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Microstructure controlled bending response in AA6016 Al alloys
Materials Science and Engineering A 528 (2011) 7068–7076
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Microstructure controlled bending response in AA6016 Al alloys Aleksandar Davidkov a,∗ , Roumen H. Petrov a,b , Peter De Smet c , Bruno Schepers c , Leo A.I. Kestens a a
Department of Material Science and Engineering, Gent University, Technologiepark 903, 9052 Gent, Belgium Department of Material Science and Engineering, Delft University of Technology, Mekelweg 2, 2628 Delft, The Netherlands c Aleris Aluminum Duffel B.V.B.A., A. Stocletlaan 87, 2570 Duffel, Belgium b
a r t i c l e
i n f o
Article history: Received 8 May 2011 Accepted 25 May 2011 Available online 31 May 2011 Keywords: Aluminium 6xxx alloys Microstructure Bendability Hemming Ductile intergranular fracture
a b s t r a c t A contemporary approach in the car weight reduction is the use of low weight and high strength Al alloys sheets for hang-on body panels production. The final step in the forming route of such panels is the attachment of the outer skin to the inner part of the panel by applying a hemming operation. This joining method is cheap, easy to perform and environment-friendly, but requires severe 180◦ bending of the edges of the outer skin which quite often results in cracking or complete tearing of the bend surface. Such kind of failure restricts the further application of the hemmed products. The microstructures after solution heat treatment and pre-aging (T4P temper state) of two grades age-hardening AA6016-type aluminium alloy sheets were studied in this work by means of optical microscopy, scanning electron microscopy and electron backscatter diffraction. The obtained results were related to the hemming response of the grades. It was found that the alloy composition is one of the main parameters controlling the bendability of these grades through the amount of the formed strengthening phases. However, the applied thermal treatment remains the key factor responsible for the favorable distribution of these phases into the microstructure. The grain size and the volume fraction of the constituent particles were found to play secondary role in forming the material bending properties and can be only used for their fine tuning. The presence of Mg2 Si (-phase) and/or Al1.9 CuMg4.1 Si3.3 (Q-phase) particles in the grain boundaries structure was recognized as a critical microstructural feature causing severe reduction in the bending ability of the sheets by promoting an intergranular fracture. The possibility of grain boundaries failure exponentially raises with the time passed due to the natural aging. © 2011 Elsevier B.V. All rights reserved.
1. Introduction The weight of automobiles continuously increases during the years with improving performance, comfort and safety features. At the same time the importance of environmental and economical issues grows and leads to raising demand of weight reduction of the cars. The favorable combination of low specific density (approximately 1/3 of the steel), good corrosion resistance, strength and formability of aluminium alloys make them attractive material for this application. As a sheet material aluminium alloys are being used mainly for production of hang-on car body panels like hoods, trunk lids, doors, etc. Currently, 6xxx heat treatable Al–Mg–Si–(Cu) alloys are used worldwide for the outer skins of the panels as they satisfy well the main requirements for high strength, dent resis-
∗ Corresponding author. Tel.: +32 9 331 04 62; fax: +32 9 264 58 33. E-mail addresses: [email protected] (A. Davidkov), [email protected] (R.H. Petrov), [email protected] (P. De Smet), [email protected] (B. Schepers), [email protected] (L.A.I. Kestens). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.05.055
tance and good formability. The major alloying element in these alloys is Si, which together with Mg and Cu forms strengthening Mg–Si or Al–Mg–Si–Cu phases. Excessive amounts of Si are usually added to improve the precipitation hardening during the short paint bake process, respectively to sufficiently raise the dent resistance of the parts [1]. Such approach enables downgauging and further weight reduction while maintaining sufficient strength for the outer panels, but however, it can have detrimental effect on the sheet’s hemming ability. Consequently, the properties of the material for such applications should be in some way compromised in order to balance between the contradictive requirements for high strength and good hemmability. On other hand, the industrial production process should be conformed to the commercial requirements for minimizing the time and cost factors during production. The hemming is an environmental friendly assembly method used in the automotive industry to join the outer skins to the inner parts of the hang-on car body panels. The requirements to the alloys subjected to hemming operation are very severe as the edge of the skin is folded over the inner part of the panel by bending to
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180◦ over a radius, equal to the half of the sheet thickness. Finally, the hem is closed by applying a pressing operation. The pre-strain already accumulated in the preceding forming operations additionally makes the sheet material more susceptible to fracture during the hemming. The aluminium alloys which were worldwide recognized as being able to satisfy these severe requirements for the car body skins are the heat treatable 6xxx series Al alloys. They are usually delivered to the car manufacturer (in some cases with a delay) in the form of coils in pre-aged T4P temper state [2]. The microstructure at this stage is quite complex and determines the properties of the material before the final forming operations. Several groups of intermetallic particles can be found embedded in the recrystallized and supersaturated matrix of these alloys [3]: (i) large-sized primary constituent particles of Al–Fe(Mn,Cr)–Si-phase (typically 1 ÷ 20 m in size), (ii) dispersoid particles of Mn–Al, Mg–Si or Al–Mg–Si–(Cu) phases (typically 0.1 ÷ 1 m in size) and (iii) Al–Mg–Si–(Cu) clusters formed during the pre-aging heat treatment which aim to preserve the formability of the material by stabilizing the microstructure and slowing down the process of natural aging and to ensure a fast bake-hardening response during the insufficiently long final artificial age-hardening process. Different number of pre-aging steps can be included into the heat treatment of these alloys, however, not carefully chosen pre-aging time or temperature can have detrimental effect on the sheet formability [2,4]. The bendability is one of the most important forming attributes of precipitation hardening 6xxx aluminium alloys. Many metallurgical, technological and processing factors in their rather complex combination can have a strong influence on it. The fracture during bending of these age-hardening aluminium alloys usually occurs through two main mechanisms. The first one, which in many cases results from an improperly applied heat treatment, is the intergranular fracture due to the presence of heterogeneously nucleated grain boundary particles [5]. Evensen et al. [6] investigated the fracture behaviour of grain boundaries as a function of the grain size and the precipitation hardening of the matrix. They found that by promoting more dispersed slip, or by grain refinement the intergranular fracture mode can be replaced by the conventional ductile rupture in which void formation and shear banding became more pronounced than grain boundary decohesion. Practically, several possible sources of intergranular fracture can be delineated: (i) reduction in the local cohesive stress due to segregation or presence of grain boundary particles, (ii) initiation, growth and coalescence of micro-voids, formed on the grain boundary particles, (iii) localization of the deformation in precipitation free zones (PFZs) adjacent to the grain boundaries, (iv) promotion of coarse slip in the bulk of the grains, and (v) coarsening of the grain size [5,7–9]. The second fracture mechanism is accepted as conventional and includes sequence of events causing localization of the strain and intense shearing. In this case it is considered that the bendability of aluminium alloy sheets is directly governed by the combined effect of micro-voids formation around large secondphase particles, strain localization and propagation of macroscopic shear bands [10–12]. Lievers et al. [13] confirmed that the shear bands and void damage development play co-operative role in promoting ductile fracture at the outer free surface during bending. They also found that the changes in porosity, which create spatial differences in constitutive behaviour needed for strain localization, are exceptionally pronounced once the strain hardening of the metal matrix has been exhausted. Hence, the bendability becomes better for materials which have higher strain hardening rate and smaller initial voids fraction. In different materials with the same amount of second-phase particles, the bendability becomes mainly dependent on the value of the strain hardening factor n, which is controlled by the properties of the aluminium matrix [14].
