DDooccooll 660000 DDPP TTWWIIPP 22 ZZTTAA WWeelldd zzoonnee

131

Fig. 4.3.20: Graphic representation of the hardness measurements performed on the TWIP 2 – DOCOL

800DP weld.

Fig. 4.3.21: Graphic representation of the hardness measurements performed on the TWIP 4 – DOCOL

600DP weld..

Fig. 4.3.22: Graphic representation of the hardness measurements performed on the TWIP 4 – DOCOL

800DP weld.

150

200

250

300

350

400

0 2 4 6 8 10 12 14 16 18 20 22 24 26

Har

dnes

s (H

V 0,

2)

TWIP 4 - 600

DDooccooll 660000 DDPP TTWWIIPP 44 ZZTTAA WWeelldd zzoonnee

150

200

250

300

350

400

450

0 2 4 6 8 10 12 14 16 18 20 22 24 26

Har

dnes

s (H

V 0,

2)

TWIP 4 - 800

DDooccooll 880000 DDPP TTWWIIPP 44 ZZTTAA WWeelldd zzoonnee

150

200

250

300

350

400

0 2 4 6 8 10 12 14 16 18 20 22 24 26

Har

dnes

s (H

V 0,

2)

TWIP 2 - 800

DDooccooll 880000 DDPP TTWWIIPP 22 ZZTTAA WWeelldd zzoonnee

132

Fig. 4.3.23: Graphic representation of the hardness measurements performed on the TWIP 5 – DOCOL

600DP weld.

Fig. 4.3.24: Graphic representation of the hardness measurements performed on the TWIP 5 – DOCOL

800DP weld.

In a general way the weld zone present hardness values similar to the TWIP steel base material values. The exception to the previous was found in TWIP 4 welds were the weld zone hardness value increases to around 350 HV. Another consistent trend was the high hardness values presented in the heat affected zone of the Docol steels. Figures 4.3.25 and 4.3.26 allow a comparative observation of the hardness measurements performed on the studied welds.

150

200

250

300

350

400

0 2 4 6 8 10 12 14 16 18 20 22 24 26

Har

dnes

s (H

V 0,

2)

TWIP 5 - 600

DDooccooll 660000 DDPP TTWWIIPP 55 ZZTTAA WWeelldd zzoonnee

150

200

250

300

350

400

450

0 2 4 6 8 10 12 14 16 18 20 22 24 26

Har

dnes

s (H

V 0,

2)

TWIP 5 - 800

DDooccooll 880000 DDPP TTWWIIPP 44 ZZTTAA WWeelldd zzoonnee

133

0

50

100

150

200

250

300

350

400

0 2 4 6 8 10 12 14 16 18 20 22 24 26

Har

dnes

s (H

V 0,

2)

TWIP 2 - 600

TWIP 3 - 600

TWIP 4 - 600

TWIP 5 - 600 0

50

100

150

200

250

300

350

400

450

0 2 4 6 8 10 12 14 16 18 20 22 24 26

Har

dnes

s (H

V 0,

2)

TWIP 2 - 800

TWIP 3 - 800

TWIP 4 - 800

TWIP 5 - 800

Fig. 4.3.25: Graphic representation of all the hardness measurements performed on Docol 600DP steel welds.

Fig. 4.3.26: Graphic representation of all the hardness measurements performed on Docol

800DP steel welds.

Conclusions Task 4.3

Based on the experimental laser welding analysis of four TWIP steel compositions, weldability of these materials was established as well as the main influencing factors. The work allowed to conclude that the joining of TWIP steels is easily achievable by laser welding. The factors that determine the good welding behaviour and quality of TWIP steels are intrinsically related to the steel composition and austenitic structure. Concerning the dissimilar steel welding all the weld and heat affected zones present acceptable hardness values which indicate that the mechanical properties of the base materials are matched or improved after laser welding. This study permitted to prove the good potentialities of TWIP materials on automotive applications where the welding ability is one the the main requirements.

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WP5: Industrial trial The industrial trial foreseen in the WP5 was not performed and for this reason the activities scheduled in WP5 have been only partially accomplished. The reasons at the basis for the delay of the cast and finally for cancelling are discussed in the following section 5.2. Anyway the results achieved in WP5 can be summarised as follows:

• Metallurgically based selection of TWIP steel to be cast (Task 5.1). • Analysis of the impact of TWIP steel production on manufacturing route.

Task 5.1 Selection of the most interesting TWIP steel variant, for automotive applications, on the basis of the previous WPs results.

Introduction The principal aims of the industrial trial heat can be summarized in the following points:

• Metallurgical characterization of the TWIP strip produced by industrial cycle; • Evaluation of TWIP steel processing ability in terms of casting, hot and cold rolling process,

etc. All the strip produced will be used exclusively for the in-depth characterization in terms of microstructure, mechanical properties, and technological properties (forming ability, welding ability, etc). Selection of the TWIP variant for the industrial trial In view of the industrial trial, to be performed at Duferco steel works, an in-depth review of the TWIP variants properties was carried out. It is worthy to remind that in the present project 5 TWIP steel compositions were selected (table 5.1.1) and the relevant properties were investigated. TWIP Variant Fe Nb Al C N Si Mn

1 Bal. - 3.02 0.01 0.0015 3.00 21.30 2 Bal. - 0.012 0.59 0.0040 0.22 22.30 3 Bal. 0.019 0.010 0.60 0.0040 0.20 17.80 4 Bal. - 1.54 0.29 0.0060 0.21 16.40 5 Bal. - 0.011 0.23 0.2 0.20 21.00

Table 5.1.1: TWIP compositions investigated in the ‘metaldesign’ project. TWIP1 and 2 were selected as reference steel. Nevertheless also for these steel grades the knowledge of metallurgical properties in terms of SFE, recrystallization kinetics, strain hardening behavior, etc , was largely lacking. In parallel, at the start of the project, the basic metallurgical parameters controlling the TWIP effect in high Mn steels were identified and investigated. The SFE (stacking fault energy) was identified as the main metallurgical parameters controlling the steel metallurgical properties. Thermodynamical models based on empirical equations (Allain, Grässel, Dumay, etc) together with Thermocalc calculations were used in order to evaluate the SFE and the effect of the main alloy elements on SFE. In addition to SFE was also recognized that the austenite stability during deformation or temperature variations (both at low and high temperature) is an other important metallurgical parameter to taken into account when different TWIP steels are compared.

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On the basis of the selected reference steels (named TWIP1 and TWIP2) three further compositions were designed aiming to achieve with different alloy contents the range of SFE able to guarantee a good TWIP effect during deformation. The reasons and the approach underlying the metallurgical design of the newer chemical compositions selected are described in the task 1.1. For the selection of the TWIP steel for the industrial trial the following aspects were focused:

• Stable austenitic microstructure revealing a TWIP effect (deformation induced twinning ) on a wide range of temperature (at least +50°C ÷ -50°C).

