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Cold-rolled complex-phase (CP) steel
grades with optimised bendability,
stretch-flangeability and anisotropy(CP-Steels)
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European Commission
Research Fund for Coal and SteelCold-rolled complex-phase (CP) steel grades
with optimised bendability,
stretch-flangeability and anisotropy
(CP-Steels)
L. RydeSwerea KIMAB
Box 55970, SE-10216 Stockholm, SWEDEN
O. Lyytinen, P. PeuraRuukki Metals Oy
Harvialantie 420, 13300 Hmeenlinna, FINLAND
M. TitovaRWTHDepartment of Ferrous Metallurgy, Inzestrae 1, 52072 Aachen, GERMANY
Y. Vilander GranbomSSAB EMEA
SE-781 84 Borlnge, SWEDEN
T. Hebesbergervoestalpine Stahl
Research & Development B3E, Materials Development, Voest-Alpine-Strae 3, Postfach 3, 4031 Linz, AUSTRIA
Contract No RFSR-CT-2006-00021
1 July 2006 to 30 June 2010
Final report
Directorate-General for Research and innovation
2012 EUR 25041 EN
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Luxembourg: Publications Office of the European Union, 2012
ISBN 978-92-79-22146-0
doi:10.2777/97678
ISSN 1831-9424
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2.6 WP4 - the relation between microstructure and formability ................................................................ 45
2.6.1 Objectives and tasks of WP4 ..................................................................................................................452.6.2 Task 4.1 Formability and Task 4.2 - Microstructures .......................................................... ................ 45
2.6.2.1 Formability and tensile testing .......................................................................................... ............. 452.6.2.2 Characterisation of microstructure and texture .............................................................................. 462.6.2.3 Properties and microstructures of the industrial material, Task 1.1 ............................................... 462.6.2.4 Properties and microstructures of samples from task 1.4 - Laboratory produced material ............ 472.6.2.5 Properties and microstructures of samples from task 3.1 - Initial laboratory annealing
simulations ........................................................................................ ............................................. 482.6.2.6 Properties and microstructures of samples from task 3.2 a wide range of microstructures......... 512.6.2.7 Properties and microstructures of samples from task 3.3 special microstructures from
dedicated lab annealing ........................................................... ....................................................... 562.6.2.8 Properties of the samples from task 5.1 - industrial trials .............................................................. 59
2.6.3 Task 4.3 Characterisation of the failure mechanism on pre-strained industrially produced materials .... 612.6.3.1.1 Fracture analyses on the fractured hole expansion specimens .................................................. 612.6.3.1.2 Analysis of the microstructure beneath the punched edge surface after HE testing ................. 62
2.6.4 Task 4.4 Critical evaluation and correlation of the formability, microstructure including textureand the detected failure mechanism, conclusions of WP4 ................................................................... ... 63
2.7 WP 5: Semi-industrial/industrial processing of optimised material ..................................................... 69
2.7.1
Objectives and tasks of WP5, industrial processing of optimised material ............................................ 69
2.7.2 Task 5.1 Processing of material with optimised microstructure on industrial scale ............................... 692.7.2.1 Initial trials of HDG at Ruukki.............................................. ......................................................... 692.7.2.2 Final on line trials of HDG at Ruukki .................................................................... ........................ 702.7.2.3 Previous trials of CAL-WQ at SSAB .................................................................. ........................... 73
2.7.2.3.1 Annealing simulation using dilatometry at KIMAB ................................................................. 732.7.2.3.2 Annealing simulation at voestalpine ........................................................................................ . 74
2.7.2.4 On line trials for CAL-WQ at SSAB ....................................................... ...................................... 742.7.2.4.1 Summary of on line trials of CAL-WQ at SSAB ............................................................... ....... 75
2.7.2.5 Final on line trials of CAL-GQ at voestalpine ............................................................................ ... 752.7.3 Task 5.2 Critical evaluation of the results from WP4 ............................................................................. 772.7.4 Task 5.3 Critical comparison of the objectives of the project and the results obtained in the project .... 792.7.5 Task 5.4 Basic characterisation of the formability based on forming limit diagrams ............................. 81
2.7.6
Task 5.5 impact of the strain rate and temperature on the mechanical properties samples fromindustrial trials ............................................... .................................................................... ..................... 852.7.6.1 Tensile testing, room temperature and at -40C and 100C ........................................................... 852.7.6.2 Dynamic tensile tests ............................................................ ......................................................... 86
2.7.7 Task 5.6 impact of the strain state on mechanical properties ..... ............................................................ 862.7.8 Task 5.8 constitutive modelling and determination of the parameters relevant for the selected
CP grades ................................................................... ................................................................... .......... 88
2.8 Conclusions ......................................................................................... ....................................................... 91
2.9 Exploitation and impact of the research results ................................................................. .................... 93
3.
LIST OF FIGURES, TABLES, ACRONYMS AND ABBREVIATIONS ................ 94
3.1 List of Figures ............................................................................................................................................94
3.2 List of tables ...............................................................................................................................................97
3.3 List of acronyms and abbreviations ......... .................................................................... ........................... 98
4. LIST OF REFERENCES ...................................................................................... 99
5. APPENDICES.................................................................................................... 100
5.1 Appendix 1 All included steels per task and annealing simulation schedules ................................... 100
5.2 Appendix 2 Results from Industrially produced material with varying texture, task 1.3. ............... 103
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5.3 Appendix 3 Results from initial laboratory simulations, task 3.1 ....................................................... 103
5.4 Appendix 4 Results from laboratory simulations of a wide range of microstructures, task 3.2 ...... 104
5.5 Appendix 5 Results from laboratory simulations of dedicated microstructures, task 3.3 ................ 105
5.6 Appendix 6 Results from final on line trials, task 5.1 ......................................................... ................. 106
5.7 Appendix 7 Results from final on line trials; Effect of strain rate and temperature, task 5.5 ......... 106
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1. FINAL SUMMARY
1.1 MAIN OBJECTIVES
The over all goal of this project was to identify microstructures with optimised balance betweenstrength, bendability and stretch-flangeability and to develop guidelines of how to produce this coldrolled sheet steel through three processing routes, i.e. continuous annealing with gas cooling,continuous annealing with quenching and the third route is hot dip galvanizing. The objectives were to
produce material via the three routes with the following properties; A tensile strength of more than800MPa, a hole expansion ratio (HE) of more than 35%, a bending angle, Ri/t
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The CCT diagrams determined show very different behaviour for the investigated steels. It seems thatthe additions of slowly diffusing elements (Mn, Mo and Cr) have a stronger impact on thetransformation kinetics than C. The maximum continuous cooling rate used in the dilatometryinvestigations was 285C/s. This was enough to describe the transformation behaviour in a CAL(continuous annealing line) and HDG (hot dip galvanizing) line, with gas quenching, but too low of acooling rate for a CAL with water quenching.
1.2.4 WP3 Laboratory annealing simulations to produce different microstructuresWP3 of this project was concerned with producing different microstructures based on WP 1, 2 and 4.The objectives of this work package has been to narrow the processing parameters in three steps tooptimise the required properties. This meant producing sheet material in industrial production andlaboratory in quantities large enough to be able to analyse the forming properties in WP4.For all simulations the same cold rolled material was used, see Table 2.
Table 2. Chemical analysis, laboratory annealingMaterial Grade Supplier Gauge
[mm]C
[wt%]Mn
[wt%]Si
[wt%]Cr
[wt%]Mo
[wt%]Nb
[wt%]V1-CR VA CP800-Mo Voest 1.50 0.11 2.1 0.15 0 0.2 0.02
R4-CR CP800-1 Ruukki 1.2 0.14 1.7 0.18 0.3 0.15
S7-CR TS 800 SSAB 1.23 0.135 1.5 0.5 0.015
1.2.5 WP4 analysing properties and structures
In work package 4 (WP4) the relation between microstructure and formability was studied whichincluded analysis of tensile behaviour, hole expansion and bendability. The main delivery from this WPwas basic know-how of the impact of the microstructure on the formability and fracture behaviour.This work has also been performed stepwise when new material has been delivered from WP3.The first demand on the material was to obtain enough strength, i.e. Rm>800MPa and when this wasachieved the next property to optimise was the hole expansion or the bending behaviour, depending onwhich was the poorest.
The experimental findings for the investigated materials allow the following hypothesis to beestablished regarding stretch-flangeability:Cleanliness of the steels is an important factor and was not intended to vary within the frame of thisproject and thus will not be considered here.The strength/hardness difference between the matrix phases in the microstructure must not be high. Thismechanical homogeneity prevents local hot-spots of stress and/or strain in the microstructureespecially at phase boundaries which act as nucleation sites for voids. Several possibilities have beentried to reach this goal for CP-grades:
Reducing the grain size of the hard martensite
Replacing hard martensite by softer tempered martensite or bainite.
Increasing the strength of the matrix by grain size reduction, other strengthening mechanisms
(solution or precipitation hardening) and /or the introduction of more bainite. Reducing the amount of soft ferrite
The experimental findings for the investigated materials regarding bending is not as obvious as for holeexpansion but the best results were achieved with the following principles:
Reducing the grain size of all constituents in the microstructure.
