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Study of deformation behavior, structure and mechanical properties of the AlSiMnFe alloy during ECAP-PBP
Micron 44 (2013) 210–217
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Study of deformation behavior, structure and mechanical properties of the AlSiMnFe alloy during ECAP-PBP A.B. Naizabekov a , V.A. Andreyachshenko a , Radim Kocich b,∗ a b
Department of Material Forming, Karaganda State Industrial University, Republic Avenue 30, 101400 Temirtau, Kazakhstan Department of Material Forming, Faculty of Metallurgy and Materials Engineering, VSˇ B TU Ostrava, 17.listopadu 15, 70833 Ostrava-Poruba, Czech Republic
a r t i c l e
i n f o
Article history: Received 30 May 2012 Received in revised form 26 June 2012 Accepted 26 June 2012 Keywords: SEM TEM Substructure ECAP-PBP Mechanical properties Microhardness
a b s t r a c t The presented article deals with the effects of equal channel angular pressing (ECAP) with a newly adjusted die geometry on the microstructure and mechanical properties of the Al–Si–Mn–Fe alloy. This alloy was subjected to two modes of heat treatment followed by the ECAP process, which led to partial back pressure (ECAP-PBP). Ultra-fine grained (UFG) structure formed through ECAP-PBP process has been studied by methods of optical as well as electron microscopy. The obtained results indicate that quenched alloys, in comparison to slowly cooled alloys, do not contain large brittle particles which subsequently initiate a premature creation of cracks. It was shown that the mechanical properties of these alloys after such processing depend first and foremost on the selected type of heat treatment and on the number of performed passes. The maximum of ultimate tensile strength (417 MPa) was obtained for quenched alloy after 3 passes. On the other hand, maximum ductility was found in slowly cooled alloy after second pass. Further passes reduced strength due to the brittle behavior of excluded particles. One of the partial findings is that there is only a small dependency of the resulting size of grains on previously applied thermal processing. The minimum grain sizes were obtained after 3 passages, where their size ranged between 0.4 and 0.8 m. The application of quick cooling after heat processing due to the occurrence of finer precipitates in the matrix seems to produce better results. © 2012 Elsevier Ltd. All rights reserved.
1. Introduction The number of researchers studying the properties and preparation of materials with improved mechanical properties keeps growing. Al and Al alloys are among the most studied in this respect. One of the reasons is the fact that these alloys have great potential for automotive as well as aerospace applications due to intrinsic properties such as low density, high specific strength, and good machinability, which make them highly suitable for such applications. Interest in these particular alloys is caused also by the possibility of increasing their strength characteristics and their relatively high stability at higher temperatures. The studied materials also include alloys in the Al–Si–Mn–Fe system, as documented by published works (Gupta, 2003a, 2003b; Zuo et al., 2004). Generally, iron exists primarily as an impurity element in Al–Si alloys. The main microstructural consequence of adding iron to Al–Si foundry alloys is the formation of intermetallics (the alloy becomes brittle due to formation of large -Al5 FeSi plates or phase ␣-Al15 (Fe,Mn)3 Si2 ). The effect of iron on
∗ Corresponding author. E-mail addresses: [email protected] (A.B. Naizabekov), [email protected] (V.A. Andreyachshenko), [email protected] (R. Kocich). 0968-4328/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.micron.2012.06.011
the mechanical properties of Al–Si alloys is well documented by Couture (1981) and Mbuya et al. (2003). However various elements can act as neutralizers of the embrittling effect of iron (e.g. Cr, Co, Mo or Mn). Manganese is probably one of most frequently used “iron-suppressing” element. As Roy et al. (1996) demonstrated the addition of manganese to iron-containing Al–Si foundry alloys can be connected with drop of the porosity in alloy. The listed Al–Si–Mn–Fe alloys may thus be prepared by conventional casting, however this requires long-term annealing at a constant temperature. An alternative is the preparation of alloys by powder metallurgy (including milling and homogenization in ball mills). This process is generally followed by forming and long-term annealing. Several studies (Gupta, 2003a, 2003b; Li et al., 2000; Marker et al., 2011; Du et al., 2008) have been published describing the phase composition or crystallographic properties resulting from a selected annealing temperature, however there is a lack of information which would describe the behavior of listed alloys with respect to a wider range of contained elements. A great deal of attention has also been given to powder metallurgy. As documented by Lee et al. (2007), powder materials may be processed by the gas atomization method and cold pressing followed by forging to create intermediate products with a fine-grain structure without residual porosity even without previous sintering. Additionally, it is necessary to note
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the better resulting mechanical properties of Al–Si–Fe alloys when omitting the sintering phase. As far as ductility is concerned, it is generally known that the plasticity or super-plasticity of these alloys is significantly affected by the applied temperature of forming. For instance, Jeong et al. (2001) provide information on the maximum elongation of the tested alloy at near-melting temperatures. It also explains the reason for this behavior, where the authors claim that the cause lies in locally melted areas which provide suitable conditions for the realization of superplastic forming. Another significant effect on plastic as well as other properties is the content of Si and Fe. The specific effect of the contents of these elements is summarized in Novák et al. (2011), which closely studies the positive effects of Si on corrosive resistance. On the other hand, an increased Si content also comes with a reduced wear resistance, while leading to higher strength of the alloy (Novák et al., 2010). Silicon may however also be contained in the alloy in the form of nano-particles at the boundaries of the Al matrix, and similarly Fe may also be localized in amorphous status in the areas between grains of Si and the Al matrix (Gaidarova, 2007). Several works (Wang et al., 2010; Fang et al., 2007; Hou et al., 2010) have mentioned the possibility of improving the mechanical properties of the alloy via modification by other elements. Aside from the primary Si, the alloy also includes eutectic ␣-iron from the primary Si, ␦-iron and eutectics (Hou et al., 2010). On the other hand, an addition of SiC into the Al–Si–Fe system leads to increased strength of the materials due to its aging (Hassan et al., 2008). Aside from the effects of chemical composition and heat treatment, the properties of the alloys are also significantly affected by the method of their forming. The use of conventional forming techniques such as, e.g. rolling or forward extrusion is widespread. Partial insight is also provided by the experiment (Srivastava et al., 2007) applying forward extrusion under heat (at temperature of 480 ◦ C) to monitor the mechanical properties of cast alloys prepared by powder metallurgy. The above-listed experiment utilized spray deposited samples in the form of disks. It is clear that defining ductility with respect to specific phase compositions is one of the necessary requirements for the subsequent practical use of these alloys. It should also be noted that, aside from conventional forming techniques, it is also desirable to monitor the effects of unconventional techniques. A significant group of unconventional forming methods is represented by methods based on the application of severe plastic deformations (SPD). Although a relatively large number of published results in the area exist, the overwhelming majority of these only apply to pure Al or other Al alloys (Reihanian et al., 2008; Alhajeri et al., 2011; El-Danaf et al., 2007). The purpose of this work is to describe the deformation behavior of AlSiMnFe alloys during the application of SPD process. To this end, the equal channel angular pressing-partial back pressure (ECAP-PBP) method was chosen without the second extrusion punch. To obtain the back pressure, a recently patented method of a construction adjustment of the SPD die was used (Naizabekov et al., 2010). The effect of ECAP-PBP on the deformation behavior of the tested alloy was studied in two different modes.