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Table 1 Chemical composition of the alloys, wt%. Grade
Si
Mg
Fe
Cu
Mn
Al
A B
1.00 0.90
0.47 0.60
0.23 0.10
0.17 0.17
0.07 0.15
Bal. Bal.
The analysis of the existing literature data show that there is still no comprehensive understanding for the influence of different microstructural features of age-hardening Al alloys on their bendability. In the current work the microstructures of two AA6016-type grades processed in the form of thin sheets were studied with the aim to find out the main microstructural features provoking the observed significant differences in their hemmability. Interrupted hemming tests in laboratory conditions were used to estimate the hemming properties of the grades and in situ bending tests were performed in a scanning electron microscope (SEM) in order to study the local deformation and the fracture development during bending. The chemical composition of the alloy, the volume fraction and the distribution of the strengthening phases as well as the grain boundary structure of the grades were considered as parameters potentially controlling the bendability. 2. Experimental Two heat-treatable AA6016-type grades in the form of 1 mm thick sheets were studied in this work after being heat-treated to T4P temper state by applying several consecutive steps, i.e. solutionizing, quenching and two pre-aging steps. During solutionizing, the grade denominated here as A was soaked for 5 s at solutionizing temperature and the other grade, denominated here as B was soaked for 55 s at the same temperature. All the other parameters of the heat-treating process were the same for both grades. Part of the sheets was kept at room temperature for several months and another part was kept at a temperature of −18 ◦ C to avoid any natural aging during the studies. The examined grades also differ in their chemical compositions but the variations are within the standard tolerances for AA6016-type aluminium alloys – see Table 1. Samples for characterization of the in-plane and the throughthickness microstructure were cut from the middle part of both sheets. Standard mechanical polishing procedure [15] was used to prepare samples for observation of the constituent particles distribution at magnification 20× with an optical microscope “Zeiss Jenavert”. Multiple digital images were acquired over an area of 0.4 mm2 and further processed by specialized software for quantitative microstructural analyses “ImageJ” [16]. The general grain structures as well as the macro-cracks paths through the microstructure were characterized using a polarized light after anodizing the mechanically polished surfaces of the samples by submerging in a 3% HFB4 water solution for a period of 1–1.5 min under an applied tension of 25 V. Samples for SEM, EDX and EBSD analysis were prepared using electrolytic polishing by means of 10% HClO4 -solution in methanol at 22 ◦ C under tension of 39 V for a time of 12 s. The characterization of the microstructures at different magnifications was done by an environmental scanning electron microscope ESEM FEI XL30 with LaB6 filament employing acceleration voltages of 5 ÷ 25 kV and working distances of 10 ÷ 20 mm. Energy Dispersive X-ray (EDX) analyses on the polished surfaces of the samples were carried out to identify the chemical composition of the matrix, constituent particles (1 ÷ 20 m) and dispersoid particles (0.1 ÷ 1 m) at a working distance of 10 mm, acceleration voltage of 25 kV and acquisition time of 100 s which correspond to approximately 1800 counts per second. Minimum 10 measurements for each type of particle were made and the results were presented as mean values. EBSD measurements in both RD × TD and RD × ND planes were carried out
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Fig. 1. Through-thickness grain structure of 1 mm thick sheets. (a) grade A, (b) grade B, RD × ND plane (RD vertical), (c) average diameter of grains (based on area fraction of grains).
and the data were collected and post-processed by TSL-OIM® version 4.6 software. Accelerating voltage of 25 kV, magnification of 180× and a working distance of 22 mm were used to scan over of an area of 1.0 mm × 1.5 mm with a step of 2 or 2.5 m in hexagonal scan grid. Data for the average grain sizes and grain size distribution were obtained and analysed. The through thickness sections of mechanically polished samples with a size of 7 mm × 12 mm were observed during 180◦ in situ bend tests with the bend line (BL) parallel to the rolling direction (RD) of the sheets. Three-point bending micro-device designed to work with FEI XL30 environmental scanning electron microscope was used for this purpose. During interrupted laboratory hemming tests rectangular samples with a size of 10 mm × 20 mm were fully folded without inserting of an intermediate sheet around axis (BL) oriented parallel or perpendicular to the sheet rolling direction. This bending test procedure was applied in three steps – to 90◦ , to 135◦ and with final pressing to 180◦ . The assessment of the hemmability was made by observing the micro- and macroscopic appearance of cracks on the outer convex surface of the samples. A comparative hemming factor (HF) ranging from 1 (very bad bendability) to 4 (very good bendability) was used to scale the hemming quality of the tested sheets. 3. Results The through-thickness microstructures of both T4P AA6016type grades consist of fine equiaxial recrystallized grains – see Fig. 1(a) and (b). However, grade B shows more heterogeneous grain distribution and larger average grain diameter in comparison to grade A, especially in the areas close to the sheet surface. The average grain diameter calculated as a weighted sum of the grain area fractions is presented in Fig. 1(c). Large ␣-Al(FeMnCr)Si constituent particles form during the solidification of these alloys and remain almost unchanged during the whole cycle of thermomechanical treatment [17]. According to the quantitative optical observations, a slightly higher volume fraction of constituent particles was found in grade A with its maximum at equivalent particle diameter between 1 and 2 m – see Fig. 2 and Table 2. The EDX analyses confirmed their Al–Fe–Si–Mn–Cr composition in both cases [18]. SEM investigations were made in order to reveal the distribution of dispersoid particles in the microstructure, especially of those lying in the grain boundaries, as they can play destructive role on
Table 2 Density and area fraction of constituent particles. Grade Density per mm Area fraction, %
2
A
B
2726 2.05
2644 1.49
the bending properties. The qualitative observations of multiple areas in both grades show higher amount of grain boundary particles in grade B. They are well aligned with the grain boundaries and in many cases quite coarse – see Fig. 3(b). Particles are also present in the grain boundaries of grade A but they are smaller and much less densely populated. Despite the compositional changes during the electrolytic polishing, the EDX analyses of these particles show close to stoichiometric amount of Mg and Si atoms for the stable Mg2 Si (-phase) (Mg/Si = 1.73 [19]) – see Fig. 4 and Table 3. The tensile properties of the grades were studied after one and two months of natural aging in pre-aged T4P temper state as well as in artificially aged T8 temper state after 2% tensile deformation and final paint-bake. Grade A has shown lower strength compared to grade B but at the same time both grades have equal ductility when expressed in terms of uniform elongation (UE), total elongation (TE) and strain hardening factor n – see Table 4. The hemming performance of the sheets was studied by observing the macro-appearance of the outer convex surface of 180◦ bent samples. Fig. 5 shows the surfaces of the samples subjected to bending in three months after the production. In the case, when the bend line (BL) was parallel to the rolling direction (RD) of the sheet (see the upper row in Fig. 5) both grades showed relatively good hemmability estimated with HF = 3.25 for grade A and HF = 2.25 Table 3 Chemical composition of grain boundary particles. Element
wt%
at%
Mg Al Si Cr Mn Fe Cu Total
1.36 97.02 0.85 0.16 0.19 0.15 0.27 100
1.51 97.30 0.82 0.08 0.09 0.07 0.11 100
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Fig. 2. Optical micrographs of mechanically polished samples in RD × TD plane (RD horizontal). (a) grade A, (b) grade B, (c) size distribution of constituent particles.
