The uniaxial compression tests were performed to assess the influence of precipitation state on flow stress, work hardening behavior and texture of modified Al-Si-Cu-Mg alloy modified with V, Zr and Ti after T6 heat treatment. A semi-empirical model was adopted and used to interpret the changes in work-hardening rate during plastic deformation. The experimental result and semi empirical model were found in good agreement. Development of crystallographic texture was also measured and made a correlation with deformation behavior of the alloy. Before and after deformation, week texture was observed in all conditions. However, before deformation, random texture was found in both states and during deformation, shear type texture components, {001}<110>, {111}<110> and {112}<110>, was dominant which was influenced the hardening of the alloy..
Keywords: Aluminum-silicon; compression; work-hardening; precipitates; texture component.
1. Introduction
During last few decades global anthropogenic emissions are growing very fast with a ground transportation being substantial contributor. Therefore, reducing vehicle mass is one of many strategies to improve fuel economy and reduce emissions that RE being widely considered [1-3]. As lightweight materials, aluminum (Al) alloys are the forefront because of their weight to strength ratio, corrosion and wear resistance, thermal, electrical conductivity and castability. Among the Al alloys, the Al-Si-Cu-Mg grades are widely used for automotive structural components [4-8].
The age-hardenable alloys change their microstructure as well as mechanical properties due to heat treatment [9]. The most common heat treatment of Al alloys is composed of solution annealing and aging. During solution stage, the alloying elements are dissolved in the matrix and the alloy behaves like a single phase material. Then, during aging, different types of precipitates are formed depending on the aging parameters and alloy chemical composition. When external load is applied, these precipitates block the movement of dislocations which increases the material strength [10-14]. There are mainly two kinds of precipitates formed during aging; coherent or semi-coherent which are sharable and incoherent which is non-sharable. The fine particles are sharable which resist the movement of the dislocation line and the line cuts through the precipitates. On the other hand, coarse particles are non-sharable and forme Orowan loops by bowing the dislocation line. Thus, presence of precipitates adds to the microstructure complexity and affects alloy mechanical properties [15-17].
The work-hardening behavior of age-hardened alloys depends on the nature of precipitates. Changing the state of precipitates from sharable to non-sharable, the deformation behavior changed and influenced the work hardening behavior during tensile[17]The influence of precipitation on work-hardening behavior of AA6111and AA7030 aluminum alloys were studied and developed a model to determine hardening behavior due to the sharable/non-sharable transition from a series of tensile test [18,19]. Another investigation showed that the precipitation in Al-Sc alloy made significant differences in work hardening behavior with aging time and temperature [20]. In addition, the effect of Ti, V and Zr addition on tensile and fatigue behavior was also investigated [21].
The orientation of the crystal is one of the key crystallographic characteristics that determine micro-plastic deformation in metals. Aluminum has a face centered cubic crystal structure where the deformation occurs by slip and slip system is composed of slip plane and direction. Since, the work hardening in polycrystalline materials is also influenced the crystallographic texture where increasing plastic deformation, the grains tend to rotate toward more stable orientations, leading to an alteration of the hardening behavior. Therefore, it is also important to understand the effect of crystallographic texture on hardening behavior of the alloys. Most of the research on texture evolution of Al alloys was performed for wrought alloys after rolling or extrusion [22-25]. Benum and Nes[26] and Kashihara and Inagaki [27] studied the effect of precipitates on texture evolution in Al alloys during rolling followed by solution treatment and aging. They reported that precipitates could control the formation of texture by pinning effect of grain boundary. The work-hardening behavior and texture formation of Al alloys during axial loading had been studied by many researchers. However, these phenomena are not yet clear for the cast alloys during deformation. Therefore, the work-hardening behavior and texture evolution of modified Al-Si-Cu-Mg alloys during compression is not clear yet. The present study was done in to two folds; first, the stress strain behavior, microstructure evolution and work-hardening behavior during compression were explained by semi-empirical model and second, the influence of precipitates on texture evolution during compression of Al-Si-Cu-Mg alloy modified with V, Zr and Ti addition was investigated.