Reduce the hardness difference between different phases Replacing hard martensite and/or soft ferrite by bainite or tempered martensite.
Ferrite is not poor for bending but the required strength level will not be reached if too much ferrite ispresent at the same time as martensite is avoided.
The experimental finding for nearly fully martensitic DP-grades TS1200M with 96.5% of temperedmartensite have also established that the homogeneous microstructure of this material provides highstrength and good hole expansion level, although the tensile elongation is low, and that this steel showsbetter HE than a dual phase microstructure with ferrite and martensite.
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Our results indicate that optimum crystallographic texture also can provide additional improvements inthe forming properties, but this was considered more difficult than manipulating the phase constituentsand was not considered in the final trials. It was shown that the best texture was achieved annealing inthe austenite-ferrite region, a processing condition that will be difficult to use in HDG or CAL-GQ.These processes normally use full austenitisation to achieve a bainitic transformation during theisothermal holding section. Without complete transformation during annealing it is likely that the yield
strength level will be too low without more additions of alloying elements.Another possibility to improve the texture is to increase the cold rolling reduction, and this was shownin laboratory trials but could not be tested in full scale.The obtained variations in crystallographic texture have been very small in all experiments except forthe specific tests performed under task 1.3.
Our findings have resulted in selected processing conditions for the industrial trials presented inWP5.
1.2.6 WP5 Industrial processing of optimised material and characterisation of the propertiesfor the end user
Semi-industrial / industrial processing of optimised material and characterisation of the material for theend user was the task for work package 5 (WP5). This included, besides the production of the material,characterisation of the formability, mechanical properties as a function of strain rate and temperatureand constitutive equations and flow curves.
The required properties from the initial project objectives; Rm>800MPa, Ri/t35% have allbeen met for all three production lines although the microstructures are different. Complex phasemicrostructures with bainite, martensite, ferrite and retained austenite were produced in the CAL-GQand HDG but it was particularly difficult to produce a bainitic microstructure in the continuouslyannealing line with water quenching. This microstructure was more of a dual phase type.
By comparing the forming limit diagrams of the initially delivered industrially produced steels and thesteels from the final industrial trials based on the results of this project, it is possible to draw some
conclusions of other forming properties: Improved biaxial balanced stretch forming, plane strain and stretch-flangeability has been
achieved for the steels produced through a hot dip galvanizing line and continuous annealingwith gas section (steels R and V). This improvement is a result of a reduction of the size of themartensite grains, elimination of the soft ferrite phase and by replacing of hard martensite bytempered martensite or bainite.
The continuously annealed and water quenched industrially produced materials have (steels S),however, better biaxial balanced stretch forming, plane strain, stretch-flangeability and uniaxialtensile properties in comparison with the results from the final industrial trails, due to a morehomogeneous microstructure and smaller martensite grain size.
The project has resulted in a new product for SSAB, a LCE1000, which has a leaner chemical
composition compared to DP1000, improved hole expansion due to less difference between the phaseconstituents and better weldability due to lower carbon equivalent.
1.2.7 WP6 Coordination and reports
Co-ordination and reports (WP6) was taken care of by the co-ordinator KIMAB.The project was delayed and granted an extension with one extra year. The reason for this was twofold:
The tasks of WP3 and WP4 could not run in parallel with WP2, as was intended in the initialplan, since the transformation kinetics were needed to set up the laboratory trials
The financial crisis during 2008-2009 made all on line trials impossible at that time.All reports were delivered on time
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1.3 CONCLUSIONS
The following conclusions can be drawn from the work performed in this project:General conclusions:
The project objectives have been met; improving the properties of steel products from threedifferent annealing lines to obtain a tensile strength of >800MPa, bending radii 35%
A new steel, SSABs LCE1000, has been the outcome of this project together with adjustmentsof the existing process parameters and compositions for the CP-steels with Rm 800MPa atRautaruukki and voestalpine
The properties and microstructures are not identical between the three products so there is stillroom for improvement
It is very difficult to obtain a bainitic structure in the CAL-WQ process due to the lack ofintermediate temperature holding zones.
About microstructure, texture and hole expansion (stretch-flangeability) or bending:
The strength/hardness difference between the matrix phases in the microstructure must not behigh. This mechanical homogeneity prevents local hot-spots of stress and/or strain in themicrostructure especially at phase boundaries which act as nucleation sites for voids. Several
possibilities have been tried to reach this goal for CP-grades:- Reducing the grain size of the hard martensite- Replacing hard martensite by softer tempered martensite or bainite.- Increasing the strength of the matrix by grain size reduction, other strengthening
mechanisms (solution or precipitation hardening) and /or the introduction of morebainite.
- Reducing the amount of soft ferrite
The experimental findings regarding bending is not as obvious as for hole expansion but thebest results were achieved with the following principles:
- Reducing the grain size of all constituents in the microstructure.
- Reduce the hardness difference between different phases- Replacing hard martensite and/or soft ferrite by bainite or tempered martensite.- Ferrite is not poor for bending but the required strength level will not be reached if too
much ferrite is present at the same time as martensite is avoided.
The texture investigation implies that a strong {111}-texture should be beneficial for the holeexpansion properties.
1.4 IMPLEMENTATION OF THE RESEARCH AND TRANSFERABILITY OF RESULTS
One of the major goals of this project has been to develop guidelines of how to produce the material, byan appropriate alloy design, and well-adjusted processing parameters through several processing routes,
i.e. continuous annealing and hot dip galvanizing, which should make the results transferablethroughout Europe. This has been successful in the lines available in this project and since the mainthree commercial annealing processes are covered, the results of this project will be able to benefit notonly the three industrial partners of this consortium, but many other steel producers in Europe. Onlyminor adjustments are intended in each type of facility to produce advanced high-strength CP material.
Another objective was to characterise the material for the end user, which will aid the user to use thefull potential of the high strength material. This has been made through FLDs and flow curves of all thenew steels and are presented in the chapter on WP5 and in the appendices.
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2. SCIENTIFIC AND TECHNICAL DESCRIPTION OF THE RESULTS
2.1 OBJECTIVES OF THE PROJECT
This project has dealt with the relationship between microstructure and formability of cold-rolledcomplex-phase steels. The over all goal of this project was to identify microstructures with optimisedbalance between strength, bendability and stretch-flangeability and to develop guidelines of how toproduce these cold rolled sheet steel through three processing routes, i.e. continuous annealing with gas
cooling, continuous annealing with quenching and the third route was hot dip galvanizing.The objectives were:
- To achieve an improved strength-formability balance with particular focus on bendability,stretch-flangeability and deep-drawing capacity. A tensile strength of more than 800MPa, ahole expansion ratio (HE) of more than 35%, a bending angle, Ri/t
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2.3 WP1 PROCESSING /DELIVERY OF DIFFERENT STEEL GRADES
2.3.1 Objectives and tasks of WP1
WP1 deals with processing and delivery of different steel grades. Some laboratory materials are alsoincluded to be able to vary parameters, e.g. processing parameters and texture, more widely than ispossible from current industrial production. The objectives were the same as the tasks of this work
package and were divided into 5 tasks listed in Table 3 below.
Table 3. Tasks of WP1Task 1.1 Delivery of fully industrially produced material from
three processing routesvoestalpine, SSAB, Ruukki
Task 1.2 Industrially as cold-rolled material Ruukki, SSAB, voestalpine
Task 1.3 Processing of material with different texture KIMAB, SSAB
Task 1.4 Lab processed material voestalpine
Task 1.5 Characterisation of all the delivered materials. Ruukki, SSAB, voestalpine
2.3.2 Description of activities and discussion of WP1
2.3.2.1 Task 1.1 Delivery of fully industrially produced material from three processingroutes
The following materials has been selected from the shelf, see Table 4. All have a tensile strength ofmore than 800MPa.
Table 4. Selected materials, (wt%)Material Supplier Gauge C Mn Si Cr Mo Nb Rp0.2
[MPa]Rm
[MPa]A80[%]
Rp0.2/Rm
R1 Ruukki 1.20 0.17 1.7 0.18 0.3 0.15 626 855 16 0.73
R2 Ruukki 2.50 0.17 1.7 0.18 0.3 0.15 574 846 14 0.68R3 Ruukki 1.20 0.17 1.7 0.18 0.3 0.16 568 836 14 0.68
V1 voest 1.50 0.11 2.1 0.15 0 0.2 0.02 720 920 11 0.78
V2 voest 1.50 0.14 2.1 0.2 0.3 0 0.02 620 790 16 0.78
S1 SSAB 1.50 0.144 1.52 0.48 0.018 640 840 15 0.76
S2 SSAB 1.80 0.172 1.59 0.51 0.017 1060 1140 6 0.93
S3 SSAB 1.00 0.170 1.63 0.47 0.017 940 1030 8 0.91
S4 SSAB 1.60 0.178 1.41 0.44 0.016 1190 1220 4 0.98
S5 SSAB 1.65 0.135 1.53 0.18 0.015 970 1200 5 0.81
A description and motivation for the selected steels are given below:
Rautaruukki has chosen one steel grade for the starting point of complex phase (CP) -steeldevelopment, a hot dip galvanized LITEC 800DPF. The steel grade represents the complex phase steelsof Rautaruukki at the time for the start of the project.Rautaruukki has processed and delivered two gauges of this CP800-steel grade: thicknesses 1.20 and2.50 mm, R1 and R2. The material was analysed and tested at Ruukki, KIMAB and RWTH. The twogauges of CP800 are produced via different galvanizing lines i.e. they have the same alloy design, buthave different annealing cycle. The composition is shown in Table 4. The same material and thicknesshas also been processed and delivered in a cold-rolled state (task 1.2). This has been used for laboratoryannealing and dilatometry simulations at the University of Oulu (subcontractor of Ruukki) and atvoestalpine (only 1.20 mm thickness). These materials were used for adjusting different microstructuresin laboratory annealing simulators in order to obtain optimal properties for CP-steels. Based on the
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transformation characteristics found in WP2, Rautaruukki developed a new coil that was included asindustrial material for further tests, the R3. This development will be described later in the report.