• Tensile properties: in particular high elongation to rupture and strain hardening coefficient. This requirement is necessary for a good forming ability and for a high energy absorption property.

• Technological properties: welding ability, forming ability, coating ability, etc. These properties are of particular importance for these steels because the main applications for TWIP steel is for automotive sector.

• Processing ability: good casting ability, hot and cold rolling ability, pickling ability.

• Lower industrial costs and impact on environment (as low Mn content as possible);

During the project almost all of these aspects were deeply investigated and now a enough clear frame of the behavior of the steel grades designed is available. Taking into consideration all the above aspects the best performing steel, as well as documented in this and previous reports, is represented by TWIP2. Nevertheless the choice of this steel grade was excluded from the start of the project because it was introduced in the project only for reference aim and is patent covered. The selection was therefore focused on the newer steel grades metallurgically designed in the present project i.e. TWIP3, TWIP4 and TWIP5. On the basis of the above metallurgical and economical aspects all the steel grades designed show several attractive features. TWIP3,4,5 steels in terms of tensile, forming ability and welding properties, even if, in some instances lower than the reference TWIP2 steel grade, revealed quite interesting properties that could find a good favor in the automotive market. Tensile elongation larger than 50% with yield stress and tensile strength of respectively 280-300MPa and >800MPa were found. The strain hardening behavior of TWIP3,4,5 is also very interesting and in the case of TWIP 3 is really very close to TWIP2 steel. From the metallurgical point of view one of the most relevant topic is represented by austenite stability during deformation both at low and high temperature. At low temperature TWIP4 and TWIP5 as well as TWIP1 revealed a significant TRIP behavior (formation of martensite during deformation even at room temperature). TWIP3 showed a higher austenite stability and a behavior closer to TWIP2 steel. At higher temperatures TWIP 3 also revealed a clear TWIP behavior even at 250°C. Finally hot torsion tests carried out on TWIP2 and TWIP3 revealed a very good intrinsic ductility for both steel. On the basis of all these metallurgical aspects, of processing ability and finally in order to reduce the extra alloy costs and impact on environment for the industrial heat was preferred the TWIP3 variant. Conclusions of task 5.1 TWIP3 is the selected TWIP grade to be cast at Duferco steelworks. TWIP3 steel grade (chemical composition reported in following table) is characterized by a markedly lower Mn content (25% less than TWIP2) that is the most expensive element in TWIP steels and, as well known, the element that could have unwished impact on environment during steel making due to Mn evaporation. The properties revealed by this steel if, in some instance, are below the TWIP2 are really promising and

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could be further improved by means of a fine tuning of the alloy design and industrial processing conditions.

Fe Mn C Si Nb min bal 16.00 0.60 0.20 0.015 max bal 17.50 0.65 0.30 0.030

Table 5.1.2 Range of elements for TWIP3 steel grade to be cast at Duferco steelworks.

Task 5.2 Coil supplying and material processing (hot and cold rolling, annealing treatments, pickling and coating). The scheduling of the TWIP industrial heat suffered the major deviation of the activities of the present project because of more than expected difficulties and technical challenges related to TWIP steel production. Two were the main issues:

1. Installation of a vacuum degassing plant due to risk of hydrogen embrittlement susceptibility of high Mn steel;

2. Definition of the steelmaking route of high Mn steel using a direct current electric arc furnace (DC-EAF).

In the following are summarized the activities carried out on these topics.

Installation of a vacuum degassing plant due to risk of hydrogen embrittlement susceptibility of high Mn steel The development of high Mn TWIP steel faces a number of technical challenges related to steelmaking, casting and strip processing. On the basis of recent literature, even if unclear, could be inferred that TWIP steels, could be sensitive to hydrogen embrittlement even in relation with the formation of martensitic phases in the microstructure of deformed parts. This phenomenon should be carefully examined in the case of the high Mn steels production without vacuum degassing facilities. Duferco La Louviere steelworks planned the installation of a VD facility aiming to achieve with this facility a significant improvement not only for hydrogen content reduction in steel but also to widen the range of feasible steel grades (low C content). For this reason the scheduling of the TWIP heat was postponed of 1 year. In fig. 5.2.1 is shown a typical kinetics of hydrogen removal curve during a VD process. As can be noted into about 20 min of VD process is possible to reduce the hydrogen content in steel from 3.5 ppm down to 1.0-1.5 ppm.

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Fig. 5.2.1 Typical kinetics of hydrogen removal curve during a VD process using different Argon pressure.

Unfortunately several causes determined a larger delay of the VD facility completion inclusive, of course, the recent steel markets crisis (2008-09) that widened the time for VD installation and commissioning work.

Evaluation of the steelmaking route of high Mn steel in a direct current electric arc furnace (DC-EAF).

The feasibility of TWIP steel on a direct current arc furnace required a in depth study of the metallurgical interactions of steel and furnace soils electrodes for the evaluation of the potential risks of furnace damagement. Whatever was the chosen operating practice at EAF (selected scrap without deP or normal scrap with deP and deslagging operations, the addition of an high amount of Mn ferroalloys in the furnace was required. The only way to perform this additions was to use a charge basket when the first baskets of scrap were melt. This operation led to the formation of a small “mountain” of Mn alloy on the bottom of the furnace, just in contact with the soils electrodes. It is well known that the melting of a ferroalloy happens for the formation of low melting points alloys between steel and the ferroalloy added, the risk was to have a heavy consumption of the soils electrodes (which are steel billets) due to the contact with Mn ferroalloys and the local high temperature in the arcing zone. The study was carried out on the basis of thermo dynamical calculation and they showed that low melting point high manganese content alloys forms during melting, but taking into account that soils electrodes are cooled and of big dimensions, it was important to evaluate the time needed to reach an homogeneous chemical condition in the liquid melt just to have an idea of possible real soils electrode’s consumption.

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Fig. 5.2.2: Fe-Mn diagram phase.

From the diagram phase Fe-Mn it can be seen that in the worst condition, the liquidus temperature of a Fe-Mn alloy, with more than 60% of Mn, is around 1250°C.

In AC furnace, the experience showed that about 15 minutes are needed to melt the whole Mn ferroalloy bulk added with the basket. The time in Duferco DC furnace should be quite similar, and taking in consideration that the soils electrode are water cooled, and that an important portion of them is at a temperature lower than 1200°C, it should be reliable that no catastrophic damages to the bottom of the furnace should occur. This hypothesis is realistic even thinking that some DC Furnaces use instead of steel billets, copper billets as soils electrodes. The melting point of copper is lower than that of whatever kind of steel, but the water cooling enables the use of this kind of soils electrodes.