Fig. 3. SEM micrographs of Mg2 Si particles allocated at the grain boundaries. (a) grade A, (b) grade B.
for grade B. In the case, when the BL was perpendicular to the RD (lower row), the hemming performance of the grades was estimated as good for grade A and as bad for grade B with a comparative hemming factors respectively HF = 2.5 and HF = 1.5.
The observation of the fracture in grade B showed that a full length primary crack spreads along the hem. The SEM investigations reveal that it propagates mainly through a ductile grain boundary fracture (at room temperatures). The micrographs in
Fig. 4. Grain boundary particle in grade B. (a) SEM micrograph, the scale bar is 500 nm, (b) EDX spectrum.
Table 4 Tensile properties of the grades. Grade
A B
One month after pre-bake
Two months after pre-bake
After paint-bake
YS (MPa)
UTS (MPa)
UE (%)
TE (%)
n (at 5%)
YS (MPa)
UTS (MPa)
UE (%)
TE (%)
n (at 5%)
YS (MPa)
124 141
241 257
22.5 22.3
26.0 25.6
0.26 0.26
133 149
250 266
22.4 22.1
26.4 26.5
0.26 0.26
235 250
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Fig. 5. Hemmability of the sheets with BL RD (upper row) and BL ⊥ RD (lower row). (a) grade A, (b) grade B.
Fig. 6. SEM micrographs of fracture surface in grade B. The arrows show specific fracture zones discussed in the text.
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Fig. 6 present the details of the fracture mechanism including grain boundaries debonding and rupture as well as strain localization and micro-voids formation at the grain boundary particles.
4. Discussion Regarding the chemical composition shown in Table 1, both grades are characterised as alloys with excess of Si, with a Mg/Si ratio respectively 0.47 for grade A and 0.67 for grade B. The excessive amount of Si is added in order to improve the precipitation hardening during the final insufficiently long paint-bake process [20]. However, the excess Si increases also the initial T4 yield strength influencing the volume fraction of the strengthening phases. Many authors [20–22] agree that the bendability deteriorates with increasing yield strength. Hence, because of the higher excess of Si in grade A, one should expect worse bending response in comparison to grade B. Actually, the volume fraction of the strengthening phases in Si-excess alloys is always limited by the level of Mg, as it is completely removed by the Si from the solution [20,23]. Thus, the volume fraction of the strengthening phases in grade B will be higher due to the higher Mg content of this alloy and the bendability will be reduced. Better hemming should be expected from the lower-Mg and lower-strength grade A. This expectation is confirmed by the results obtained from the hemming tests – cf. Fig. 5. The Fe content in these alloys influences the volume fraction of the coarse non-deformable constituent particles which are considered as strongly involved into the fracture process during bending by their role in voids formation [24,25]. Part of the excess Si is also consumed by those particles, but however the sufficiently high excess of Si in both alloys provides the further formation of strengthening phases. In order to diminish the volume fraction of the constituent particles and to improve the bendability in grade B, the Fe content was dropped to 0.10 wt%. According to the microstructural study, as expected, lower volume fraction of constituent particles was found in grade B in comparison to grade A – see Fig. 2 and Table 2. However, in spite of the reduced volume fraction of constituent particles, the bendability of grade B does not improve as compared to the higher-Fe-containing grade A. This fact gives a clue that the fracture mechanisms involved in this grade during bending does not necessary include formation of voids around constituent particles and their further interaction with the propagating shear bands. The EDX analyses of those particles show higher content of Mn in grade B than in grade A [18] which is in agreement with the chemical composition of these alloys – cf. Table 1. However, the reduced amount of dissolved Mn reduces the possibility for formation of incoherent, non-deformable ␣-AlMnSi dispersoid particles (with a size around 0.1 m) during the homogenization annealing. Such dispersoids are considered as improving the bendability by promoting of grain refinement, homogenization of the slip distribution (decreasing the slip band spacing and broadening the slip band width) and in this way reducing the tendency of intergranular fracture [26]. The soak time during the solutionizing of grade B was much longer in comparison to grade A. Hence, its matrix was enriched of alloying elements to a greater extend and its precipitation potential was increased. The large grain boundary particles found in grade B – cf. Fig. 3(b) – were formed by precipitation of the same alloying elements (Mg, Si and also Cu) which form the strengthening phases in this alloy. It is logical to conclude that if these large particles are formed in the grain boundaries, the adjacent zones will be depleted from alloying elements, i.e. the strengthening ability of the adjacent to the grain boundaries areas would be lower than the bulk of the grains. At the same time the cohesive bonds between the grains are weakened due to the presence of large amount of grain
Fig. 7. Deterioration of hemmability during natural aging. The steeper line of grade B shows faster deterioration. Samples bent along RD (BL ⊥ RD).
boundary particles. The continuously increasing difference in the strength between the interior of the grains and the grain boundary areas, taking place for instance during natural aging in T4P temper state, will automatically lead to increased possibility of grain boundaries fracture followed by severe deterioration of the bending properties of grade B. Such a fracture mechanism is mentioned in [27] and gives one very possible explanation for the observed failure of the grain boundaries during the bending in grade B. The scanning electron microscope fractographs of the cracked samples of grade B are presented in Fig. 6 and further discussed in details in the next paragraph. The increment of strength in T4P temper state due to natural aging is observed during the storage period at ambient temperatures in both grades. The results from the tensile tests, given in Table 3 show the usual dependence of YS and UTS on the time elapsed after solutionizing. For example the YS rises with about 7% for grade A and with 5% for grade B in one month of natural aging. The UTS rises respectively with about 5% and 3% for the same period of time. In both cases a relatively uniform increase of YS and UTS is observed. When tracing the changes in the hemming behaviour of the grades – see Fig. 7 – in spite of the subjectivity in determining the HF, which is obtained by visual observation of the formed hem, it seems that the hemmability of the grades deteriorates not as uniform as the increase of the strength for the same period of natural aging. The better hemmability of grade A one month after solutionizing, assessed with HF = 3.0 is reduced within one month with about 8% to HF = 2.75 and the worse hemmability of grade B assessed with HF = 2.5 is reduced for the same period with 20% to HF = 2.0. This more pronounced deterioration in the hemmability of grade B could confirm the abovementioned hypothesis for the possible continuous increase of the strength difference between the grain boundaries and the bulk of the grains during the natural aging, which could determine the prevailing intergranular mechanism of fracture observed in this grade. This strength difference is supposed to be the highest for the grades aged to peak yield strength [28]. Such mechanism of failure is possible in grade A too, as its microstructure also shows particles embedded in the grain boundaries – cf. Fig. 3(a). Fig. 8 presents some initial features of grain boundary fracture observed in grade A. However propagation of cracks did not occur until the end of the 180◦ bend. It seems that the sparse population of grain boundary particles in combination with the lower strength of the matrix in grade A make this fracture mechanism less possible. The study of the fracture surfaces in grade B (cf. Fig. 6) supposed that the phenomenon of grain boundary ductile fracture (GBDF) is involved into the failure of this grade during hemming [5]. Fig. 6(a) shows that the propagation of the primary crack follows the grain boundaries because the polycrystalline structure of the material is visible on the fracture surface. The intact convex
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Fig. 8. Grain boundaries decohesion observed on RD × ND surface of grade A (a) at 50◦ bending, (b) at 180◦ bending.