2. Experimental
A hypoeutectic Al-7%Si-1%Cu-0.5%Mg alloy modified with additions of transition metals V, Zr, and Ti was used as a test material in the present study. The nominal chemical composition is shown in Table 1. The alloy was cast from a temperature of 730 ï‚°C into ASTM B108 steel mold and then subjected to T6 heat treatment consisting of two-step solution annealing at 501ï‚°C for 0.5 h and at 525ï‚°C for 4.5 h, followed by artificial aging at 150ï‚°C for 100 h.
The compression tests were performed following ASTM E9M-09 standard at room temperature (25°C) on as-cast (AC), supper saturated solid solution (SSSS) and T6 heat treated samples using the computerized United tensile testing machine. A cylindrical type specimen with a dimension of 5Ã-8 mm was used for compression test at a strain rate of 10-3s-1. The flow stress contribution was calculated using semi-empirical model proposed by Sharma et al. [17] and Cheng et al. [18] . After the compression testing, the samples were subjected to metallographic examinations using light microscope (Nikon EPIPHOT) equipped with CLEMEX quantitative image analysis software.
To study the crystallographic orientation in the grain, distribution of textures were determined by measuring pole figures (PF) using line scan X-ray diffraction technique and PANalytical X-ray diffractometer. The experiment was conducted by measuring incomplete pole figures between Ψ = 0° and 75° in the back-reflection mode using Cu Kα radiation at 45 kV and 40 mA [28]. During X-ray diffraction, the loading axis of the sample was placed along the x-axis of the stage of diffractometer. Based on measured pole figures (PF), orientation distribution function (ODF) was calculated using MTEX software.
3. Results and Discussion
3.1 Influence of flow stress on microstructure
The typical stress strain behaviors of the studied alloy during compression in as-cast, SSSS and T6 heat treated conditions are given in Fig. 1. The obtained yield strengths from the experiments are 107, 164 and 318 MPa for SSSS, as-cast and T6 heat treated alloy, respectively. It is clear that SSSS sample has lower strength than AC and T6 samples due to dissolution of phases and absence of precipitates. Fig. 2 shows the micrograph of the alloy before and after deformation in as-cast and T6 heat treated condition. The as-cast and T6 heat treated both samples contained α-Al (#1), eutectic silicon (#2) and intermetallics (#3) phases. In the as-cast condition, eutectic silicon phase was needle or plate like shape (Fig. 2a) where almost globular or spherical eutectic silicon phase was found in T6 heat treated condition (Fig. 2c). The micrographs show that the major cracks, found in secondary phases, were almost parallel to the compressive loading axis (Fig 2b & d). Agarwal et al. conducted uniaxial compression test of 6061 (T651) alloy and found that the cracks were formed in Fe-rich intermetalics [29, 30]. Dighe et al. [31] and Poole and Dowdle , [32] observed the damage evolution of Si particles in Al-Si-Mg alloys and reported that the cracks in Si particles were appeared parallel to the loading axis.
In multiphase alloys, stress arises due to different crystal structure, misfit between phases and their shape. When the interfacial stress exceeds the strength of the interface or the stress exceeds the fracture stress of the particle, damage or debonding occurs; hence, the load transfer from matrix to particles is important. If the particles are spherical in shape, the interfacial stress generated around the particles, are not enough to fracture or debond the interface [33]. In the as-cast condition (Fig 2a), Si particles were plate like in shape which increased the interfacial stress and stress generated at the tip of the plate contributing to frequent cracking. On the contrary, T6 heat treated alloy exhibited less damage in Si particles which were almost spherical in shape (Fig. 2). As a result, the strength of as-cast alloy was lower than the T6 heat-treated alloy (Fig. 1).
3.2 Work hardening
Fig. 3(a) represents the work-hardening behavior of AC, SSSS and T6 heat treated alloy by plotting the instantaneous work hardening rate (θ=dσ/dε) versus flow stress increments (σ-σy) which is well known as Kocks-Mecking (KM) plot. As shown in Fig. 3(b), the typical KM plot for FCC single crystal shows three stages. The stages II to IV represent respectively: linear hardening rate, falling of hardening rate and finally saturation. However, the hardening stage III & IV are generally found in polycrystalline materials during compression [34, 35].
