SMALL CREEP DEFORMATIONS AND RESIDUAL STRESS RELAXATION IN WELDED AND SHOT PEENED AL MG SI CU-ALUMINUM ALLOYS

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1 477 SMALL CREEP DEFORMATIONS AND RESIDUAL STRESS RELAXATION IN WELDED AND SHOT PEENED AL MG SI CU-ALUMINUM ALLOYS T. Hirsch, S. Daichendt IWT - Stiftung Institut für Werkstofftechnik, Bremen, Germany ABSTRACT The stability of compressive residual stresses and the stability of strain hardening states are seen to have high technical importance. For a laser beam welded AlMgSiCu-alloy, results of applied stress and residual stress relaxation at elevated temperatures 100<T<150 C are presented. These experiments prove the existence of back stresses. The level of these stresses can be estimated to be in the range of Orowan stresses for metastable precipitates of the age hardened Aluminum alloy. Pure thermal relaxation reduces residual stresses for temperatures >100 C, and only small differences between welding zones and base material occur. Shot peening induced residual stresses decrease to a large extent in the weld seam areas yet remain stable to a higher extent in the base material after stress relaxation tests (and also after creep tests, low cycle alternating thermal or thermal mechanical tests). Calculations of residual stress relaxation based on applied stress relaxation tests underestimates the relaxation behavior due to the different nature of applied stresses and residual stresses and the higher relaxation of surface residual stresses. INTRODUCTION Precipitation-hardened wrought Aluminum alloys are construction materials for the automotive and aircraft industries. The copper alloyed AA6056 is used for its high corrosion resistance, good weldability, high strength as well as high damage tolerance [1]. Laser beam welding of this alloy causes characteristic microstructures and inferior mechanical properties in the weld seam [2, 3]. Surface treatments such as shot peening are appropriate methods to increase the fatigue strength of these welded materials [4]. It is well known, that residual stresses at the surface and/or in surface layers determine the fatigue and corrosion resistance.. However there is a lack of knowledge about the long-term stability of induced residual stresses and strain hardening states regarding complex superimpositions of cyclic thermal and mechanical service conditions. Service temperatures can be as high as around C, the onset of creep deformation in AlMgSi-Aluminum Alloys. The relaxation behavior of applied stresses and residual stresses was examined for the base material and welded and shot peened conditions at temperatures from 100 C to 150 C.

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3 478 INVESTIGATED MATERIAL The chemical composition is given in Table 1. The received T4 aged sheet material was 3.1 mm thick. Without solution annealing these sheets were age-hardened for 4 hours at 190 C to reach condition T6. The different types of experiments were executed with standard flat tensile test specimens of 150 mm in length (gauge length 50 mm and width 10 mm). Table 1: Chemical composition of the alloy AA 6056 in wt-% (bulk analysis by optical emission spectrometry, OES) AA 6056 Si Mg Cu Mn Zn Fe Al OES [wt-%] Bulk EXPERIMENTAL PROCEDURE Laser beam welding was executed transverse to the rolling direction with a CO 2 -laser TLF turbo and a filler wire of Al-12%Si. From the base material and welded sheet tensile test specimens were manufactured according to DIN The weld seam was milled to remove an approximately 100µm thick surface layer on both sides of the samples. Applied stress relaxation tests were carried out in a temperature range 100 < T < 150 C. Initial applied stresses were selected to be between 50 and 70% of the Yield strength at testing temperature (base material). Tests were executed with an electronically controlled material testing machine Z020 (manufacturer Zwick-Roell company, Germany) equipped with a suitable furnace. Strains were measured with a clip-on extensometer of initial gauge length of 20 mm, and stress changes of the straincontrolled tests recorded continuously by the instruments software. Residual stress evaluation used the conventional sin 2 χ method with a conventional Bragg-Brentano Geometry (Bruker-AXS D8) equipped with Vanadium filtered Cr-Kα radiation. A primary beam aperture of 1 mm diameter was selected for the local analysis over the weld seam. The diffracted beam slit was 2 mm. The {311}-lattice plane of α-aluminum at 2θ0= was chosen for the analysis, the measured range 137 < 2θ <142 in steps of 0.05 with measurement time 20 sec. The sin²χmeasurements utilized 17 χ-angles equidistant in sin²χ (χmax= ±60 ). X-ray elastic constants of ½ s 2 = *10-6 MPa -1 and s 1 = *10-6 MPa -1 were used. RESULTS AND DISCUSSION Welding produced a typical residual stress profile with a local minimum of about 0 MPa in the weld center and maximum tensile residual stresses in the heat affected zones. Hardness was measured on the cross section. In the fusion zone (FZ) a minimum hardness of 80 HV is detected followed by values of about 100 HV in the HAZ. The base material shows values of about 138 HV. The shot peening was executed with shot S230 in an air blasting machine. Shot pressure was 1.74 bar and coverage 200 %. This treatment resulted in an Almen intensity of 0.22 mma. After