voestalpine has selected two CP-grades with a minimum tensile strength of 800MPa. The chemicalcompositions of these two grades are given in Table 4.Generally, CP-grades have a complex microstructure consisting of bainite, martensite, temperedmartensite, and minor polygonal ferrite. In comparison to DP-grades, such a microstructure is
characterised by a high yield ratio and a lower uniform and total elongation in the tensile test. However,CP-grades show excellent bendability and stretch-flangeability.One main task for producing such microstructures is to suppress or at least minimise the formation ofpolygonal ferrite.Molybdenum is very effective in avoiding the formation of polygonal ferrite and therefore it is oftenfound in CP-grades (V1, VA_CP800_Mo). However, molybdenum has also its disadvantages as itincreases the hot rolling forces significantly and it is also quite expensive. The second steel grade (V2,VA_CP800) is a try to compensate molybdenum with chromium and higher carbon content.
Four different steel grades of a DP-type have been chosen for the project by SSAB,see Table 4.These grades have a tensile strength between 800 and 1200 MPa after continuous annealing.For a tensile strength of 1000 MPa, two coils with different thickness have been processed.For a tensile strength of 1200 MPa, two coils with different properties have been processed. One is afully martensitic steel and the other has 10 to 15 % second phase.
Below, in Table 5, you find the annealing temperatures for the chosen steel grades.
Table 5. Annealing temperaturesMaterial Grade Thickness
[mm]Annealing
temperature [C]Overageing
temperature [C]
S1 TS 800 1.50 760 400
S2 TS 1000-1 1.80 810 400
S3 TS 1000-2 1.00 790 400S4 TS 1200 (M) 1.60 850 400
S5 TS 1200 (DP) 1.65 850 200
2.3.2.2 Task 1.2 Industrially as cold-rolled material
The industrially as cold rolled materials selected are all aiming for the same strength level, with atensile strength of more than 800 MPa. These were the steels used in most laboratory simulations, seeTable 6. The only exception is the final task 3.3 where new steels were added. Those steels aredescribed in the chapter 2.5.2.3.
Table 6. Chemical composition of the industrially cold rolled materials (wt%)Material Grade Supplier Gauge
[mm]C Mn Si Cr Mo Nb
V1-CR VA CP800-Mo voest 1.50 0.11 2.1 0.15 0 0.2 0.02
R4-CR CP800-1 Ruukki 1.2 0.14 1.7 0.18 0.3 0.15
S7-CR TS 800 SSAB 1.23 0.135 1.5 0.5 0.015
2.3.2.3 Task 1.3 Processing of material with different texture
Task 1.3 includes the processing of material with different texture both in laboratory and in thecontinuous annealing line by KIMAB and SSAB. voestalpine has assisted with annealing larger piecesfor mechanical properties and RWTH did hole expansion testing on those. The variations in texturewere obtained through variations in the rolling schedule and coiling temperature, see Table 7.
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Table 7. Task 1.3 Texture programme, chemical composition, coiling and cold reduction
Grade Material Supplier Gauge(mm)
C(%)
Mn(%)
Si(%)
Nb(%)
Condition
S6 TS 800 SSAB 1.5 0.13 1.51 0.21 0.015 Contin. An.
S1-HR TS 800 LAB SSAB 3.47 0.144 1.52 0.48 0.018 Hot rolled
Grade Material Coiling temperature[C]
Cold rolling[%]
S6-A TS 800-CT1 515 56%
S6-B1,S6-B2
TS 800-CT2 600 56%,69%
S6-C TS 800-CT3 660 56%
S1-Tex1S1-Tex2S1-Tex3S1-Tex4S1-Tex5
TS 800 LAB 600 50%,60%,70%,(77%,84%)
The first three grades are industrially processed by SSAB and their properties and structures will bediscussed in the following chapter.The laboratory material was delivered after hot rolling and coiling and was cold rolled to differentreductions in the laboratory. The highest cold rolling reductions, i.e. more than 70%, was not possible toachieve in the available semi sized laboratory cold rolling mill for the widths necessary for formingtests and further attempts to find an alternative mill was not successful in the available time frame. Thehighest reductions, 77% and 84%, were only achieved on samples too narrow to be tested for holeexpansion or tensile testing.Initial annealing simulations of small pieces for texture analysis and microstructure characterisationhave been performed on samples between 50% reduction and 84% in a quenching dilatometer. Basedon the results from these tests, larger samples were then annealed in voestalpines annealing simulator
with principally the same annealing schedules.The purpose of the laboratory tests was to find heating routes to vary the strength of the crystallographictextures while maintaining the martensite content at approximately 20-25% and thereby the requiredstrength level. The performed annealing schedules are shown in Figure 1. The annealing simulation wasmade in a Bhr 805 Dilatometer to ensure that the time-temperature schedule was controlled with highaccuracy. The dilatometry tests results in rather small samples, fit only for texture and microstructureanalysis.
Figure 1. Two temperature schedules were selected; one with annealing in the two phaseregion and one in the austenite region
0
100
200
300
400
500
600
700
800
900
1000
0 100 200 300 400 500 600
Time, s
Temperature,C
quenching
austenitised
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The annealing schedules were selected to simulate continuous annealing to produce a dual phase steelwith a strength level of approximately 800 MPa through:
1. annealing in the ferrite-austenite region and natural cooling in the gasjet section beforequenching at a relatively high temperature, above AC1
2. completely austenitise the material and then cool down with gas to a low temperature toallow for ferrite to transform and then quench to maintain approximately 25% martensite.
The overageing temperature was selected deliberately high, 350C, to be able to analyse the martensitetexture with EBSD. A change in the overageing temperature is not expected to alter the texture ormartensite content unless much higher temperatures are applied than are used here. In the largersamples, annealed by voestalpine, this is corrected for to achieve a correct strength. For the intercriticalanneal, called quenching the overageing temperature was set to 200C and for the austenitisingschedule the top temperature was increased to 900C to ensure complete phase transformation. For thelatter an overageing temperature of 400C was selected.
The textures of the investigated samples shown as ODF sections with 2=45 are found in Figure 2.The texture is weaker if the material is annealed above AR3 (labelled austenitised)compared to in thetwo phase region (AR1). The texture is also weaker at lower cold rolling reductions, as expected.
Another notable feature in the crystallographic textures of the samples annealed in the two phaseregion (labelled quench) is the strong -fibre, particularly in the 50% cold rolled material. This
indicates that there may be unrecrystallised regions in the material. For deep drawing, a strong -fibre
is beneficial and a strong -fibre is detrimental, but it is not certain how the texture will affect otherforming properties. The relationship between deep drawing and texture is well established for singlephase material but it is likely that multi phase material will behave somewhat different and theseresults show a promising range of textures to test this.
Austenitised50%CR
Austenitised60%CR
Austenitised70%CR
Austenitised77%CR
Austenitised84%CR
Quench 50%CR Quench 60%CR Quench 70%CR Quench 77%CR Quench 84%CR
Figure 2. The texture is weaker if the material is annealed in theaustenite region or if the cold rolling reduction is reduced.
ODF sections with 2=45
The microstructure of four samples (50% and 70% cold rolling and austenitised and quenchedrespectively) has been studied with EBSD to analyse the approximate amount of martensite and theresults are shown in Figure 3. The level of martensite after annealing in the austenite region ismeasured to minimum 22% and maximum 28%, which is promising in order to achieve the correctstrength level. However, the investigate area is rather small, so the results may be somewhat uncertain.The suspicion of partly recrystallised structured as indicated previously from the crystallographic
texture showed to be correct; the EBSD mapping of the 50% cold rolled sample annealed in the twophase region showed 30% unrecrystallised grains. There was no evidence of unrecrystallised areas inthe 70% cold rolled sample.