The instantaneous hardening rate of AC, SSSS and T6 heat treated alloy and plastic stress plot reveal only stage of III and IV as KM hardening. The nature of the curves indicates that the hardening rate is higher in T6 heat treated alloy compared to as-cast and SSSS states. At the beginning of stage III, forest hardening occurred due to dislocation interaction and accumulation, and dynamic recovery of dislocation which caused linear reduction of hardening rate. When extrapolating the stage III to zero hardening rate (θ=0), the flow stress at that point would be in steady state condition which is known as Voce stress, σv (Fig. 3b) where the mobile dislocations are generated, stored and recovered very rapidly. The evaluated Voce stresses of SSSS, AC and T6 heat treated alloys were found 188, 256 and 376 MPa (Fig. 4). However, in practice, stage III is interrupted by stage IV before it goes to zero at severe plastic deformation where dï±/dï¥ is zero [35].
3.3 Effect of precipitate on flow stress
The contribution of precipitates to flow stresses for age harden alloy can be expressed using the Eq. (1) which was suggested by Kocks and adapted by Sharma et al. [16] and Cheng et al.[17] in the following form:
, (1)
Where; is total stress, is intrinsic stress, is solid solution strengthening, is strengthening by eutectic phase, is strengthening by dislocation and is strengthening by precipitates, n is the parameter which varies between 1 to 2, depending the average obstacle density. Since, the studied alloy was peak aged at 150°C for 100 h hence, n = 1.5 was used here.
For a condition of supper saturated solid solution, when plastic strain; then the yield strength can be expressed as follows:
, (2)
where, is the yield strength of the materials.
For an age hardened alloy, when plastic strain; then the yield strength can be expressed as follows:
, (3)
The flow stress contributed by the eutectic phase can be written in the following form as follows
, (4)
Again, the flow strength contribution by solid solution depends on the solubility (Css) of alloying elements in solvent being α-Al which can be expressed using following equation:
, (5)
Here, is the critical resolved shear stress for solid solution contribution, and m are the materials constant where value is considered 24 MPa per atomic percent and m = 2/3 .
The yield strength in SSSS condition calculated from the empirical Eq. (5), was 82 MPa where Css was considered as 2.67. But, the experimentally obtained yield strength was 107 MPa which is 29 point higher than the calculated value. Therefore, the empirical Eq (5) was modified by Sharma et al. [16], by incorporating two constants A and B.
, (6)
The unknown constant A = 42 and B = 29 were used in the present study to calculate the σss in ssss condition
Mecking and Kocks[37] and Estrin [38] proposed the changes of dislocation density approach with plastic strain using following equation:
, (7)
where, Ï is the dislocation density, k1 = 6.86 ï‚´ 108 m-1 is the parameter related to statistical storage of dislocations, fk2 = 61 parameter related to dynamic recovery of dislocation and kD = 3.32 ï‚´ 1016 [17] is the parameter related to the storage of dislocation geometrically due to the nonshareable precipitates and forest dislocations.
Taylor proposed that the flow stress contribution by dislocation can be expressed as follows:
, (8)
where = 0.3 is a constant, b = 0.284 nm is Burgers vector, G = 27 GPa is the shear modulus, M = 3.06 is the Taylor factor for fcc metals.
From Eq. (7) Sharma et al. [17] proposed that for shareable precipitates, the term kD = 0 and simplified solution is:
, (9)
They also proposed for nonshareable precipitates, the parameter related to statistical storage of dislocations is k1 = 0. Then, the solution comes out as [16]:
, (10)
From Eq. (8) & Eq. (9) the hardening due to shareable precipitates can be expressed using the following equation:
, (11)
From Eq. (8) & Eq. (10) the hardening due to nonshareable precipitate can be express using the following equation:
, (12)
where hardening due to shareable precipitates and hardening due to nonshareable precipitates.