surface next to the primary crack of the same sample is shown in Fig. 6(b). Relative sliding of the grains to each-other happened during increased deformation and leaded to height changes and formation of rough surface relief. Moreover, microcracks opened following the shear parallel to the grain boundaries – see the regions, pointed with arrows 1 in Fig. 6(b). Slip line markings and deformation steps, characteristic for inhomogeneous planar slip in the bulk of the grains are visible on the fracture surfaces of these microcracks. Such steps are also clearly seen in large amount on the primary crack fracture surface shown with the arrows 1 in Fig. 6(d)–(f). Their presence can be explained with decohesion of the grain boundaries which happens before the planar slip in the grains, i.e. the slip steps on the grain’s surfaces are formed after the fracture of the grain boundaries, when the adjacent grains no longer inhibit the dislocation motion in the neighboring grain. A fracture mechanism, involving reduction in the local cohesive stresses of the grain boundaries seems to be responsible for such type of fracture [6]. In confirmation of this mechanism, such decohesions of the grain boundaries were observed at quite early stages during the bending of grade B. Fig. 9(a) shows a SEM-micrograph of a polished RD × ND sample surface of this grade bent to 50◦ . The grain boundary debonding is already clearly seen to occur before any significant amount of deformation was accumulated in the bulk of the grains. The above observation supposes that the grain boundary debonding is rather triggered by the presence of high amount of grain boundary particles than because of the stresses arising from the dislocation pile-ups in the end of the slip lines at the grain boundaries dur-
ing the planar slip [5]. Even at the end of the bending when the hem is closed (180◦ ) the amount of accommodated strain is low and the grains are almost completely separated from each other – see Fig. 9(b). After 50◦ bending of grade A the height changes on the free sample surface are also present – see Fig. 10(a). However, larger accommodated strain expressed as surface slip lines and deformation steps exists in the interior of the grains as compared to grade B. The grain-to-grain compatibility is generally kept to the end of the bend – see Fig. 10(b) – predominantly due to the low amount of grain boundary particles in combination with the lower strength of the matrix in this grade. The grain boundary fracture in grade B involves also another mechanism triggered by strain localization in the weak grain boundary areas. The micro-dimples observed on the fracture surfaces – see the arrows 2 in Fig. 6(c), (e) and (f) – are indication that the heterogeneously nucleated grain boundary particles visible in Fig. 3(b) lead to formation, growth and coalescence of micro-voids, which finally results in grain boundaries ductile fracture (GBDF). This type of fracture is very common in precipitation-hardening aluminium alloys especially after aging in T6 temper state [5]. The GBDF is promoted by large amounts of grain boundary particles created by precipitation of the same alloying elements, forming the matrix precipitates (in this case Mg, Si and Cu) and is facilitated by the surrounding precipitation free zones (PFZs) that inevitable form along the grain boundaries during the aging process. According to the observations of the fracture behaviour during bending of grade B and considering the two above described
Fig. 9. Grain boundaries decohesion observed on RD × ND surface of grade B (a) at 50◦ bending, (b) at 180◦ bending.
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Fig. 10. Slip lines and deformation steps observed on RD × ND surface of grade A (a) at 50◦ bending, (b) at 180◦ bending.
Fig. 11. Through thickness optical micrograph of middle-width RD × ND plane of hemmed grade B sample. (a) Intergranular crack propagation in zone A and transgranular in zone B. The scale bar is 100 m. (b) Insert with higher magnification of zone A.
fracture mechanisms one can conclude that the presence of Mg2 Si grain boundary particles observed in the microstructure of grade B unhesitatingly contributes to the bad hemming response of this high-strength aluminium alloy grade. It seems that the high tensile stresses in the outer layers of the bent sample as well as the suppressed sliding due to the geometrical restrictions during bending, in combination with the weak grain boundary structure and the high strength of the matrix favour the propagation of predominantly intergranular type of fracture. Fig. 11 presents a through-thickness optical micrograph of the middlewidth RD × ND plane of a hemmed sample. The primary crack which fractures the hem propagates intergranularly in the tensile zone A. Approaching the sample’s neutral line, the tensile stresses
Table 5 Compositional and microstructural factors, influencing the bendability of AA6016type grades. Factors
Influence
Chemical composition (Mg and Si content) Volume fraction of strengthening phases Volume fraction of constituent particles Volume fraction of grain boundary particles Grain size
Strong Strong Medium Strong Medium
weaken and change to compressive, which alters the type of fracture to transgranular – see zone B in Fig. 11. 5. Summary and conclusions The hemming response of two industrial grades age hardening AA6016-type aluminium alloy sheets in T4P temper state was studied in this work and discussed as a function of the chemical composition and the applied heat treatment. The results confirmed that the volume fractions of the strengthening phases as well as their distribution in the microstructure are the main parameters controlling the bendability of these alloys. While the absolute amount of Mg–Si phases is a function of the chemical composition of the alloys, their distribution into the microstructure depends mainly on the applied heat-treatment. The microstructural observations in grade B showed dense population of Mg2 Si grain boundary particles heterogeneously formed by precipitation of the alloying elements in the grain boundaries. These particles, in combination with the high strength of the matrix in this grade are considered as the main reasons for the observed bad hemming response. The lower strength of grade A, resulting from the applied shorter solutionizing combined with a weak population of grain boundary particles determines its relatively good hemming properties. Other microstructural parameters, like grain
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size distribution and volume fraction of constituent particles play minor role and can be only used for fine tuning of the strength and the bending properties of these grades. Table 5 summarizes the influencing compositional and microstructural factors and the extent to which they affect the hemmability of these grades. The geometrical limitations of the bending deformation during hemming favour the intergranular fracture mechanism when the grain boundaries are weakened due to the presence of large amounts of particles. The fracture mechanism, including voids formation in the vicinity of the coarse (1 ÷ 20 m) constituent particles and their interaction with the propagating shear bands becomes not determinative for the fracture process. Thus, the fluctuations of the iron content of the grades in the studied limits (0.1 ÷ 0.25% Fe) and the slight differences in the volume fraction of the constituent particles do not play primary role in the hemming ability of these grades. According to the obtained in this study experimental results and their discussion, the following conclusions can be drawn: 1. Deterioration of the hemming ability of high strength AA6016type aluminium alloy sheets is observed in microstructures containing weakened grain boundaries due to the formation of large population of Mg2 Si particles during the solutionizing heattreatment of the grades. 2. The combination of strong interior of the grains and weak boundary areas triggers propagation of intergranular cracks which catastrophically decrease the bending ability. 3. The intergranular mechanism of fracture becomes more pronounced after natural aging of the grades due to the increased strength of the matrix determining the larger relative difference in strength with the remained weak grain boundary areas. 4. Two mechanisms of intergranular fracture are involved (i) grain boundaries debonding at the early stages of bending, due to the weak adhesion of the grains and (ii) grain boundary ductile fracture at later stages, promoted by strain localization and nucleation, growth and coalescence of voids around the grain boundary particles.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure controlled bending response in AA6016 Al alloys Aleksandar Davidkov a,∗ , Roumen H. Petrov a,b , Peter De Smet c , Bruno Schepers c , Leo A.I. Kestens a a