2. Materials and methods 2.1. Alloy and processing The experiment material was the Al–Si–Mn–Fe alloy with following chemical composition (wt.%): Al – base, Si – 4.0%, Mn – 1.12%, Fe – 0.5%, Ca – 0.34%, Cu – 0.32%. The initial samples were defined by a square 30 mm × 30 mm, 70 mm long section. The prepared samples were subsequently divided into two groups
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Fig. 1. Geometry of the die used for ECAP-PBP process.
(samples A and samples Q). Samples A were annealed at a temperature of 550 ◦ C for 1 h and then slowly cooled in the furnace. The second group (Q) was annealed at 550 ◦ C for 1 h and cooled in water. Group A was then subjected to ECAP-PBP at room temperature, while group Q to ECAP-PBP at 200 ◦ C. The extrusion was carried out using a die whose shape is depicted in Fig. 1. It is clear that the channel in the die comprises two main parts. The first part of the channel has constant cross-section and a geometry defined by angles ϕ = 135◦ and = 40◦ . This is then followed by the second part of the channel, defined by a smaller cross-section in one direction and connected to the first part at an angle of 45◦ . It needs to be said that the whole extrusion process is based on using a larger number of dies, whereas each die linearly reduces the cross-section in one direction. This method of deformation is thus based on a similar principle as the recently proposed IE-ECAP (integrated extrusion and equal channel angular pressing) process (Orlov et al., 2011). While in IE-ECAP forward extrusion is preceded by angular extrusion, ECAP-PBP reverses the order of these phases. Additionally, the reduction of the cross-section only occurs in one direction of the material. Palm oil was selected as the lubricant. Samples were deformed by 1–3 passages through the selected deformation route C (rotation of samples between each passage by 180◦ around the sample’s axis) (Valiev and Langdon, 2006). The reduction of the cross-section in the specific direction is 1 mm from each side. The dimensions of the cross-section in the orthogonal direction remain constant during the whole process. The scheme (Fig. 1) indicates that this process differs from the classical ECAP process. The used configuration allows the deformation area to be spread out across three zones. The material undergoes the shear strain in zone I (channel bend). On the other hand, zone II is characteristic by a back pressure created from the reduced channel cross-section. The third zone is then the area between both parts of the channel, where the main deformation will be mostly generated in the peripheral layers of the processed sample (i.e. there will be a reduction of the sample’s cross-section in one direction). This means that during the whole process (in all 3 zones) the extruded material is affected mostly by the pressure stress caused by the die geometry. It is generally known that compressive stresses support the ductility of the material, which may predict a favorable course of the experiment. All the evaluated samples came from the central areas of the formed samples (elimination of external layers of samples). Optical microscopy (OM) was executed with LEICA-DMIRM machine equipped with a microhardness indenter Anton Paar MHT10. This
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Fig. 2. Microstructure of the alloy after heat treatment: sample A (a) and sample Q (b).
method was used to evaluate especially the primary particles and the eutectic phase. Scanning electron microscopy (SEM) on a JEOL JSM 5910 microscope was also used, for instance for the evaluation of fracture areas after the performed tension tests. The transmission electron microscope (TEM) JEOL JEM 2100 was used to monitor the substructural changes during heat treatment and subsequent extrusion. Mechanical properties were evaluated via tension tests carried out on MI40KU tensile machine. The cross-head speed of the tension machine was 0.5 mm/min. This value corresponds with the value of strain rate 0.56 × 10−3 s−1 .
3. Results 3.1. OM analysis Due to the content of elements in the alloys, which is greater than their maximum dissolubility in a solid solution, these elements are present in the alloy as particles or in the form of excluded phases with other elements. The created eutectics are mostly defined as Al–Si, whereas it is known that these contain ∼1.65 wt.% Si at a temperature of around 577 ◦ C (Murray and McAlister, 1984). The attached images indicate that the sample Q (Fig. 2b) had a larger amount of elements dissolved during thermal processing than sample A (Fig. 2a). It is clear that the microstructure of sample A indicates the presence of relatively unequally spread coagulated particles with sizes ranging between 2 and 10 m. These particles occur in the form of clusters, which however break down during continued deformation and also gradually become more homogenously spread in
the matrix. The attached images show the borders of grains of the solid solution, the size of which is 150–200 m in the length and 30–50 m orthogonally. Additionally, the particles of secondary phases distributed across the whole cross-section of the sample are also visible. Sample Q (Fig. 2b), on the other hand, does not display any clear and sharp grain edges. It includes relatively homogenously distributed particles of sizes below 1 m. A certain difference is clear regarding the occurrence of groups of 3–5 m particles, which are not homogenously spread in the matrix. The situation after the performed ECAP-PBP (3 passes) is depicted in Fig. 3, which shows both groups of samples (A and Q). It is clear that during ECAP no significant changes of the size or shape of large particles in the structure have occurred. On the contrary, smaller particles are more clustered together. Additionally from the above-listed, the total number of particles has increased, confirming the process of dynamic aging caused by the plastic deformation. Particles are present inside grains as well as on their boundaries. The alloy in status Q differs especially by the fact that during ECAP-PBP the number of fine particles increases; this is due to annealing followed by the subsequent hardening. On the other hand, larger particles do not undergo changes in their size or distribution during the deformation. 3.2. TEM analysis Individual changes in the structure were monitored via transmission microscopy. Images were taken after heat treatment and after ECAP-PBP to allow comparison. Samples after three passes were processed.