From Eq. (1) to Eq. (12) the total flow stress contribution can be calculated as follows:
, (13)
Using Eq. (13) the flow stress increment was calculated and plotted against the plastic strain considering the hardening due to shareable, nonshareable, and both precipitates and compared with the experimental results as shown in Fig. 5(a). It shows that the experimental results make a good agreement with the model where both types of precipitates were considered. So, the studied alloy was hardened during compression not only the shareable precipitates but also by the nonshareable precipitates. However, the flow stress was affected by shareable precipitates rather than the non-shareable precipitates. Such finding indicates that the alloy contained more fine precipitates.
Sharma et al. [16] and Cheng et al. [17] studied the work hardening behavior during tensile loading of AA2219, AA6111 and AA7030 alloys and reported that shareable and non-shareable, precipitates were present in the age hardened alloys and the hardening behavior was affected by both precipitates. At the beginning of deformation, dislocations cut off precipitates without forming loops. When the plastic deformation reaches higher level, the dislocations pile up and make Orowan loops around the non-shareable precipitates. So, at the beginning, hardening was controlled by shareable precipitates and at higher level of plastic deformation, the deformation was affected by shareable with non-shareable precipitates as seen in Fig 5(b).
3.4 Effect of precipitates on texture
The crystallographic texture depends on the alloy chemical compositions, microstructures, crystallographic structure as well as processing conditions [39, 40]. Fig. 6 shows the (111), (100) and (110) PF of as-cast and T6 heat treated samples before and after 15% deformation. In general, all the pole figures show weak texture with maximum intensity 2.4 to 4.6 mrd. Before deformation, the maximum intensity in the (111), (100) and (110) PFs of as-cast samples were 2.4, 2.3 and 1.7 mrd, respectively. After deformation the intensity changed to 2.6, 2.7 and 2.5 mrd and concentrated to the center of the (111) and (100) PF, while for (110) PF, it moved parallel to loading direction (LD). It was also noticed that in T6 heat treated condition, the intensity of (111), (100) and (110) PFs were 2.6, 2.7 and 1.5 mrd, respectively. But, after deformation, the intensity changed to 4.6 and 3 mrd for (111) and (100), respectively, while for (110) PF, the value of intensity was remained the same. However, after deformation, the position of maximum intensity of (111) and (100) PFs for both conditions were same. For PF of (110) in T6 condition, no changed is occurred while in as-cast condition the highest intensity was moved parallel to loading direction.
The orientation distribution function (ODF) of ϕ2=0 and ϕ2=45 sections for the alloy in as-cast, SSSS and T6 heat treated conditions is given in Fig. 7. The density function, f(g) was calculated at the fixed Eular angle of (35 45 0), (55 35 65), (90 30 45), (0 0 45), (0 55 45) and (0 35 45) for their respective typical deformed texture components of {011}<211>, {123}<634>, {112}<111>, {001}<110>, {111}<110> and {112}<110> known as Bs, S, Cu, shear1, shear2 and shear3 as given in Table 2. It is evident that there was no ideal fiber texture before deformation. But, after deformation, typical rolling texture with a well-developed β fiber was noticed. Before deformation, the values of f(g) in as-cast sample were found to 0.83, 1.72, 0.53, 1.52, 5.22 and 2.25 mrd for deformed texture components of {011}<211>, {123}<634>, {112}<111>, {001}<110>, {111}<110> and {112}<110>, respectively. But, in T6 condition, the value of f(g) for similar texture components was 0.14, 1.76, 0.07, 0.2, 0.92 and 0.43 mrd. After deformation of as-cast sample, the value of f(g) of Bs, S, shear2 and shear3 texture components decreased from 0.83 to 0.42, 1.72 to 0.91, 5.22 to 2.42 and 2.25 to 0.45 mrd, respectively, while the Cu and shear 1 increased from 0.53 to 1.52 and 2.46 to 6.53 mrd. In contrary, after deformation of T6 heat treated sample, the value of f(g) of S, Cu, shear 1, shear2 and shear3 texture components increased from 01.76 to 1.82, 0.07 to 1.51, 0.2 to 5.60, 0.92 to 4.12 and 0.43 to 0.74 mrd while Bs texture component decreased to 0.14 to 0.11 mrd. The difference can readily be attributed to the value of density function of texture components of as-cast and T6 heat treated samples after deformation. The intensity in T6 condition increased for all texture components except Bs while in as-cast condition Cu and shear 1 increased and rest of the components intensity decreased with deformation. Similar rolling texture components were reported by Deng et al. for hot rolled plate of 7050 Al alloy [41].
