EFFECT OF QUASI-STATIC PRESTRAIN ON SUBSEQUENT DYNAMIC TENSILE CURVES

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1 EFFECT OF QUASI-STATIC PRESTRAIN ON SUBSEQUENT DYNAMIC TENSILE CURVES L. Durrenberger 1, A. Rusinek 1, A. Molinari 1, D. Cornette 2 1 Laboratory of Physics and Mechanics of Materials, UMR CNRS 75-54, University of Metz, Ile du Saulcy, Metz cedex, France 2 Arcelor Research, Voie Romaine, BP 30320, F Maizières les Metz, France durrenberger@lpmm.univ-metz.fr ABSTRACT The effects of a quasi-static prestrain on subsuent dynamic tensile curves have been investigated by using interrupted tests where the specimen was completely unloaded before reloading at high strain rate. The reloading flow curve was compared with results obtained at constant loading rate (same strain rate as during the previous reloading). The behaviour of three cold rolled steels produced by ARCELOR was analyzed in this study. Introduction Ecological and safety preoccupations ruire the development of Ultra High Strength Steels in the automotive industry. Thanks to the very high mechanical characteristics of these steels, the safety performance is improved without increasing car weight. It has been shown in [1] that a prestrain process improves the crash behaviour of a crash-box structure. For an efficient numerical simulation of industrial applications, a good knowledge of material behaviour is needed. During the quasi-static manufacturing of crash structure, a microstructure evolution is observed which leads to a modification of the material response at high strain rate. In this study, the effects of strain rate history have been investigated using an interrupted test in which the specimen is loaded at quasi-static strain rate, unloaded, and then reloaded at high strain rate. A convincing demonstration of history effects by means of a rapid change in strain rate fruently ruires a change of more than three orders of magnitude in the value of strain rate [2]. Bake hardening steels, dual phase steels and TRansformation Induced by Plasticity (TRIP) steels are fruently used in automotive industries. A grade of each of these families (BH260, DP and TRIP800) is analyzed in this paper. Experimental procedure The thicknesses of the sheet steels are respectively 1.643mm, 1.585mm and 1.335mm for BH260, DP and TRIP800 steels. The material characterization was performed applying two experimental methods: the quasi-static behavior ( ) was determined using a screw controlled machine whereas the high strain rates (~1000 s -1 ) was obtained by a tension split Hopkinson bar. The chemical compositions of the investigated steels are listed in Table 1. Material C Mn Si Al P BH DP TRIP Table 1. Chemical compositions of the investigated steels

2 Micrograph observations of non-prestrained materials have been performed with an optical microscope. The analysis of the BH260 micrograph, after a Dino etching, reveals an average grain size of about 15 µm, Figure 1.a. For the case of DP and TRIP800 steels, Lepera and Klemm etchings have been respectively performed. In Figure 1.b, the different phases of the DP can be observed. The ferrite appears in clear brown, the martensite in white and carbureted phases (cementite+perlite) in dark brown. The volume fractions of the above phases are respectively 85%, 5% and 10%. In Figure 1.c, the ferrite phase of the TRIP steel appears in brown, whereas the residual austenite and the bainite phases appear in white. The initial volume fraction of austenite is 19%. a) b) c) Figure 1. Micrograph of the investigated steels a) BH260 b) DP c) TRIP800 The prestrain process has been performed using sheet steels with dimensions of 350*460mm. The tensile specimens were then machined using the previous sheet steel only in the homogenous strain zone [1]. All the tests were performed perpendicularly to the rolling direction. At the end of the quasi-static process, the strain tensor is the following, ε = ε (1) Schmitt et al [3] have proposed a parameter α =(D 1:D 2)/( D 1. D 2 ) to characterize a two-stage strain path. In this expression, D 1 and D 2 represent the strain rate tensor during the prestrain and the subsuent deformation, respectively, and D is the norm of D defined by D =(D ij D ij ) 0.5. The so-called quasi-monotonic, quasi-bauschinger, and orthogonal strain-path changes are defined by α = 1, -1 and 0 respectively. In our case the subsuent loading corresponds to a quasi-monotonic loading because α = Results Klepaczko et al [2,4] have already shown that the effects of strain rate history on the flow stress depend of the microstructure of the material. For the case of BCC metals, a dynamic tensile loading after an initial quasi-static deformation leads to an increase of the flow stress compared to the dynamic loading curve without quasi-static prestraining. Conversely, for FCC metals, a dynamic tensile loading after an initial quasi-static deformation leads to a decrease of the flow stress compared to the dynamic loading curve. Von-Mises uivalent plastic strains p p ε = 2 / 3 ε ij ε are used to compare the behaviour before and after the prestrain process. ij The bake hardening steels, which are mild steels with solid solutions, show a typical behavior of BCC metal, Figure 3.a. Uenishi [5] had performed TEM observations on specimens of solution-hardened steels deformed at high (1000 s -1 ) and low (0.001 s -1 ) strain rates. It has been shown that an increase of strain rates leads to an increase of the dislocation density, especially at low strains, and delays the onset of dislocation organization (Figure 2). During a jump of strain rate, the dislocation density increases inside an organized structure (Figure 2.c), leading to a high macroscopic work-hardening rate θ = σ ε just after the jump.