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Austenitised 50%CR, 22% martensite Austenitised 70%CR, 27% martensite
Quench 50%CR, 25% martensite (black),25% unrecrystallised structure(red) Quench 70%CR,28% martensite
Figure 3. The microstructure is showing approximately the same level of martensite (here indark grey or black) and there are unrecrystallised structures in the sample with50% cold rolling and annealed in the two phase region
Figure 4. Hole expansion ratio of annealed SSAB-materials with different texturesThe results of the hole expansion tests, see Figure 4, show that with an increase in the cold rollingreduction the hole expansion value increases. The hole expansion values of SSAB-grades, which havebeen annealed in the two phase region, are a little higher than the hole expansion of the samples whichhave been annealed in the austenitic range. This is an indication that the same textures that arefavourable for deep drawing may also be beneficial for a high hole expansion value, i.e. a strong {111}-ND texture is preferable. It is however, necessary to point out that the samples have different thicknessdue to the different cold rolling levels and that this also has an influence on the forming properties. Incomparison with the industrially produced TS800, the hole expansion values are substantially better forthe laboratory samples, 48% as minimum for the 50% cold rolled material annealed in the two phaseregion, whereas the industrially produced had a hole expansion ratio of 35%.
2.3.2.3.1 Industrially produced material with varying texture
SSAB provided material with different textures through variation of coiling temperature after hotrolling and cold rolling reduction as described in Table 7.
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The mechanical behaviour of the different TS800 texture materials were investigated both in the as hotrolled condition and after cold rolling and annealing, i.e. the finished DP-product. For the as hot rolledmaterials, the true stress-true strain curves are presented in Figure 5 and the values of the yield andtensile strength and elongation in Table 8 and micrographs of the structures (optical micrographs) areshown in Figure 6. As can be seen from graphs and tables, an additional coiling temperature of 470C isincluded. This material was only added for comparison here and is not tested for hole expansion.
Figure 5. True stress-true strain of the as hot rolled materials.From the curves above we see that the lower the coiling temperature the higher the yield strength. Wealso achieve a pronounced Lders strain for the higher coiling temperature as well as a larger fractureelongation. The differences are pronounced.
Table 8. As hot rolled condition:C.T. [C] YS [MPa] TS [MPa] A80 [%]
660 444/442 551/550 24/23
600 513/520 615/616 16/14
515 574/572 667/665 14/8*
470 585/590 680/679 11/11*) Fracture outside the gauge length
The microstructures of the texture materials in the as hot rolled condition are presented in Figure 6.A more homogenous microstructure is achieved when the material is coiled at lower temperatures.Higher coiling temperatures (660C and 600C) result in a banded structure where the dark bandsconsist of pearlite. The structure of the lower coiling temperatures is best described as bainitic.
Figure 6.
Microstructure of the as hot rolledmaterials.
300
400
500
600
700
800
0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14 0.16Strain [mm]
Stress[MPa]
C.T. 470CC.T. 515C
C.T. 600C
C.T. 660C
C.T . 600C C.T. 515C C.T. 470CC.T . 660C C.T . 600C C.T. 515C C.T. 470CC.T . 660C C.T . 600C C.T. 515C C.T. 470CC.T . 660C
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After cold rolling and annealing, the difference in mechanical properties is smaller; see Figure 7 andTable 9. The sheet with the lowest coiling temperatures (470C and 515C) exhibit higher yield andtensile properties than the ones with the higher (660C and 600C) coiling temperatures.
Figure 7. True stress-true strain of the cold rolled and annealed samples.Table 9. Mechanical properties of the finished products (DP-steels)C.T. [C] YS [MPa] TS [MPa] A80 [%]
660 552/553 854/857 16/14
600 553/549 839/839 13/14
515 578/571 882/877 15/13
470 566/567 873/874 13/14
The microstructure after cold rolling and annealing seem to be reversed in fineness and bandingcompared to the hot rolled structures; Coiling at temperatures above 600C seems to result in thesmallest and the most homogenous ferrite grain size, as well as the most finely dispersed martensiteparticles, see the micrographs in Figure 8.
Figure 8. Microstructure of the DP texture materials after cold rolling and annealing.An explanation to this behaviour may be found in the effect of Nb. The precipitation of Nb-carbides isknown to be most efficient around 600C, and may be incomplete at lower. It is likely that highercoiling temperatures will allow the particles to grow somewhat.
Besides varying the coiling temperature, the cold rolling reduction was also varied; a TS800 strip wassubjected to two cold rolling reductions; one at low reduction (L.R. 56%) and one with high reduction
(H.R. 69%). Looking at the true stress-true strain curves presented in Figure 9, there does not seem tobe a significant difference in mechanical properties after annealing as a result of the different rollingreductions.
400
500
600
700
800
900
1000
0 0.02 0.04 0.06 0.08 0.1 0.12
Strain [mm/mm]
Stress[MPa]
C.T. 660C
C.T. 600C
C.T. 515CC.T. 470C
C.T 600C C.T 515C C.T 470CC.T 660C C.T 600C C.T 515C C.T 470CC.T 660C
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for the coil with the higher cold rolling reduction. A higher reduction should increase the driving forcefor recrystallisation, which is why only a shorter annealing or a lower temperature could explain theresults.
CR 56%, CT 515CCR
56%, CT 600C CR 56%, CT 660C
CR 69%, CT 600C
Figure 11. The crystallographic texture of the industrially produced coils with varying texture.ODF sections with 2=45
EBSD confirmed that there were deformed areas in the microstructure of the 69% cold rolled material
and only recrystallised grains with the typical -fibre orientation in the material that was cold rolled56% prior to annealing, see Figure 12. The orientations are coloured relative to the end orientation of
the -fibre, {001} and allowing 7 spread. Grains with an orientation that deviates more than 7are grey.
a) b)
Deviation from {001}
Figure 12. EBSD maps of the a) 56% cold rolled and b) 69% cold rolled steel. Coilingtemperature after hot rolling of 600C. The orientations are coloured relative to
the end orientation of the -fibre, {001}.
Tensile data and hole expansion results are also shown in appendix 2, chapter 5.1.The results from thehole expansion tests of the industrially produced textured material are shown in Figure 13. The besthole expansion value was found for material coiled at 600C and cold rolled to 56%. This materialshowed a hole expansion value of 51% and this result is much better than the industrially producedTS800, that only had 35%. The main difference between these two materials seems to be a 100MPalower yield stress for the new coil (CT600C-CR56%), which may explain the better hole expansion.The explanation for the poor result of the higher cold rolling (69%) is probably that the unrecrystallised
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regions, which were found with EBSD, will reduce the forming properties. The poor result for thematerial with the high coiling temperature, CT 660C, is probably also caused by unrecrystallised
remaining areas, as indicated by the strong -fibre in the texture. The pronounced banding may alsodecrease the HE results.
Figure 13. Hole expansion tests of the various cold rolling reductions and coilingtemperatures2.3.2.3.2 Conclusions regarding texture and hole expansion from task 1.3
The conclusions from the textured materials are that:
A strong {111}-ND texture seems to be beneficial to obtain a high hole expansion ratio. Thistexture is also known to result in high r-values.
The best production route to achieve a strong {111}-ND texture and is to use a coilingtemperature around 600C, as high cold rolling reduction as possible, anneal the material inthe two phase region BUT ensure that the material is completely recrystallised.
2.3.2.4 Task 1.4 Lab processed materialFor the processing of different microstructures based on lab processed material, Rautaruukki studied theeffect of microalloying on the structure and mechanical properties of CP-steel. At the same time thepossibility of reducing the amount of Mo was studied. Two laboratory heats, Ru-4 and Ru-5, wereprepared with 0,015 wt% Nb additions without or with Mo (Table 11).
Table 11. Chemical compositions of the investigated steels in wt%.Material C Si Mn Cr Mo Nb
Ru-1 (normal CP800) 0.17 0.18 1.7 0.3 0.15 -
Ru-4 0.17 0.18 1.7 0.3 - 0.015
Ru-5 0.17 0.18 1.7 0.3 0.15 0.015
The heats were cast by using a vacuum induction furnace. After cooling, the ingots were wrapped instainless steel foil heated to 1200C and held for 24 hours for homogenizing. After removing the scaleand cutting the ingots to suitable sizes the pieces were reheated in Argon shielding gas to 1240C forone hour and hot rolled. The finish rolling temperature was 750 - 800C and samples were cooled down(at 1-2C/s) to a temperature of 500 - 600C, where they were held to simulate coiling. After this, thesamples were cold rolled with 50 % reduction. The cold rolled samples were heat treated according tovarious temperature schedules in a Gleeble 1500 simulator to simulate the galvanizing line. Theannealing and holding zone temperatures were varied. Two tests were carried out for each simulation
cycle. One was used for testing the RT-tensile properties and the other specimen for microstructureanalysis employing the LePera's etchant, see chapter 2.6.2.4 in WP4.
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2.3.2.5 Task 1.5 Characterisation of all the delivered materials.
The microstructure analysis resulted in the following phases in the respective material, see Table 12. Anadditional coil was also included from Ruukki; this coil was produced based on the phasetransformation results under WP2 and should belong to the tasks under WP3 (Processing of materialwith different microstructure and texture) but is included here since both the microstructure andproperties of the industrial materials are compared here.