4 479 peening compressive residual stresses of 70 to 100 MPa were measured at the surface. A typical depth profile, obtained by electrochemical layer removal, revealed maximum residual stresses of 200 MPa in 100 µm surface distance and zero values at 300 µm as shown in Fig. 1 (right). Pronounced strain hardening occurred and was measured by the integral breadth (IB) of diffraction lines. The fusion zone always gave lowest values of residual stresses and strain hardening followed by values of BM and HAZ. Residual stresses [MPa] BM HAZ FZ HAZ BM 2,0 IB RS 1,6 1,2 0,8 0, , IB [ ] Residual stresses [MPa] RS 2,0 IB 1,6 FZ HAZ BM 1,2 0,8 0, , Distance from weld center [mm] Distance from surface [µm] IB [ ] Figure1: Surface values (left) and depth profiles (right) of residual stresses (RS) and integral breadth (IB) of X-Ray diffraction lines of shot peened surfaces. Nominal stress [MPa] T=150 C=const ,5 1,0 1,5 2,0 2,5 Time [*10 5 s] Applied Stress: 219 MPa Applied Stress: 157 MPa Figure 2. Relaxation behavior of base material at 150 C with indicated initial applied stresses.

5 480 In figure 2 the base material s relaxation behaviour is plotted against time. An initially strong relaxation is observed and the amount of relaxed stresses increases with initial applied stress values. With increasing time almost constant stress levels occur. Figure 3 gives a similar plot for welded and shot-peened specimens. The initial decrease of stresses is higher, but the stresses level off at a slower rate compared to the base material. 240 T=150 C=const. Nominal stress [MPa] Applied stresses: 220 MPa 200 MPa 181 MPa Figure 3. Relaxation behavior of welded and peened samples at 150 C with indicated initial applied stresses ,5 1,0 1,5 2,0 2,5 Time [*10 5 s] The results indicate the existence of high internal back stresses, as results of creep tests with the same material states do [5]. Plastic deformation of crystalline solids at high temperatures does not take place under the action of the whole applied stress σ, but only under a part of it [6]. This behavior is represented by the difference of applied stress σ and back σ B (s eff = σ σ B ) [6]. Back stresses are caused by other dislocations by a dislocation substructure and by obstacles like strengthening particles. If back stresses do not depend on applied stresses they are defined as threshold stresses. Introducing threshold stresses for creep tests results in a reduction of high stress sensitivity parameters to reasonable values. Back stresses are reported for precipitationhardened alloys [7, 8], for Aluminum metal matrix composites [9], for Aluminum alloys manufactured by powder metallurgy [10] and for dispersion-hardened alloys [11, 12]. With applied stresses equal to back stresses, the effective stresses for the transformation of elastic into plastic strains during relaxation is equal to zero as can be seen from Figures 2 and 3. Table 2 summarizes back stresses evaluated from relaxation tests. The observed values decrease with temperature and generally give lower values in the welded condition. Contrary to the result after creep tests [5], some dependence from applied initial stresses can be observed. If only precipitations or dispersoids are responsible for the back stresses these values should be independent of applied stresses. If compared with creep tests, during stress relaxation tests recovery of dislocation struc-

6 481 tures is somewhat restricted due to the short testing times. As a consequence the dislocation structure also contributes to σ B [9]. On the other hand, the dislocation structure depends on applied stresses, which explains the stress-dependence of σ B. Table 2: Back stresses for different test temperatures Base material Welded, shot-peened Initial applied stress σ 0 [MPa] Back stress σ B B [MPa] Initial applied stress σ 0 [MPa] Back stress σ B B [MPa] 100 C C C Three deformation models are proposed for the origin and magnitude of back stresses (for example see [10]): a) stresses required to cause a dislocation bowing between incoherent particles (the Orowan stress), b) extra back stresses required to create an additional dislocation line length as dislocation segments climb over a particle and c) stresses required to detach the dislocation from the incoherent particle after climb is completed. Estimates of Orowan stresses can be made if partially coherent or incoherent precipitates must be bypassed [13]. Careful Metallographic inspection revealed that the investigated alloy contains dispersoid particles with a size of 0.1 to 0.2 µm and distances between 2 and 10 µm (see also [2]). This would result in back stresses of some MPa only and obviously is only a small contribution to the total value of the back stresses. In the case of partially coherent β precipitates of the age-hardened condition, Transmission Electron Microscopy revealed lengths of 10 to 20 nm and distances from 25 to 50 nm [2]. For simplification in the calculation these precipitates were thought to be spherical. Threshold stresses then vary from 25 to 180 MPa. This last value agrees quite well with experimental data in table 2. Thus the assumption of a deformation mechanism by dislocation climb for bypassing the partially coherent ß -precipitates with increasing temperature is supported. Figure 4 summarizes results of residual stress measurements after relaxation tests at 150 C. The test parameters are given in figure 4. Open squares characterize the initially existing values of the shot peening process. As already mentioned the fusion (FZ) and heat affected zones (HAZ) give lower residual stress values due to differences in microstructure and strength. Almost independent of the initial applied stress a pronounced relaxation of surface residual stresses occurred after 40 to 60 hours testing time. This relaxation is highest in the FZ and HAZ (30% of initial values remain) whereas the base material shows a relaxation of residual stresses from 100 MPa to 60 to 70 MPa (70% of initial values remain). Pure thermal relaxation of shot peening induces residual stresses after 150 C isothermal annealing for some 15 hours resulted in 24% remaining residual stress values