Moreover, when the bulk or ferroalloy is quite melt, the argon stirring performed by bottom porous plugs should enhance the chemical homogenization, bringing rapidly to non dangerous conditions (the Fe-Mn alloy at 20% of Mn has a melting point which is similar to some stainless steels, which can be produced by DC furnaces). As the chemical and thermal homogenization of melt, is function even of steel recirculation inside the furnace, induced by the use of porous plugs which are present on the bottom of Duferco La Louviere’s EAF, fluid dynamic simulations, by the use of numerical and physical modeling (water models) should be necessary to evaluate the real industrial risk and to optimize the production route. Laboratory tests were moreover required to validate the theoretical study. It was not possible to perform such kind of laboratory tests due to lack of time so the heat was not performed, even if it is quite clear that, determining the necessary right conditions, it should be possible to perform the production of high Mn steels even in a DC Furnace.

Economical impact of high Mn TWIP steel grade production In the present section will be summarised the results of an indepth analysis aiming to evaluate for all the processes involved in the industrial manufacturing route the impact of the TWIP steel production due to the very peculiar characteristics of this grade in terms of chemical composition, steel physical

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properties, etc. First of all it is worthy to separate the aspects related to different process area of manufacturing route: 1) Steelmaking. 2) Continuous casting. 3) Hot rolling. 4) Cold rolling and final annealing. For what concern the first two points (primary area) the impact represented by TWIP production is mainly related to:

• Steelmaking process: The impact of TWIP steel production in the primary area can be envisaged both in terms of process and environment. In terms of steelmaking process the issue related to hydrogen embrittlement susceptibility resulted with an importance stronger than expected in the technical annex. This issue has suggested to delay the industrial heat scheduling in order to evaluate the risk of hydrogen embrittlement for TWIP steels and the need to include a vacuum degassing facility (VD) for industrial TWIP steel manufacturing. The total time for EAF+LF process is strictly related to the specific steelmaking practice adopted. This means that different practice (in terms of type of Fe alloy additions) and different facilities adopted (LF/VD, deslagging), could yield to dramatically different total working time. As discussed in the previous section the specific steelmaking route of high Mn steel has to be tailored according if DC or AC EAF process is adopted. The total steelmaking process time is expected longer than conventional carbon steel and this could result in higher energy consumption and higher refractory wear (EAF, ladle) due to high MnO content in the slag. In terms of enviromental impact the high Mn content of TWIP steels determines during both steelmaking and continuous casting large amount of smoke constituted mainly of Manganese oxides. The management of so large amount of smoke constitute a environmental problem that could require investments in terms of facilities necessary to avoid pollution. This problem is not limited only to steelmaking process but also during continuous casting process in the transitory of tundish filling (without covering powder) smoke can be developed. The implementation of a tundish cover could be a solution.

• Continuous casting process: TWIP steels are characterized by lower liquidus temperature with respect to low-carbon steel and also the solidification range (liquidus-solidus) is wide (84°C for TWIP2 steel). These peculiarities have a strong impact on the casting process. It is worthy to mention, among others, the following topics:

o Set-up/design of tailored casting powder for high Mn steel: the low liquidus temperature and high Mn content does not allow the use of the same casting powders used for low-C steels. Lubrication unevenness could be a cause of break-out and other casting problems (cracks, depressions).

o Optimization of slab secondary cooling: Due to TWIP steel high temperature strength the adoption of not correct practice of slab secondary cooling (below mould) could damage seriously the bending and strengthening rolls of casting machine after only one sequence cast.

If all these aspects are not optimized additional costs in terms of production loss and maintenance costs could occur. Concerning the hot rolling process the most critical aspect is represented by insertion of relatively small number of TWIP slabs within the charge of walking beam furnace production. Due to quite different physical properties of TWIP steel with respect to conventional carbon steel the adoption of hot charging

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of slabs is necessary because is not recommended to change the working conditions (in terms of soaking temperature and time) of the walking beam furnace. Of course this means that suitable facilities for hot slab stacking as to be considered. The impact of hot charging can be assumed mainly on the level of management of production planning and in case of stop at the hot rolling mill. Finally for what concern the cold rolling and annealing process the impact of TWIP steel production is mainly related to

• Cold rolling process: TWIP steels work harden very quickly during cold deformation. The work hardening of TWIP steels is significantly higher than stainless steels as shown by plain strain compression tests. At a true strain of 0.5 the difference in terms of stress is about 38%. The above result implies difficulties in cold rolling in terms of loads, number of passes or reduction ratio at each stage. Depending on the hot strip thickness, an intermediate annealing treatment (two step cold rolling process) could be necessary for obtaining the aimed final cold strip thickness (<2.0 mm).

• Annealing process: As stressed in previous WP3, TWIP steels tensile properties are seriously

affected by decarburization. The choice of a suitable annealing furnace atmosphere is of paramount importance to avoid the occurrence of decarburization.

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Objectives of the project

The main objectives of this project are:

1. To complete the understanding of the TWIP steel metallurgy particularly for what regard the following fields: • Solidification microstructure; • Recrystallization behaviour (dynamic and static) and texture formation. • Tensile properties and work hardening ability, precipitation behaviour as a function of

the steel chemical composition and its influence on the strip properties. • Influence of steel chemical composition on microstructure, mechanical properties,

strain hardening behaviour as well as on application properties such as formability, weldability and coatability.

2. To design a metallurgically based manufacturing route (hot/cold rolling process, annealing

treatment) taking into account the specific capability of the DUFERCO plants. Comparison of initially planned activities and work accomplished The major deviation of the activities of the present project was caused by more than expected difficulties and technical challenges related to steelmaking of TWIP steel. Two are the main issues:

• On the basis of recent developments, one particular challenge is represented by the fact that TWIP steels, could be sensitive to hydrogen embrittlement even in relation with the presence of martensitic phases in the microstructure of deformed parts. This phenomenon should be carefully examined in the case of the high Mn steels production without vacuum degassing facilities.

• Evaluation of the steelmaking route of high Mn steel in a direct current electric arc furnace (DC-EAF).

The above issue determined the delay of the project of 1,5 years and the solution was pointed out in the installation at Duferco La Louviere of a vacuum degassing facility (VD) in order to produce a significantly lower H content in steel.