Table 12. Summary of the phase constituents (%) and properties of the different industriallyproduced steels
Phase\Material R1 R2 R3 S1 S2 S3 S4 S5 V1 V2
Thickness, mm 1.2 2.5 1.25 1.5 1.8 1 1.6 1.65 1.5 1.5
Ferrite(F) 37.0 37.0 15 76 20 30 3.5 14 42 36
Martensite(M) 13.0 13.0 12.0 4 7
TemperedMartensite(TM)
2.6 2.6 24 80 70 96.5 86 15
Carbide freeBainite(CfB)
42.0 42.0 29 27 57
Bainite (B) 5.4 5.4 44 12
Rm, [MPa] 855 846 820 840 1140 1030 1220 1200 920 790
Ri/t 1.2 3 1.2 0.7 3.5 1 3.5 1 3 1.7
HE [%] 19 17 56 35 30 71 44 40 47 58
2.3.3 Conclusions from WP1, delivery of material
All included industrial materials have tensile strengths above 800 MPa which is the target of thisproject. It is clear that this strength level can be obtained through a large variation in phase constituentswhich allows for freedom to design a suitable microstructure for the required propertiesIt seems that the bending properties need to be improved to meet the required level, except for a few of
these steels. The best bending properties are found for the continuously annealed and quenched steels,whereas the bending properties of the CP-structures after hot dip galvanizing or continuously annealedwith gas cooling, needs to be better.The claimed good hole expansion for a complex phase structure is found for the steels of voestalpine,however the first steels delivered by Ruukki, with a CP-structure, showed poor HE. The explanation forthis lies in the high amount of soft ferrite in combination with a large fraction of untempered martensite.Since the hot dip galvanizing line does not allow any overageing, tempering of the martensite is notrealistic and the hardness difference between the phases must be reduced by removing the soft ferrite. Inthe third material that Ruukki included, much of the soft ferrite is replaced by bainite, and the HE-valueis dramatically enhanced. The improved forming properties are likely to be related to the morehomogeneous structure, and the reduced difference in hardness between the phases. The hole expansionvalues of the steels with very high levels of martensite also strengthens this theory. These HE values are
around 40%, which implies that strength alone does not deteriorate the hole expansion, but rather thedifference in hardness between soft and hard phases, and good HE can be achieved when the fraction ofthe hard phase is large enough to make it continuous.
Furthermore, the hole expansion value does not only depend on the phase fraction of material, but alsoon the thickness of the sheets. The thicker TS1000-1 shows a much poorer HE than TS1000-2, eventhough the phase constituents are nearly the same. The steels from Ruukki are also delivered in twogauges, 1.25mm and 2.5mm respectively, but the thickness effect is not obvious here, since the HEvalues are very poor for both steels for other reasons discussed earlier.
The texture investigation implies that a strong {111}-texture should be beneficial for the hole expansionproperties.
Very promising results were obtained in the laboratory tests by modifying the chemical compositionusing Nb microalloying. Some additional study is still needed, in order to reduce the amount of Mo.
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2.4 WP2 TRANSFORMATION BEHAVIOUR DURING ANNEALING
2.4.1 Objectives and tasks of WP2
Transformation behaviour during annealing was analysed in order to find processing windows, i.e.chemical composition, time and temperatures, to produce the desired balance between selected phases
in the microstructure. The main work was performed using interrupted annealing trials, dilatometry andThermoCalc calculations and by analysing the microstructures in SEM and TEM. Dictra calculationswere also used but with limited usefulness. The objectives were the same as the tasks and the work wasdivided into four tasks, see Table 13:
Table 13. Tasks of WP2Task 2.1 Kinetics of the cementite dissolution and austenite
formation in the two phase region and austenite rangeKIMAB, voestalpine,Ruukki
Task 2.2 Transformation kinetics during cooling Ruukki, voestalpine
Task 2.3 Calculation of the transformation behaviour duringheating, soaking and cooling
KIMAB, voestalpine
Task 2.4 Delivery of know-how KIMAB, voestalpine,Ruukki
Task 2.1:Kinetics of the cementite dissolution and austenite formation in the two phase region andaustenite range has been investigated for three steels by interrupted annealing trials and dilatometricinvestigations and microstructure investigations.Task 2.2:Transformation kinetics during cooling has been studied for the same three steels with aGleeble and CCT diagrams for these are presented in the report. The microstructures of the sampleswere also compared with the Gleeble results.Task 2.3:Calculation of the transformation behaviour during heating, soaking and cooling based onthermodynamic (ThermoCalc) and kinetic (Dictra) modelling has been made, but the outcome of thekinetic calculations have not shown very useful and are extracted from this final report. The reason for
this is that the necessary assumptions of the allowed diffusion distances influence the result to such adramatic extent that the predictions were of no practical use; a small change in diffusion distances,based on microstructural observations, would change the predicted dissolutions time from too short totoo long, compared to those used in the annealing lines. The calculations were able to predict that themore alloyed steels were likely to need longer soaking times for dissolution compared to the leaner, butthis was also found in the laboratory studies. Therefore, it was better to rely on the laboratorysimulations and limit further use of Dictra calculations.The results from Tasks 1 and 2 are more valuable in determining new processing parameters formodifying the microstructure. Equilibrium calculations by ThermoCalc are helpful in identifyingdifferences in transformation temperatures, but since only a few of the studied annealing processesreach equilibrium and have high heating and cooling rates, this can only be used as a rough estimate.The used calculations are indicated in the graphs of task 2.1.Task 2.4:delivery of know-how is the summary of all results in WP1.
2.4.2 Description of activities and discussion of WP2
2.4.2.1 Task 2.1 Kinetics of the cementite dissolution and austenite formation in the twophase region and austenite range
The formation of austenite is of fundamental interest for CP grades, as bainite, martensite, and temperedmartensite typical constituents of a CP microstructure are products of the transformation fromaustenite during cooling. Therefore, experiments on the annealing simulator and in the dilatometer weredone by voestalpine. Three industrially produced grades were investigated. The basic difference
between the three compositions is the alloying levels of C, Mn and Mo, even though small fractions ofCr (Ruukki) and Nb are also present in (SSAB and voestalpine) the steels. The chemical compositionsof the grades are given in Table 14.
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Table 14. Chemical composition of the industrially produced CP-grades (wt%)Material Supplier C Mn Cr Mo Si Nb
V1 voestalpine 0.11 2.1 0 0.2 0.15 0.02
S7 SSAB 0.14 1.5 0 0 0.2 0.015
R4 Ruukki 0.16 1.75 0.3 0.15 0.18 0
The materials were heat treated with different parameters for time and temperature. The annealingtemperature (Tan) varied from 760C to 840C, the heating rate (HR) varied from 1/s to 80/s, and theholding time at (tan) was changed from 0 seconds to 80 seconds.After simulating the different time temperature cycles, the microstructure and the mechanical propertiesof the samples were investigated with optical microscopy and tensile testing. Carbide dissolution withFEG SEM was also studied after either etching with Picral (picric acid in ethanol) or in light electrolyticetch (3 V-10 s) in Chromium-acid (water solution of CrO3).
2.4.2.1.1 Carbide dissolution and microstructure change during annealingThe basic raw material was available as cold rolled (V1, S7, R4) and as hot rolled (V1, R4 only)In the Ruukki and SSAB materials the carbide phase appeared in the cold rolled condition in longpearlitic areas that were 100m long or more, but only one m or a few m wide. The distance betweenthese areas were somewhere between a few m and over 10m. Within these areas the carbide particlesoccurred as plates with thickness around 0.1m or rounded with diameters from 0.1- 0.25m.The voestalpine cold rolled material differed somewhat from the previous two. The pearlite area lengthwas from 50- 100m and the width from few m to 10m. There were also more rounded carbides,with diameters from 0.25m to 1 m.An example of the microstructure of the studied materials as cold rolled and as heat treated is shown inFigure 14. The two etching techniques used are also illustrated.
a) b)Figure 14. In a) an example of cold rolled Ruukki material and in b) after annealing,
HR=20/s, Tan=760C, tan=0s. In a) Picral was used as etchant and the carbidesappear white. In b) Cr-acid was used as etchant and carbides appear black.
Complete carbide dissolution is a prerequisite for the full utilization of the alloying elementsstrengthening potential. This is not only since carbon is the most important alloying element but alsobecause other important elements may be locked in the carbides; e.g. it is found that carbides areenriched with Mn, the second most important alloying element.The rate of carbide dissolution varies between the three steels, where the fastest to dissolve its carbidesis the steel from voestalpine. The slowest dissolution is shown in the steel from Ruukki, probably due to
the Cr-content of this material. The results show that carbide dissolution rate may be the factor thatcontrols the amount of austenite during the annealing if the soaking time is short or if the temperature isnot high enough.
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A summary of the results from the carbide dissolution analysis is shown in Figure 15 for the heatingrate of 20C/s, which is selected since this is the most relevant heating rate for the industrial annealingschedules involved. Complete dissolution of carbides is achieved at the indicated dotted lines.
Figure 15. Dissolution time of carbides at different temperatures for the three steels, V1, R4and S7 (voestalpine, Ruukki and SSAB).
2.4.2.1.2 Transformation of ferrite to austenite during heating and soaking.
Austenite formation was studied using dilatometry and microstructure studies. Dilatometry producedthe phase transformation curves and microstructure studies of the annealing simulator samples revealedthe amount of martensite in the steel, see Figure 16. The data presented here were collected from 0s and80s of annealing at various temperatures, i.e. quenched immediately after reaching the desiredtemperature or after a hold that corresponds to a normal annealing time in an industrial line. The
heating rate up to the selected temperatures was 20 C/s.