3 a) b) c) Figure 2. Microstructure observed in a hot-rolled IF steel [5] a) after 17% at low strain rate b) after 15% at high strain rate c) after 4% in high strain rate following a quasi-static prestrain of 17% In Dual Phase steels, the dislocation distribution is initially non uniform. Dislocation cells are arranged around the martensitic islets because the ferrite phase has to accommodate the volume variation of the martensite during the cooling phase of the manufacturing of the steel [6]. This accommodation is performed by plastic deformation of the ferrite phase, which explains that the dislocations cells are created at the interface. The plastic deformation process depends on the dimensions of the martensitic islands [7,8], which are very hard metallurgical elements (mean ~ 60 HRC) [9]. All phases are first deformed elastically. Then, the martensitic phase continue to be elastically deformed whereas the ferrite phase is deformed plastically; the dislocation cells are propagating from the interface to the grain interior of the ferrite. If martensitic islands are smalls, they undergo no plastic deformation and can be described as rigid particules dispersed in a ferritic matrix, whereas if their dimensions are bigger a plastic deformation can occur after excessive elongation of the ferrite phase [10]. It appears that the presence of the martensite plates explains the high stress level of these steels but the hardening is mainly controlled by the evolution of the ferrite phase. As a result, the effect of strain rate history on DP steels, Figure 3.b, is typical of the behaviour of BCC metal. BH s s DP 819 s s s s a) b) Figure 3. Strain rate history effect a) BH260 b) DP For the case of the TRIP800 steel considered in this study, dynamic tensile loading following initial quasi-static deformation leads to a decrease of the flow stress compared to the pure dynamic loading curve, Figure 4.a.

4 1 TRIP dε/dt = s s s a) b) Figure 4. TRIP800 steel a) Strain rate history effect b) Evolution of the residual austenite volume fraction as a function of the plastic strain (dots are experimental data) TRIP steel s microstructure is composed of soft ferrite grains with bainite (or martensite) and retained austenite. The retained austenite (mean ~ 30 HRC [9]) transforms into martensite during deformation. The evolution of the ferrite microstructure and the growth of the martensite volume fraction f M are mainly responsible of the increasing stress during deformation. The Figure 4.b provides the evolution of the volume fraction of austenite f A in terms of the plastic strain for a tensile quasi-static loading. No experimental data are available at high strain rate because all specimens were fractured during the test with tension split Hopkinson bars. The evolution of the volume fraction of martensite is given by: f M = f A 0 + f M 0 - f A where f A 0 and f M 0 are respectively the initial volume fractions of austenite and martensite. It has been shown in early studies that the volume fraction of martensite depends on strain, strain rate and temperature [12] and also on the strain path [13]. In fact, for a given value of the plastic strain, the volume fraction increases with strain rate and stress triaxiality. Tomita et al [12] have proposed a formulation to model the strain-induced martensitic transformation kinetics. In the proposed expression, the volume fraction of martensite f M is considered to be dependent upon plastic strain, strain rate and temperature. Using this formalism, a schematic evolution of the volume fraction of martensite in terms of the plastic strain is shown in Figure 5. Quasi-static loading Dynamic loading Jump of strain rate Plastic strain Figure 5. Schematic illustration of the evolution of the martensite volume fraction with the plastic strain