Department of Material Science and Engineering, Gent University, Technologiepark 903, 9052 Gent, Belgium Department of Material Science and Engineering, Delft University of Technology, Mekelweg 2, 2628 Delft, The Netherlands c Aleris Aluminum Duffel B.V.B.A., A. Stocletlaan 87, 2570 Duffel, Belgium b
a r t i c l e
i n f o
Article history: Received 8 May 2011 Accepted 25 May 2011 Available online 31 May 2011 Keywords: Aluminium 6xxx alloys Microstructure Bendability Hemming Ductile intergranular fracture
a b s t r a c t A contemporary approach in the car weight reduction is the use of low weight and high strength Al alloys sheets for hang-on body panels production. The final step in the forming route of such panels is the attachment of the outer skin to the inner part of the panel by applying a hemming operation. This joining method is cheap, easy to perform and environment-friendly, but requires severe 180◦ bending of the edges of the outer skin which quite often results in cracking or complete tearing of the bend surface. Such kind of failure restricts the further application of the hemmed products. The microstructures after solution heat treatment and pre-aging (T4P temper state) of two grades age-hardening AA6016-type aluminium alloy sheets were studied in this work by means of optical microscopy, scanning electron microscopy and electron backscatter diffraction. The obtained results were related to the hemming response of the grades. It was found that the alloy composition is one of the main parameters controlling the bendability of these grades through the amount of the formed strengthening phases. However, the applied thermal treatment remains the key factor responsible for the favorable distribution of these phases into the microstructure. The grain size and the volume fraction of the constituent particles were found to play secondary role in forming the material bending properties and can be only used for their fine tuning. The presence of Mg2 Si (-phase) and/or Al1.9 CuMg4.1 Si3.3 (Q-phase) particles in the grain boundaries structure was recognized as a critical microstructural feature causing severe reduction in the bending ability of the sheets by promoting an intergranular fracture. The possibility of grain boundaries failure exponentially raises with the time passed due to the natural aging. © 2011 Elsevier B.V. All rights reserved.
1. Introduction The weight of automobiles continuously increases during the years with improving performance, comfort and safety features. At the same time the importance of environmental and economical issues grows and leads to raising demand of weight reduction of the cars. The favorable combination of low specific density (approximately 1/3 of the steel), good corrosion resistance, strength and formability of aluminium alloys make them attractive material for this application. As a sheet material aluminium alloys are being used mainly for production of hang-on car body panels like hoods, trunk lids, doors, etc. Currently, 6xxx heat treatable Al–Mg–Si–(Cu) alloys are used worldwide for the outer skins of the panels as they satisfy well the main requirements for high strength, dent resis-
∗ Corresponding author. Tel.: +32 9 331 04 62; fax: +32 9 264 58 33. E-mail addresses: [email protected] (A. Davidkov), [email protected] (R.H. Petrov), [email protected] (P. De Smet), [email protected] (B. Schepers), [email protected] (L.A.I. Kestens). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.05.055
tance and good formability. The major alloying element in these alloys is Si, which together with Mg and Cu forms strengthening Mg–Si or Al–Mg–Si–Cu phases. Excessive amounts of Si are usually added to improve the precipitation hardening during the short paint bake process, respectively to sufficiently raise the dent resistance of the parts [1]. Such approach enables downgauging and further weight reduction while maintaining sufficient strength for the outer panels, but however, it can have detrimental effect on the sheet’s hemming ability. Consequently, the properties of the material for such applications should be in some way compromised in order to balance between the contradictive requirements for high strength and good hemmability. On other hand, the industrial production process should be conformed to the commercial requirements for minimizing the time and cost factors during production. The hemming is an environmental friendly assembly method used in the automotive industry to join the outer skins to the inner parts of the hang-on car body panels. The requirements to the alloys subjected to hemming operation are very severe as the edge of the skin is folded over the inner part of the panel by bending to
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180◦ over a radius, equal to the half of the sheet thickness. Finally, the hem is closed by applying a pressing operation. The pre-strain already accumulated in the preceding forming operations additionally makes the sheet material more susceptible to fracture during the hemming. The aluminium alloys which were worldwide recognized as being able to satisfy these severe requirements for the car body skins are the heat treatable 6xxx series Al alloys. They are usually delivered to the car manufacturer (in some cases with a delay) in the form of coils in pre-aged T4P temper state [2]. The microstructure at this stage is quite complex and determines the properties of the material before the final forming operations. Several groups of intermetallic particles can be found embedded in the recrystallized and supersaturated matrix of these alloys [3]: (i) large-sized primary constituent particles of Al–Fe(Mn,Cr)–Si-phase (typically 1 ÷ 20 m in size), (ii) dispersoid particles of Mn–Al, Mg–Si or Al–Mg–Si–(Cu) phases (typically 0.1 ÷ 1 m in size) and (iii) Al–Mg–Si–(Cu) clusters formed during the pre-aging heat treatment which aim to preserve the formability of the material by stabilizing the microstructure and slowing down the process of natural aging and to ensure a fast bake-hardening response during the insufficiently long final artificial age-hardening process. Different number of pre-aging steps can be included into the heat treatment of these alloys, however, not carefully chosen pre-aging time or temperature can have detrimental effect on the sheet formability [2,4]. The bendability is one of the most important forming attributes of precipitation hardening 6xxx aluminium alloys. Many metallurgical, technological and processing factors in their rather complex combination can have a strong influence on it. The fracture during bending of these age-hardening aluminium alloys usually occurs through two main mechanisms. The first one, which in many cases results from an improperly applied heat treatment, is the intergranular fracture due to the presence of heterogeneously nucleated grain boundary particles [5]. Evensen et al. [6] investigated the fracture behaviour of grain boundaries as a function of the grain size and the precipitation hardening of the matrix. They found that by promoting more dispersed slip, or by grain refinement the intergranular fracture mode can be replaced by the conventional ductile rupture in which void formation and shear banding became more pronounced than grain boundary decohesion. Practically, several possible sources of intergranular fracture can be delineated: (i) reduction in the local cohesive stress due to segregation or presence of grain boundary particles, (ii) initiation, growth and coalescence of micro-voids, formed on the grain boundary particles, (iii) localization of the deformation in precipitation free zones (PFZs) adjacent to the grain boundaries, (iv) promotion of coarse slip in the bulk of the grains, and (v) coarsening of the grain size [5,7–9]. The second fracture mechanism is accepted as conventional and includes sequence of events causing localization of the strain and intense shearing. In this case it is considered that the bendability of aluminium alloy sheets is directly governed by the combined effect of micro-voids formation around large secondphase particles, strain localization and propagation of macroscopic shear bands [10–12]. Lievers et al. [13] confirmed that the shear bands and void damage development play co-operative role in promoting ductile fracture at the outer free surface during bending. They also found that the changes in porosity, which create spatial differences in constitutive behaviour needed for strain localization, are exceptionally pronounced once the strain hardening of the metal matrix has been exhausted. Hence, the bendability becomes better for materials which have higher strain hardening rate and smaller initial voids fraction. In different materials with the same amount of second-phase particles, the bendability becomes mainly dependent on the value of the strain hardening factor n, which is controlled by the properties of the aluminium matrix [14].