The austenite fractions from dilatometry data, the martensite fractions from the annealing simulationsand phase equilibrium calculated with ThermoCalc for various annealing temperatures at annealingtimes of 0s and 80s respectively for the three steels are compared in Figure 17.
A large fraction of austenite is already present immediately after heating up to the soaking temperatureand thus transform during heating. The results also shows that 80s of annealing is usually very close toequilibrium for all the investigated steels.There is a slight discrepancy between the results from the carbide dissolution studies and these datasince both the steel from Ruukki and SSAB seem to be closer to equilibrium than the steel fromvoestalpine, whereas the carbide dissolution study showed the Ruukki steel to be the slowest to
dissolve. It is however, possible that shorter annealing times would have given a different result. Insome of the graphs it seems that the volume fraction of austenite (or martensite) is higher thanequilibrium, which of course is impossible, but this is due to the precision when evaluating dilatometrycurves or optical micrographs of these fine structures.An example of the changes in microstructure is shown in Figure 16, along with the tensile strength as afunction of the annealing temperature and time.
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Figure 17. A Comparison between ThermoCalc calculations and results from dilatometry andmicrostructures after annealing simulations show that 80s of annealing is usuallyvery close to equilibrium for all the investigated steel.
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2.4.2.2 Task 2.2 Transformation kinetics during cooling
For the investigation of transformation kinetics during cooling the same three steel grades were selectedas in the previous tasks. The chemical composition of the three steels is given in Table 14. In the teststhe effect of process temperatures on the mechanical properties and microstructure was studied andCCT-diagrams were produced from the dilatometry data.
2.4.2.2.1 CCT-diagrams for transformation kinetics during coolingCCT diagrams were determined using the data acquired from the dilatation measurements using aGleeble 1500 simulator at University of Oulu. The specimens were heated at 20C/s to 860C(austenitic range), held for 40s prior to cooling at different rates. The target range for the cooling rates
was from 1C/s to 300-400C/s. For cooling rates of 100C/s, specimens were thinned to achieve thedesired cooling rates. The cooling rates of greatest interest from the point of view of development of CPsteels were quite different depending on the steel composition in question. The variety of differentcooling rates was included in order to achieve more complete CCT diagrams. Higher cooling ratesbeyond about 250C/s were not easily accessible on the Gleeble simulator, even for the thinnedspecimens. Samples can be thinned only to a certain level, until thinning will lead to a disturbingly largeslipping and/or tilting of the C-strain gauge.Austenite decomposition into ferrite, bainite, martensite or a mixture of these phases can be ascertained
by the inflexion point in the shape of the dilatation curves. The effect of cooling rate and the reheating(annealing) temperature on the phase transformation characteristics is clearly revealed in the dilatationcurves.
2.4.2.2.1.1 CCT diagram for the SSAB steel, S7The CCT diagram of SSAB steel is presented in Figure 18, which shows a large ferrite phase field(As/Af) extending over at least until 200C/s, beyond which it is possible to achieve bainite and/or
martensite by continuous cooling at 285C/s. It is possible that a very small fraction of high Cmartensite (M2) forms at lower temperatures, owing to C-enrichment of remaining austenite followingnearly complete ferrite transformation at higher temperatures. Even though C content is about 0.14% inthis steel, but in the absence of B and any strong austenite stabilizer except 1.5%Mn, the steel has
shown poor hardenability until at least a cooling rate of 100C/s. Reheating at a different annealingtemperature at 780C and cooling at 20C/s (broken cooling curve in Figure 18) did not show anysignificant effect on the phase transformation temperatures.
Figure 18. CCT diagram for the SSAB steel, S7
0
100
200
300
400
500
600
700
800
900
1000
1 10 100 1000
Time, s
Temperature,C
1C/s
5C/s
20C/s
50C/s
75C/s
100C/s
200C/s
285C/s
As
Af/Bf
M2
20C/s_RH780C
As_RH780C
Af_RH780C
F
M2
F+B
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2.4.2.2.1.2 CCT diagram for the voestalpine steel, V1The CCT diagram for the voestalpine steel is shown in Figure 19. It also revealed a large ferrite field,but a bainite nose appeared beyond about 50C/s cooling rate suggesting the possibility of achievingbainite and/or martensite beyond this cooling rate and nearly complete martensite beyond about100C/s by continuous cooling. This steel showed relatively higher hardenability obviously due to thepresence of Mo (0.2%) and higher Mn content (2.17%), despite the fact that C content was lowest(about 0.11%), if compared with that of other steels (Table 14). Mo is known to delay ferrite formation,though it is a strong carbide and/or ferrite former. As a result, the ferrite nose appears at about 50C/scooling rate, even though higher Mn too has a strong effect. Both Mn and Mo are known to delaybainite formation as well [6] and hence, move the bainite nose to right, so that it is possible to getcompletely displacive transformation of austenite to martensite by continuous cooling at rates over100C/s. Both the start (Ms) and end (Mf) of martensite temperatures could be easily identified fromthe dilatation curves and the Ms temperature is quite close to that predicted by equations given byStuhlmann [7]. At lower cooling rates, it is possible to achieve a very small fraction of high C-martensite (M2) as revealed by the inflexion in the dilatation curves and shown in the CCT diagram.Changing the annealing temperature from 860C to 780C prior to continuous cooling at 20C/s didnot, however, show any marked difference, even though the austenite fractions and their C contents canbe slightly different at the two reheating (annealing) temperatures.
It is noteworthy that both the steels from SSAB and VA have a small fraction of Nb, which is alsoexpected to increase the hardness and improve the hardenability, provided the cooling rate andcomposition are such that polygonal ferrite is not formed during cooling [8]. However, the data here isnot adequate to explain the effect of Nb in these CCT diagrams and also, the boron-alloyed steels usedby Somani et al. [8] were somewhat different, essentially meant for direct quenching.
Figure 19. CCT diagram for the voestalpine steel, V1
2.4.2.2.1.3 CCT diagram for the Ruukki steel, R4The CCT diagram for the Ruukki Steel is shown in Figure 20. The phase transformation behaviour inthis steel was somewhat between the two steels from SSAB and voestalpine. Even though the ferritephase field is large and nearly similar as seen in the CCT diagram of voestalpine steel (Figure 19), thebainite nose has extended beyond about 285C/s, thus depicting the fact that bainite will be obtained inaddition to martensite (fractions depending on cooling rate) by continuous cooling at all cooling ratesbetween 50-285C/s. It is expected that nearly complete martensite can be achieved by fast cooling
0
100
200
300
400
500
600
700
800
900
1000
1 10 100 1000
Time, s
Temperature,
C
1C/s
5C/s
20C/s
50C/s
75C/s
100C/s
200C/s
285C/s
As
Af
Bs
Bf
Ms
Mf
M2
20C/s_RH780C
As_RH780C
M2_RH780C
F
B
M
M2
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beyond about 300C/s, even though there are no data to show that. The Ms and Mf temperatures can beconveniently computed from the inflexion in the dilatation curves and the prediction from Stuhlmannsequation [7] for Ms temperature is also close to those achieved here at different cooling rates, althoughthe C content can significantly vary owing to the formation of different fractions of bainite prior to theformation of martensite. There appears to be a bay between ferrite and bainite phase fields extendingbetween 1-20C/s cooling rates. Inflexion in the curves at very low temperatures at 1-5C/s coolingrates suggest formation of a very small fraction of M2 martensite, but needs to be checked by
microscopy more closely. There seems to be some effect of reheating temperature at 780C, instead of860C normally used for all the experiments here. The ferrite transformation temperature is not verydifferent (720C), but bainite start (Bs) temperature is relatively lower by about 50 and no martensitecould be achieved. Presence of a very small fraction of M2 martensite cannot be ruled out, as marked inthe CCT diagram.Even though the steel from Ruukki has relatively higher C (0.16%C), it showed relatively lowerhardenability compared to that of voestalpine steel (0.11%C) with the bainite nose extending to veryhigh cooling rates up to at least 285C/s, possibly due to somewhat lower Mn and Mo contents andabsence of Nb, although there is a small fraction of Cr (0.3%) in Ruukkis steel. C, Mn, Nb, Cr and Moare all known to retard the bainite formation [6], but their individual effects are not easy to ascertain orcompute and also, the prior thermal history including the cooling rate can strongly influence thetransformation characteristics.
Figure 20. CCT diagrams for the Ruukki Steel, R4.
2.4.2.2.1.4 Correlation between microstructure and CCT-diagrams for all steelsSelected specimens from the dilatation curves were examined for microstructures by optical microscopyto correlate to the CCT diagrams (Figure 21-Figure 24). The specimens were etched with LePera toreveal the distinction between martensite, bainite and ferrite phases. Microstructures of samplesannealed in the austenitic region (860C) with different cooling rates (1, 20, 100 and 200C/s) areshown in Figure 21 for the SSAB steel grade, in Figure 22 for the steel grade of voestalpine and inFigure 23 for the steel grade of Ruukki. Microstructures of all the three steel grades annealed in theintercritical temperature (780C) for 40 seconds and cooled 20C/s are shown in Figure 24.