5 If the evolution of the volume fractions of austenite and martensite are considered as frozen, the overall plastic response of the TRIP steel is solely governed by the deformation of the (BCC) ferrite phase. Indeed, the martensite phase is very hard and does not contribute to the total deformation. Therefore, under the hypothesis of fixed volume fractions, the macroscopic response of the TRIP steel to a strain rate jump would be typical of a material with a BCC structure as in Figure 3. To have a full understanding of flow stress evolution, the additional effect due to the growth of the martensitic phase must be analysed. The evolution of f M for different strain rate histories, as illustrated in Figure 5, shows that the growth of the martensitic phase leads to a flow stress evolution similar to that of a FCC structure. The global response of a TRIP steel being the result of the interplay of the deformation of the BCC ferrite and of the growth of the martensitic phase, strain rate history effects can be viewed as a combination of BCC and FCC responses. It is worth noting that the TRIP800 steel considered in this study shows a typical behaviour of FCC metals, while experiments made by Bleck et col [11] on a TRIP700 with 11% of initial volume fraction of austenite show that the strain rate history effects are identical to those of BCC metals. These different responses can be attributed to the initial volume fraction of austenite which is larger for TRIP800 than for TRIP700. Thus the potential of martensitic transformation is larger for TRIP800, which favors FCC like effects. Conclusion It has been shown in earlier studies that the effect of strain rate history on the flow curve depends on the material microstructure. The evolution of the ferrite microstructure, which depends on strain, temperature and strain rate, governs the evolution of the flow stress in BCC steels. It has been shown in this study that DP steels exhibit the same behaviour as BCC metals. Since TRIP steels are composed of different phases, strain rate history effects are the result of the competition between the microstructure evolution of the ferrite phase and the austenito-martensitic transformation. References 1. Durrenberger, L., Even, D., Molinari, A., Rusinek, A., Influence of the strain path on crash properties of a crash-box structure by experimental and numerical approaches. J. Phys. IV (2006). 2. Klepaczko, J.R., Duffy, J., History effects in polycrystalline BCC metals and steel subjected to rapid changes in strain rate and temperature. Arch. Mech. 34, Warszawa (1982). 3. Schmitt; J-H., Shen, E.L., Raphanel, J.L., A parameter for measuring the magnitude of a change of strain path : Validation and comparison with experiments on low steel. Int. Jour. of Plasticity 10(5) (1994). 4. Klepaczko, J.R., Frantz, R.A., Duffy, J., History effects in polycristalline FCC metals subjected to rapid changes in strain rate and temperature. Rozprawy inzynierskie, Engineering transactions 25, ,(1977). 5. Uenishi, A., Thermodynamical behaviour of solution-hardened interstitial free steels at high strain rates. Ph. D. Thesis, Université Paris 13, Korzekwa, D.A., Matlock, D.K., Krauss, G., Dislocation substructure as a function of strain in DP steel. Met. Trans. 15 A, (1984). 7. Byun, T.S., Kim, I.S., Tensile properties and inhomogeneous deformation of Ferrite-Martensite dual-phase steels. Mater. Sci (1993). 8. Bag, A., Ray, K.K., Dwarakadasa, E.S., Influence of martensite content and morphology on tensile and impact properties of high-martensite dual-phase steels. Metall. Mater. Trans. 30 (A) (1999). 9. Hildenwall, B., Prediction of the residual stresses created during quenching, Ph. D. Thesis, Linkoping University (1979). 10. Rashid, M.S., Cprek, E.R., Relationship between microstructure and formability in two high-strength, low-alloy steels. In: Niemeier, B.A., Schmeider, A.K., Newby, J.R. (Eds.), Formability Topics-Metallic Materials, ASTM STP 647. American Society for Testing and Materials, (1978). 11. Bleck, W. et col., Crash relevant properties and dynamic denting of pre-strained or pre-loaded high strength steels for automotive parts, Research Programme of the Research Fund for Coal and Steel Final report, (2005). 12. Tomita, Y., Iwamoto, T., Constitutive modeling of TRIP steel and its application to the improvement of mechanical properties. Int. J. Mech. Sci. 37, No (1995). 13. Andersson, R., Deformation characteristics of stainless steels. Doctoral thesis - division of manufacturing systems engineering, (2005).