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Table 1 Chemical composition of the alloys, wt%. Grade
Si
Mg
Fe
Cu
Mn
Al
A B
1.00 0.90
0.47 0.60
0.23 0.10
0.17 0.17
0.07 0.15
Bal. Bal.
The analysis of the existing literature data show that there is still no comprehensive understanding for the influence of different microstructural features of age-hardening Al alloys on their bendability. In the current work the microstructures of two AA6016-type grades processed in the form of thin sheets were studied with the aim to find out the main microstructural features provoking the observed significant differences in their hemmability. Interrupted hemming tests in laboratory conditions were used to estimate the hemming properties of the grades and in situ bending tests were performed in a scanning electron microscope (SEM) in order to study the local deformation and the fracture development during bending. The chemical composition of the alloy, the volume fraction and the distribution of the strengthening phases as well as the grain boundary structure of the grades were considered as parameters potentially controlling the bendability. 2. Experimental Two heat-treatable AA6016-type grades in the form of 1 mm thick sheets were studied in this work after being heat-treated to T4P temper state by applying several consecutive steps, i.e. solutionizing, quenching and two pre-aging steps. During solutionizing, the grade denominated here as A was soaked for 5 s at solutionizing temperature and the other grade, denominated here as B was soaked for 55 s at the same temperature. All the other parameters of the heat-treating process were the same for both grades. Part of the sheets was kept at room temperature for several months and another part was kept at a temperature of −18 ◦ C to avoid any natural aging during the studies. The examined grades also differ in their chemical compositions but the variations are within the standard tolerances for AA6016-type aluminium alloys – see Table 1. Samples for characterization of the in-plane and the throughthickness microstructure were cut from the middle part of both sheets. Standard mechanical polishing procedure [15] was used to prepare samples for observation of the constituent particles distribution at magnification 20× with an optical microscope “Zeiss Jenavert”. Multiple digital images were acquired over an area of 0.4 mm2 and further processed by specialized software for quantitative microstructural analyses “ImageJ” [16]. The general grain structures as well as the macro-cracks paths through the microstructure were characterized using a polarized light after anodizing the mechanically polished surfaces of the samples by submerging in a 3% HFB4 water solution for a period of 1–1.5 min under an applied tension of 25 V. Samples for SEM, EDX and EBSD analysis were prepared using electrolytic polishing by means of 10% HClO4 -solution in methanol at 22 ◦ C under tension of 39 V for a time of 12 s. The characterization of the microstructures at different magnifications was done by an environmental scanning electron microscope ESEM FEI XL30 with LaB6 filament employing acceleration voltages of 5 ÷ 25 kV and working distances of 10 ÷ 20 mm. Energy Dispersive X-ray (EDX) analyses on the polished surfaces of the samples were carried out to identify the chemical composition of the matrix, constituent particles (1 ÷ 20 m) and dispersoid particles (0.1 ÷ 1 m) at a working distance of 10 mm, acceleration voltage of 25 kV and acquisition time of 100 s which correspond to approximately 1800 counts per second. Minimum 10 measurements for each type of particle were made and the results were presented as mean values. EBSD measurements in both RD × TD and RD × ND planes were carried out
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Fig. 1. Through-thickness grain structure of 1 mm thick sheets. (a) grade A, (b) grade B, RD × ND plane (RD vertical), (c) average diameter of grains (based on area fraction of grains).
and the data were collected and post-processed by TSL-OIM® version 4.6 software. Accelerating voltage of 25 kV, magnification of 180× and a working distance of 22 mm were used to scan over of an area of 1.0 mm × 1.5 mm with a step of 2 or 2.5 m in hexagonal scan grid. Data for the average grain sizes and grain size distribution were obtained and analysed. The through thickness sections of mechanically polished samples with a size of 7 mm × 12 mm were observed during 180◦ in situ bend tests with the bend line (BL) parallel to the rolling direction (RD) of the sheets. Three-point bending micro-device designed to work with FEI XL30 environmental scanning electron microscope was used for this purpose. During interrupted laboratory hemming tests rectangular samples with a size of 10 mm × 20 mm were fully folded without inserting of an intermediate sheet around axis (BL) oriented parallel or perpendicular to the sheet rolling direction. This bending test procedure was applied in three steps – to 90◦ , to 135◦ and with final pressing to 180◦ . The assessment of the hemmability was made by observing the micro- and macroscopic appearance of cracks on the outer convex surface of the samples. A comparative hemming factor (HF) ranging from 1 (very bad bendability) to 4 (very good bendability) was used to scale the hemming quality of the tested sheets. 3. Results The through-thickness microstructures of both T4P AA6016type grades consist of fine equiaxial recrystallized grains – see Fig. 1(a) and (b). However, grade B shows more heterogeneous grain distribution and larger average grain diameter in comparison to grade A, especially in the areas close to the sheet surface. The average grain diameter calculated as a weighted sum of the grain area fractions is presented in Fig. 1(c). Large ␣-Al(FeMnCr)Si constituent particles form during the solidification of these alloys and remain almost unchanged during the whole cycle of thermomechanical treatment [17]. According to the quantitative optical observations, a slightly higher volume fraction of constituent particles was found in grade A with its maximum at equivalent particle diameter between 1 and 2 m – see Fig. 2 and Table 2. The EDX analyses confirmed their Al–Fe–Si–Mn–Cr composition in both cases [18]. SEM investigations were made in order to reveal the distribution of dispersoid particles in the microstructure, especially of those lying in the grain boundaries, as they can play destructive role on
Table 2 Density and area fraction of constituent particles. Grade Density per mm Area fraction, %
2
A
B
2726 2.05
2644 1.49
the bending properties. The qualitative observations of multiple areas in both grades show higher amount of grain boundary particles in grade B. They are well aligned with the grain boundaries and in many cases quite coarse – see Fig. 3(b). Particles are also present in the grain boundaries of grade A but they are smaller and much less densely populated. Despite the compositional changes during the electrolytic polishing, the EDX analyses of these particles show close to stoichiometric amount of Mg and Si atoms for the stable Mg2 Si (-phase) (Mg/Si = 1.73 [19]) – see Fig. 4 and Table 3. The tensile properties of the grades were studied after one and two months of natural aging in pre-aged T4P temper state as well as in artificially aged T8 temper state after 2% tensile deformation and final paint-bake. Grade A has shown lower strength compared to grade B but at the same time both grades have equal ductility when expressed in terms of uniform elongation (UE), total elongation (TE) and strain hardening factor n – see Table 4. The hemming performance of the sheets was studied by observing the macro-appearance of the outer convex surface of 180◦ bent samples. Fig. 5 shows the surfaces of the samples subjected to bending in three months after the production. In the case, when the bend line (BL) was parallel to the rolling direction (RD) of the sheet (see the upper row in Fig. 5) both grades showed relatively good hemmability estimated with HF = 3.25 for grade A and HF = 2.25 Table 3 Chemical composition of grain boundary particles. Element
wt%
at%
Mg Al Si Cr Mn Fe Cu Total
1.36 97.02 0.85 0.16 0.19 0.15 0.27 100
1.51 97.30 0.82 0.08 0.09 0.07 0.11 100
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Fig. 2. Optical micrographs of mechanically polished samples in RD × TD plane (RD horizontal). (a) grade A, (b) grade B, (c) size distribution of constituent particles.