The second issue is related to the fact that the world-wide experience on TWIP steelmaking is only on alternate current electric arc furnace. The feasibility of TWIP steel on a direct current arc furnace required a in depth study of the metallurgical interactions of steel and furnace soils electrodes for the evaluation of the potential risks of furnace damagement. Whatever was the choosen operating practice at EAF (selected scrap without deP or normal scrap with deP and deslagging operations, the addition of an high amount of Mn ferroalloys in the furnace was required. The only way to perform this additions was to use a charge basket when the first baskets of scrap were melt. This operation led to the formation of a small “mountain” of Mn alloy on the bottom of the furnace, just in contact with the soils electrodes. It is well known that the melting of a ferroalloy happens for the formation of low melting points alloys between steel and the ferroalloy added, the risk was to have a heavy consumption of the soils electrodes (which are steel bars) due to the contact with Mn ferroalloys and the local high temperature in the arcing zone. The study was carried out on the basis of thermo dynamical calculation and they showed that low melting point high manganese content alloys forms during melting, but taking into account that soils electrodes are cooled and of big dimensions, it was important to evaluate the time needed to reach an homogeneous chemical condition in the liquid melt just to have an idea of possible real soils electrode’s consumption. As the chemical and thermal homogenization of melt, is function even of steel recirculation inside the furnace, induced by the use of porous plugs which are present on the bottom of Duferco La Louviere’s EAF, fluid dynamic simulations, by the use of numerical and physical modeling (water models) should be necessary to evaluate the real industrial risk and to optimize the production

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route. Laboratory tests were moreover required to validate the theoretical study. It was not possible to perform such kind of laboratory tests due to lack of time so the heat was not performed, even if it is quite clear that, determining the necessary right conditions, it should be possible to perform the production of high Mn steels even in a DC Furnace. The study was carried out on the basis of thermodinamical calculation and also taking into account the steelmaking experience of other high Mn steels nevertheless laboratory tests to validate the theoretical study were needed and the time was not enough. The missed realization of the industrial heat has determined the impossibility to perform the Task 5.3 and Task 5.4 mainly relevant to characterization of the industrial hot and cold rolled strip.

Conclusions Five TWIP steel grades (Fe-Mn-C-N and Fe-Mn-Al-Si-C) were selected on the basis of literature review and thermodynamical calculations adopting as main characteristics for alloy design the SFE and austenite phase stability. The chemical composition are reported in the following table.

Table 5.1.3: Selected TWIP grades chemical analysis.

The selected TWIP grades were characterised in terms on SFE, mechanical properties, microstructure evolution during deformation, recrystallization, precipitation behaviour. The results revealed that TWIP steel have really excellent tensile properties that make them extremely attractive for automobile and for structural reinforcement. In addition, the positive strain rate sensitivity make TWIP steels suitable for applications devoted to energy absorption (crashworthiness). The performances analysis revealed that only TWIP2 and TWIP3 clearly behave as expected from TWIP steels, with the typical hardening stage associated to profuse deformation twinning The TWIP steel variants with significant Al and Si content (TWIP1 and TWIP4) require a additional tuning of steel chemistry to better balance the effect of Mn and C and avoid the formation of second phases (ferrite or deformation induced martensitic phases). Al seems to be effective on suppress the γ ε-martensite transformation but the formation of α-martensite has been detected on strained samples. At same way the variant TWIP5 with high nitrogen and low C requires an additional set-up of the steel chemical composition in order to increase the SFE and improve the austenite stability during deformation. This grade could be interesting for a significantly lower carbides precipitation in the range 500-600°C. The results achieved in the project allowed to define the suitable manufacturing route to produce the TWIP steel. Concerning the strip annealing process the avoidance of strip surface decarburization was stressed. This means that a proper atmosphere control during annealing at high temperature must be taken. Even if a technical work relevant to the feasibility study has been carried out, due to the steel market crisis, the proposed industrial TWIP steel heats was not carried out for two main reasons: too much increased time and costs for installation of a vacuum degassing facility (VD), compared to the RFCS scheduled duration of the project and the industrial available resources, and the not completely in time explored comprehension of the process (thermodynamical modellization and calculations) for the production of high Mn steel by direct current electric furnace even with further expertise of external consultants. The criticality of the industrial heat is at least in part balanced by an indepth not foreseen work aimed on focus to theroretical (diffusion evaluation inside austenitic lattice steel) and experimental (tensile tests on hydrogen charged specimens, TDA) study on the hydrogen embrittlement for TWIP steel and at the same time a not foreseen work done on the decarburization effect of TWIP tensile properties during the annealing step of the production process.

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Exploitation and impact of the research results The present project consisted in the metallurgically based design of TWIP steels assuming as main key parameters the stacking fault energy (SFE) and the austenite stability. The performances of the TWIP variants studied, reveal that also small difference in chemical composition yield to quite different proprierties in terms of deformation mechanisms and austenite stability. The approach used and the quite large amount of results gathered, among which data on recrystallization behaviour, precipitation, tensile properties at low and high strain rate, high temperature flow stress, will be useful to outline the guideline for design of TWIP steel manufacturing route. For the development and application of high Mn steels a further continuation of the investigations on hydrogen embrittlement of TWIP steels seems to be necessary. In particular the activities should be focused on fracture mechanics tests to measure critical stress intensity factor, crack growth rate, delayed failure tests and slow strain rate tests (SSRT). Fundamental aspects concerning the main trapping sites of diffusible hydrogen and activation energies for its desorption, and thermal desorption analysis (TDA), has to be clarified. Similarly the effect of process conditions (for example annealing, galvanization) on embrittlement of TWIP steel has to be investigated. Another field of investigation completely lacking at moment is represented by characterization of creep behaviour of TWIP steel. This investigation could open new applications fields for TWIP steels. The investigated TWIP steel have shown excellent tensile properties that could open new innovative application fields in addition to utilize them as extremely attractive for automobile, even for the positive strain rate sensitivity make TWIP steels suitable for applications devoted to energy absorption (crashworthiness) and for structural reinforcement as for body in white. New activities could be therefore directed, to fill up an existing gap in particular for the automotive industry, on an innovative metallurgically based design of TWIP steel, on a dedicated manufacturing route, to obtain a new product combining high-strength, improved crash resistence and reduction of gas emissions through lightweighting. The following pubblications were carried out using the results achieved in the present project: 1) Fatigue Behavior of four high-Mn TWIP steels, Metallurgical and Materials Transactions A, 41A, 2010, 1102-1108. 2) Fatigue behavior of high-Mn TWIP steels, Materials Science and Engineering A, 517, 2009, 68-77. 3) High-Cycle Fatigue Behavior of Ultrafine-Grained Austenitic Stainless and TWIP steels, submitted to Materials Science and Engineering A. 4)Kinetics of recrystallization and grain growth of cold rolled TWIP steel, Advanced Materials Research Vols. 89-91 (2010) pp 153-158 5) Fatigue Behavior of Four High-Mn Twinning Induced Plasticity Effect Steels, 1102—VOLUME 41A, MAY 2010 METALLURGICAL AND MATERIALS TRANSACTIONS A. 6) Javier Gil Sevillano, “An alternative model for the strain hardening of FCC alloys that twin, validated for twinning-induced plasticity steel”, Scripta Materialia, 60 (2009) 336-339. 7) “Kinetics of recrystallization and grain growth of cold rolled TWIP steel”, Advanced Materials Research, 89-91 (2010) 153-158 8) ”Hall-Petch relationship of a TWIP steel”, Key Engineering Materials, 423 (2010) 147-152. 9) F. de las Cuevas, M. Reis, A. Ferraiuolo, G. Pratolongo, L.P.Karjalainen, J. Alkorta, J. Gil Sevillano, ”Efecto Hall-Petch en un acero TWIP”, XI Congreso Nacional de Propiedades Mecánicas de los Sólidos PMS2008, Cádiz, Spain, 2008. 10) F. de las Cuevas, a, M. Reis, A. Ferraiuolo, G. Pratolongo, L.P. Karjalainen, V. García Navas, J. Gil Sevillano, “Kinetics of recrystallization and grain growth of cold rolled TWIP steel”, THERMEC 09, Berlin (Germany), 2009. List of figures and tables WP1: Testing material supply and basic metallurgical characterisation of cast materials Fig.1.2.1.1: CSM VIM facility.