Fig. 3. Microstructure of the alloy after ECAP-PBP: sample A (a) and sample Q (b).
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Fig. 4. TEM image of AlSiMnFe alloy after ECAP-PBP: (a) after 1 pass (bright field) and (b) after 2 passes (dark field).
The application of a single pass meant that the original grains became finer, as documented by Fig. 4a. However these created grains are not yet equiaxed, and the grains are clearly elongated in the direction of extrusion. As is visible, grains sized 1–2 m (in the orthogonal direction to the direction of extrusion) and 3–5 m (in the direction of extrusion) are created after the first pass. Excluded particles are distributed unevenly in the matrix and it is clear that they are not ordered in larger clusters. A large number of present particles contribute to the grains being finer, whereas these may often serve as nuclei for the creation of new grains (Lukáˇc et al., 2007). Fig. 4b shows one particular particle inside the area of the boundary of two adjacent grains, which has very probably played a role in making the structure finer. On the other hand, the second pass leads to the creation of finer grains due to shear strain – the grains range within 0.8–1.2 m. After two passes the
structure of grains is already homogenous. It should be noted that the second pass also disintegrated the original elongated grains which were visible in the structure after the first pass and turned of some grains. Generally, it may be said that after the second pass we can see the creation of a homogenous equiaxed ultra-fine substructure with certain grains smaller than 0.5 m. The third pass led to the re-creation of grains with a characteristic elongation in the direction of extrusion. Their size ranged between 0.4 and 0.8 m. This finding is very probably related to the reverse nature of the deformation when using deformation path C (Valiev and Langdon, 2006). The increase of the imposed strain after 3 passages also meant an increased activity of dislocations, where the images clearly display tangles of dislocations in some grains. The refinement of grains in this passage is less intense than in the first two passes. It should however be mentioned that this deformation leads
Fig. 5. Mechanical properties of the AlSiMnFe alloy after heat treatment and ECAP-PBP: (a) ultimate tensile strength, (b) elongation to failure and (c) microhardness.
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Fig. 6. SEM images of fracture areas: (a) sample A before ECAP-PBP, (b) sample Q before ECAP-PBP, (c) sample A after ECAP-PBP (3 passes), and (d) sample Q after ECAP-PBP (3 passes).
to a larger amount of small low-angle particle fragments. Additionally, there is also a clear influence of the imposed strain on the size of particles, and this may be considered one of the causes for the creation of microcracks. 3.3. Mechanical properties Tensile tests were carried out with the specimens whose shape is shown in Fig. 5a. Dimensions of the working part of the sample were L = 15 mm and Ø d = 3 mm. Fig. 5a also demonstrates the course of the ultimate tensile strength (UTS) for samples after heat treatment and subsequent ECAP-PBP for both groups of samples. It is clear that samples A exhibit a significant increase of UTS already after the first pass. However, the strength drops after the second pass and the third passage does not significantly affect the UTS of ECAPed materials. The ductility (elongation to failure) of materials also drops after the first pass (Fig. 5b). However, the alloy’s ductility increases significantly after the second pass. Although the UTS of the alloy after the third performed passage is almost identical as after the second passage, its ductility drops drastically (∼50%). In contrast to group A, samples in group Q demonstrate a reduction of UTS after the first pass, however subsequent passes increase UTS. The maximum ductility values are obtained by the alloy after the second pass. This behavior may be explained by changes in the microstructure, redistribution of particles and dynamic aging. In relation to the abovementioned, it is also necessary to mention the probable process of restoration processes (recovery) which cause the redistribution of dislocations. The size and density of precipitated particles significantly affects the process of these softening processes, and these may be used as efficient barriers against the occurrence of softening (Humphreys and Hatherly, 2002).
The measured microhardness values have also been evaluated to supplement the results (Fig. 5c). It is clear that during ECAPPBP, each passage increases the microhardness. Higher values were found for the slowly cooled samples (A). This means that the strengthening effect in this alloy was higher when cooled slowly in the furnace than when quenched in water (Q). 3.4. SEM analysis The analysis of fracture areas targeted the creation of cracks, failures and their subsequent propagation. Images of fractures are presented in Fig. 6, where it becomes clear that fracture areas exhibit the plastic type of cracks. It is possible to notice common properties between images from various states. A specific example could be the presence of the same phases after heat treatment (Fig. 6a and b) and after the following deformation (ECAP-PBP) (Fig. 6c and d). This means that inclusions after the performed heat treatment detected by OM in the microstructure are also visible on fracture areas even after ECAP-PBP. The attached images indicate that cracks and their subsequent propagation occur via two mechanisms. The first is along larger particles, and the second is their propagation around smaller particles. For samples A, we confirmed the occurrence of globular particles with sizes ranging within 5–15 m with a failure propagating around them. Additionally, the alloy contained larger particles (30–40 m) whereas the propagation of failure can be described by brittle fracture. On the other hand, samples Q clearly contained smaller particles on fracture areas, although the occurrence of particles remains bimodal. Globular particles sized 5–10 m together with particles sized up to 20 m also affected the fracture, similarly as in the previous case.
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Fig. 7. Schematic diagram of the initiation of cracks during ECAP-PBP processing.