7 482 Residual stresses [MPa] BM HAZ FZ HAZ BM 0 T= 150 C= const Distance from the weld center [mm] 220MPa,2358 min 200MPa,3800 min 180MPa,2424 min as peened Figure 4: Residual stresses across the weld seam after applied stress relaxation tests with indicated parameters (initial applied stress, testing time at end of test). (t)/σ 0 Ratio stress/initial stress σ 1,2 1,0 σ 0 = 160 MPa 0,8 0,6 0,4 Creep tests Relaxation tests 0,2 Temp.-changes Temp.-stress-changes Log time [min] Figure 5: Calculated relaxation of applied stress normalized by the initial applied σ 0 = 250 MPa stress and normalized residual stresses after interruption of tests at the indicated time (T= 150 C). Finally the results from applied stress relaxation tests are compared with the relaxation of residual stresses in these tests and after alternating thermal and mechanical loads to detect similarities or differences. Details of the latter experiments can be found elsewhere [5]. Figure 5 displays calculated relaxation of applied stresses normalized by the initial applied stresses and normalized residual stresses after interruption of tests at the given time. Residual stress values from the fu-

8 483 sion zone are plotted in figure 5 as they show the highest amount of relaxation. The initial applied stress of 250 MPa used in the calculation of a relaxation test was chosen as residual stress depth profiles after shot peening show maximum compressive residual stresses of about this value. The lower initial applied stress of 160 MPa is in the range of back stresses and consequently almost no load induced stress relaxation should occur. A higher amount of residual stresses is expected to relax at the surface due to the high dislocation density induced by the shot peening process. As proved by figure 5 residual stresses after the different testing procedures indeed relax to a far higher amount as expected even from the maximum residual stress value after shot peening. The difference between applied stresses and residual stresses is obvious. CONCLUSION Relaxation tests for temperatures 100<T<150 C prove the existence of high internal back stresses. The level of these back stresses can be estimated to be in the range of Orowan stresses of metastable Al-based precipitates. With the knowledge of back stresses experimental results could be analytically modeled. Shot peening induced residual stresses after applied stress relaxation tests (and also after creep tests, low cycle alternating thermal or thermal mechanical tests) decrease to a large extent in the weld seam areas and remain stable to a higher extent in the base material. Calculations of residual stress relaxation based on applied stress relaxation tests underestimate the relaxation behavior due to the higher relaxation of residual stresses at the surface and the different nature of applied stresses and residual stresses. The effect of back stresses seems to be more pronounced during applied stress relaxation compared to residual stress relaxation. ACKNOWLEDGEMENTS The authors would like to thank AIF (German Department of Economy) for financial support. REFERENCES [1] M. Tanaka, T. Warner, La Revue de Métallurgie-CIT, May (2003), p [2] B. C. Meyer, H. Doyen, D. Emanowski, G. Tempus, T. Hirsch, P. Mayr, Metallurgical and Materials Transactions A, Vol. 31A (2000), p [3] B. C Meyer, Ph. D. Thesis, University of Bremen (1999) [4] T. Hirsch, P. Starker, E. Macherauch, Haerterei Technische Mitteilung (1982), p. 100 [5] S. Daichendt, T. Hirsch, Final report 13680N, Arbeitsgemeinschaft industrieller Forschungsvereinigungen, Koeln, Germany (2007) [6] J. Cadek, Creep in Metallic Materials, Materials Science Monographs, 48 (1988) [7] P. Zhang, Scripta Materialia 52 (2005), p. 277 [8] J. Wang, X. Wu, K. Xia, Materials Science and Engineering A (1997), p. 287 [9] G. Kausträter, B. Skrotzki, G. Eggeler, Materials Science and Engineering A (2001), p. 716

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