Fig. 3. SEM micrographs of Mg2 Si particles allocated at the grain boundaries. (a) grade A, (b) grade B.
for grade B. In the case, when the BL was perpendicular to the RD (lower row), the hemming performance of the grades was estimated as good for grade A and as bad for grade B with a comparative hemming factors respectively HF = 2.5 and HF = 1.5.
The observation of the fracture in grade B showed that a full length primary crack spreads along the hem. The SEM investigations reveal that it propagates mainly through a ductile grain boundary fracture (at room temperatures). The micrographs in
Fig. 4. Grain boundary particle in grade B. (a) SEM micrograph, the scale bar is 500 nm, (b) EDX spectrum.
Table 4 Tensile properties of the grades. Grade
A B
One month after pre-bake
Two months after pre-bake
After paint-bake
YS (MPa)
UTS (MPa)
UE (%)
TE (%)
n (at 5%)
YS (MPa)
UTS (MPa)
UE (%)
TE (%)
n (at 5%)
YS (MPa)
124 141
241 257
22.5 22.3
26.0 25.6
0.26 0.26
133 149
250 266
22.4 22.1
26.4 26.5
0.26 0.26
235 250
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Fig. 5. Hemmability of the sheets with BL RD (upper row) and BL ⊥ RD (lower row). (a) grade A, (b) grade B.
Fig. 6. SEM micrographs of fracture surface in grade B. The arrows show specific fracture zones discussed in the text.
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Fig. 6 present the details of the fracture mechanism including grain boundaries debonding and rupture as well as strain localization and micro-voids formation at the grain boundary particles.
4. Discussion Regarding the chemical composition shown in Table 1, both grades are characterised as alloys with excess of Si, with a Mg/Si ratio respectively 0.47 for grade A and 0.67 for grade B. The excessive amount of Si is added in order to improve the precipitation hardening during the final insufficiently long paint-bake process [20]. However, the excess Si increases also the initial T4 yield strength influencing the volume fraction of the strengthening phases. Many authors [20–22] agree that the bendability deteriorates with increasing yield strength. Hence, because of the higher excess of Si in grade A, one should expect worse bending response in comparison to grade B. Actually, the volume fraction of the strengthening phases in Si-excess alloys is always limited by the level of Mg, as it is completely removed by the Si from the solution [20,23]. Thus, the volume fraction of the strengthening phases in grade B will be higher due to the higher Mg content of this alloy and the bendability will be reduced. Better hemming should be expected from the lower-Mg and lower-strength grade A. This expectation is confirmed by the results obtained from the hemming tests – cf. Fig. 5. The Fe content in these alloys influences the volume fraction of the coarse non-deformable constituent particles which are considered as strongly involved into the fracture process during bending by their role in voids formation [24,25]. Part of the excess Si is also consumed by those particles, but however the sufficiently high excess of Si in both alloys provides the further formation of strengthening phases. In order to diminish the volume fraction of the constituent particles and to improve the bendability in grade B, the Fe content was dropped to 0.10 wt%. According to the microstructural study, as expected, lower volume fraction of constituent particles was found in grade B in comparison to grade A – see Fig. 2 and Table 2. However, in spite of the reduced volume fraction of constituent particles, the bendability of grade B does not improve as compared to the higher-Fe-containing grade A. This fact gives a clue that the fracture mechanisms involved in this grade during bending does not necessary include formation of voids around constituent particles and their further interaction with the propagating shear bands. The EDX analyses of those particles show higher content of Mn in grade B than in grade A [18] which is in agreement with the chemical composition of these alloys – cf. Table 1. However, the reduced amount of dissolved Mn reduces the possibility for formation of incoherent, non-deformable ␣-AlMnSi dispersoid particles (with a size around 0.1 m) during the homogenization annealing. Such dispersoids are considered as improving the bendability by promoting of grain refinement, homogenization of the slip distribution (decreasing the slip band spacing and broadening the slip band width) and in this way reducing the tendency of intergranular fracture [26]. The soak time during the solutionizing of grade B was much longer in comparison to grade A. Hence, its matrix was enriched of alloying elements to a greater extend and its precipitation potential was increased. The large grain boundary particles found in grade B – cf. Fig. 3(b) – were formed by precipitation of the same alloying elements (Mg, Si and also Cu) which form the strengthening phases in this alloy. It is logical to conclude that if these large particles are formed in the grain boundaries, the adjacent zones will be depleted from alloying elements, i.e. the strengthening ability of the adjacent to the grain boundaries areas would be lower than the bulk of the grains. At the same time the cohesive bonds between the grains are weakened due to the presence of large amount of grain
Fig. 7. Deterioration of hemmability during natural aging. The steeper line of grade B shows faster deterioration. Samples bent along RD (BL ⊥ RD).
boundary particles. The continuously increasing difference in the strength between the interior of the grains and the grain boundary areas, taking place for instance during natural aging in T4P temper state, will automatically lead to increased possibility of grain boundaries fracture followed by severe deterioration of the bending properties of grade B. Such a fracture mechanism is mentioned in [27] and gives one very possible explanation for the observed failure of the grain boundaries during the bending in grade B. The scanning electron microscope fractographs of the cracked samples of grade B are presented in Fig. 6 and further discussed in details in the next paragraph. The increment of strength in T4P temper state due to natural aging is observed during the storage period at ambient temperatures in both grades. The results from the tensile tests, given in Table 3 show the usual dependence of YS and UTS on the time elapsed after solutionizing. For example the YS rises with about 7% for grade A and with 5% for grade B in one month of natural aging. The UTS rises respectively with about 5% and 3% for the same period of time. In both cases a relatively uniform increase of YS and UTS is observed. When tracing the changes in the hemming behaviour of the grades – see Fig. 7 – in spite of the subjectivity in determining the HF, which is obtained by visual observation of the formed hem, it seems that the hemmability of the grades deteriorates not as uniform as the increase of the strength for the same period of natural aging. The better hemmability of grade A one month after solutionizing, assessed with HF = 3.0 is reduced within one month with about 8% to HF = 2.75 and the worse hemmability of grade B assessed with HF = 2.5 is reduced for the same period with 20% to HF = 2.0. This more pronounced deterioration in the hemmability of grade B could confirm the abovementioned hypothesis for the possible continuous increase of the strength difference between the grain boundaries and the bulk of the grains during the natural aging, which could determine the prevailing intergranular mechanism of fracture observed in this grade. This strength difference is supposed to be the highest for the grades aged to peak yield strength [28]. Such mechanism of failure is possible in grade A too, as its microstructure also shows particles embedded in the grain boundaries – cf. Fig. 3(a). Fig. 8 presents some initial features of grain boundary fracture observed in grade A. However propagation of cracks did not occur until the end of the 180◦ bend. It seems that the sparse population of grain boundary particles in combination with the lower strength of the matrix in grade A make this fracture mechanism less possible. The study of the fracture surfaces in grade B (cf. Fig. 6) supposed that the phenomenon of grain boundary ductile fracture (GBDF) is involved into the failure of this grade during hemming [5]. Fig. 6(a) shows that the propagation of the primary crack follows the grain boundaries because the polycrystalline structure of the material is visible on the fracture surface. The intact convex
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Fig. 8. Grain boundaries decohesion observed on RD × ND surface of grade A (a) at 50◦ bending, (b) at 180◦ bending.