The conducted thermal processing was to ensure the possibility of controlling the amount and size of particles distributed in the structure. As is clear from the listed results, the selected type of annealing at 550 ◦ C for 1 h provided suitable conditions for the exclusion of elements from the alloy as individual particles. Quick cooling after annealing meant a reduction of the size of excluded particles as well as the distance between them. Simultaneously it also led to the occurrence of a large number of heterogeneously distributed fine particles. It is clear that the mechanical properties of AlFeSiMn alloys are affected not only by the contents of elements in the solid solution but also by plastic deformation.
meant the occurrence of large (10 m) as well as small particles, homogenously distributed across the cross-section of the sample. The solid solution is no longer oversaturated after the performed annealing, which already leads to conditions for the creation of the aforementioned areas after the first pass. It is clear that during the first passages the plastic deformation is first applied in the solid solution, which is accompanied by the movement of large particles. After two and more passes the generated stresses in the metal also lead to deformation within individual particles, which leads to their fragmentation and a reduction of their mechanical properties. Additionally, the stresses led to the diffusion of atoms of elements in the solid solution to particles, leading to their increase. This process is sketched in Fig. 7.
4.1. Influence of heat treatment on microstructure
4.2. Development of mechanical properties
Both cases (A and Q) exhibited a significant refinement of grains after performing ECAP-PBP. The final size of grains only differed slightly between both selected procedures. However, the distribution and size of particles present in the structure depended first and foremost on the selected type of heat treatment. As was shown, an oversaturated solid solution exhibits signs of strain aging after applied plastic deformations. This means that the first pass already leads to the exclusion of soft particles in the whole volume of the sample. This exclusion however occurs practically without any mutual interaction between such particles. This is also confirmed by the small number of large particles occurring in the structure during the whole deformation cycle. During subsequent passes, the small particles localize to isolated areas due to better thermodynamic conditions. The concentration of these particles in such areas is higher than the concentration in the solid solution adjacent to them. This increases their density during plastic deformation due to the influence of stresses during the movement of material layers. A growing number of cycles make the movement of layers centralized. Heat treatment followed by slow cooling in the furnace
After 4 and more passes, the deformed samples lose their cohesion. As follows from microstructural observations, it is clear that this loss is caused by the small cohesion of the created particles. The hardness of particles and secondary phases also determines the resulting mechanical properties of the alloy. This is also related to the strengthening effect of the deformed alloy. The presence of a large number of particles leads to the intense inhibition of the movement of dislocations during deformation in ECAP-PBP, which increases hardness and reduces the ductility of the alloy. The acting stresses during process led to the redistribution of particles, whereas fine particles are separated into individual regions. This leads to the localization of strain primarily into areas with lower concentrations of these particles. The result is a reduction of the hardness and increase of the ductility of alloys. The back pressure factor needs to be noted, as documented also by published works (Mckenzie and Lapovok, 2010), as this positively affects the deformation behavior of Al alloys. The experiment confirmed this positive effect is also capable of exhibiting partial back pressure in the proposed die. This may be documented by a higher number
4. Discussion
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of passages in comparison to using a classical die with an identical curvature angle of the channel. However, a higher number of passes mean that coarser particles cannot resist the stresses and cracks begin to occur. This finding also affects ductility. For alloys after HT and subsequent quenching in water, the high stress lead to massive exclusion of particles from the solid solution, reducing the resulting ductility of the alloy. The subsequent ECAP-PBP leads to a manifestation of the strain aging of the material, which causes increased hardness while maintaining a relatively satisfactory level of ductility. Although samples Q also exhibited brittle fracture of large particles, this does not negatively affect its mechanical properties as opposed to the case of samples A. It should be mentioned that microhardness grew relatively linearly with the number of passages for both groups A and Q. This finding may be explained by an increased occurrence of particles in the structure and also by the decreasing distance between them. Although both groups of samples exhibit a similar trend of increasing microhardness, certain differences do exist. Samples A had higher microhardness values than group Q. The probable cause for this is the difference in the processing of both groups. Samples A were air-cooled and deformed at room temperature, which led to the creation of a certain volume of particles due to HT and further particles were then created during the deformation itself, whereas the increased density of dislocations led to an increase of hardness. The microhardness of group Q increased especially due to the generation of particles caused by imposed strain. For group Q, which was quenched in water and subsequently deformed at a temperature of 200 ◦ C, this has led to suitable conditions for a more gradual increase of hardness. 5. Conclusions The work describes an experiment to determine the deformation behavior of alloys in the AlSiFeMn system during the newly designed ECAP-PBP process. Another goal of the work was to describe the behavior of this alloy after selected modes of heat treatment. The main conclusions from the conducted experiment are: (1) An ECAP-PBP is a useful technology that leads to a significant refinement of grains to the submicrometer range (down to 0.5–0.8 m) in the alloy, both when applied at room temperature and at 200 ◦ C. (2) A relatively equiaxed substructure with an average size of grains ranging in 0.8–1.2 m was already detected after the two performed passes. (3) The strain imposed after heat treatment and subsequent quenching leads to a precipitation of small particles from the oversaturated solid solution. Their subsequent clustering increases the mechanical properties of the alloy. (4) During ECAP-PBP we first see an increase of the strength properties followed by their decrease caused by the brittleness of excluded particles. Maximum ultimate tensile strength (417 MPa) was obtained after the third passage. (5) Heat treatment followed by material quenching seems to be a very suitable method before the deformation of AlSiFeMn alloys, especially due to the possibility of obtaining fine particles. Both tested cases exhibited a gradual increase in the microhardness of the deformed alloy with a growing number of passages. (6) The fracture areas demonstrated the effects of the presence of coarse particles as well as localized groups of fine particles, which played a major role in the creation and propagation of cracks.
Acknowledgements This paper was created under the Programme 120 “Grant financing the scientific studies” Committee of the science project no. 1099 financed by Ministry of Education and science of the Republic of Kazkhstan and under the project no. CZ.1.05/2.1.00/01.0040 “Regional Materials Science and Technology Centre” within the frame of the operation program “Research and Development for Innovations” financed by the Structural Funds and from the state budget of the Czech Republic. ˇ Dr. P. Jonsta (VSB-TU Ostrava, Czech Republic) is greatly appreciated for his help in the mechanical properties measurements.