0
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Time, s
T
emperature,
C
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285C/s
As
Af
Bs
Bf
Ms
Mf
M2
20C/s_RH780C
As_RH780C
Bs_RH780C
Bf_RH780C
M2_RH780C
F
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M2
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a) b) c) d)Figure 21. Microstructures of the SSAB steel S7, temperature 860C, cooling rate a) 1C/s b)
20C/s c) 100C/s d) 200C/s, holding time 40s, heating rate 20C/s, etched withLePera.
The microstructure of the steel grade from SSAB was quite coarse and banded with lower cooling ratesand a fine microstructure was only achieved with the highest cooling rates. The coarse grain structurewas probably due to the lack of Mo in alloying combined with relatively high C content.The microstructure of CP-steel of voestalpine was the finest of all the three CP-steels for all the coolingrates. The microstructure refinement was most likely due to alloying elements Mo and Nb [9].
a) b) c) d)Figure 22. Microstructures of the voestalpine steel V1, temperature 860C, cooling rate a)
1C/s b) 20C/s c) 100C/s d) 200C/s s, holding time 40s, heating rate 20C/s,etched with LePera
The microstructure of the steel of Ruukki was somewhere in the middle of the two other steels grades:banded with the low cooling rate (1 C/s) and finer with the higher cooling rates which is probably dueto the effect of Mo on the microstructure refinement.
The microstructures resulting from the intercritical annealing differed somewhat from the other results,see Figure 24. The steel grade of Ruukki seemed most banded, but the difference between the steelgrades was quite small. The steel grade of Ruukki has the highest C content and quite high Mn content,so the banded structure could be a result of segregation of these elements. All the microstructures fromthe intercritical annealing seemed to be only partly recrystallised.
a) b) c) d)Figure 23. Microstructures of the Ruukki steel R4, temperature 860C, cooling rate a) 1C/s b)
20C/s c) 100C/s d) 200C/s s, holding time 40s, heating rate 20C/s, etched withLePera
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a) b) c)Figure 24. Microstructures of CP800 of a) SSAB b) voestalpine and c) Ruukki, temperature
780C, cooling rate 20C/s, holding time 40s, heating rate 20C/s, etched withLePera
2.4.2.3 Task 2.4 Conclusions from transformation behaviour during annealing andcooling
The carbide dissolution kinetics will control the rate of the phase transformation if the annealingtemperature is low or the soaking time is short. The kinetics varies with the chemical composition of thesteel and it was shown that a Cr-addition will slow down the dissolution rate. A phase transformationclose to equilibrium was usually achieved at a soaking time of 80 s or more, although much of the phasetransformation is done during the heating, before the actual annealing temperature is reached.The CCT diagrams, determined in the dilatometry investigations, show very different transformationbehaviour for the investigated steels. It seems that the additions of slowly diffusing elements (Mn, Moand Cr) have a stronger impact on the transformation kinetics than C. The maximum continuous coolingrate used in the dilatometry investigations was 285C/s. This was enough to describe the transformationbehaviour in a CAL (continuous annealing line) and HDG (hot dip galvanizing) line, with gasquenching, but too low of a cooling rate for a CAL with water quenching.According to the dilatometry investigations, the steel from SSAB has the poorest hardenability, due to
its lean alloy composition. It requires cooling rates of more than 285C/s to transform to martensiteand/or bainite. The steel from voestalpine has the highest hardenability of the three investigated steels.This is due to the high Mn-content and the Mo-addition. Bainite can be achieved at cooling ratesbetween 50C/s and 100C/s and the material will transform to martensite at rates over 100C/s. Thesteel from Ruukki has an intermediate hardenability. Full martensite transformation will only beachieved at rates above 285C/s, but bainite is formed between 50C/s and 285C/s. The hardenabilitycan be explained by the chemical composition of the Ruukki steel, which is leaner, compared to thesteel from voestalpine, although the Cr-content may be responsible for the extended bainite field.
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2.5 WP3-DIFFERENT MICROSTRUCTURES
2.5.1 Objectives and tasks of WP3
WP3 of this project was concerned with producing different microstructures based on WP 1, 2 and 4.WP3 consists of task 3.1, 3.2, and 3.3 (Table 15). The objectives of this work package has been tonarrow the processing parameters in three steps to optimise the required properties. Details on the
annealing simulations are found in appendix 1.
Table 15. Tasks of WP3Task 3.1 Processing of different microstructures and textures
based on industrial as cold-rolled material or labprocessed material
Ruukki and voestalpine
Task 3.2 Processing of a wide range of special microstructuresand textures based on suggestions from WP 4
Ruukki and voestalpine
Task 3.3 Based on the work in WP 4 and WP2 microstructuresand textures resulting in particular excellentformability will be produced
Ruukki and voestalpine
2.5.2 Description of activities and discussion of results in WP3
As described previously, WP 3 consisted of three loops of annealing simulations in order to narrow theprocessing parameters stepwise. For all simulations the same cold rolled material was used,see Table 16.
Table 16. Chemical analysis, laboratory annealingMaterial Grade Supplier Gauge
[mm]C Mn Si Cr Mo Nb
V1-CR VA CP800-Mo voest 1.50 0.11 2.1 0.15 0 0.2 0.02
R4-CR CP800-1 Ruukki 1.2 0.14 1.7 0.18 0.3 0.15
S7-CR TS 800 SSAB 1.23 0.135 1.5 0.5 0.0152.5.2.1 Task 3.1: Processing of different microstructures based on industrial as cold-
rolled material
The first laboratory annealing performed in the annealing simulator were the following for the differentannealing cycles and steels, see Figure 25 to Figure 27. These were selected in order to investigate theeffect of varying processing parameters for each individual line. It was made using the steel of eachindividual partner; i.e. the V1 steel was used for the CAL with gas quenching, the R4 steel was used forthe hot dip galvanizing simulation and the S7 was used for the CAL process with water quenching.The schedule for a continuous annealing line with gas quenching is shown in Figure 25. The parametersthat can be adjusted in this process are the temperature of annealing (Tan), the quenching temperature(TQ) and temperature of overageing (TOA). It is also possible to adjust the times during the process but
this is done by changing the line speed and will affect all other parameters, which is true for all threeinvestigated processes. In this project we were interested in producing a product with a complex phasemicrostructure which includes different variants of bainite. The selected parameters were therefore:annealing in the fully austenitic range, gas quenching from one selected temperature, and finallyoverageing at different temperatures. The main part of the phase transformation from austenite tobainite is supposed to happen during the overageing stage, if ferrite formation during cooling can beavoided. The CCT diagram of the V1 steel, see Figure 19, indicates that the cooling rate must be morethan 50C/s and that the martensite start temperature is approximately 420C. The overageingtemperatures selected lies between 300C and 500C, see Figure 25, to produce various bainites withdifferent amounts of martensite. Some ferrite is normally formed during the cooling stages, whichenriches the austenite in carbon, which in turn decrease the Ms temperature. It is therefore not expectedthat any of the selected schedules would produce a fully martensitic microstructure.
The schedule for the annealing stage of a continuous annealing line with water quenching is shown inFigure 26. There is also a section with overageing in the line but that was not considered in this initialexperiment. The parameters that were adjusted in this experiment are the temperature of annealing (Tan)
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and the quenching temperature (TQ). The CCT diagram of this steel, see Figure 18, indicate that bainiteformation is very difficult to achieve. A wide range of annealing temperatures, both in the intercriticalrange and fully austenitic, followed by very different quenching temperatures were selected to evaluatethe possibilities.The selected schedules for the hot dip galvanizing process, see Figure 27, were focused annealingtemperature, since the CCT diagram, see Figure 20, indicated that annealing in the intercritical rangewould change the transformation kinetics, and on the overageing time, i.e. the time that is needed for
the bainitic transformation.
Cycle Tan[C]
TQ[C]
TOA[C]
1 830 700 500
2 830 700 450
3 830 700 400
4 830 700 3505 830 700 300
Figure 25. Schematic time-temperature cycle of the lab anneals simulating a continuousannealing line with gas quenching for steel V1.
Cycle T2[C]
T3[C]
CR1[C/s]
T4[C]
6 760 780 -2.5 730
7 740 760 -2.5 710
8 780 800 -20 4009 820 840 -8 680
10 820 840 -22 400
Figure 26. Schematic time-temperature cycle of the lab anneals simulating a continuousannealing line with water quenching for steel S7.
Cycle Tan[C]
TQ[C]
toa[s]
11 830 720 48
12 830 720 95
13 780 700 48
14 780 700 95
15 805 710 95
Figure 27. Schematic time-temperature cycle of the hot dip galvanizing laboratory cycles forsteel R4
The mechanical properties and the hole expansion (HE) values achieved in the first laboratory annealedsamples are presented in Table 17.