Benum and Nes [23], and Engler and Hirsch [26] performed review on the texture control by thermomechanical processing of Al-Mg-Si alloy sheet during cold rolling and reported that particle size larger than 1µm can stimulate at the deformation zones which interact with slip dislocations which weaken the texture formation. The studied alloy contains large eutectic Si phase in as-cast and T6 condition which also restricted the movement of dislocations resulted weak texture (Fig. 6 and Fig. 7). In addition, the T6 heat treated alloy contains more precipitates than as-cast and the flow stress in T6 condition was contributed by dislocation strengthening as well as strengthening by precipitates (Fig. 5). Therefore, the weakening effect was more prominent in T6 condition, showing the lowest intensity value of 5.60. It is also evident that the accumulation of higher dislocation density caused by the induced strain and supression of dynamic recovery during deformation (Fig. 3) which formed cell structure resulted weak texture (Fig. 6 and Fig. 7). Further, with increasing plastic deformation, the grains tend to rotate toward more stable orientations, leading to an alteration of the work hardening behavior. As plastic deformation continues, dislocations are generated and obstacle by precipitates. But, due to the presence of shareable and non-shareable precipitates, the rapid work hardening behaviour was attributed in T6 heat treated sample compared to as-cast sample, which hindered the movement of dislocation as well as rotation of the crystal. Therefore, in T6 heat treated sample was gained weak texture.
Conclusion
The influence of the precipitation state on the work hardening behavior of Al-7%Si-1%Cu-0.5%Mg alloy with additions of V, Zr and Ti in as-cast, solid solution and T6 heat treated condition has been studied through uniaxial compression deformation. The microstructural observations were made through LOM to understand the phases present before and after deformation. The initial work-hardening rate is found to increase with an increase in the yield strength and exhibiting a minimum at the SSSS condition. The slope of the θ(dθ/dσ)-( σ-σy) curves gradually decreases from the SSSS to the T6 condition. The calculated flow stress from semi empirical model vs plastic strain curve are found to be in good agreement with the experimental data. With this model, it is also possible to provide the contributions of shearable and nonshearable precipitates for the dislocation hardening during compression.
Texture evolutions in the Al-7%Si-1%Cu-0.5%Mg alloy with additions of V, Zr and Ti in as-cast, and T6 heat treated condition are investigated using XRD diffraction. The pole figures in both states, i.e., as-cast, and T6 heat treated, of the alloy show week texture. Before deformation, the orientation distribution function shows random texture. However, after deformation, shear1, {001}<110> and shear 2, {111}<110> texture components were dominant during deformation. Hence, the density function of {001}<110> texture component increased in both states, while it decreased for {111}<110> component in as-cast state and increased in T6 heat treated which influenced the work hardening.
Acknowledgements
The authors would like to acknowledge the financial support of the Advanced Structural Materials for Next Generation Vehicles (ASM-NGV) Program of Natural Resources Canada, and the Natural Sciences and Engineering Research Council of Canada (NSERC) and AUTO21 Network of Centers of Excellence for providing financial support. One of the authors (D.L. Chen) is grateful for the financial support by NSERC-DAS Award, Premier's Research Excellence Award (PREA), Canada Foundation for Innovation (CFI), and Ryerson Research Chair (RRC) program. The authors would also like to thank Q. Li, A. Machin, J. Amankrah, R. Churaman, F. Comert and P. Lele for their assistance in the experiments. The authors also thank Professor S. Bhole for the helpful discussion as well as P. Newcombe, H. Webster and D. Salesh from CANMET-MTL for the casting of the test bars.