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Study of deformation behavior, structure and mechanical properties of the AlSiMnFe alloy during ECAP-PBP A.B. Naizabekov a , V.A. Andreyachshenko a , Radim Kocich b,∗ a b
Department of Material Forming, Karaganda State Industrial University, Republic Avenue 30, 101400 Temirtau, Kazakhstan Department of Material Forming, Faculty of Metallurgy and Materials Engineering, VSˇ B TU Ostrava, 17.listopadu 15, 70833 Ostrava-Poruba, Czech Republic
a r t i c l e
i n f o
Article history: Received 30 May 2012 Received in revised form 26 June 2012 Accepted 26 June 2012 Keywords: SEM TEM Substructure ECAP-PBP Mechanical properties Microhardness
a b s t r a c t The presented article deals with the effects of equal channel angular pressing (ECAP) with a newly adjusted die geometry on the microstructure and mechanical properties of the Al–Si–Mn–Fe alloy. This alloy was subjected to two modes of heat treatment followed by the ECAP process, which led to partial back pressure (ECAP-PBP). Ultra-fine grained (UFG) structure formed through ECAP-PBP process has been studied by methods of optical as well as electron microscopy. The obtained results indicate that quenched alloys, in comparison to slowly cooled alloys, do not contain large brittle particles which subsequently initiate a premature creation of cracks. It was shown that the mechanical properties of these alloys after such processing depend first and foremost on the selected type of heat treatment and on the number of performed passes. The maximum of ultimate tensile strength (417 MPa) was obtained for quenched alloy after 3 passes. On the other hand, maximum ductility was found in slowly cooled alloy after second pass. Further passes reduced strength due to the brittle behavior of excluded particles. One of the partial findings is that there is only a small dependency of the resulting size of grains on previously applied thermal processing. The minimum grain sizes were obtained after 3 passages, where their size ranged between 0.4 and 0.8 m. The application of quick cooling after heat processing due to the occurrence of finer precipitates in the matrix seems to produce better results. © 2012 Elsevier Ltd. All rights reserved.
1. Introduction The number of researchers studying the properties and preparation of materials with improved mechanical properties keeps growing. Al and Al alloys are among the most studied in this respect. One of the reasons is the fact that these alloys have great potential for automotive as well as aerospace applications due to intrinsic properties such as low density, high specific strength, and good machinability, which make them highly suitable for such applications. Interest in these particular alloys is caused also by the possibility of increasing their strength characteristics and their relatively high stability at higher temperatures. The studied materials also include alloys in the Al–Si–Mn–Fe system, as documented by published works (Gupta, 2003a, 2003b; Zuo et al., 2004). Generally, iron exists primarily as an impurity element in Al–Si alloys. The main microstructural consequence of adding iron to Al–Si foundry alloys is the formation of intermetallics (the alloy becomes brittle due to formation of large -Al5 FeSi plates or phase ␣-Al15 (Fe,Mn)3 Si2 ). The effect of iron on
∗ Corresponding author. E-mail addresses: [email protected] (A.B. Naizabekov), [email protected] (V.A. Andreyachshenko), [email protected] (R. Kocich). 0968-4328/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.micron.2012.06.011
the mechanical properties of Al–Si alloys is well documented by Couture (1981) and Mbuya et al. (2003). However various elements can act as neutralizers of the embrittling effect of iron (e.g. Cr, Co, Mo or Mn). Manganese is probably one of most frequently used “iron-suppressing” element. As Roy et al. (1996) demonstrated the addition of manganese to iron-containing Al–Si foundry alloys can be connected with drop of the porosity in alloy. The listed Al–Si–Mn–Fe alloys may thus be prepared by conventional casting, however this requires long-term annealing at a constant temperature. An alternative is the preparation of alloys by powder metallurgy (including milling and homogenization in ball mills). This process is generally followed by forming and long-term annealing. Several studies (Gupta, 2003a, 2003b; Li et al., 2000; Marker et al., 2011; Du et al., 2008) have been published describing the phase composition or crystallographic properties resulting from a selected annealing temperature, however there is a lack of information which would describe the behavior of listed alloys with respect to a wider range of contained elements. A great deal of attention has also been given to powder metallurgy. As documented by Lee et al. (2007), powder materials may be processed by the gas atomization method and cold pressing followed by forging to create intermediate products with a fine-grain structure without residual porosity even without previous sintering. Additionally, it is necessary to note
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the better resulting mechanical properties of Al–Si–Fe alloys when omitting the sintering phase. As far as ductility is concerned, it is generally known that the plasticity or super-plasticity of these alloys is significantly affected by the applied temperature of forming. For instance, Jeong et al. (2001) provide information on the maximum elongation of the tested alloy at near-melting temperatures. It also explains the reason for this behavior, where the authors claim that the cause lies in locally melted areas which provide suitable conditions for the realization of superplastic forming. Another significant effect on plastic as well as other properties is the content of Si and Fe. The specific effect of the contents of these elements is summarized in Novák et al. (2011), which closely studies the positive effects of Si on corrosive resistance. On the other hand, an increased Si content also comes with a reduced wear resistance, while leading to higher strength of the alloy (Novák et al., 2010). Silicon may however also be contained in the alloy in the form of nano-particles at the boundaries of the Al matrix, and similarly Fe may also be localized in amorphous status in the areas between grains of Si and the Al matrix (Gaidarova, 2007). Several works (Wang et al., 2010; Fang et al., 2007; Hou et al., 2010) have mentioned the possibility of improving the mechanical properties of the alloy via modification by other elements. Aside from the primary Si, the alloy also includes eutectic ␣-iron from the primary Si, ␦-iron and eutectics (Hou et al., 2010). On the other hand, an addition of SiC into the Al–Si–Fe system leads to increased strength of the materials due to its aging (Hassan et al., 2008). Aside from the effects of chemical composition and heat treatment, the properties of the alloys are also significantly affected by the method of their forming. The use of conventional forming techniques such as, e.g. rolling or forward extrusion is widespread. Partial insight is also provided by the experiment (Srivastava et al., 2007) applying forward extrusion under heat (at temperature of 480 ◦ C) to monitor the mechanical properties of cast alloys prepared by powder metallurgy. The above-listed experiment utilized spray deposited samples in the form of disks. It is clear that defining ductility with respect to specific phase compositions is one of the necessary requirements for the subsequent practical use of these alloys. It should also be noted that, aside from conventional forming techniques, it is also desirable to monitor the effects of unconventional techniques. A significant group of unconventional forming methods is represented by methods based on the application of severe plastic deformations (SPD). Although a relatively large number of published results in the area exist, the overwhelming majority of these only apply to pure Al or other Al alloys (Reihanian et al., 2008; Alhajeri et al., 2011; El-Danaf et al., 2007). The purpose of this work is to describe the deformation behavior of AlSiMnFe alloys during the application of SPD process. To this end, the equal channel angular pressing-partial back pressure (ECAP-PBP) method was chosen without the second extrusion punch. To obtain the back pressure, a recently patented method of a construction adjustment of the SPD die was used (Naizabekov et al., 2010). The effect of ECAP-PBP on the deformation behavior of the tested alloy was studied in two different modes.