surface next to the primary crack of the same sample is shown in Fig. 6(b). Relative sliding of the grains to each-other happened during increased deformation and leaded to height changes and formation of rough surface relief. Moreover, microcracks opened following the shear parallel to the grain boundaries – see the regions, pointed with arrows 1 in Fig. 6(b). Slip line markings and deformation steps, characteristic for inhomogeneous planar slip in the bulk of the grains are visible on the fracture surfaces of these microcracks. Such steps are also clearly seen in large amount on the primary crack fracture surface shown with the arrows 1 in Fig. 6(d)–(f). Their presence can be explained with decohesion of the grain boundaries which happens before the planar slip in the grains, i.e. the slip steps on the grain’s surfaces are formed after the fracture of the grain boundaries, when the adjacent grains no longer inhibit the dislocation motion in the neighboring grain. A fracture mechanism, involving reduction in the local cohesive stresses of the grain boundaries seems to be responsible for such type of fracture [6]. In confirmation of this mechanism, such decohesions of the grain boundaries were observed at quite early stages during the bending of grade B. Fig. 9(a) shows a SEM-micrograph of a polished RD × ND sample surface of this grade bent to 50◦ . The grain boundary debonding is already clearly seen to occur before any significant amount of deformation was accumulated in the bulk of the grains. The above observation supposes that the grain boundary debonding is rather triggered by the presence of high amount of grain boundary particles than because of the stresses arising from the dislocation pile-ups in the end of the slip lines at the grain boundaries dur-
ing the planar slip [5]. Even at the end of the bending when the hem is closed (180◦ ) the amount of accommodated strain is low and the grains are almost completely separated from each other – see Fig. 9(b). After 50◦ bending of grade A the height changes on the free sample surface are also present – see Fig. 10(a). However, larger accommodated strain expressed as surface slip lines and deformation steps exists in the interior of the grains as compared to grade B. The grain-to-grain compatibility is generally kept to the end of the bend – see Fig. 10(b) – predominantly due to the low amount of grain boundary particles in combination with the lower strength of the matrix in this grade. The grain boundary fracture in grade B involves also another mechanism triggered by strain localization in the weak grain boundary areas. The micro-dimples observed on the fracture surfaces – see the arrows 2 in Fig. 6(c), (e) and (f) – are indication that the heterogeneously nucleated grain boundary particles visible in Fig. 3(b) lead to formation, growth and coalescence of micro-voids, which finally results in grain boundaries ductile fracture (GBDF). This type of fracture is very common in precipitation-hardening aluminium alloys especially after aging in T6 temper state [5]. The GBDF is promoted by large amounts of grain boundary particles created by precipitation of the same alloying elements, forming the matrix precipitates (in this case Mg, Si and Cu) and is facilitated by the surrounding precipitation free zones (PFZs) that inevitable form along the grain boundaries during the aging process. According to the observations of the fracture behaviour during bending of grade B and considering the two above described
Fig. 9. Grain boundaries decohesion observed on RD × ND surface of grade B (a) at 50◦ bending, (b) at 180◦ bending.
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Fig. 10. Slip lines and deformation steps observed on RD × ND surface of grade A (a) at 50◦ bending, (b) at 180◦ bending.
Fig. 11. Through thickness optical micrograph of middle-width RD × ND plane of hemmed grade B sample. (a) Intergranular crack propagation in zone A and transgranular in zone B. The scale bar is 100 m. (b) Insert with higher magnification of zone A.
fracture mechanisms one can conclude that the presence of Mg2 Si grain boundary particles observed in the microstructure of grade B unhesitatingly contributes to the bad hemming response of this high-strength aluminium alloy grade. It seems that the high tensile stresses in the outer layers of the bent sample as well as the suppressed sliding due to the geometrical restrictions during bending, in combination with the weak grain boundary structure and the high strength of the matrix favour the propagation of predominantly intergranular type of fracture. Fig. 11 presents a through-thickness optical micrograph of the middlewidth RD × ND plane of a hemmed sample. The primary crack which fractures the hem propagates intergranularly in the tensile zone A. Approaching the sample’s neutral line, the tensile stresses
Table 5 Compositional and microstructural factors, influencing the bendability of AA6016type grades. Factors
Influence
Chemical composition (Mg and Si content) Volume fraction of strengthening phases Volume fraction of constituent particles Volume fraction of grain boundary particles Grain size
Strong Strong Medium Strong Medium
weaken and change to compressive, which alters the type of fracture to transgranular – see zone B in Fig. 11. 5. Summary and conclusions The hemming response of two industrial grades age hardening AA6016-type aluminium alloy sheets in T4P temper state was studied in this work and discussed as a function of the chemical composition and the applied heat treatment. The results confirmed that the volume fractions of the strengthening phases as well as their distribution in the microstructure are the main parameters controlling the bendability of these alloys. While the absolute amount of Mg–Si phases is a function of the chemical composition of the alloys, their distribution into the microstructure depends mainly on the applied heat-treatment. The microstructural observations in grade B showed dense population of Mg2 Si grain boundary particles heterogeneously formed by precipitation of the alloying elements in the grain boundaries. These particles, in combination with the high strength of the matrix in this grade are considered as the main reasons for the observed bad hemming response. The lower strength of grade A, resulting from the applied shorter solutionizing combined with a weak population of grain boundary particles determines its relatively good hemming properties. Other microstructural parameters, like grain
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size distribution and volume fraction of constituent particles play minor role and can be only used for fine tuning of the strength and the bending properties of these grades. Table 5 summarizes the influencing compositional and microstructural factors and the extent to which they affect the hemmability of these grades. The geometrical limitations of the bending deformation during hemming favour the intergranular fracture mechanism when the grain boundaries are weakened due to the presence of large amounts of particles. The fracture mechanism, including voids formation in the vicinity of the coarse (1 ÷ 20 m) constituent particles and their interaction with the propagating shear bands becomes not determinative for the fracture process. Thus, the fluctuations of the iron content of the grades in the studied limits (0.1 ÷ 0.25% Fe) and the slight differences in the volume fraction of the constituent particles do not play primary role in the hemming ability of these grades. According to the obtained in this study experimental results and their discussion, the following conclusions can be drawn: 1. Deterioration of the hemming ability of high strength AA6016type aluminium alloy sheets is observed in microstructures containing weakened grain boundaries due to the formation of large population of Mg2 Si particles during the solutionizing heattreatment of the grades. 2. The combination of strong interior of the grains and weak boundary areas triggers propagation of intergranular cracks which catastrophically decrease the bending ability. 3. The intergranular mechanism of fracture becomes more pronounced after natural aging of the grades due to the increased strength of the matrix determining the larger relative difference in strength with the remained weak grain boundary areas. 4. Two mechanisms of intergranular fracture are involved (i) grain boundaries debonding at the early stages of bending, due to the weak adhesion of the grains and (ii) grain boundary ductile fracture at later stages, promoted by strain localization and nucleation, growth and coalescence of voids around the grain boundary particles.
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