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Fig. 1.3.1: Images of TWIP1 ingot. a) Macrostructure of the TWIP1 ingot sample, polished and Nital 2% etched b) SEM image of structure 2% Nital etch. Fig. 1.3.2: a) Optical microstructures of TWIP1 in the cast state of lathy (black) and vermicular ferrite; b) XRD of TWIP-1 in the cast state showing the presence of duplex structure. Fig. 1.3.3: The element distribution (X-ray map) in the austenite and ferrite phases in cast TWIP1, (a) SEM image, (b) Al distribution, and (c) Mn distribution. Fig. 1.3.4: Isothermal cross-sections of Fe-Mn-Al system. (a) 1200°C, (b) 1000°C [2]. Fig. 1.3.5: SEM-EDS analysis locations on the cast TWIP1. Fig.1.3.6: OM photo of the microstructure of the cast TWIP2. Fig. 1.3.7: Presence of alfa-martensite in the cast TWIP3. Fig. 1.3.8: OM photo of TWIP4 in the cast state. Fig. 1.3.9: OM photos of the microstructure of the cast TWIP-5. WP2: Fundamental investigations on the physical metallurgy of TWIP steels Fig 2.1.1.1: TWIP2 microstructure after annealing treatments at 1000°Cx5 min (average grain size 32 μm). Fig. 2.1.1.2: Extended nodes detected on a TWIP2 sample. Fig. 2.1.2.1: Measure of hcp-martensite fraction on sample deformed at 0.30 at low temperature. Fig. 2.2.1.1: Schedule used in double-hit compression tests. Fig. 2.2.1.2: SEM-EBSD photo of TWIP-2 steel as heated at 1200°C for 2 min and water quenched. Red lines reveal high-angle (>15°) grain boundaries. Fig. 2.2.1.3: Typical true stress-strain curves of TWIP-2 (Fe-22Mn-0.6C) steel at high temperatures and constant strain rate of 0.1 s-1. Curves for the low-carbon steel are included for comparison. constant strain rate of 0.1 s-1. Curves for the low-carbon steel are included for comparison. Fig. 2.2.1.4: Dependence of peak stress (σp) and peak strain (εp) on the inverse temperature at the strain rate of 0.1 s-1 for TWIP-2 and low-C steels. Fig. 2.2.1.5: Plot of ln(εp) vs ln Z for the low-C and TWIP-2 (22Mn-0.6C) steels. Fig. 2.2.1.6: SRX rates of TWIP-2 at constant strain rate of 0.1 s-1. Double-hit compression data and fitted curves. Times for 50% recrystallization, t50, and the Avrami exponents are listed. Fig. 2.2.7: Typical true stress-strain curves of TWIP-3 (Fe-18Mn-0.6C-0.02Nb) steel at high temperatures and the constant strain rate of 0.1 s-1. A curve for Fe-26Mn-0.14C steel is included for comparison. Fig. 2.2.1.8: SRX of TWIP-3 (0.2 strain at constant strain rate 0.1 s-1). Double compression data and fitted curves. Fig. 2.2.1.9: SRX of TWIP-3 determined in a relaxation test at 1000°C/0.2/0.1 s-1. Fig. 2.2.1.10: Typical true stress-strain curves of TWIP-4 (Fe-16Mn-1.5Al-0.3C) at the constant strain rate of 0.1 s-1. Flow stress curves for 25Mn1Al [9] steel are included for comparison. Fig. 2.2.1.11: SRX rates of TWIP-4 at the constant strain rate of 0.1 s-1. Double-hit compression data and fitted curves. Data for 25Mn1Al at 1000°C from [9]. Fig. 2.2.1.12: True stress-strain curves of TWIP-5 (Fe-22Mn-0.2C-0.2N) steel at high temperatures and at the constant strain rate of 0.1 s-1. Curve for TWIP-2 (Fe-22Mn-0.6C) is included for comparison. Fig. 2.2.1.13: SRX of TWIP-5 at the constant strain rate of 0.1 s-1. Double-hit compression data and fitted curves. Fig. 2.2.1.14: SRX rates (strain 0.2, strain rate 0.1 s-1) for the 25Mn1Al steel. Double-hit compression data and fitted curves. Fig. 2.2.1.15: Comparison of SRX kinetics of 25Mn1Al to low-carbon, Nb and Type 304 steels. Data for the Nb-steel from [2]. Fig. 2.2.1.16: Plot of t50 vs strain to calculate Fig. 2.2.1.17: Plot of t50 vs strain to calculate the power of strain for 25Mn1Al. Fig. 2.2.1.18: Time t50 of SRX for 25Mn1Al vs the inverse absolute temperature. Fig. 2.2.2.1: OIM image of hot rolled TWIP2 sample (inverse pole figure and grain maps, rolling plane normal

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Fig. 2.2.2.2: Hall-Petch plot relationship for TWIP 2 steel with composition 22% Mn- 0.6% C, of the hardness, H, as a function of the inverse of the square root of the grain size, D, for structures after fully recrystallization or after fully recrystallization and grain growth. The grain size was determined as the mean linear intercept using EBSP-OIM images, counting both, twin boundaries together with other high-angle boundaries with misorientation higher of 12º. Fig. 2.2.2.3: Recrystallization kinetics for all reductions derived from the softening by annealing temperature of 700ºC as function of time. TWIP 2 steel with composition 22% Mn- 0.6% C Fig. 2.2.2.4: Grain growth equation. D* is calculated from mean linear intercept method, twin boundaries not counted as grain boundaries. TWIP 2 steel with composition 22% Mn- 0.6% C. Fig.2.2.2.5 Macro-textures, pole figures.a) TWIP steel 60% cold rolled with the typical brass-type texture {011}<211>.b) Recrystallization texture, D ≤ 2 μm.c) Grain growth texture with D ≈ 9 μm. Fig. 2.2.3.1: Comparison of model 1 (black line) with experimental data for a cold reduction of 60% with an annealing temperature of 600°C. Fig. 2.2.3.2: Comparison of model 2 (red line) with (corrected) experimental data for a cold reduction of 60% with an annealing temperature of 600°C. Fig. 2.2.3.3: Comparison of model 3 (blue line) with (corrected) experimental data for a cold reduction of 60% with an annealing temperature of 600°C. Fig. 2.2.3.4: Comparison of model 3 with a range of experimental data, for a constant annealing temperature of 600°C. A) 40% reduction B) 50% reduction C) 60% reduction and D) 70% reduction. Fig.2.3.1: Microstructure of the TWIP variant after soaking at 800°C for 10h. Fig. 2.3.2: TWIP1 - 700°Cx10h Fig. 2.3.3: SEM image of the TWIP2 sample after soaking at 700°C for 10h Fig. 2.3.4: SEM image of the TWIP3 sample after soaking at 700°C for 10h. Fig. 2.3.5: SEM image of the TWIP2 sample after soaking at 600°C for 10h. Fig. 2.3.6: TWIP3: SEM + EDS image of the cementite precipitated on the austenitic grain boundaries after a soaking at 600°C for 10h. Fig. 2.3.7: TWIP4 at 700°C: within some grains is detected a precipitation of ferrite + carbides Fig. 2.3.8: SEM + EDS image of the TWIP 4 samples after a soaking at 600°C for 10h. Fig. 2.3.9: SEM + EDS image of the TWIP 4 samples after a soaking at 500°C for 10h. Fig. 2.3.10: TWIP 5 specimen after soaking for 10h at 500°C Fig. 2.3.11: TWIP1-700°C: Low magnification image Fig. 2.3.12: TWIP1-700°C: Carbide of cementite type (FeMn)3C, the analysis is reported in fig. 2.3.11. Fig. 2.3.13: (Fe,Mn)3C carbide image with relevant diffraction Fig. 2.3.14: Vanadium carbo-nitride image with relevant diffraction. Fig. 2.3.15: Nb content of the precipitates selected on extraction replica. Fig. 2.3.16: TWIP3 – 600°C: (Fe,Mn)3C carbide image Fig. 2.3.17: TWIP5 at 700°Cx10h - (Ti, Al) carbo-nitrides Fig. 2.4.1: Typical slab temperature in conventional route. Fig. 2.4.2: Typical slab cooling evolution adopting the hot charge route (piling up in a box). Fig 2.4.3: Comparison of the MFS of TWIP2 and a conventional low-C steel. Fig 2.4.4: TWIP1 scale analysis Fig 2.4.7 TWIP4 scale analysis Fig 2.4.5: TWIP4 scale analysis Fig. 2.4.6: Scale microstructure and main phases detected on: a) TWIP2 and; b) TWIP3 steel samples. Fig. 2.4.7: Scale microstructure and main phases detected on TWIP5 WP3 Study of the deformation mechanisms and strain hardening behaviour Fig. 3.1.1.1a): tensile tests results carried out at room temperature and higher temperature 150-350°C. Fig. 3.1.1.1b): tensile tests results carried out at room temperature and higher temperature 150-350°C. Fig. 3.1.1.2: The strain rate sensitivity of the five TWIP steel grades was evaluated collecting all the data, torsion, tensile tests and Hopkinson bar tests as well. Fig. 3.1.1.3: Calculated work hardening curves for each TWIP grade. Fig. 3.1.1.4: True stress-true strain curves for TWIP3-4.

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Fig. 3.1.1.4: TWIP 2 x500 a) head region etchant I, b) strain = 0.4, etchant I Fig. 3.1.1.5: TWIP 3 x500 a) head region etchant I, b) strain = 0.4, etchant I. Fig. 3.1.1.6: TWIP 4 x500 a) head region (strain=0) etchant I, b) strain = 0.4 etchant II. Fig. 3.1.1.7: TWIP 5 x500 a) head region etchant I, b) strain = 0.4 etchant II. Fig. 3.1.1.8: Calculated volume fractions of constituents in TWIP steels during tensile straining. A) TWIP 1, B) TWIP 2, C) TWIP 3, D) TWIP 4 and E) TWIP 5. Fig. 3.1.1.9: Ferritoscope measurements of α’bcc martensite on strained TWIP steels samples (deformation 10%). Fig. 3.1.1.10: Ferritoscope measurements of α’bcc martensite on strained TWIP steels samples (deformation 30%). Fig 3.1.1.11: TWIP1: deformed at 250°C (strain 0.3 - strain rate = 0.01). Fig.3.1.1.12: TWIP3: microstructure deformed at: a) 250°C; and b) 350°C (strain 0.3 - strain rate = 0.01). Fig.3.1.1.13: TWIP4: deformed structure at: a) 250°C; b) 350°C (strain 0.3). Fig.3.1.1.14: TWIP5: deformed at 250°C (strain 0.3 - strain rate = 0.01). Fig. 3.1.1.3.1a): Longitudinal cross section of a TWIP3 tensile specimen close the fracture surface. Fig. 3.1.1.3.1b): Subsurface microstructure after tensile test (TWIP3) – Klem II etching Fig. 3.1.1.3.2: TWIP2 microstructure after annealing at 1000°C x 300s Fig. 3.1.1.3.3: TWIP2 microstructure after annealing at 1200°C x 1100s. Fig. 3.1.1.3.4: Comparison of the decarburization depth of TWIP steels and C45 steels at 1000°C and 1200°C. Fig. 3.1.1.3.5: GDOES Carbon and manganese concentration profiles on TWIP2,3,4. Fig. 3.1.1.3.6: Tensile test curve of hydrogen charged TWIP3 sample. Fig. 3.1.1.3.7: Results of tensile tests carried out on sample with different hydrogen content. Fig. 3.1.1.3.8: DTA plot on blank and hydrogen pre-charged sample. Fig. 3.1.2.1a): Comparison between the quasi-static mechanical properties, yield strength Rp0.2, tensile strength Rm, and fracture elongation A of the present TWIP steels [15] and those of TRIP700, DP600 and H340LAD [17] at RT (the strain rate ≈ 10-3 s-1). Fig. 3.1.2.1 b): Comparison between the dynamic mechanical properties, yield strength Rp0.2, tensile strength Rm, and fracture elongation A of the present TWIP steels and those of TRIP700, DP600 and H340LAD [17] at RT (the strain rate of order of 103 s-1). Fig. 3.1.2.2: Comparison between the stress ratio in quasi-static and dynamic tensile tests of the present TWIP steels and those of TRIP700, DP600 and H340LAD [17]. Fig. 3.1.2.3: Strain-hardening coefficient (n-value) from the Hollomon equation vs. true plastic strain for TWIP2 tested at 1210 s-1 and RT. Fig. 3.1.3.1: Torsion shear stress-plastic shear strain curves at =ε 1.4·10-3 s-1 and corresponding Kocks-Mecking plots (strain rate vs. flow stress). Fig. 3.1.3.2: Torsion shear stress-plastic shear strain curves at =ε 0.113 s-1 and corresponding Kocks-Mecking plots (strain rate vs. flow stress). Fig. 3.1.3.3: Torsion shear stress-plastic shear strain curves at =ε 3.73 s-1 and corresponding Kocks-Mecking plots (strain rate vs. flow stress). Fig. 3.1.3.4: Torsion shear stress for three shear strain levels, 0.173, 0.520 and 1.040, as a function of temperature, for the three strain rates tested (TWIP2, TWIP3). Fig. 3.1.3.5. Torsion shear stress for three shear strain levels, 0.173, 0.520 and 1.040, as a function of temperature, for the three strain rates tested (TWIP4, TWIP5). Fig. 3.1.3.6: TWIP steels, maximum shear stress in torsion as a function of temperature for different strain rates. Fig. 3.1.3.7: Torsion shear failure strain of the four TWIP steels as a function of temperature and strain rate. Fig. 3.1.4.1: Hot ductility curve for TWIP1 (Fe-22Mn-3Al-3Si). Fig. 3.1.4.2: Hot ductility curve of TWIP2. Curve of AISI 304 is included for comparison.

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Fig. 3.1.4.3: Hot ductility curve of TWIP3. Curve of TWIP2 is included for comparison. Fig. 3.1.4.4: Hot ductility curve of TWIP4. Curve of TWIP2 is included for comparison. Fig. 3.1.4.5: Hot ductility curve of TWIP5. Curve of TWIP2 is included for comparison. Fig. 3.1.5.1: Hall-Petch plot of the tensile yield stress at 10-3 s-1 (flow stress for 0.002 tensile plastic strain) of steel TWIP1 and TWIP2. Fig. 3.1.5.2: Hall-Petch slopes for the flow stress at fixed true plastic tensile strain values, steels TWIP1 and TWIP3 as a function of plastic true tensile strain, up to ε = 0.15. Fig. 3.1.5.3: Hall-Petch slopes of steel TWIP2 at room temperature and three tensile strain rates, as a function of tensile plastic strain. Fig. 3.1.5.4: Work hardening rate vs. true flow stress of the TWIP 2 steel with composition 22% Mn-0.6% C. Tensile tests at room temperature and 10-3 s-1 for a equiaxed grain size in the range 1.5 μm < D < 50 μm. Fig. 3.1.5.5a: D* grain size (mean linear intercept) without considering twin boundaries . Fig. 3.1.5.5b: D id., with twin boundaries counted as grain boundaries. Fig. 3.1.5.6. Twin Boundary Fraction vs. Grain size D* (twin boundaries not counted). Fig. 3.2.1. S-N curves of the investigated TWIP steels and those of 301LN and 316L and high-strength TRIP steel are included for comparison [1,2]. Fig. 3.2.2: The final fatigued structure of TWIP (Fe-22Mn-0.6C) steel after failure. Fig. 3.2.3: Effect of the grain size on the fatigue life of the TWIP steel (Fe-22Mn-0.6C). All tests were carried out at the amplitude of 500 MPa. Fig. 3.3.1: Charpy tests results for the five TWIP compositions. Fig. 3.4.1: Plain strain compression tests results on TWIP2 variant. WP 4 – Basic characterisation of application properties: formability, weldability and coating ability Fig. 4.1.1.1 – Graphic representation of the Erichesen tests results for the five TWIP compositions. Fig. 4.1.2.1 - Comparison between the Erichsen index (IE) in high-speed (and conventional) testing and the total elongation (A) from quasi-static tensile tests [3] and that of AISI 304 [5]. Fig. 4.2.1: Sample TWIP 2 DM . Fig. 4.2.2: Sample TWIP 2 DQ. Fig. 4.2.3: Sample TWIP 5 DM (200X). Fig. 4.2.4: Sample TWIP 2 DM (200X). Fig. 4.2.5: Sample TWIP 2 DQ (500X). Fig. 4.2.6: Sample TWIP 5 DM (200X). Fig. 4.2.7: Sample TWIP 5 DQ (500X). Fig. 4.3.1: Visual observation of the front surface of the TWIP3 diode laser weld. Fig. 4.3.2: Visual observation of the root surface of the TWIP3 diode laser weld Fig. 4.3.3: Low-magnification light optical micrograph from TWIP2 steel showing a cross section of the welded joint Fig. 4.3.4: Low-magnification light optical micrograph from TWIP3 steel showing a cross section of the welded joint Fig. 4.3.5: Low-magnification light optical micrograph from TWIP4 steel showing a cross section of the welded joint Fig. 4.3.6: Micrography of TWIP2 steel base material (500X). Fig. 4.3.7: Micrography of TWIP3 welded sample showing the transition base material- welded material (200X). Fig. 4.3.8: Micrography of TWIP3 welded joint (200X). Fig. 4.3.9: Micrography of TWIP2 welded material showing the presence of ε-martensite (200X). Fig. 4.3.10: Vickers hardness measurements Fig. 4.3.11: Docol 800DP steel microstructure taken in a scanning microscope (X500). Fig. 4.3.12: Tailored blank weld configuration used in the welding tests. Fig. 4.3.13: Visual observation of the front surface of the TWIP4 – Docol 600DP CO2 laser weld.

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Fig. 4.3.14: Low-magnification light optical micrograph showing a cross section of the welded joint - TWIP2 and Docol 600DP steels Fig. 4.3.15: Low-magnification light optical micrograph showing a cross section of the welded joint - TWIP2 and Docol 800DP steels. Fig. 4.3.16: Low-magnification light optical micrograph showing a cross section of the welded Fig. 4.3.17: Low-magnification light optical micrograph showing a cross section of the welded joint - TWIP3 and Docol 600DP steels. Fig. 4.3.18: Hardness distribution curve of a Docol 800 DP laser weld Figure 4.3.21 – Low-magnification light optical micrograph showing a cross section of the welded joint - TWIP3 and Docol 600DP steels. Fig. 4.3.19: Graphic representation of the hardness measurements performed on the TWIP 2 – DOCOL 600DP weld. Fig. 4.3.20: Graphic representation of the hardness measurements performed on the TWIP 2 – DOCOL 800DP weld. Fig. 4.3.21: Graphic representation of the hardness measurements performed on the TWIP 4 – DOCOL 600DP weld. Fig. 4.3.22: Graphic representation of the hardness measurements performed on the TWIP 4 – DOCOL 800DP weld. Fig. 4.3.23: Graphic representation of the hardness measurements performed on the TWIP 5 – DOCOL 600DP weld. Fig. 4.3.24: Graphic representation of the hardness measurements performed on the TWIP 5 – DOCOL 800DP weld. Fig. 4.3.25: Graphic representation of all the hardness measurements performed on Docol 600DP steel welds. Fig. 4.3.26: Graphic representation of all the hardness measurements performed on Docol 800DP steel welds WP5 Industrial trial Fig. 5.2.1 Typical kinetics of hydrogen removal curve during a VD process using different Argon pressure. Fig. 5.2.2: Fe-Mn diagram phase List of table WP1 table Table 1.1.1.1 Selected TWIP steel compositions (Fe balance). Table 1.1.2.1: Numerical values and functions used for the calculations of SFE in two models Table 1.1.2.2: Calculated SFE for TWIP steels at 298 K. Table 1.2.1.1 Ingots chemical analysis. Table 1.3.1: Partition ratios PD of alloying elements between austenite and ferrite in TWIP 1. Table 1.3.2: Tsolidus and Tliquidus of TWIP steels. Table 1.3.3 Ingots chemical analysis. WP2 table Table 2.1.1.1: SFE results achieved on TWIP 2 Table 2.1.1.2: Model parameters for calculation of SFE in Dumay et al model [4]. Table 2.1.1.3: Calculated SFE for TWIP steels at 298 K. Table 2.1.2.1: Md30γ ε evaluation from fig.2.1.2.2. Table 2.2.1.1: Peak stress at 1000°C of different TWIP steels. Table 2.2.1.2: Values of t50 at 1000°C and 1100°C relevant to TWIP steels variants. Table 2.2.2.1. Annealing treatments Table 2.2.2.2. Values of ksoft and B obtained from Avrami fittings to the recrystallization results obtained for different reductions and annealing temperatures. Table 2.2.2.3. Annealing processes map of the TWIP steel in T-t space. Not explored regions in blank.

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Table 2.3.1: thermal treatments for precipitation study. Table 2.4.1: Bending after a sag test with 3h soaking at 1200°C. Table 2.4.2: Calculated Tsolidus and Tliquidus of TWIP steels. Table 2.4.3: HRM Rolling schedule - TWIP 22Mn 0.6C steel- slab format 1000x250x28 mm Table 2.4.4: Calculated TWIP steel finishing rolling schedule for strip format: 1000x28x4 mm. Table 2.4.5: Results of the pickling tests. WP3 table Table 3.1.1.1: Tensile test at 20°C results (cold rolled and annealed at 1000°Cx5min) Table 3.1.1.2: Tensile test results at different temperature with mechanically polished surface Table 3.1.1.3: Values of physical parameters used in modelling Table 3.1.1.4: Obtained model fitting parameters Table 3.1.1.5: Microstructural constituents/phases before and after testing of TWIP steels at room temperature. Table 3.1.1.6 Mechanical twins occurrence during deformation at high temperature (strain rate 0.01s-1) Table 3.1.1.3.1: Annealing tests carried out to study the strip surface decarburization. Table 3.1.1.3.2: Comparison of Vickers hardness (HV200g) average values on bulk and decarburized zone. Table 3.1.1.3.3 Table 3.1.1.3.4 Low strain rate tensile tests reference and hydrogen charged TWIP3 sample. Table 3.1.1.3.5 LECO RC-412 main features. Table 3.1.5: TWIP2, numerical results of tensile tests. Table 3.3.1 Charpy tests results. Table 3.4.1: Stress after a true strain of 0.5 WP4 Table Table 4.1.1.1 Erichsen tests results. Table 4.1.1.2 Erichsen testing parameters. Table 4.1.1.3 Characteristics of the TWIP_1 (Fe-22Mn-3Al-3Si) samples Table 4.1.1.4 Results of conventional Erichsen tests performed at ISQ. Table 4.1.1.5 Characters of the TWIP_2 samples after the high-speed Erichsen tests. Table 4.1.1.6 Characteristics of the TWIP_3 samples after the high-speed Erichsen tests. Table 4.1.1.7 Characteristics of the TWIP_4 specimens after the high-speed Erichsen tests. Table 4.1.1.9 Hardness (HV5) of the TWIP steels before and after high-speed Erichsen testing. Table 4.2.1 Thickness of coating layers. Table 4.2.2 EDS analysis. Table 4.3.1 Optimal laser welding parameters for each “TWIP” steel composition Table 4.3.2 ASTM grain size measurements using the circular interception technique according with the ASTM E 112 standard. Table 4.3.3 Docol DP chemical composition. Table 4.3.4 Docol DP mechanical properties. WP5: Table Table 5.1.1: TWIP compositions investigated in the ‘metaldesign’ project. Table 5.1.2 Range of elements for TWIP3 steel grade to be cast at Duferco steelworks. List of References Task 1.1.2 1. G.B. Olsen and M. Cohen, Metall. Trans. A, 7A (1976) 1897 2. O. Grässel, G. Frommeyer, C. Derder and H. Hofmann, J. Phys. IV, 7 (1997) C5- 3. S. Allain, J.P. Chateau, O. Bouaziz, S. Migot and N. Guelton, Mat. Sci. Eng. A, 387-389 (2004)

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European Commission

eUR 25063 — Metallurgical design of high-strength austenitic Fe-C-Mn steels with excellent formability (Metaldesign)

A. Ferraiuolo, A. Smith, J. G. Sevillano, F. de las Cuevas, P. Karjalainen, G. Pratolongo, H. Gouveia, M. Mendes Rodrigues

Luxembourg: Publications Office of the European Union

2012 — 155 pp. — 21 × 29.7 cm

Research Fund for Coal and Steel series

ISBN 978-92-79-22205-4

doi:10.2777/14874

ISSN 1831-9424

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Five TWIP steel grades (Fe-Mn-C-N and Fe-Mn-Al-Si-C) were selected, adopting stack-ing fault energy (SFE) and austenite phase stability as the main metallurgical character-istics for alloy design. The selected TWIP grades were characterised in terms of SFE, mechanical properties, microstructure, recrystallisation, welding and forming ability.

The performance analysis revealed that only TWIP2 and TWIP3 clearly behave as expected from TWIP steels, with the typical hardening stage associated with profuse deformation twinning. The variants with lower Mn and C content (TWIP1, TWIP4, TWIP5) require a further set-up of steel chemistry to better balance the SFE and avoid the for-mation of second phases (ferrite or deformation-induced martensitic phases). The results achieved in the project allowed to define the suitable industrial manufacturing route to produce the TWIP steel. However two main issues resulted with an importance stronger than expected in the technical annex: hydrogen embrittlement susceptibility and strip surface decarburization. The first issue has suggested to delay the industrial heat sched-uling to evaluate the risk of hydrogen embrittlement for TWIP steels and the need to include a vacuum degassing facility for industrial TWIP steel manufacturing. The decar-burization issue revealed that during the annealing process a proper atmosphere control is necessary to avoid within the decarburised layer the formation of deformation induced martensitic phases (α’+ε).The activity results revealed that the investigated TWIP steels have excellent tensile properties together with good welding and galvanising ability that make them extremely attractive for automotive applications both for structural reinforce-ment (body in white) and for energy absorption.