2. Materials and methods 2.1. Alloy and processing The experiment material was the Al–Si–Mn–Fe alloy with following chemical composition (wt.%): Al – base, Si – 4.0%, Mn – 1.12%, Fe – 0.5%, Ca – 0.34%, Cu – 0.32%. The initial samples were defined by a square 30 mm × 30 mm, 70 mm long section. The prepared samples were subsequently divided into two groups
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Fig. 1. Geometry of the die used for ECAP-PBP process.
(samples A and samples Q). Samples A were annealed at a temperature of 550 ◦ C for 1 h and then slowly cooled in the furnace. The second group (Q) was annealed at 550 ◦ C for 1 h and cooled in water. Group A was then subjected to ECAP-PBP at room temperature, while group Q to ECAP-PBP at 200 ◦ C. The extrusion was carried out using a die whose shape is depicted in Fig. 1. It is clear that the channel in the die comprises two main parts. The first part of the channel has constant cross-section and a geometry defined by angles ϕ = 135◦ and = 40◦ . This is then followed by the second part of the channel, defined by a smaller cross-section in one direction and connected to the first part at an angle of 45◦ . It needs to be said that the whole extrusion process is based on using a larger number of dies, whereas each die linearly reduces the cross-section in one direction. This method of deformation is thus based on a similar principle as the recently proposed IE-ECAP (integrated extrusion and equal channel angular pressing) process (Orlov et al., 2011). While in IE-ECAP forward extrusion is preceded by angular extrusion, ECAP-PBP reverses the order of these phases. Additionally, the reduction of the cross-section only occurs in one direction of the material. Palm oil was selected as the lubricant. Samples were deformed by 1–3 passages through the selected deformation route C (rotation of samples between each passage by 180◦ around the sample’s axis) (Valiev and Langdon, 2006). The reduction of the cross-section in the specific direction is 1 mm from each side. The dimensions of the cross-section in the orthogonal direction remain constant during the whole process. The scheme (Fig. 1) indicates that this process differs from the classical ECAP process. The used configuration allows the deformation area to be spread out across three zones. The material undergoes the shear strain in zone I (channel bend). On the other hand, zone II is characteristic by a back pressure created from the reduced channel cross-section. The third zone is then the area between both parts of the channel, where the main deformation will be mostly generated in the peripheral layers of the processed sample (i.e. there will be a reduction of the sample’s cross-section in one direction). This means that during the whole process (in all 3 zones) the extruded material is affected mostly by the pressure stress caused by the die geometry. It is generally known that compressive stresses support the ductility of the material, which may predict a favorable course of the experiment. All the evaluated samples came from the central areas of the formed samples (elimination of external layers of samples). Optical microscopy (OM) was executed with LEICA-DMIRM machine equipped with a microhardness indenter Anton Paar MHT10. This
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Fig. 2. Microstructure of the alloy after heat treatment: sample A (a) and sample Q (b).
method was used to evaluate especially the primary particles and the eutectic phase. Scanning electron microscopy (SEM) on a JEOL JSM 5910 microscope was also used, for instance for the evaluation of fracture areas after the performed tension tests. The transmission electron microscope (TEM) JEOL JEM 2100 was used to monitor the substructural changes during heat treatment and subsequent extrusion. Mechanical properties were evaluated via tension tests carried out on MI40KU tensile machine. The cross-head speed of the tension machine was 0.5 mm/min. This value corresponds with the value of strain rate 0.56 × 10−3 s−1 .
3. Results 3.1. OM analysis Due to the content of elements in the alloys, which is greater than their maximum dissolubility in a solid solution, these elements are present in the alloy as particles or in the form of excluded phases with other elements. The created eutectics are mostly defined as Al–Si, whereas it is known that these contain ∼1.65 wt.% Si at a temperature of around 577 ◦ C (Murray and McAlister, 1984). The attached images indicate that the sample Q (Fig. 2b) had a larger amount of elements dissolved during thermal processing than sample A (Fig. 2a). It is clear that the microstructure of sample A indicates the presence of relatively unequally spread coagulated particles with sizes ranging between 2 and 10 m. These particles occur in the form of clusters, which however break down during continued deformation and also gradually become more homogenously spread in
the matrix. The attached images show the borders of grains of the solid solution, the size of which is 150–200 m in the length and 30–50 m orthogonally. Additionally, the particles of secondary phases distributed across the whole cross-section of the sample are also visible. Sample Q (Fig. 2b), on the other hand, does not display any clear and sharp grain edges. It includes relatively homogenously distributed particles of sizes below 1 m. A certain difference is clear regarding the occurrence of groups of 3–5 m particles, which are not homogenously spread in the matrix. The situation after the performed ECAP-PBP (3 passes) is depicted in Fig. 3, which shows both groups of samples (A and Q). It is clear that during ECAP no significant changes of the size or shape of large particles in the structure have occurred. On the contrary, smaller particles are more clustered together. Additionally from the above-listed, the total number of particles has increased, confirming the process of dynamic aging caused by the plastic deformation. Particles are present inside grains as well as on their boundaries. The alloy in status Q differs especially by the fact that during ECAP-PBP the number of fine particles increases; this is due to annealing followed by the subsequent hardening. On the other hand, larger particles do not undergo changes in their size or distribution during the deformation. 3.2. TEM analysis Individual changes in the structure were monitored via transmission microscopy. Images were taken after heat treatment and after ECAP-PBP to allow comparison. Samples after three passes were processed.
Fig. 3. Microstructure of the alloy after ECAP-PBP: sample A (a) and sample Q (b).
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Fig. 4. TEM image of AlSiMnFe alloy after ECAP-PBP: (a) after 1 pass (bright field) and (b) after 2 passes (dark field).
The application of a single pass meant that the original grains became finer, as documented by Fig. 4a. However these created grains are not yet equiaxed, and the grains are clearly elongated in the direction of extrusion. As is visible, grains sized 1–2 m (in the orthogonal direction to the direction of extrusion) and 3–5 m (in the direction of extrusion) are created after the first pass. Excluded particles are distributed unevenly in the matrix and it is clear that they are not ordered in larger clusters. A large number of present particles contribute to the grains being finer, whereas these may often serve as nuclei for the creation of new grains (Lukáˇc et al., 2007). Fig. 4b shows one particular particle inside the area of the boundary of two adjacent grains, which has very probably played a role in making the structure finer. On the other hand, the second pass leads to the creation of finer grains due to shear strain – the grains range within 0.8–1.2 m. After two passes the
structure of grains is already homogenous. It should be noted that the second pass also disintegrated the original elongated grains which were visible in the structure after the first pass and turned of some grains. Generally, it may be said that after the second pass we can see the creation of a homogenous equiaxed ultra-fine substructure with certain grains smaller than 0.5 m. The third pass led to the re-creation of grains with a characteristic elongation in the direction of extrusion. Their size ranged between 0.4 and 0.8 m. This finding is very probably related to the reverse nature of the deformation when using deformation path C (Valiev and Langdon, 2006). The increase of the imposed strain after 3 passages also meant an increased activity of dislocations, where the images clearly display tangles of dislocations in some grains. The refinement of grains in this passage is less intense than in the first two passes. It should however be mentioned that this deformation leads
Fig. 5. Mechanical properties of the AlSiMnFe alloy after heat treatment and ECAP-PBP: (a) ultimate tensile strength, (b) elongation to failure and (c) microhardness.
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Fig. 6. SEM images of fracture areas: (a) sample A before ECAP-PBP, (b) sample Q before ECAP-PBP, (c) sample A after ECAP-PBP (3 passes), and (d) sample Q after ECAP-PBP (3 passes).
to a larger amount of small low-angle particle fragments. Additionally, there is also a clear influence of the imposed strain on the size of particles, and this may be considered one of the causes for the creation of microcracks. 3.3. Mechanical properties Tensile tests were carried out with the specimens whose shape is shown in Fig. 5a. Dimensions of the working part of the sample were L = 15 mm and Ø d = 3 mm. Fig. 5a also demonstrates the course of the ultimate tensile strength (UTS) for samples after heat treatment and subsequent ECAP-PBP for both groups of samples. It is clear that samples A exhibit a significant increase of UTS already after the first pass. However, the strength drops after the second pass and the third passage does not significantly affect the UTS of ECAPed materials. The ductility (elongation to failure) of materials also drops after the first pass (Fig. 5b). However, the alloy’s ductility increases significantly after the second pass. Although the UTS of the alloy after the third performed passage is almost identical as after the second passage, its ductility drops drastically (∼50%). In contrast to group A, samples in group Q demonstrate a reduction of UTS after the first pass, however subsequent passes increase UTS. The maximum ductility values are obtained by the alloy after the second pass. This behavior may be explained by changes in the microstructure, redistribution of particles and dynamic aging. In relation to the abovementioned, it is also necessary to mention the probable process of restoration processes (recovery) which cause the redistribution of dislocations. The size and density of precipitated particles significantly affects the process of these softening processes, and these may be used as efficient barriers against the occurrence of softening (Humphreys and Hatherly, 2002).
The measured microhardness values have also been evaluated to supplement the results (Fig. 5c). It is clear that during ECAPPBP, each passage increases the microhardness. Higher values were found for the slowly cooled samples (A). This means that the strengthening effect in this alloy was higher when cooled slowly in the furnace than when quenched in water (Q). 3.4. SEM analysis The analysis of fracture areas targeted the creation of cracks, failures and their subsequent propagation. Images of fractures are presented in Fig. 6, where it becomes clear that fracture areas exhibit the plastic type of cracks. It is possible to notice common properties between images from various states. A specific example could be the presence of the same phases after heat treatment (Fig. 6a and b) and after the following deformation (ECAP-PBP) (Fig. 6c and d). This means that inclusions after the performed heat treatment detected by OM in the microstructure are also visible on fracture areas even after ECAP-PBP. The attached images indicate that cracks and their subsequent propagation occur via two mechanisms. The first is along larger particles, and the second is their propagation around smaller particles. For samples A, we confirmed the occurrence of globular particles with sizes ranging within 5–15 m with a failure propagating around them. Additionally, the alloy contained larger particles (30–40 m) whereas the propagation of failure can be described by brittle fracture. On the other hand, samples Q clearly contained smaller particles on fracture areas, although the occurrence of particles remains bimodal. Globular particles sized 5–10 m together with particles sized up to 20 m also affected the fracture, similarly as in the previous case.
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Fig. 7. Schematic diagram of the initiation of cracks during ECAP-PBP processing.
The conducted thermal processing was to ensure the possibility of controlling the amount and size of particles distributed in the structure. As is clear from the listed results, the selected type of annealing at 550 ◦ C for 1 h provided suitable conditions for the exclusion of elements from the alloy as individual particles. Quick cooling after annealing meant a reduction of the size of excluded particles as well as the distance between them. Simultaneously it also led to the occurrence of a large number of heterogeneously distributed fine particles. It is clear that the mechanical properties of AlFeSiMn alloys are affected not only by the contents of elements in the solid solution but also by plastic deformation.
meant the occurrence of large (10 m) as well as small particles, homogenously distributed across the cross-section of the sample. The solid solution is no longer oversaturated after the performed annealing, which already leads to conditions for the creation of the aforementioned areas after the first pass. It is clear that during the first passages the plastic deformation is first applied in the solid solution, which is accompanied by the movement of large particles. After two and more passes the generated stresses in the metal also lead to deformation within individual particles, which leads to their fragmentation and a reduction of their mechanical properties. Additionally, the stresses led to the diffusion of atoms of elements in the solid solution to particles, leading to their increase. This process is sketched in Fig. 7.
4.1. Influence of heat treatment on microstructure
4.2. Development of mechanical properties
Both cases (A and Q) exhibited a significant refinement of grains after performing ECAP-PBP. The final size of grains only differed slightly between both selected procedures. However, the distribution and size of particles present in the structure depended first and foremost on the selected type of heat treatment. As was shown, an oversaturated solid solution exhibits signs of strain aging after applied plastic deformations. This means that the first pass already leads to the exclusion of soft particles in the whole volume of the sample. This exclusion however occurs practically without any mutual interaction between such particles. This is also confirmed by the small number of large particles occurring in the structure during the whole deformation cycle. During subsequent passes, the small particles localize to isolated areas due to better thermodynamic conditions. The concentration of these particles in such areas is higher than the concentration in the solid solution adjacent to them. This increases their density during plastic deformation due to the influence of stresses during the movement of material layers. A growing number of cycles make the movement of layers centralized. Heat treatment followed by slow cooling in the furnace
After 4 and more passes, the deformed samples lose their cohesion. As follows from microstructural observations, it is clear that this loss is caused by the small cohesion of the created particles. The hardness of particles and secondary phases also determines the resulting mechanical properties of the alloy. This is also related to the strengthening effect of the deformed alloy. The presence of a large number of particles leads to the intense inhibition of the movement of dislocations during deformation in ECAP-PBP, which increases hardness and reduces the ductility of the alloy. The acting stresses during process led to the redistribution of particles, whereas fine particles are separated into individual regions. This leads to the localization of strain primarily into areas with lower concentrations of these particles. The result is a reduction of the hardness and increase of the ductility of alloys. The back pressure factor needs to be noted, as documented also by published works (Mckenzie and Lapovok, 2010), as this positively affects the deformation behavior of Al alloys. The experiment confirmed this positive effect is also capable of exhibiting partial back pressure in the proposed die. This may be documented by a higher number
4. Discussion
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of passages in comparison to using a classical die with an identical curvature angle of the channel. However, a higher number of passes mean that coarser particles cannot resist the stresses and cracks begin to occur. This finding also affects ductility. For alloys after HT and subsequent quenching in water, the high stress lead to massive exclusion of particles from the solid solution, reducing the resulting ductility of the alloy. The subsequent ECAP-PBP leads to a manifestation of the strain aging of the material, which causes increased hardness while maintaining a relatively satisfactory level of ductility. Although samples Q also exhibited brittle fracture of large particles, this does not negatively affect its mechanical properties as opposed to the case of samples A. It should be mentioned that microhardness grew relatively linearly with the number of passages for both groups A and Q. This finding may be explained by an increased occurrence of particles in the structure and also by the decreasing distance between them. Although both groups of samples exhibit a similar trend of increasing microhardness, certain differences do exist. Samples A had higher microhardness values than group Q. The probable cause for this is the difference in the processing of both groups. Samples A were air-cooled and deformed at room temperature, which led to the creation of a certain volume of particles due to HT and further particles were then created during the deformation itself, whereas the increased density of dislocations led to an increase of hardness. The microhardness of group Q increased especially due to the generation of particles caused by imposed strain. For group Q, which was quenched in water and subsequently deformed at a temperature of 200 ◦ C, this has led to suitable conditions for a more gradual increase of hardness. 5. Conclusions The work describes an experiment to determine the deformation behavior of alloys in the AlSiFeMn system during the newly designed ECAP-PBP process. Another goal of the work was to describe the behavior of this alloy after selected modes of heat treatment. The main conclusions from the conducted experiment are: (1) An ECAP-PBP is a useful technology that leads to a significant refinement of grains to the submicrometer range (down to 0.5–0.8 m) in the alloy, both when applied at room temperature and at 200 ◦ C. (2) A relatively equiaxed substructure with an average size of grains ranging in 0.8–1.2 m was already detected after the two performed passes. (3) The strain imposed after heat treatment and subsequent quenching leads to a precipitation of small particles from the oversaturated solid solution. Their subsequent clustering increases the mechanical properties of the alloy. (4) During ECAP-PBP we first see an increase of the strength properties followed by their decrease caused by the brittleness of excluded particles. Maximum ultimate tensile strength (417 MPa) was obtained after the third passage. (5) Heat treatment followed by material quenching seems to be a very suitable method before the deformation of AlSiFeMn alloys, especially due to the possibility of obtaining fine particles. Both tested cases exhibited a gradual increase in the microhardness of the deformed alloy with a growing number of passages. (6) The fracture areas demonstrated the effects of the presence of coarse particles as well as localized groups of fine particles, which played a major role in the creation and propagation of cracks.
Acknowledgements This paper was created under the Programme 120 “Grant financing the scientific studies” Committee of the science project no. 1099 financed by Ministry of Education and science of the Republic of Kazkhstan and under the project no. CZ.1.05/2.1.00/01.0040 “Regional Materials Science and Technology Centre” within the frame of the operation program “Research and Development for Innovations” financed by the Structural Funds and from the state budget of the Czech Republic. ˇ Dr. P. Jonsta (VSB-TU Ostrava, Czech Republic) is greatly appreciated for his help in the mechanical properties measurements.
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