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Table 17. Properties of the first laboratory annealed samples
Material Cycle Gauge[mm]
Rp0.2[MPa]
Rm[MPa]
Ag[%]
A80[%]
HE[%]
V1 1 1.52 509 898 8.3 11.7 49
V1 2 1.52 411 766 12.7 18.0 64
V1 3 1.51 437 597 13.0 18.3 73
V1 4 1.52 466 806 10.2 14.5 71
V1 5 1.53 478 851 9.1 12.2 55
S7 6 1.24 576 1143 4.9 5.6 16
S7 7 1.23 541 1116 7.5 8.7 14
S7 8 1.23 463 959 12.0 14.5 18
S7 9 1.23 685 1078 2.3 3.5 10
S7 10 1.23 388 790 16.2 21.8 29
R4 11 1.19 530 812 9.3 13.3 59
R4 12 1.18 529 801 8.3 11.2 60
R4 13 1.18 438 821 8.1 11.6 36
R4 14 1.18 422 807 7.1 10.3 32
R4 15 1.18 498 801 7.3 11.1 64
2.5.2.2 Task 3.2: Processing of a wide range of special microstructures based onsuggestions from WP 4
In Task 3.2 ten different time-temperature schedules were tested by voestalpine in the annealingsimulator, with the aim to produce various combinations of different phases in the microstructure. The
same three materials, which were used in previous annealing simulations, were selected (Table 16).These ten annealing cycles are presented in Figure 28-Figure 30. The performed simulations of acontinuous annealing line with water quenching and separate overageing is shown in Figure 28 .Annealing both in the austenite region as well as in the intercritical range was performed and variousoverageing temperatures were investigated. Figure 29 shows how the simulations of a continuousannealing line with water quenching, soft cooling in the gas jet section and overageing was made.Annealing was performed only in the austenite region and the overageing temperatures were set to200C.Figure 30 illustrates the simulations of a continuous annealing line with gas quenching or hot dipgalvanizing line. Annealing was performed in the austenite region and various overageing temperatureswere tested.The detailed analysis of the annealed samples can be found in WP4. A short summery of the results of
this analysis is depicted in Table 18.The best combination of HE values and bending properties were found for cycle 9; which consists ofannealing in the fully austenitic region at 880C, gas cooling to 450C and overageing at thistemperature for 40s before cooling to room temperature.
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Figure 28. Simulations of a continuous annealing line with water quenching and separate
overageing. Annealing both in the austenite region as well as in the intercriticalrange was performed and various overageing temperatures were tested.
Figure 29. Simulations of a continuous annealing line with water quenching, soft cooling inthe gas jet section and overageing. Annealing was performed in the austenite regionand the overageing temperatures were set to 200C.
Figure 30. Simulations of a continuous annealing line with gas quenching or hot dipgalvanizing line. Annealing was performed in the austenite region and variousoverageing temperatures were tested.
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Table 18. Properties and microstructures obtained in task 3.2
Rp 0,2 Rm Ag A25 Ri/t HE Phase constituents
Cycle 1 [MPa] [MPa] [%] [%] [%] [%]
V1-C1 1169 1172 0,2 2,4 6,7 52 100%M
S7-C1 938 973 3,6 8,7 5,3 67 80%M 20%F
R4-C1 1216 1245 2,2 5 6,7 88 100%MCycle 2
V1-C2 859 898 1,4 3,5 4,7 73 100%M
S7-C2 908 848 5,2 7,3 4 51 94%M 6%F
R4-C2 1230 1373 3,2 7,4 5 87 100%M
Cycle 3
V1-C3 1141 1145 0,9 4,3 2,7 72 100%M
S7-C3 1181 1383 1,4 1,5 3 33 95%M 5%F
R4-C3 1211 1450 2,3 5,8 4,2 69 100%M
Cycle 4
V1-C4 910 947 1,6 4,1 5 42 70%M 30%F
S7-C4 699 833 5,8 11,1 1,1 33 60%M 40%FR4-C4 882 5,2 6,5 6,7 83 82%M 18%F
Cycle 5
V1-C5 900 950 0,2 0,2 4 33 67%M 33%F
S7-C5 872 1137 6,1 10,7 1,3 26 61%M 39%F
R4-C5 1285 1376 2,5 2,9 5 73 86%M 14%F
Cycle 6
V1-C6 907 1145 4,8 9,4 3 29 83%M 17%F
S7-C6 551 962 9,5 10,8 1,7 23 66%M 34%F
R4-C6 1241 1488 2,6 3,2 4,2 44 94%M 6%F
Cycle 7
V1-C7 676 948 3,3 5,6 0,8 38 33%M 60%F 7%CFBS7-C7 446 779 14,6 19,2 0,7 24 19%M 79%F 2%CFB
R4-C7 705 1122 7,1 8,5 1,7 22 47%M 37%F 16%CFB
Cycle 8
V1-C8 602 985 9,2 12,1 1,3 24 36%M 62%F 2%UB
S7-C8 499 603 11,3 13 0,5 84 86%F 14%P
R4-C8 608 1046 10 16,1 1,7 16 41%M 54%F 5%UB
Cycle 9
V1-C9 597 788 11,8 18,5 0,8 71 10%M 40%F 20%UB 30%CFB
S7-C9 484 664 5,5 6,2 0,5 48 1%M 84%F 15%P
R4-C9 658 841 8,3 13,4 1,25 59 10%M 90%UB
Cycle 10V1-C10 605 856 8,3 12,8 0,8 46 7%M 30%F 33%LB 30%CFB
S7-C10 391 665 5,1 5,1 0,7 28 10%M 80%F 10%P
R4-C10 690 920 5,7 10,5 1 49 5%M 5%UB 80%LB 10%CFBM=Martensite, F=Ferrite, UB=Upper bainite, LB=Lower bainite, CFB= Carbide free bainite, P=Pearlite
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2.5.2.3 Task 3.3 Processing of special microstructures based on suggestions from WP 4
In task 3.3, the main task was to perform laboratory annealing trials in order to find optimal processingparameters for the industrial trials.Each industrial partner has suggested five time temperature cycles corresponding to their productionfacilities; SSAB: continuous annealing line with water quench; voestalpine: continuous annealing linewith gas jet cooling; Ruukki: hot dip galvanizing line. The selected time-temperature schedules are
based on the previous results from phase transformation data from WP2, task 3.1, task 3.2 and on otherknowledge about the specific production route.In order to get enough material for a systematic investigation, each time temperature cycle wasperformed for four samples with a size of 450x250mm. After annealing, the samples were pickled andthe edges were cut off since these were not annealed. Then three samples were sent to RWTH Aachenfor mechanical testing and one sample was sent to KIMAB for microstructural investigation.The results of these investigations (see WP4) were the basis for selecting the proper time temperaturecycles for the industrial trials.
2.5.2.3.1 Annealing trials for a continuous annealing line with gas cooling, CAL-GQ
2.5.2.3.1.1 Material:The chemical composition of the material selected for these annealing trials is given in Table 19. TheCP800 grade without molybdenum was chosen as the results of this annealing simulation should be thebasis for industrial trials. Molybdenum, despite being a very interesting alloying element, has to beavoided in the industrial production as much as possible due to economic reasons because of its veryhigh price.
2.5.2.3.1.2 Time temperature cycles:The annealing parameters (Tan: annealing temperature, TQ: quenching temperature; Toa: overageingtemperature) for the five time temperature cycles can be found in Table 20. Additionally, cycle 1 to 3and cycle 4 & 5 are depicted in Figure 31 and Figure 32, respectively. For all time temperature cyclesan annealing temperature of 840C was applied. The quenching temperature was varied between 680Cand 750C and the overageing temperature was varied in the range of 350C to 400C.
Table 19. Chemical composition of the lab annealed samples (wt%)Material Grade C Mn Cr Mo Si Nb
V2-CR VA_CP800 0.14 2.1 0.3 0 0.2 0.02
Table 20. Annealing parameters for the five time temperature cycles
Cycle
Tan
[C]
TQ
[C]
Toa
[C]1 840 680 350
2 840 680 375
3 840 680 400
4 840 750 350
5 840 750 400
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Figure 31. Time-temperature cycle 1, 2 and 3 for the continuous annealing line CALVAS
Figure 32. Time-temperature cycle 4 and 5 for the continuous annealing line CALVAS2.5.2.3.1.3 Results:The results obtained by RWTH Aachen reveal, that for all five annealing cycles excellent holeexpansion ratios (>70%!) and bending results could be obtained. The results will be shown in detail inWP4. Also the total elongation is very high (clearly above 10%) for all cycles.The results of cycle 1 to 3 show that with increasing overageing temperature the yield strength increasesand the tensile stress decreases. This is a very typical behaviour as with an increasing overageingtemperature a bainite is formed during overageing that is softer but has a higher yield ratio.A comparison of the results of cycle 1 and 3 with cycle 4 and 5 indicates that by decreasing thequenching temperature also slightly higher yield strength can be obtained. Due to a higher quenchingtemperature less ferrite is formed during soft cooling, whereas more bainite can be formed during
overageing, leading to higher yield strength.
Cycle 1, 2, 3
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900
0 200 400 600 800 1000
t [s]
T[C]
Cycle 4, 5
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T[C]
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2.5.2.3.1.4 Industrial trials:Based on these results, two industrial trials were planned. As for all cycles in the laboratory experimentexcellent hole expansion ratios and bending results have been obtained, the main goal for the industrialtrials is to produce two CP800 grades with different yield strength levels. That can be achieved byapplying different quenching temperatures (680C and 750C). In order to get a tensile strength clearlyabove 800MPa, an overageing temperature of 350C has to be chosen. Therefore, cycle 1 and