Effect of pre-deformation on microstructure and mechanical properties of WE43 ma

Y.H. Kng, Z.H. Hung, S.C. Wng, H. Yn, R.S. Chen,∗, J.C. Hung

a Guangdong-Hong Kong Joint Research and Development Center on Advanced Manufacturing Technology for Light Alloys, Guangdong Institute of Materials and Processing, Guangzhou 510650, China

b Institute of Metal Research Chinese Academy of Sciences, Shenyang 110016, China

c Hong Kong Institute for Advanced Study; Department of Materials Science and Engineering, City University of Hong Kong, Hong Kong, China

Abstract In this work, the microstructural evolution and mechanical properties of a pre-deformed WE43 magnesium alloy when aged at 250 and 300 °C were further investigated. It is found that the abundant deformation twins introduced by pre-deformation were maintained within the alloy during the aging treatment. Second particles formed at the twin boundaries and coarsened with aging time, especially at 300 °C.When peak-aged at 250 °C, the fine metastable β′′′and β′precipitates formed in the un-deformed alloy have been transformed into relatively large β1 and β precipitates by the pre-deformation. While peak-aged at 300 °C, the pre-deformation obviously refined the β precipitates.Mechanical properties indicate that pre-deformation can increase the yield strength by 19MPa and 54MPa for the peak-aged alloy at 250 °C and 300 °C, respectively, and will not obviously deteriorate the tensile elongations.

Keywords: WE43 magnesium alloy; Pre-deformation; Aging; Microstructure; Mechanical properties.

1. Introduction

The precipitation strengthened WE43 alloy (Mg–Y–Nd system) owns a high specific strength, creep resistance and corrosion resistance, and has become the most successful commercial creep-resistant Mg alloy for aerospace and transportation industries [1–4]. The high performance is achieved by an artificial aging treatment to form a uniformly dispersed fine precipitates with good thermal stability within the Mg matrix [5–9]. The precipitation behavior of the Mg–Y–Nd system, including the Mg–Y and Mg–Nd system alloys has been extensively investigated [5–7,10–13]. Particularly,the early stage of precipitation has become a hot research topic recently in order to further optimize the microstructure and mechanical properties of these alloys by heat treatment[14–17]. Furthermore, various plastic deformation processes have been developed for the WE43 alloy to further improve the mechanical properties (e.g. strength and ductility) by refining grain size, and thus to expand its applications as a high performance wrought Mg alloy [18–28].

Fig. 1. Schematic diagram of the cold impact forging (i.e. pre-deformation): the forging direction was changed in the sequence (i.e. x→y→z) and a deformation strain of 5% at each direction was employed.

It is well known for precipitation hardened Al alloys like Al–Cu and Al–Zn–Mg system that a cold work (i.e. predeformation) prior to aging would exert noteworthy effect on the precipitation behavior and relevant mechanical properties[29–32]. This is attributed to the generation of dislocations,which could enhance both the nucleation and diffusivities of precipitation-forming element in alloy[33–35].It has been reported for an Al–Cu–Li–Mg–Ag alloy that a pre-stretch step promoted the formation of T1precipitate at the expense of δ’,θ’, and S’ precipitates, resulting in an enhancement of precipitation strengthening [30,36]. For an Mg–Gd–Nd–Zr alloy,it has also been found that a pre-stretch step promoted the nucleation of β1phase at 200 °C [37]. The similar promotion of β1nucleation by a cold work prior to aging at 200–250°C has also been observed in the WE54 and WE43 alloys[7,38]. It is worth noting that deformation twinning is readily activated for Mg alloys due to their hexagonal close-packed(HCP) structure, different from Al alloys with face-centered cubic (FCC) structure [39,40]. In ductile Mg–Gd and Mg–Zn alloys, the Gd and Zn solute atoms were observed to segregate at twin boundaries (TBs) [41]. Even some second phases can be formed at TBs in pre-deformed Mg–Nd–Zn–Zr alloy[42]. These TBs decorated with solute atoms and precipitates were suggested to effectively hinder migration of these TBs and glide of dislocations, and thus leaded to a twin boundary hardening [41,43,44].

Recently, it is reported that cold impact forging in multidirections with a small deformation strain is efficient to produce deformation twins (DTs) and dislocations in Mg alloys[45,46]. Therefore, in our previous work [38], we have investigated the effects of the cold impact forging (i.e. predeformation) on the microstructure and mechanical properties of the WE43 alloy when aged at 200 °C. It has been found that a high density of DTs and dislocations was produced within the alloy. After peak-aged at 200 °C, the alloy strength was significantly improved due to the addition of dislocation/twin boundary strengthening and the enhancement of precipitation hardening. However, the tensile elongation was largely decreased. Thus, in present work we intend to further explore the microstructural evolution and mechanical properties of the pre-deformed WE43 alloy when aged at 250 °C and 300 °C in order to optimize the mechanical properties.

2. Experimental procedures

The current experimental WE43 alloy has a chemical composition of Mg–4.38Y–2.72Nd–1.10Gd–0.56Zr (wt.%, or Mg–1.27Y–0.49Nd–0.18Gd–0.16Zr (at %)), which is analyzed by inductively coupled plasma atomic emission spectroscopy(ICP-AES).The alloy melt was prepared by pure Mg(>99.95%), Y (>99%), Nd (99%), Gd (99%), and the Mg–30Zr master alloy (all in wt.%) in an electric resistance furnace with a stainless steel crucible at about 780 °C under the protection of an anti-oxidizing flux.After holding the melt for 0.5h, the melt was poured into a preheated mild steel mold to obtain the as-cast alloy ingot. In an electric furnace under Ar blowing, the ingot was solution treated at 525 °C for 8h and then quenched into an 80 °C hot water (i.e. the T4 condition).After solution treatment, a dimension of 50×50×50mm of cubic block specimen was machined from the T4 alloy ingot and subjected to a cold impact forging (i.e. pre-deformation).The process was performed by one impact on various three directions with 5% deformation strain per direction as the schematic diagram shown in Fig. 1. The total forging strain is 15%. The condition after the pre-deformation was assigned as the T3 condition. Then, small samples were taken from the center section of the T3 specimen for microstructure observation, artificial aging treatment, Vickers hardness and tensile tests. These samples were taken from locations 5mm apart from the surface. Both the pre-deformed T3 and un-deformed T4 alloy were aged at 250 °C and 300 °C in an electric furnace under Ar blowing.

The alloy microstructures were examined by optical microscope (OM), scanning electron microscope (SEM, