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AZ91 composite sheets
Materials Science & Engineering A 763 (2019) 138118
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The microstructure, mechanical properties and tensile deformation mechanism of rolled AlN/AZ91 composite sheets
T
Bin Zhanga, Changlin Yanga,*, Yunxia Suna, Xinlin Lib, Feng Liua,** a b
State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, Shaanxi, 710072, PR China College of Materials Science and Chemical Engineering, Harbin Engineering University, Harbin, Heilongjiang, 150001, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Composites Rolling Microstructure Mechanical properties Deformation mechanism
Due to the low ductility and poor formability, few studies were reported about the rolling behavior of magnesium matrix composites. In this study, in-situ AlN/AZ91 composite sheets were rolled at 400 °C with the total thickness reduction of 80 %. Compared with other rolled magnesium matrix composites, rolled AlN/AZ91 composite sheets in this study showed the good surface quality and better strength-ductility combination, which is expected to extend its application in industry. Although fully recrystallized α-Mg grains exhibited a basal rolling texture, there was no significant difference in mechanical properties along the rolling direction and transverse direction due to the weakening effect of AlN ceramic particles on the basal texture. During rolling, the high density of dislocations were formed around AlN ceramic particles owing to the residual plastic strain between AlN ceramic particles and AZ91 matrix, which promoted the continuous precipitation of γ-Mg17Al12 phase, so densely distributed nano-precipitates were obtained in the interior of the grains. TEM analysis reveals in-situ AlN ceramic particles not acted as the cracking source during rolling, which had a strong interfacial bonding with magnesium matrix, contributing to the good rolling capacity. Strengthening mechanism analysis suggests the high strength increment was mainly attributed to the dislocation strengthening and precipitation strengthening. For tensile axis along rolling direction and transverse direction, the dominated deformation mechanism at room temperature was pyramidal plane slip, leading to a relative high ductility.
1. Introduction Magnesium matrix composites have attracted increasing attentions in aerospace and automobile fields because of their excellent properties, such as the low density, good wear resistance, high specific strength and modulus, etc [1,2]. So far, various reinforcements such as SiC [3,4], TiC [5], TiB2 [1,6], GNPs [7] and AlN [2,8,9] reinforced magnesium matrix composites have been fabricated successfully by the ex-situ and in-situ fabrication techniques. As is known, the mechanical properties of magnesium matrix composites are very dependent on the size of the reinforcement. In general, the micro- and submicro-sized reinforcements enhance the strength of the matrix alloy at the expense of the ductility [10]. Recent years, some researchers found that the nano-sized reinforcements can improve the strength and ductility of the matrix alloys simultaneously [1,2,4,9,11]. Although the mechanical properties of magnesium alloys can be improved by the reinforcements, the low mechanical properties of ascast magnesium matrix composites can not meet the mechanical
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property requirement for using in structural components. In order to extend the application of magnesium matrix composites in industries, the secondary processing after casting is often used to further improve the mechanical properties of magnesium matrix composites. As one of the most common metal processing technique in practice, rolling has been widely used in magnesium alloys for acquiring the high performance plates. However, due to the low ductility and poor formability of magnesium matrix composites, so far there have been few studies on the rolling behavior of magnesium matrix composites. Furthermore, the small number of previous studies all focused on the rolling behavior of particle reinforced AZ31 magnesium matrix composites. To the authors’ knowledge, to date the rolling behaviors of AZ91 magnesium matrix composites have not been studied and reported. Previous studies of authors [9] have demonstrated that the ductility of as-cast magnesium alloy AZ91 also can acquire a significant increase (from 3.9 ± 0.5 % to 20 ± 3 %) due to the formation of in-situ AlN ceramic particles, as well as the increase of strength, which made AZ91 alloy possess the great potential for deformation. In this work, for
Corresponding author. Corresponding author. E-mail addresses: [email protected] (B. Zhang), [email protected] (C. Yang), [email protected] (F. Liu).
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https://doi.org/10.1016/j.msea.2019.138118 Received 21 May 2019; Received in revised form 4 July 2019; Accepted 5 July 2019 Available online 09 July 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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possessed the excellent rolling formability. Fig. 1 (b) shows the engineering stress-strain curves of the as-cast and rolled samples along the rolling direction (RD) and transverse direction (TD). The corresponding yield strength, ultimate tensile strength and elongation to fracture are listed in Table 1. It can be seen from Fig. 1 (b) that both the yield strength and ultimate tensile strength of AlN/AZ91 composites achieved a great improvement by multi-pass rolling. As shown in Table 1, the yield strength and ultimate tensile strength of AlN/AZ91 composites were increased from 103 ± 6 MPa and 240 ± 10 MPa to 338 ± 5 MPa (RD), 351 ± 6 MPa (TD) and 387 ± 7 MPa (RD), 405 ± 9 MPa (TD) respectively. While the elongation to fracture was decreased after rolling, the rolled AlN/AZ91 composite sheets still maintained a relative high ductility. Fig. 1 (c) compared the yield strength and elongation to fracture of various rolled magnesium matrix composites [12–15]. Evidently, the rolled AlN/AZ91 sheets in this study realized a better strength-ductility combination among them. Before rolling, the as-cast AlN/AZ91 sheets were first SST at 415 °C for 24 h. The SST microstructure was shown in Fig. 2 (a). It can be seen in Fig. 2 (a) that the SST microstructure was consist of α-Mg grains and some black particles. A majority of granular black particles were undissolved Al8Mn5 phase, and little mass-like and fiber-like black particles were in-situ formed AlN clusters. Fig. 2 (b) shows the enlarged SEM image of the rectangular region in Fig. 2 (a). The inset is the EDS analysis result of the particle marked with cross. Combined with the EDS analysis result, it can be concluded that the AlN clusters were consist of many nano- and submicro-sized AlN ceramic particles. In addition, it can be found in Fig. 2 (a) that only little macro-segregation occurrence in SST AlN/AZ91 sheets and most of AlN ceramic particles were uniformly distributed in AZ91 matrix. After multi-pass rolling, the microstructure of AlN/AZ91 sheets is shown in Fig. 2 (c). By comparing Fig. 2 (a) and (c), it can be found that the average size of AlN clusters was decreased after multi-pass rolling, which means the distribution of AlN ceramic particles got improvement after multi-pass rolling. Additionally, by contrast, it also can be concluded the rolled AlN/AZ91 sheet in Fig. 2 (c) has a full dynamic recrystallization microstructure with the grain size significantly refined. Fig. 2 (d) exhibits the grain size variation of SST and rolled AlN/AZ91 sheets. As shown in Fig. 2 (d), after multi-pass rolling, the average grain size of α-Mg was decreased from 75 ± 2 μm to 25 ± 3 μm. The anisotropic mechanical properties of rolled AlN/AZ91 composite sheets along RD and TD (see Table 1) implies that the recrystallized α-Mg grains may produce a preferred orientation during rolling. Hence, the macro-texture of rolled AlN/AZ91 composite sheets was measured by XRD in this study. Fig. 3 displays the (0001) pole figure and inverse pole figure of rolled AlN/AZ91 sheets. As shown in Fig. 3 (a), the rolled AlN/AZ91 sheets exhibits a basal rolling texture. The c-axes of α-Mg grains were slightly tilted away from the normal direction (ND) and lied on the ND-RD planes of the rolled AlN/AZ91 composite sheets, thereby producing an ellipsoidal intensity distribution in the basal plane pole figure. In Fig. 3 (b), the intensity distribution is focused on the center of the circle arc, which means the ⟨10-10⟩ direction of α-Mg grains was tilted toward RD during multi-pass rolling. However, due to in-situ formed AlN particles can weaken the basal texture effectively [16], so the basal rolling texture in this study had the lower maximum texture intensity, which was also confirmed in the higher ratio of the yield strength along RD and TD (≈0.96). To further investigate the microstructural evolution during rolling, EBSD test was performed for rolled AlN/AZ91 sheets. Fig. 4 (a) shows the inverse pole figure (IPF) of rolled AlN/AZ91 sheets. It can be seen from Fig. 4 (a) that the coarse grains and fine grains coexisted in the rolled microstructure. The corresponding orientation imaging map is shown in Fig. 4 (b). The low angle boundaries (LABs, 2°–15°) and high angle boundaries (HABs, > 15°) were marked with red lines and black lines, respectively. As shown in Fig. 4 (b), the grain boundaries of these coarse grains and fine grains all belong to HABs. It further demonstrated that α-Mg grains occurred fully recrystallization during rolling,
further improving the mechanical properties and extending its engineering application, the AlN/AZ91 composite was rolled at 400 °C and a total reduction of 80 % was achieved. The microstructure, mechanical properties and tensile deformation mechanism of rolled AlN/ AZ91 composite sheets were investigated systematically. 2. Experiment In this work, commercial AZ91 casting magnesium alloy was chosen as the matrix alloy. The confirmed chemical composition was Mg9.13Al-0.74Zn-0.27Mn (wt. %). Nitrogen gas with high purity (99.999 %) was used as the reaction gas, and argon gas with high purity (99.999 %) was used to maintain an inert atmosphere in the reaction chamber. In-situ AlN/AZ91 composites with the volume fraction of 1.5 % were fabricated by the liquid nitriding method. The detailed fabrication processing of the liquid niriding method were described in Ref. [9]. Lastly, the melt was poured into the permanent mould and cast into the plates with a length of 100 mm, a width of 60 mm and a thickness of 20 mm. Before rolling, the as-cast AlN/AZ91 composites were firstly solid solution treated (SST) at 415 °C for 24 h. Then the SST AlN/AZ91 composite sheets were used for repeating rolling and annealing. The AlN/AZ91 composite sheets were rolled from 13 mm into 2.6 mm with the total thickness reduction of 80 %. The reduction of the first pass was 5 %, and the reduction of other pass was kept at 20 %. Before each rolling pass, the sheets were firstly annealed at 400 °C for 15 min and then rolled at the cold roller. The nitrogen content of rolled AlN/AZ91 composite sheets was measured by the oxygen nitrogen hydrogen analyzer (LECO ONH836) and the alloying element content was measured by the inductively coupled plasma optical emission spectrometer (Thermo Scientific iCAP 7000). To ensure the reliability, at least five samples were tested. The average composition of rolled AlN/AZ91 composite sheets was Mg1.25N-9.02Al-0.65Zn-0.23Mn (wt. %). The rolled specimens were mechanically ground, subsequently polished and etched in acetic picric solution (5 g picric acid, 5 ml acetic acid, 10 ml distilled water and 80 ml ethanol) at the room temperature for 12s. Optical microscopy (OM), high resolution scanning electron microscopy (HRSEM, FEI Nova 450 SEM) equipped with an energy dispersive spectrum (EDS) and electron backscattered diffraction (EBSD) were employed to observe the microstructure. The EBSD samples were electrochemical polished in an AC2 electrolyte at −30 °C for 90 s. The dislocations and dislocation/ particle interactions in rolled AlN/AZ91 composite sheets were also analyzed using transmission electron microscopy (TEM, FEI Talos F200X) in addition to the size and morphology of the particles. TEM foils were prepared by mechanically grinding them down to the thickness of 30 μm and further thinned using a precision ion polishing system (Gatan 691, USA). The tensile stress and strain of the specimens with thickness of 2 mm, width of 6 mm and gauge length of 25 mm, were measured in the SHIMADZU AG-X Tester at a strain rate 10−3s−1. The tensile specimens were prepared at the angles of 0° and 90° to the rolling direction by wire-cutting. For the tensile test at room temperature, five specimens were tested and the displacement was measured with an Epsilon axial extensometer. Macro-texture measurements of tensile specimens with different deformation were carried out by Bruker D8 Discover diffractometer with a high resolution area detector, operating at 35 kV and 40 mA. The pole figures were calculated using the quantitative texture analysis software MULTEX 3. 3. Results Fig. 1 (a) displays the size and surface quality of AlN/AZ91 composite sheets before and after rolling. It is obvious that the AlN/AZ91 composite sheets grew longer after 80 % rolling deformation and the width of AlN/AZ91 sheet had no significant increase. More importantly, as shown in Fig. 1 (a), there was no edge crack observed on the surface of rolled sheets, which proved that in-situ AlN/AZ91 composite sheets 2
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Fig.1. (a) The size and surface quality of AlN/AZ91 composite sheets before and after rolling, (b) the engineering stress-strain curves of as-cast and rolled AlN/AZ91 sheets, and (c) the comparison of the yield strength vs. Elongation to failure of various rolled magnesium matrix composites.
Table 1 Tensile properties of as-cast and rolled AlN/AZ91 composites. Materials As-cast AlN/AZ91 Rolled AlN/AZ91
Direction
σ0.2 (MPa)
σb (MPa)
ε (%)
TD RD
103 ± 6 351 ± 6 338 ± 5
240 ± 10 405 ± 9 387 ± 7
20 ± 3 10.5 ± 0.6 8.1 ± 0.3
and these coarse grains and fine grains in Fig. 4 (a) all were the dynamic recrystallized grains. In addition, it also can be found from Fig. 4 (b) that a lot of LABs were formed within α-Mg grains, which suggests that high density of dislocation substructures were produced during rolling. The number fractions of the LABs and HABs were estimated, as listed in Fig. 4 (c). Meanwhile, two dominate peaks can be observed in
Fig. 3. The (0001) pole figure (a) and inverse pole figure (b) analyzed from the upper surface of the rolled sample.
Fig. 2. The microstructures (a, c) and grain size distributions (d) of solid solution treated and rolled AlN/AZ91 composites, respectively, (b) the corresponding HRSEM graph of the rectangular region in (a). 3
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Fig. 4. (a) IPF map of rolled AlN/AZ91 composite sheets, (b) the corresponding orientation imaging map, and (c) the misorientation distribution calculated from EBSD in rolled AlN/AZ91 composite sheets.
Fig. 4 (c). The first high peak around 5° represents the high density of dislocation substructures and the second lower peak around 30° corresponded to the recrystallized grain boundaries [17]. The LABs and HABs were further analyzed by means of TEM. Fig. 5 (a) confirms that a lot of dislocation substructures were formed during multi-pass rolling, such as dislocation tangles and dislocation walls, etc. On the one hand, high density of dislocations were formed and tangled
around the uniform distributed in-situ AlN ceramic particles during rolling, while few dislocations were observed in the area without AlN ceramic particles. On the other hand, during rolling, by the annihilation and rearrangement of dislocations, the number of dislocations in the cell interiors had decreased and the cells became subgrains. The regular dislocation walls evolved into low angle grain boundaries (subgrain boundaries). Fig. 5 (b) shows the bright field image of HABs in rolled
Fig. 5. TEM images of LABs and HABs in rolled AlN/AZ91 composite sheets: (a) various dislocation substructures and (b) the migration of recrystallized grain boundary. 4
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Fig. 6. (a–c) FESEM images of γ-Mg17Al12 precipitates in rolled AlN/AZ91 composite sheets, the EDS analysis in (c) corresponding to continuous precipitates, (d) the size distributions of sub-micron γ-Mg17Al12 precipitates estimated by FESEM images.
cross. Combined with the SAED pattern, the black particles with submicro- and nano-size in Fig. 7 (a) were identified as the in-situ formed AlN ceramic particles. Similarly to the microstructure in Fig. 5 (a), high densities of dislocations were produced around in-situ AlN ceramic particles in Fig. 7 (a). More interestingly, it can be found in Fig. 7 (a) that a lot of nano-precipitates (marked with yellow arrows) were formed in the region of high densities of dislocations. Fig. 7(c–d) shows the TEM image of nano-precipitates in the vicinity of AlN ceramic particles and the corresponding SAED pattern. According to analysis of the diffraction rings in Fig. 7 (d), it can be identified that these densely distributed particles around in-situ AlN ceramic particles in Fig. 7 (c) were the γ-Mg17Al12 nano-precipitates. The size and density of γMg17Al12 nano-precipitates were estimated from the TEM images (at least five TEM images were used for analyzing). The average diameter of γ-Mg17Al12 nano-precipitates was 23.5 nm and the volume fraction of the γ-Mg17Al12 nano-precipitates was about 9.49 %. Fig. 8 (a) presents the bright field TEM image of rolled AlN/AZ91 composite sheets. γ-Mg17Al12 precipitates and in-situ AlN ceramic particles were dispersed uniformly in the magnesium matrix. The AlN/ Mg interface from the yellow rectangular region surrounded by dash lines in Fig. 8 (a) was characterized by HRTEM. Combined the corresponding fast Fourier transformed (FFT) SAED patterns, it can be concluded from Fig. 8 (b) that in-situ AlN ceramic particles and magnesium matrix possessed a strong interfacial bonding in AlN/AZ91 composite sheets. No crack initiation and propagation was detected in the AlN ceramic particles or at the AlN/Mg interface, which indicates dislocations can cut or bypass in-situ AlN particles instead of piling up during rolling, and the load can transfer from the soft magnesium matrix to the hard AlN ceramic particles effectively. The tensile fracture surfaces of the rolled AlN/AZ91 composite sheets along TD and RD were observed by SEM, as shown in Fig. 9. It is obvious that both samples along TD and RD exhibited the brittle-ductile mixed fractured feature, i.e. the cleavage planes and dimples were observed at the same time in the tensile samples along TD and RD.
AlN/AZ91 composite sheets. It can be found clearly in Fig. 5 (b) that the recrystallized grain boundaries were migrated from the dislocation free region to the dislocation rich region by grain boundary bulging during rolling. Before rolling, as shown in Fig. 2 (a), those coarse eutectic βMg17Al12 phases distributed along the grain boundaries [9] had been fully dissolved into α-Mg matrix. As shown in Fig. 6 (a), after multi-pass rolling, densely distributed precipitates were observed in AlN/AZ91 composite sheets again as well as little AlN clusters, which suggests that γ-Mg17Al12 phase was dynamically precipitated from the matrix during rolling. The morphology, size and precipitated type of γ-Mg17Al12 phase had a great difference from that of as-cast. Firstly, it can be seen from Fig. 6 (b) that there were two types of γ-Mg17Al12 precipitates: discontinuous precipitation and continuous precipitation. While most of the γ-Mg17Al12 precipitates were observed within α-Mg grains and only little γ-Mg17Al12 particles were distributed along the grain boundaries. The number fraction of discontinuous γ-Mg17Al12 particles precipitated on the grain boundaries was estimated by the SEM images as about 7 %. So the γ-Mg17Al12 particles precipitated during rolling in this study were mainly continuous precipitation. The enlarge FESEM image of these continuous precipitated γ-Mg17Al12 particles in α-Mg grains is shown in Fig. 6 (c). Evidently, the continuous precipitated γ-Mg17Al12 particles had a spherical morphology and a sub-micron size. In this study, the size of the continuous precipitated sub-micron γ-Mg17Al12 particles was estimated under SEM (at least five SEM images were used for estimating). The size distribution of the continuous precipitated γMg17Al12 particles is shown in Fig. 6 (d). It can be seen from Fig. 6 (d) that the size of the continuous precipitated γ-Mg17Al12 particles was mainly ranged from 300 nm to 800 nm and its average size was about 487.5 nm. In addition, the volume fraction of γ-Mg17Al12 precipitates was also estimated by the SEM images, as 6.36 %. Fig. 7 (a) exhibits the bright field TEM image of rolled AlN/AZ91 composite sheets. Fig. 7 (b) is the corresponding selected area electron diffraction (SAED) pattern of the black particle marked with the white 5
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Fig. 7. (a) Dislocations tangled by AlN ceramic particles and precipitated nano-Mg17Al12 particles around the AlN ceramic particles, (b) the SAED pattern of AlN ceramic particles, (c) TEM image of nano-Mg17Al12 precipitates and the corresponding SAED pattern (d).
cracking of the rolled sheets. For AZ91 alloys hot rolled at 400 °C, edge cracks began to occur when the accumulated reduction was about 50% and edge cracking became more severe with the accumulated reduction further increasing [18]. In this study, the AlN/AZ91 composite sheets were rolled at 400 °C with a total thickness reduction of 80 %. Importantly, no surface crack and edge crack was formed on the surface of rolled AlN/AZ91 composite sheets (Fig. 1a), which means that in-situ AlN ceramic particles can improve the rolling capability of AZ91 alloys significantly. Based on previous study [9], the ductility of as-cast AZ91 alloy acquired a remarkable enhancement due to the formation of insitu AlN ceramic particles, and this made AlN/AZ91 composites possess large deformation ability. So the rolling capacity improvement of AZ91 alloys may originate from its ductility enhancement due to the formation of in-situ AlN ceramic particles. Additionally, the interfacial bonding between AlN ceramic particles and magnesium matrix plays the critical role in the rolling capability. The poor interface easily acts as the source of cracking, resulting in the crack initiation in the early of rolling. In this case, the rolled sheets will fracture with a small thickness reduction. In this study, the AlN/Mg interfaces had a strong bonding (Figs. 8 and 9), which well realized the load transfer from the matrix to
However, by contrast, it can be found that all the amount, size and depth of dimples on the fracture surface along TD are much larger than that along RD. It suggests the rolled AlN/AZ91 composite sheets along TD have a better ductility than that along RD, which is consistent with the mechanical properties in Table 1. In addition, some particles with diameter of sub-micron, which were identified as γ-Mg17Al12 precipitates by the EDS analysis (the inset in Fig. 9a), were observed inside the dimples as indicated by the yellow arrows in Fig. 9 (a). Thus, it can be concluded that cracks were initiated at the γ-Mg17Al12 precipitates rather than in-situ AlN ceramic particles, which is agree well with the result of Fig. 8. 4. Discussion 4.1. Effect of in-situ AlN ceramic particles on the rolling capability of AZ91 alloys The rolling capability of materials can be characterized by the degree of edge cracking. In general, increasing the rolled thickness reduction and decreasing the rolling temperature will promote the edge 6
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Fig. 8. (a) The bright field TEM image of rolled AlN/AZ91 composite sheets, (b) the HRTEM image of the AlN/Mg interface in rolled AlN/AZ91 composite sheets and the corresponding fast Fourier transformed SAED patterns from the rectangular regions.
Fig. 9. Fracture surfaces of rolled AlN/AZ91 composite tensile sample along TD (a) and RD (b), the EDS analysis result corresponding to the particles in the dimples (yellow arrows). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
And continuous precipitates nucleate in the interior of grains, which depends on the volume diffusion [19]. In other words, if the volume diffusion is accelerated, the continuous precipitation will be promoted. In this study, due to the formation of in-situ AlN ceramic particles, high densities of dislocations were produced and tangled around the AlN ceramic particles during rolling (Figs. 5 and 7a). As shown in Figs. 5, 7a and 8a, most of in-situ formed AlN ceramic particles were uniformly dispersed in the interior of α-Mg grains. So these high densities of dislocations around AlN ceramic particles were located in the interior of α-Mg grains. On the one hand, these high densities of dislocations can accelerate the diffusion of aluminum atoms and promote the precipitation of γ-Mg17Al12 phase. On the other hand, dislocations can act as the heterogeneous nucleation sites for γ-Mg17Al12 precipitates [20] and promote its precipitation. Hence, most of γ-Mg17Al12 precipitates formed during rolling were continuous precipitates in this study.
Table 2 The volume fractions and average sizes of submicro- and nano-AlN ceramic particles.
Submicro-AlN Nano-AlN
Average size (nm)
Volume fraction (%)
231.2 23.7
0.62 0.71
AlN ceramic particles. As a result, no crack initiated at the AlN/Mg interfaces (Fig. 8) during rolling, which also contributed to the large rolling reduction and good rolling capability.
4.2. Effect of in-situ AlN ceramic particles on the precipitation behavior of AZ91 alloys For AZ91 alloys, most of γ-Mg17Al12 precipitates produced during hot rolling at 400 °C were distributed along the grain boundaries [18]. However, in this study, about 93 % (number fraction) γ-Mg17Al12 precipitates were continuous precipitates (Fig. 6b). It demonstrated that in-situ AlN ceramic particles promoted the continuous precipitation of γ-Mg17Al12 phase during multi-pass rolling. As is known, discontinuous precipitates nucleate on the grain boundaries, so the discontinuous precipitation is favored when the grain boundary diffusion is dominant.
4.3. Strengthening mechanisms analysis In this study, after multi-pass rolling, the yield strength of as-cast AlN/AZ91 composite achieved a significant improvement (Fig. 1b and Table 1). The significant increase in yield strength was caused by grain boundary strengthening, precipitation strengthening and dislocation strengthening. Firstly, during rolling, the grain size of α-Mg was refined owing to 7
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Fig. 10. (a) The increment comparison of yield strength (σ0.2) in AlN/AZ91 composites due to grain boundary strengthening, precipitation strengthening and dislocation strengthening, respectively, (b) the theoretical and experimental values of the increased yield strength in AlN/AZ91 composites during rolling.
AlN ceramic particles due to the residual plastic strain between AlN ceramic particles and AZ91 matrix during rolling [26], which can effectively increase the yield strength of AlN/AZ91 composites. The increased yield strength Δσd due to dislocation strengthening can be calculated by the following equation [26]:
the dynamic recrystallization. For the grain boundary strengthening, the Hall-Petch relationship is applied and the strength increment Δσg is calculated as follows: −1
−1
2 Δσg = k (drolled − das −2cast )
(1) 1/2
for AZ91 Where k is the Hall-Petch coefficient and is 0.13 MPa m alloy, drolled and das − cast are the average grain size of the rolled (25 μm) and as-cast (50 μm [9]) AlN/AZ91 composites, respectively. According to calculating, the increased yield strength due to grain boundary strengthening is about 7.62 MPa. Secondly, densely distributed γ-Mg17Al12 precipitates formed during rolling can give rise to significant strengthening effect by Orowan mechanism. In this study, most of γ-Mg17Al12 precipitates belong to continuous precipitation. Compared with discontinuous precipitation, continuous precipitated γ-Mg17Al12 precipitates inside the grains are more effective for inhibiting the dislocation motion [19–22], which resulted in the high strength contribution of the precipitation strengthening. The increased stress for the slip dislocations to bypass γMg17Al12 precipitates is described by the Orowan equation [23,24]:
Δσp =
dp Gb ln r0 2πλ 1 − ν
Δσd =
Gb 2π 1 − ν (
0.779 f
− 0.785) dt
(2)
ln
0.785dt b
12ΔTΔαV bd
(4)
where β is the strengthening coefficient (is 1.25), ΔT is the temperature difference between rolling processing and mechanical tests, Δα is the difference in thermal expansion coefficient between AlN ceramic and AZ91 matrix (is 26.3 × 10−6 K−1 [9]), V is the volume fraction of AlN ceramic particles and d is the average size of AlN ceramic particles. In this study, although in-situ formed AlN ceramic particles had a multiscale size, the fractions of submicro- and nano-AlN ceramic particles were much more than that of micro-AlN ceramic particles [9]. The volume fractions and average sizes of submicro- and nano-AlN ceramic particles in AlN/AZ91 composites were estimated from the TEM and SEM images, which are listed in Table 2. Accordingly, the increased yield strength due to dislocation strengthening from submicro- and nano-AlN ceramic particles are 32.73 MPa and 109.38 MPa, respectively. The micro-sized AlN ceramic particles had the lower volume fraction (~0.17 %) and the larger size, which resulted in a weak effect of dislocation strengthening. Thus, the dislocation strengthening from the micro-AlN ceramic particles can be ignored in this study. Above all, the increased yield strength due to dislocation strengthening from insitu AlN ceramic particles is 142.11 MPa. Fig. 10 (a) shows a comparison of the contributions of grain boundary strengthening, precipitation strengthening and dislocation strengthening to the yield strength increment of rolled AlN/AZ91 composites. Evidently, the increased yield strength of AlN/AZ91 composites after rolling is mainly attributed to the dislocation strengthening (~142.11 MPa) and the precipitation strengthening (~95.47 MPa), which means in-situ AlN ceramic particles effectively improved the strength during rolling via increasing the dislocation density and promoting the continuous precipitation of γ-Mg17Al12 phase. While the continuous precipitated γ-Mg17Al12 particles within grains can not pin the grain boundary, so the dynamically recrystallized grains had a larger size, resulting in the low contributions for yield strength increment (~7.62 MPa). Furthermore, the increased yield strength during rolling can be predicted by the following equation:
Where Δσp is the increased stress due to the precipitation strengthening, G is the shear modulus (is 15 GPa for magnesium), b is magnitude of the Burgers vector of the slip dislocations (is 0.32 nm), ν is the Poisson's ratio (is 0.29 for magnesium), λ is the effective planar inter-precipitate spacing, dp is the mean planar diameter of the precipitates and r0 is the dislocation core radius (commonly, the assumption r0 = b is used). For the spherical precipitates (Fig. 7c), the effective planar inter-precipitate spacing λ and the mean planar diameter of the precipitates dp can be calculated according to the method of J-F. Nie et al. [25], and Eq. (2) can be written as:
Δσp =
3 βGb
(3)
where f is the volume fraction of γ-Mg17Al12 precipitates and dt is the average diameter of the spherical γ-Mg17Al12 precipitates. In this study, dynamic precipitated γ-Mg17Al12 precipitates had nano and sub-micron two sizes. Firstly, for γ-Mg17Al12 submicro-precipitates, the average diameter and the volume fraction were 487.5 nm and 6.36 %, respectively. According to calculating, the Orowan strengthening from submicro-precipitates is 5.72 MPa. Secondly, dynamic precipitated nanosized γ-Mg17Al12 precipitates had a finer size (dt = 23.5 nm) and higher content ( f = 9.49 %). In this case, the Orowan strengthening from γMg17Al12 nano-precipitates is 89.75 MPa. Thus, the increased yield strength due to γ-Mg17Al12 precipitates is 95.47 MPa. Thirdly, high densities of dislocations were formed around in-situ
ΔσTheory = Δσg + Δσp + Δσd
(5)
where ΔσTheory is the theoretical predicted value of the yield strength increment during rolling. By adding up the contributions of grain boundary strengthening, precipitation strengthening and dislocation strengthening, the theoretical predicted value ΔσTheory is 245.2 MPa. The theoretical predicted value ΔσTheory and the experimental obtained 8
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Fig. 11. (0001), (10-10), (10–11) and (11–22) pole figures of rolled AlN/AZ91 samples after 0 % (a, d g, j), 3%(b, e, h, k) and 10 % (c, f, i, l) tensile deformation along TD.
and pyramidal II plane (10–12) pole figures of rolled AlN/AZ91 composite samples along TD and RD with different tensile strain, respectively. It can be seen from the basal pole figures in Figs. 11 and 12 that the basal plane texture exhibits an ellipsoidal intensity distribution around the pole figure center and no intensity distribution at the north and south poles, which means that the major deformation mechanism of rolled AlN/AZ91 composite sheets at room temperature is dislocation slip rather than twinning [27]. Generally, the twinning behavior is influenced by grain size, texture and second-phase. Firstly, some studies reported that twinning has the positive grain size effect [28,29], i.e. twinning forms more difficult in finer grains. Secondly, C.M. CepedaJiménez et al. [30] confirmed that the twinning behavior during inplane tension of rolled magnesium alloys was due to the local stress accommodation rather than texture effect. Besides, X. Li and J.B. Clarke et al. [31,32] found a high level of Mg17Al12 precipitates could suppress the twinning behavior during deformation. In this study, the grain size
values along TD and RD are compared in Fig. 6 (b). The theoretical predicted values ΔσTheory (245.2 MPa) is consistent well with the experiment value along TD ΔσExperiment − TD (248 MPa) and RD ΔσExperiment − RD (235 MPa), and exactly predict the increased yield strength of AlN/AZ91 composites during rolling. 4.4. Tensile deformation mechanism analysis As shown in Fig. 1 (c), the rolled AlN/AZ91 composite sheets possessed a better ductility among the various rolled magnesium matrix composite sheets. The elongation to fracture is even much higher than that of AZ31 (typical wrought magnesium alloy) matrix composite sheets. To study the reasons for high ductility of rolled AlN/AZ91 composite sheets, the tensile deformation mechanism at room temperature was studied in this section. Fig. 11 and Fig. 12 display the basal plane (0001), prismatic plane (10-10), pyramidal I plane (10–11) 9
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Fig. 12. (0001), (10-10), (10–11) and (11–22) pole figures of rolled AlN/AZ91 samples after 2.5 % (a, d g, j), 4 % (b, e, h, k) and 7 % (c, f, i, l) tensile deformation along RD.
was decreased effectively (Fig. 2d) and a high density of γ-Mg17Al12 precipitates (Figs. 6b and 7c) were obtained during rolling. Based on above analysis, it can be concluded that the weak twinning mode of rolled AlN/AZ91 composite sheets during tensile deformation may mainly attributed to the finer grain size and high density of γ-Mg17Al12 precipitates. During tensile deformation, the slip planes tend to rotate to lie parallel to the tensile axis, so the activated slip systems can be identified indirectly by analyzing the evolution of texture intensity during deformation. For the tensile axis along TD, as shown in Fig. 11, the maximum texture intensity of the basal plane increased slightly from 3.58 to 3.64 and 3.69 with increasing the tensile strain to 3 % and 10 %, which increased by 1.7 % and 3.1 % respectively. It suggests that a small number of basal planes were initiated and rotated to parallel with
the tensile axis during tension. In contrast, the maximum intensity of the prismatic plane decreased by 15.8 % and 18.1 % with increasing the tensile strain, which means that the prismatic plane slip was suppressed and the prismatic planes were gradually far away from the tensile axis during tension. Surprisingly, it can be found in Fig. 11 that both the maximum intensities of pyramidal I plane and pyramidal II acquired a significant improvement. The maximum intensities of the pyramidal I plane and pyramidal II plane increased by 7.4 % and 5.2 % when the tensile strain increased to 3 %. With further increasing the tensile strain to 10 %, the maximum intensities of the pyramidal I plane and pyramidal II increased by 21 % and 14 %, which must be the result of pyramidal dislocation activity in favorably oriented grains. It demonstrates that profuse pyramidal plane slip systems were activated when the rolled AlN/AZ91 composite sheets tension along TD. 10
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Fig. 13. The variation of the maximum texture intensities as a function of tensile strain along TD (a) and RD (b).
Similarly, for the tensile axis along RD, the pole figures of basal plane, prismatic plane and pyramidal plane with different tensile strain are shown in Fig. 12. According to analyzing the evolution of the maximum intensities during tension, similar conclusions were obtained for the tension along RD. With increasing the tensile strain, the maximum intensities of the basal plane and prismatic plane decreased gradually, which means the basal plane slip and prismatic plane slip were suppressed. In contrary, the maximum intensities of the pyramidal I plane and pyramidal II got a larger increase. Thus, the tensile deformation along RD was mainly attributed to the pyramidal slip. In general, the dominate slip systems at room temperature in magnesium are the basal plane slip system [33]. However, the basal plane slip only provide two independent slip systems [27], which fails to satisfy the Taylor criterion [34], resulting in the poor ductility of magnesium alloys and magnesium matrix composites. Fig. 13 shows the maximum intensities variation of different slip planes with the tensile strain along TD and RD. It can be seen from Fig. 13 that no matter the tensile axis along TD or RD, only the maximum intensities of pyramidal plane (pyramidal I and pyramidal II) increased significantly with the tensile strain. This suggests that pyramidal slip dominated the tensile deformation of rolled AlN/AZ91 composite at room temperature. Similar result was also obtained in rolled Mg–Li alloys by analyzing texture evolution using elasto-plastic self-consistent (EPSC) model [35]. The activity of pyramidal plane slip increased the number of slip systems and made the Taylor criterion get satisfied [27]. Hence, a high elongation to fracture was obtained in the rolled AlN/AZ91 composite sheets (Fig. 1c). Furthermore, it should be noted in Fig. 13 (a) that the maximum intensities of pyramidal I plane and pyramidal II plane almost increased at constant rate with increasing the tensile strain. However, as shown in Fig. 13 (b), with increasing the tensile strain, the maximum intensities of pyramidal I plane and pyramidal II plane first increased and then almost remained unchanged. This suggests that the number of favorably oriented grains for activating pyramidal slip along RD was much less than that along TD during later deformation, which may lead to the less ductility in RD (Table 1). The limited activity of basal plane slip and prismatic plane slip was mainly attributed to the unfavorable grain orientation. In rolled AlN/ AZ91 composite sheets, the grains were orientated such that their caxes were nearly normal to the sheet plane with a tilt along RD, and the prismatic planes exhibited a dominate orientation normal to TD (Fig. 3), which lead to the grain in a unfavorable orientation and formed the lower Schmid factors for basal plane slip and prismatic plane slip. But, due to the low critical resolved shear stress (CRSS) for basal plane slip, basal plane slip may be activated in few grains with special orientation, so the maximum intensities of basal plane had a
slight increase in Fig. 11. Since no non-basal slip was observed in rolled AZ91 alloys during tension at room temperature [18], it can be safely concluded that the pyramidal plane slip systems activity in rolled AlN/ AZ91 composite sheets is due to the formation of in-situ AlN ceramic particles. The formation of in-situ AlN ceramic particles may reduce the CRSS of pyramidal plane slip or the intrinsic stacking fault I1 energy, resulting in the activity of pyramidal plane slip systems [36,37]. Further research works need to be carried on for exploring the underlying mechanisms in future. 5. Conclusions In this work, in-situ AlN/AZ91 composite sheets fabricated by liquid nitridation method were subjected to multi-pass rolling with the aim of improving mechanical properties and extend its industrial application. The microstructure, mechanical properties and tensile deformation mechanism of rolled AlN/AZ91 composite sheets were investigated, and the obtained main conclusions were summarized as follows: (1) In-situ AlN/AZ91 composites had excellent rolling capacity. In situ AlN/AZ91 composite sheets can be rolled at 400 °C to the 80 % thickness reduction and no crack was formed on its surface. The excellent rolling capacity was mainly attributed to the high as-cast ductility and good interfacial bonding; (2) In this study, dynamically precipitated γ-Mg17Al12 particles were mainly belong to continuous precipitation. During rolling, high density of dislocations were formed near AlN ceramic particles, which distributed uniformly within grains, promoting the continuous precipitation of γ-Mg17Al12 phase; (3) By rolling, the mechanical properties of AlN/AZ91 composites realized a better strength-ductility combination compared to other rolled magnesium matrix composites. The increased strength was mainly attributed to the dislocation strengthening and precipitation strengthening; (4) The dominant tensile deformation mechanism of rolled AlN/AZ91 composite sheets was the pyramidal plane slip. The formation of insitu AlN ceramic particles activated the pyramidal dislocations, leading to a higher ductility in rolled AlN/AZ91 composite sheets. Acknowledge The authors are grateful for financial support for this research by the National Natural Science Foundation of China (Grant No. 51771151), the National Key R&D Program of China (Grant No. 2017YFB0305100), the Research Fund of the State Key Laboratory of Solidification Processing, China (Grant No. 2019-TS-02) and the Innovation 11
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Foundation for Doctor Dissertation of Northwestern Polytechnical University, China (Grant No. CX201905).
Mater. 64 (2011) 986–989. [18] L. Wang, Y.M. Kim, J. Lee, B.S. You, Improvement in rollability of AZ91 magnesium alloy by carbon addition, Mater. Sci. Eng. A 528 (2011) 943–949. [19] K.N. Braszczyńska-Malik, Discontinuous and continuous precipitation of magnesium-aluminum type alloys, J. Alloy. Comp. 477 (2009) 870–876. [20] C.L. Mendis, K. Hono, Understanding precipitation processes in magnesium alloys, in: M.O. Pekguleryuz, K.U. Kainer, A.A. Kaya (Eds.), Fundamentals of Magnesium Alloy Metallurgy, Woodhead Publishing, Cambridge, 2013, pp. 125–151. [21] C.P. Wang, R.L. Xin, D.R. Li, B. Song, M.Y. Wu, Q. Liu, Enhancing the age hardening response of rolled AZ80 alloy by pre-twinning deformation, Mater. Sci. Eng. A 680 (2017) 152–156. [22] L.Y. Jiang, D.F. Zhang, X.W. Fan, F. Guo, G.S. Hu, H.S. Xue, F.S. Pan, The effect of Sn addition on aging behavior and mechanical properties of wrought AZ80 magnesium alloy, J. Alloy. Comp. 620 (2015) 368–375. [23] J.D. Embury, D.J. Lloyd, Aluminium alloys: contemporary research and applications, in: A.K. Vasudevan, R.D. Doherty (Eds.), Treatise on Materials Science and Technology, vol. 31, Academic Press, New York, 1989, p. 579. [24] J.D. Robson, N. Stanford, M.R. Barnett, Effect of particles in promoting twin nucleation in a Mg-5 wt.% Zn alloy, Scr.Mater. 63 (2010) 823–826. [25] J.-F. Nie, Effects of precipitate shape and orientation on dispersion strengthening in magnesium alloys, Scripta Mater. 48 (2003) 1009–1015. [26] Y.C. Guo, K.B. Nie, X.K. Kang, K.K. Deng, J.G. Han, Z.H. Zhu, Achieving highstrength magnesium matrix nanocomposite through synergistical effect of external hybrid (SiC+TiC) nanoparticles and dynamic precipitated phase, J. Alloy. Comp. 771 (2019) 847–856. [27] A.S. Khan, A. Pandey, T. Gnäupel-Herold, R.K. Mishra, Mechanical response and texture evolution of AZ31 alloy at large strains for different strain rate and temperatures, Int. J. Plast. 27 (2011) 688–706. [28] S.R. Agnew, C.N. Tome, D.W. Brown, T.M. Holden, S.C. Vogel, Study of slip mechanisms in a magnesium alloy by neutron diffraction and modeling, Scripta Mater. 48 (2003) 1003. [29] M.R. Barnett, Z. Keshavarz, A.G. Beer, D. Atwell, Influence of grain size on the compressive deformation of wrought Mg-3Al-1Zn, Acta Mater. 52 (2004) 5093. [30] C.M. Cepeda-Jiménez, J.M. Molina-Aldareguia, M.T. Pérez-Prado, Effect of grain size on slip activity in pure magnesium polycrystals, Acta Mater. 84 (2015) 443–456. [31] X. Li, F. Jiao, T. Al-Samman, S. Ghosh Chowdhury, Influence of second-phase precipitates on the texture evolution of Mg-Al-Zn alloys during hot deformation, Scripta Mater. 66 (2012) 159–162. [32] J.B. Clarke, Age hardening in a Mg-9 wt.% alloy, Acta Mater. 16 (1968) 141. [33] N. Dixit, K.Y. Xie, K.J. Hemker, K.T. Ramesh, Microstructural evolution of pure magnesium under high strain rate loading, Acta Mater. 87 (2015) 56–67. [34] G.I. Taylor, Plastic strain in metals, J. Inst. Met. 62 (1938) 307–338. [35] M. Zecevic, I.J. Beyerlein, M. Knezevic, Activity of pyramidal I and II < c+a > slip in Mg alloys as revealed by texture development, J. Mech. Phys. Solids 111 (2018) 290–307. [36] R.K. Sabat, A.P. Brahme, R.K. Mishra, K. Inal, S. Suwas, Ductility enhancement in Mg-0.2%Ce alloy, Acta Mater. 161 (2018) 246–257. [37] S. Sandlobes, M. Friak, S. Zaefferer, A. Dick, S. Yi, D. Letzig, Z. Pei, L.-F. Zhu, J. Neugebauer, D. Raabe, The relation between ductility and stacking fault energies in Mg and Mg-Y alloys, Acta Mater. 60 (2012) 3011–3021.
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The microstructure, mechanical properties and tensile deformation mechanism of rolled AlN/AZ91 composite sheets
T
Bin Zhanga, Changlin Yanga,*, Yunxia Suna, Xinlin Lib, Feng Liua,** a b
State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, Shaanxi, 710072, PR China College of Materials Science and Chemical Engineering, Harbin Engineering University, Harbin, Heilongjiang, 150001, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Composites Rolling Microstructure Mechanical properties Deformation mechanism
Due to the low ductility and poor formability, few studies were reported about the rolling behavior of magnesium matrix composites. In this study, in-situ AlN/AZ91 composite sheets were rolled at 400 °C with the total thickness reduction of 80 %. Compared with other rolled magnesium matrix composites, rolled AlN/AZ91 composite sheets in this study showed the good surface quality and better strength-ductility combination, which is expected to extend its application in industry. Although fully recrystallized α-Mg grains exhibited a basal rolling texture, there was no significant difference in mechanical properties along the rolling direction and transverse direction due to the weakening effect of AlN ceramic particles on the basal texture. During rolling, the high density of dislocations were formed around AlN ceramic particles owing to the residual plastic strain between AlN ceramic particles and AZ91 matrix, which promoted the continuous precipitation of γ-Mg17Al12 phase, so densely distributed nano-precipitates were obtained in the interior of the grains. TEM analysis reveals in-situ AlN ceramic particles not acted as the cracking source during rolling, which had a strong interfacial bonding with magnesium matrix, contributing to the good rolling capacity. Strengthening mechanism analysis suggests the high strength increment was mainly attributed to the dislocation strengthening and precipitation strengthening. For tensile axis along rolling direction and transverse direction, the dominated deformation mechanism at room temperature was pyramidal plane slip, leading to a relative high ductility.
1. Introduction Magnesium matrix composites have attracted increasing attentions in aerospace and automobile fields because of their excellent properties, such as the low density, good wear resistance, high specific strength and modulus, etc [1,2]. So far, various reinforcements such as SiC [3,4], TiC [5], TiB2 [1,6], GNPs [7] and AlN [2,8,9] reinforced magnesium matrix composites have been fabricated successfully by the ex-situ and in-situ fabrication techniques. As is known, the mechanical properties of magnesium matrix composites are very dependent on the size of the reinforcement. In general, the micro- and submicro-sized reinforcements enhance the strength of the matrix alloy at the expense of the ductility [10]. Recent years, some researchers found that the nano-sized reinforcements can improve the strength and ductility of the matrix alloys simultaneously [1,2,4,9,11]. Although the mechanical properties of magnesium alloys can be improved by the reinforcements, the low mechanical properties of ascast magnesium matrix composites can not meet the mechanical
*
property requirement for using in structural components. In order to extend the application of magnesium matrix composites in industries, the secondary processing after casting is often used to further improve the mechanical properties of magnesium matrix composites. As one of the most common metal processing technique in practice, rolling has been widely used in magnesium alloys for acquiring the high performance plates. However, due to the low ductility and poor formability of magnesium matrix composites, so far there have been few studies on the rolling behavior of magnesium matrix composites. Furthermore, the small number of previous studies all focused on the rolling behavior of particle reinforced AZ31 magnesium matrix composites. To the authors’ knowledge, to date the rolling behaviors of AZ91 magnesium matrix composites have not been studied and reported. Previous studies of authors [9] have demonstrated that the ductility of as-cast magnesium alloy AZ91 also can acquire a significant increase (from 3.9 ± 0.5 % to 20 ± 3 %) due to the formation of in-situ AlN ceramic particles, as well as the increase of strength, which made AZ91 alloy possess the great potential for deformation. In this work, for
Corresponding author. Corresponding author. E-mail addresses: [email protected] (B. Zhang), [email protected] (C. Yang), [email protected] (F. Liu).
**
https://doi.org/10.1016/j.msea.2019.138118 Received 21 May 2019; Received in revised form 4 July 2019; Accepted 5 July 2019 Available online 09 July 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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possessed the excellent rolling formability. Fig. 1 (b) shows the engineering stress-strain curves of the as-cast and rolled samples along the rolling direction (RD) and transverse direction (TD). The corresponding yield strength, ultimate tensile strength and elongation to fracture are listed in Table 1. It can be seen from Fig. 1 (b) that both the yield strength and ultimate tensile strength of AlN/AZ91 composites achieved a great improvement by multi-pass rolling. As shown in Table 1, the yield strength and ultimate tensile strength of AlN/AZ91 composites were increased from 103 ± 6 MPa and 240 ± 10 MPa to 338 ± 5 MPa (RD), 351 ± 6 MPa (TD) and 387 ± 7 MPa (RD), 405 ± 9 MPa (TD) respectively. While the elongation to fracture was decreased after rolling, the rolled AlN/AZ91 composite sheets still maintained a relative high ductility. Fig. 1 (c) compared the yield strength and elongation to fracture of various rolled magnesium matrix composites [12–15]. Evidently, the rolled AlN/AZ91 sheets in this study realized a better strength-ductility combination among them. Before rolling, the as-cast AlN/AZ91 sheets were first SST at 415 °C for 24 h. The SST microstructure was shown in Fig. 2 (a). It can be seen in Fig. 2 (a) that the SST microstructure was consist of α-Mg grains and some black particles. A majority of granular black particles were undissolved Al8Mn5 phase, and little mass-like and fiber-like black particles were in-situ formed AlN clusters. Fig. 2 (b) shows the enlarged SEM image of the rectangular region in Fig. 2 (a). The inset is the EDS analysis result of the particle marked with cross. Combined with the EDS analysis result, it can be concluded that the AlN clusters were consist of many nano- and submicro-sized AlN ceramic particles. In addition, it can be found in Fig. 2 (a) that only little macro-segregation occurrence in SST AlN/AZ91 sheets and most of AlN ceramic particles were uniformly distributed in AZ91 matrix. After multi-pass rolling, the microstructure of AlN/AZ91 sheets is shown in Fig. 2 (c). By comparing Fig. 2 (a) and (c), it can be found that the average size of AlN clusters was decreased after multi-pass rolling, which means the distribution of AlN ceramic particles got improvement after multi-pass rolling. Additionally, by contrast, it also can be concluded the rolled AlN/AZ91 sheet in Fig. 2 (c) has a full dynamic recrystallization microstructure with the grain size significantly refined. Fig. 2 (d) exhibits the grain size variation of SST and rolled AlN/AZ91 sheets. As shown in Fig. 2 (d), after multi-pass rolling, the average grain size of α-Mg was decreased from 75 ± 2 μm to 25 ± 3 μm. The anisotropic mechanical properties of rolled AlN/AZ91 composite sheets along RD and TD (see Table 1) implies that the recrystallized α-Mg grains may produce a preferred orientation during rolling. Hence, the macro-texture of rolled AlN/AZ91 composite sheets was measured by XRD in this study. Fig. 3 displays the (0001) pole figure and inverse pole figure of rolled AlN/AZ91 sheets. As shown in Fig. 3 (a), the rolled AlN/AZ91 sheets exhibits a basal rolling texture. The c-axes of α-Mg grains were slightly tilted away from the normal direction (ND) and lied on the ND-RD planes of the rolled AlN/AZ91 composite sheets, thereby producing an ellipsoidal intensity distribution in the basal plane pole figure. In Fig. 3 (b), the intensity distribution is focused on the center of the circle arc, which means the ⟨10-10⟩ direction of α-Mg grains was tilted toward RD during multi-pass rolling. However, due to in-situ formed AlN particles can weaken the basal texture effectively [16], so the basal rolling texture in this study had the lower maximum texture intensity, which was also confirmed in the higher ratio of the yield strength along RD and TD (≈0.96). To further investigate the microstructural evolution during rolling, EBSD test was performed for rolled AlN/AZ91 sheets. Fig. 4 (a) shows the inverse pole figure (IPF) of rolled AlN/AZ91 sheets. It can be seen from Fig. 4 (a) that the coarse grains and fine grains coexisted in the rolled microstructure. The corresponding orientation imaging map is shown in Fig. 4 (b). The low angle boundaries (LABs, 2°–15°) and high angle boundaries (HABs, > 15°) were marked with red lines and black lines, respectively. As shown in Fig. 4 (b), the grain boundaries of these coarse grains and fine grains all belong to HABs. It further demonstrated that α-Mg grains occurred fully recrystallization during rolling,
further improving the mechanical properties and extending its engineering application, the AlN/AZ91 composite was rolled at 400 °C and a total reduction of 80 % was achieved. The microstructure, mechanical properties and tensile deformation mechanism of rolled AlN/ AZ91 composite sheets were investigated systematically. 2. Experiment In this work, commercial AZ91 casting magnesium alloy was chosen as the matrix alloy. The confirmed chemical composition was Mg9.13Al-0.74Zn-0.27Mn (wt. %). Nitrogen gas with high purity (99.999 %) was used as the reaction gas, and argon gas with high purity (99.999 %) was used to maintain an inert atmosphere in the reaction chamber. In-situ AlN/AZ91 composites with the volume fraction of 1.5 % were fabricated by the liquid nitriding method. The detailed fabrication processing of the liquid niriding method were described in Ref. [9]. Lastly, the melt was poured into the permanent mould and cast into the plates with a length of 100 mm, a width of 60 mm and a thickness of 20 mm. Before rolling, the as-cast AlN/AZ91 composites were firstly solid solution treated (SST) at 415 °C for 24 h. Then the SST AlN/AZ91 composite sheets were used for repeating rolling and annealing. The AlN/AZ91 composite sheets were rolled from 13 mm into 2.6 mm with the total thickness reduction of 80 %. The reduction of the first pass was 5 %, and the reduction of other pass was kept at 20 %. Before each rolling pass, the sheets were firstly annealed at 400 °C for 15 min and then rolled at the cold roller. The nitrogen content of rolled AlN/AZ91 composite sheets was measured by the oxygen nitrogen hydrogen analyzer (LECO ONH836) and the alloying element content was measured by the inductively coupled plasma optical emission spectrometer (Thermo Scientific iCAP 7000). To ensure the reliability, at least five samples were tested. The average composition of rolled AlN/AZ91 composite sheets was Mg1.25N-9.02Al-0.65Zn-0.23Mn (wt. %). The rolled specimens were mechanically ground, subsequently polished and etched in acetic picric solution (5 g picric acid, 5 ml acetic acid, 10 ml distilled water and 80 ml ethanol) at the room temperature for 12s. Optical microscopy (OM), high resolution scanning electron microscopy (HRSEM, FEI Nova 450 SEM) equipped with an energy dispersive spectrum (EDS) and electron backscattered diffraction (EBSD) were employed to observe the microstructure. The EBSD samples were electrochemical polished in an AC2 electrolyte at −30 °C for 90 s. The dislocations and dislocation/ particle interactions in rolled AlN/AZ91 composite sheets were also analyzed using transmission electron microscopy (TEM, FEI Talos F200X) in addition to the size and morphology of the particles. TEM foils were prepared by mechanically grinding them down to the thickness of 30 μm and further thinned using a precision ion polishing system (Gatan 691, USA). The tensile stress and strain of the specimens with thickness of 2 mm, width of 6 mm and gauge length of 25 mm, were measured in the SHIMADZU AG-X Tester at a strain rate 10−3s−1. The tensile specimens were prepared at the angles of 0° and 90° to the rolling direction by wire-cutting. For the tensile test at room temperature, five specimens were tested and the displacement was measured with an Epsilon axial extensometer. Macro-texture measurements of tensile specimens with different deformation were carried out by Bruker D8 Discover diffractometer with a high resolution area detector, operating at 35 kV and 40 mA. The pole figures were calculated using the quantitative texture analysis software MULTEX 3. 3. Results Fig. 1 (a) displays the size and surface quality of AlN/AZ91 composite sheets before and after rolling. It is obvious that the AlN/AZ91 composite sheets grew longer after 80 % rolling deformation and the width of AlN/AZ91 sheet had no significant increase. More importantly, as shown in Fig. 1 (a), there was no edge crack observed on the surface of rolled sheets, which proved that in-situ AlN/AZ91 composite sheets 2
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Fig.1. (a) The size and surface quality of AlN/AZ91 composite sheets before and after rolling, (b) the engineering stress-strain curves of as-cast and rolled AlN/AZ91 sheets, and (c) the comparison of the yield strength vs. Elongation to failure of various rolled magnesium matrix composites.
Table 1 Tensile properties of as-cast and rolled AlN/AZ91 composites. Materials As-cast AlN/AZ91 Rolled AlN/AZ91
Direction
σ0.2 (MPa)
σb (MPa)
ε (%)
TD RD
103 ± 6 351 ± 6 338 ± 5
240 ± 10 405 ± 9 387 ± 7
20 ± 3 10.5 ± 0.6 8.1 ± 0.3
and these coarse grains and fine grains in Fig. 4 (a) all were the dynamic recrystallized grains. In addition, it also can be found from Fig. 4 (b) that a lot of LABs were formed within α-Mg grains, which suggests that high density of dislocation substructures were produced during rolling. The number fractions of the LABs and HABs were estimated, as listed in Fig. 4 (c). Meanwhile, two dominate peaks can be observed in
Fig. 3. The (0001) pole figure (a) and inverse pole figure (b) analyzed from the upper surface of the rolled sample.
Fig. 2. The microstructures (a, c) and grain size distributions (d) of solid solution treated and rolled AlN/AZ91 composites, respectively, (b) the corresponding HRSEM graph of the rectangular region in (a). 3
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Fig. 4. (a) IPF map of rolled AlN/AZ91 composite sheets, (b) the corresponding orientation imaging map, and (c) the misorientation distribution calculated from EBSD in rolled AlN/AZ91 composite sheets.
Fig. 4 (c). The first high peak around 5° represents the high density of dislocation substructures and the second lower peak around 30° corresponded to the recrystallized grain boundaries [17]. The LABs and HABs were further analyzed by means of TEM. Fig. 5 (a) confirms that a lot of dislocation substructures were formed during multi-pass rolling, such as dislocation tangles and dislocation walls, etc. On the one hand, high density of dislocations were formed and tangled
around the uniform distributed in-situ AlN ceramic particles during rolling, while few dislocations were observed in the area without AlN ceramic particles. On the other hand, during rolling, by the annihilation and rearrangement of dislocations, the number of dislocations in the cell interiors had decreased and the cells became subgrains. The regular dislocation walls evolved into low angle grain boundaries (subgrain boundaries). Fig. 5 (b) shows the bright field image of HABs in rolled
Fig. 5. TEM images of LABs and HABs in rolled AlN/AZ91 composite sheets: (a) various dislocation substructures and (b) the migration of recrystallized grain boundary. 4
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Fig. 6. (a–c) FESEM images of γ-Mg17Al12 precipitates in rolled AlN/AZ91 composite sheets, the EDS analysis in (c) corresponding to continuous precipitates, (d) the size distributions of sub-micron γ-Mg17Al12 precipitates estimated by FESEM images.
cross. Combined with the SAED pattern, the black particles with submicro- and nano-size in Fig. 7 (a) were identified as the in-situ formed AlN ceramic particles. Similarly to the microstructure in Fig. 5 (a), high densities of dislocations were produced around in-situ AlN ceramic particles in Fig. 7 (a). More interestingly, it can be found in Fig. 7 (a) that a lot of nano-precipitates (marked with yellow arrows) were formed in the region of high densities of dislocations. Fig. 7(c–d) shows the TEM image of nano-precipitates in the vicinity of AlN ceramic particles and the corresponding SAED pattern. According to analysis of the diffraction rings in Fig. 7 (d), it can be identified that these densely distributed particles around in-situ AlN ceramic particles in Fig. 7 (c) were the γ-Mg17Al12 nano-precipitates. The size and density of γMg17Al12 nano-precipitates were estimated from the TEM images (at least five TEM images were used for analyzing). The average diameter of γ-Mg17Al12 nano-precipitates was 23.5 nm and the volume fraction of the γ-Mg17Al12 nano-precipitates was about 9.49 %. Fig. 8 (a) presents the bright field TEM image of rolled AlN/AZ91 composite sheets. γ-Mg17Al12 precipitates and in-situ AlN ceramic particles were dispersed uniformly in the magnesium matrix. The AlN/ Mg interface from the yellow rectangular region surrounded by dash lines in Fig. 8 (a) was characterized by HRTEM. Combined the corresponding fast Fourier transformed (FFT) SAED patterns, it can be concluded from Fig. 8 (b) that in-situ AlN ceramic particles and magnesium matrix possessed a strong interfacial bonding in AlN/AZ91 composite sheets. No crack initiation and propagation was detected in the AlN ceramic particles or at the AlN/Mg interface, which indicates dislocations can cut or bypass in-situ AlN particles instead of piling up during rolling, and the load can transfer from the soft magnesium matrix to the hard AlN ceramic particles effectively. The tensile fracture surfaces of the rolled AlN/AZ91 composite sheets along TD and RD were observed by SEM, as shown in Fig. 9. It is obvious that both samples along TD and RD exhibited the brittle-ductile mixed fractured feature, i.e. the cleavage planes and dimples were observed at the same time in the tensile samples along TD and RD.
AlN/AZ91 composite sheets. It can be found clearly in Fig. 5 (b) that the recrystallized grain boundaries were migrated from the dislocation free region to the dislocation rich region by grain boundary bulging during rolling. Before rolling, as shown in Fig. 2 (a), those coarse eutectic βMg17Al12 phases distributed along the grain boundaries [9] had been fully dissolved into α-Mg matrix. As shown in Fig. 6 (a), after multi-pass rolling, densely distributed precipitates were observed in AlN/AZ91 composite sheets again as well as little AlN clusters, which suggests that γ-Mg17Al12 phase was dynamically precipitated from the matrix during rolling. The morphology, size and precipitated type of γ-Mg17Al12 phase had a great difference from that of as-cast. Firstly, it can be seen from Fig. 6 (b) that there were two types of γ-Mg17Al12 precipitates: discontinuous precipitation and continuous precipitation. While most of the γ-Mg17Al12 precipitates were observed within α-Mg grains and only little γ-Mg17Al12 particles were distributed along the grain boundaries. The number fraction of discontinuous γ-Mg17Al12 particles precipitated on the grain boundaries was estimated by the SEM images as about 7 %. So the γ-Mg17Al12 particles precipitated during rolling in this study were mainly continuous precipitation. The enlarge FESEM image of these continuous precipitated γ-Mg17Al12 particles in α-Mg grains is shown in Fig. 6 (c). Evidently, the continuous precipitated γ-Mg17Al12 particles had a spherical morphology and a sub-micron size. In this study, the size of the continuous precipitated sub-micron γ-Mg17Al12 particles was estimated under SEM (at least five SEM images were used for estimating). The size distribution of the continuous precipitated γMg17Al12 particles is shown in Fig. 6 (d). It can be seen from Fig. 6 (d) that the size of the continuous precipitated γ-Mg17Al12 particles was mainly ranged from 300 nm to 800 nm and its average size was about 487.5 nm. In addition, the volume fraction of γ-Mg17Al12 precipitates was also estimated by the SEM images, as 6.36 %. Fig. 7 (a) exhibits the bright field TEM image of rolled AlN/AZ91 composite sheets. Fig. 7 (b) is the corresponding selected area electron diffraction (SAED) pattern of the black particle marked with the white 5
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Fig. 7. (a) Dislocations tangled by AlN ceramic particles and precipitated nano-Mg17Al12 particles around the AlN ceramic particles, (b) the SAED pattern of AlN ceramic particles, (c) TEM image of nano-Mg17Al12 precipitates and the corresponding SAED pattern (d).
cracking of the rolled sheets. For AZ91 alloys hot rolled at 400 °C, edge cracks began to occur when the accumulated reduction was about 50% and edge cracking became more severe with the accumulated reduction further increasing [18]. In this study, the AlN/AZ91 composite sheets were rolled at 400 °C with a total thickness reduction of 80 %. Importantly, no surface crack and edge crack was formed on the surface of rolled AlN/AZ91 composite sheets (Fig. 1a), which means that in-situ AlN ceramic particles can improve the rolling capability of AZ91 alloys significantly. Based on previous study [9], the ductility of as-cast AZ91 alloy acquired a remarkable enhancement due to the formation of insitu AlN ceramic particles, and this made AlN/AZ91 composites possess large deformation ability. So the rolling capacity improvement of AZ91 alloys may originate from its ductility enhancement due to the formation of in-situ AlN ceramic particles. Additionally, the interfacial bonding between AlN ceramic particles and magnesium matrix plays the critical role in the rolling capability. The poor interface easily acts as the source of cracking, resulting in the crack initiation in the early of rolling. In this case, the rolled sheets will fracture with a small thickness reduction. In this study, the AlN/Mg interfaces had a strong bonding (Figs. 8 and 9), which well realized the load transfer from the matrix to
However, by contrast, it can be found that all the amount, size and depth of dimples on the fracture surface along TD are much larger than that along RD. It suggests the rolled AlN/AZ91 composite sheets along TD have a better ductility than that along RD, which is consistent with the mechanical properties in Table 1. In addition, some particles with diameter of sub-micron, which were identified as γ-Mg17Al12 precipitates by the EDS analysis (the inset in Fig. 9a), were observed inside the dimples as indicated by the yellow arrows in Fig. 9 (a). Thus, it can be concluded that cracks were initiated at the γ-Mg17Al12 precipitates rather than in-situ AlN ceramic particles, which is agree well with the result of Fig. 8. 4. Discussion 4.1. Effect of in-situ AlN ceramic particles on the rolling capability of AZ91 alloys The rolling capability of materials can be characterized by the degree of edge cracking. In general, increasing the rolled thickness reduction and decreasing the rolling temperature will promote the edge 6
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Fig. 8. (a) The bright field TEM image of rolled AlN/AZ91 composite sheets, (b) the HRTEM image of the AlN/Mg interface in rolled AlN/AZ91 composite sheets and the corresponding fast Fourier transformed SAED patterns from the rectangular regions.
Fig. 9. Fracture surfaces of rolled AlN/AZ91 composite tensile sample along TD (a) and RD (b), the EDS analysis result corresponding to the particles in the dimples (yellow arrows). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
And continuous precipitates nucleate in the interior of grains, which depends on the volume diffusion [19]. In other words, if the volume diffusion is accelerated, the continuous precipitation will be promoted. In this study, due to the formation of in-situ AlN ceramic particles, high densities of dislocations were produced and tangled around the AlN ceramic particles during rolling (Figs. 5 and 7a). As shown in Figs. 5, 7a and 8a, most of in-situ formed AlN ceramic particles were uniformly dispersed in the interior of α-Mg grains. So these high densities of dislocations around AlN ceramic particles were located in the interior of α-Mg grains. On the one hand, these high densities of dislocations can accelerate the diffusion of aluminum atoms and promote the precipitation of γ-Mg17Al12 phase. On the other hand, dislocations can act as the heterogeneous nucleation sites for γ-Mg17Al12 precipitates [20] and promote its precipitation. Hence, most of γ-Mg17Al12 precipitates formed during rolling were continuous precipitates in this study.
Table 2 The volume fractions and average sizes of submicro- and nano-AlN ceramic particles.
Submicro-AlN Nano-AlN
Average size (nm)
Volume fraction (%)
231.2 23.7
0.62 0.71
AlN ceramic particles. As a result, no crack initiated at the AlN/Mg interfaces (Fig. 8) during rolling, which also contributed to the large rolling reduction and good rolling capability.
4.2. Effect of in-situ AlN ceramic particles on the precipitation behavior of AZ91 alloys For AZ91 alloys, most of γ-Mg17Al12 precipitates produced during hot rolling at 400 °C were distributed along the grain boundaries [18]. However, in this study, about 93 % (number fraction) γ-Mg17Al12 precipitates were continuous precipitates (Fig. 6b). It demonstrated that in-situ AlN ceramic particles promoted the continuous precipitation of γ-Mg17Al12 phase during multi-pass rolling. As is known, discontinuous precipitates nucleate on the grain boundaries, so the discontinuous precipitation is favored when the grain boundary diffusion is dominant.
4.3. Strengthening mechanisms analysis In this study, after multi-pass rolling, the yield strength of as-cast AlN/AZ91 composite achieved a significant improvement (Fig. 1b and Table 1). The significant increase in yield strength was caused by grain boundary strengthening, precipitation strengthening and dislocation strengthening. Firstly, during rolling, the grain size of α-Mg was refined owing to 7
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Fig. 10. (a) The increment comparison of yield strength (σ0.2) in AlN/AZ91 composites due to grain boundary strengthening, precipitation strengthening and dislocation strengthening, respectively, (b) the theoretical and experimental values of the increased yield strength in AlN/AZ91 composites during rolling.
AlN ceramic particles due to the residual plastic strain between AlN ceramic particles and AZ91 matrix during rolling [26], which can effectively increase the yield strength of AlN/AZ91 composites. The increased yield strength Δσd due to dislocation strengthening can be calculated by the following equation [26]:
the dynamic recrystallization. For the grain boundary strengthening, the Hall-Petch relationship is applied and the strength increment Δσg is calculated as follows: −1
−1
2 Δσg = k (drolled − das −2cast )
(1) 1/2
for AZ91 Where k is the Hall-Petch coefficient and is 0.13 MPa m alloy, drolled and das − cast are the average grain size of the rolled (25 μm) and as-cast (50 μm [9]) AlN/AZ91 composites, respectively. According to calculating, the increased yield strength due to grain boundary strengthening is about 7.62 MPa. Secondly, densely distributed γ-Mg17Al12 precipitates formed during rolling can give rise to significant strengthening effect by Orowan mechanism. In this study, most of γ-Mg17Al12 precipitates belong to continuous precipitation. Compared with discontinuous precipitation, continuous precipitated γ-Mg17Al12 precipitates inside the grains are more effective for inhibiting the dislocation motion [19–22], which resulted in the high strength contribution of the precipitation strengthening. The increased stress for the slip dislocations to bypass γMg17Al12 precipitates is described by the Orowan equation [23,24]:
Δσp =
dp Gb ln r0 2πλ 1 − ν
Δσd =
Gb 2π 1 − ν (
0.779 f
− 0.785) dt
(2)
ln
0.785dt b
12ΔTΔαV bd
(4)
where β is the strengthening coefficient (is 1.25), ΔT is the temperature difference between rolling processing and mechanical tests, Δα is the difference in thermal expansion coefficient between AlN ceramic and AZ91 matrix (is 26.3 × 10−6 K−1 [9]), V is the volume fraction of AlN ceramic particles and d is the average size of AlN ceramic particles. In this study, although in-situ formed AlN ceramic particles had a multiscale size, the fractions of submicro- and nano-AlN ceramic particles were much more than that of micro-AlN ceramic particles [9]. The volume fractions and average sizes of submicro- and nano-AlN ceramic particles in AlN/AZ91 composites were estimated from the TEM and SEM images, which are listed in Table 2. Accordingly, the increased yield strength due to dislocation strengthening from submicro- and nano-AlN ceramic particles are 32.73 MPa and 109.38 MPa, respectively. The micro-sized AlN ceramic particles had the lower volume fraction (~0.17 %) and the larger size, which resulted in a weak effect of dislocation strengthening. Thus, the dislocation strengthening from the micro-AlN ceramic particles can be ignored in this study. Above all, the increased yield strength due to dislocation strengthening from insitu AlN ceramic particles is 142.11 MPa. Fig. 10 (a) shows a comparison of the contributions of grain boundary strengthening, precipitation strengthening and dislocation strengthening to the yield strength increment of rolled AlN/AZ91 composites. Evidently, the increased yield strength of AlN/AZ91 composites after rolling is mainly attributed to the dislocation strengthening (~142.11 MPa) and the precipitation strengthening (~95.47 MPa), which means in-situ AlN ceramic particles effectively improved the strength during rolling via increasing the dislocation density and promoting the continuous precipitation of γ-Mg17Al12 phase. While the continuous precipitated γ-Mg17Al12 particles within grains can not pin the grain boundary, so the dynamically recrystallized grains had a larger size, resulting in the low contributions for yield strength increment (~7.62 MPa). Furthermore, the increased yield strength during rolling can be predicted by the following equation:
Where Δσp is the increased stress due to the precipitation strengthening, G is the shear modulus (is 15 GPa for magnesium), b is magnitude of the Burgers vector of the slip dislocations (is 0.32 nm), ν is the Poisson's ratio (is 0.29 for magnesium), λ is the effective planar inter-precipitate spacing, dp is the mean planar diameter of the precipitates and r0 is the dislocation core radius (commonly, the assumption r0 = b is used). For the spherical precipitates (Fig. 7c), the effective planar inter-precipitate spacing λ and the mean planar diameter of the precipitates dp can be calculated according to the method of J-F. Nie et al. [25], and Eq. (2) can be written as:
Δσp =
3 βGb
(3)
where f is the volume fraction of γ-Mg17Al12 precipitates and dt is the average diameter of the spherical γ-Mg17Al12 precipitates. In this study, dynamic precipitated γ-Mg17Al12 precipitates had nano and sub-micron two sizes. Firstly, for γ-Mg17Al12 submicro-precipitates, the average diameter and the volume fraction were 487.5 nm and 6.36 %, respectively. According to calculating, the Orowan strengthening from submicro-precipitates is 5.72 MPa. Secondly, dynamic precipitated nanosized γ-Mg17Al12 precipitates had a finer size (dt = 23.5 nm) and higher content ( f = 9.49 %). In this case, the Orowan strengthening from γMg17Al12 nano-precipitates is 89.75 MPa. Thus, the increased yield strength due to γ-Mg17Al12 precipitates is 95.47 MPa. Thirdly, high densities of dislocations were formed around in-situ
ΔσTheory = Δσg + Δσp + Δσd
(5)
where ΔσTheory is the theoretical predicted value of the yield strength increment during rolling. By adding up the contributions of grain boundary strengthening, precipitation strengthening and dislocation strengthening, the theoretical predicted value ΔσTheory is 245.2 MPa. The theoretical predicted value ΔσTheory and the experimental obtained 8
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Fig. 11. (0001), (10-10), (10–11) and (11–22) pole figures of rolled AlN/AZ91 samples after 0 % (a, d g, j), 3%(b, e, h, k) and 10 % (c, f, i, l) tensile deformation along TD.
and pyramidal II plane (10–12) pole figures of rolled AlN/AZ91 composite samples along TD and RD with different tensile strain, respectively. It can be seen from the basal pole figures in Figs. 11 and 12 that the basal plane texture exhibits an ellipsoidal intensity distribution around the pole figure center and no intensity distribution at the north and south poles, which means that the major deformation mechanism of rolled AlN/AZ91 composite sheets at room temperature is dislocation slip rather than twinning [27]. Generally, the twinning behavior is influenced by grain size, texture and second-phase. Firstly, some studies reported that twinning has the positive grain size effect [28,29], i.e. twinning forms more difficult in finer grains. Secondly, C.M. CepedaJiménez et al. [30] confirmed that the twinning behavior during inplane tension of rolled magnesium alloys was due to the local stress accommodation rather than texture effect. Besides, X. Li and J.B. Clarke et al. [31,32] found a high level of Mg17Al12 precipitates could suppress the twinning behavior during deformation. In this study, the grain size
values along TD and RD are compared in Fig. 6 (b). The theoretical predicted values ΔσTheory (245.2 MPa) is consistent well with the experiment value along TD ΔσExperiment − TD (248 MPa) and RD ΔσExperiment − RD (235 MPa), and exactly predict the increased yield strength of AlN/AZ91 composites during rolling. 4.4. Tensile deformation mechanism analysis As shown in Fig. 1 (c), the rolled AlN/AZ91 composite sheets possessed a better ductility among the various rolled magnesium matrix composite sheets. The elongation to fracture is even much higher than that of AZ31 (typical wrought magnesium alloy) matrix composite sheets. To study the reasons for high ductility of rolled AlN/AZ91 composite sheets, the tensile deformation mechanism at room temperature was studied in this section. Fig. 11 and Fig. 12 display the basal plane (0001), prismatic plane (10-10), pyramidal I plane (10–11) 9
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Fig. 12. (0001), (10-10), (10–11) and (11–22) pole figures of rolled AlN/AZ91 samples after 2.5 % (a, d g, j), 4 % (b, e, h, k) and 7 % (c, f, i, l) tensile deformation along RD.
was decreased effectively (Fig. 2d) and a high density of γ-Mg17Al12 precipitates (Figs. 6b and 7c) were obtained during rolling. Based on above analysis, it can be concluded that the weak twinning mode of rolled AlN/AZ91 composite sheets during tensile deformation may mainly attributed to the finer grain size and high density of γ-Mg17Al12 precipitates. During tensile deformation, the slip planes tend to rotate to lie parallel to the tensile axis, so the activated slip systems can be identified indirectly by analyzing the evolution of texture intensity during deformation. For the tensile axis along TD, as shown in Fig. 11, the maximum texture intensity of the basal plane increased slightly from 3.58 to 3.64 and 3.69 with increasing the tensile strain to 3 % and 10 %, which increased by 1.7 % and 3.1 % respectively. It suggests that a small number of basal planes were initiated and rotated to parallel with
the tensile axis during tension. In contrast, the maximum intensity of the prismatic plane decreased by 15.8 % and 18.1 % with increasing the tensile strain, which means that the prismatic plane slip was suppressed and the prismatic planes were gradually far away from the tensile axis during tension. Surprisingly, it can be found in Fig. 11 that both the maximum intensities of pyramidal I plane and pyramidal II acquired a significant improvement. The maximum intensities of the pyramidal I plane and pyramidal II plane increased by 7.4 % and 5.2 % when the tensile strain increased to 3 %. With further increasing the tensile strain to 10 %, the maximum intensities of the pyramidal I plane and pyramidal II increased by 21 % and 14 %, which must be the result of pyramidal dislocation activity in favorably oriented grains. It demonstrates that profuse pyramidal plane slip systems were activated when the rolled AlN/AZ91 composite sheets tension along TD. 10
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Fig. 13. The variation of the maximum texture intensities as a function of tensile strain along TD (a) and RD (b).
Similarly, for the tensile axis along RD, the pole figures of basal plane, prismatic plane and pyramidal plane with different tensile strain are shown in Fig. 12. According to analyzing the evolution of the maximum intensities during tension, similar conclusions were obtained for the tension along RD. With increasing the tensile strain, the maximum intensities of the basal plane and prismatic plane decreased gradually, which means the basal plane slip and prismatic plane slip were suppressed. In contrary, the maximum intensities of the pyramidal I plane and pyramidal II got a larger increase. Thus, the tensile deformation along RD was mainly attributed to the pyramidal slip. In general, the dominate slip systems at room temperature in magnesium are the basal plane slip system [33]. However, the basal plane slip only provide two independent slip systems [27], which fails to satisfy the Taylor criterion [34], resulting in the poor ductility of magnesium alloys and magnesium matrix composites. Fig. 13 shows the maximum intensities variation of different slip planes with the tensile strain along TD and RD. It can be seen from Fig. 13 that no matter the tensile axis along TD or RD, only the maximum intensities of pyramidal plane (pyramidal I and pyramidal II) increased significantly with the tensile strain. This suggests that pyramidal slip dominated the tensile deformation of rolled AlN/AZ91 composite at room temperature. Similar result was also obtained in rolled Mg–Li alloys by analyzing texture evolution using elasto-plastic self-consistent (EPSC) model [35]. The activity of pyramidal plane slip increased the number of slip systems and made the Taylor criterion get satisfied [27]. Hence, a high elongation to fracture was obtained in the rolled AlN/AZ91 composite sheets (Fig. 1c). Furthermore, it should be noted in Fig. 13 (a) that the maximum intensities of pyramidal I plane and pyramidal II plane almost increased at constant rate with increasing the tensile strain. However, as shown in Fig. 13 (b), with increasing the tensile strain, the maximum intensities of pyramidal I plane and pyramidal II plane first increased and then almost remained unchanged. This suggests that the number of favorably oriented grains for activating pyramidal slip along RD was much less than that along TD during later deformation, which may lead to the less ductility in RD (Table 1). The limited activity of basal plane slip and prismatic plane slip was mainly attributed to the unfavorable grain orientation. In rolled AlN/ AZ91 composite sheets, the grains were orientated such that their caxes were nearly normal to the sheet plane with a tilt along RD, and the prismatic planes exhibited a dominate orientation normal to TD (Fig. 3), which lead to the grain in a unfavorable orientation and formed the lower Schmid factors for basal plane slip and prismatic plane slip. But, due to the low critical resolved shear stress (CRSS) for basal plane slip, basal plane slip may be activated in few grains with special orientation, so the maximum intensities of basal plane had a
slight increase in Fig. 11. Since no non-basal slip was observed in rolled AZ91 alloys during tension at room temperature [18], it can be safely concluded that the pyramidal plane slip systems activity in rolled AlN/ AZ91 composite sheets is due to the formation of in-situ AlN ceramic particles. The formation of in-situ AlN ceramic particles may reduce the CRSS of pyramidal plane slip or the intrinsic stacking fault I1 energy, resulting in the activity of pyramidal plane slip systems [36,37]. Further research works need to be carried on for exploring the underlying mechanisms in future. 5. Conclusions In this work, in-situ AlN/AZ91 composite sheets fabricated by liquid nitridation method were subjected to multi-pass rolling with the aim of improving mechanical properties and extend its industrial application. The microstructure, mechanical properties and tensile deformation mechanism of rolled AlN/AZ91 composite sheets were investigated, and the obtained main conclusions were summarized as follows: (1) In-situ AlN/AZ91 composites had excellent rolling capacity. In situ AlN/AZ91 composite sheets can be rolled at 400 °C to the 80 % thickness reduction and no crack was formed on its surface. The excellent rolling capacity was mainly attributed to the high as-cast ductility and good interfacial bonding; (2) In this study, dynamically precipitated γ-Mg17Al12 particles were mainly belong to continuous precipitation. During rolling, high density of dislocations were formed near AlN ceramic particles, which distributed uniformly within grains, promoting the continuous precipitation of γ-Mg17Al12 phase; (3) By rolling, the mechanical properties of AlN/AZ91 composites realized a better strength-ductility combination compared to other rolled magnesium matrix composites. The increased strength was mainly attributed to the dislocation strengthening and precipitation strengthening; (4) The dominant tensile deformation mechanism of rolled AlN/AZ91 composite sheets was the pyramidal plane slip. The formation of insitu AlN ceramic particles activated the pyramidal dislocations, leading to a higher ductility in rolled AlN/AZ91 composite sheets. Acknowledge The authors are grateful for financial support for this research by the National Natural Science Foundation of China (Grant No. 51771151), the National Key R&D Program of China (Grant No. 2017YFB0305100), the Research Fund of the State Key Laboratory of Solidification Processing, China (Grant No. 2019-TS-02) and the Innovation 11
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Foundation for Doctor Dissertation of Northwestern Polytechnical University, China (Grant No. CX201905).
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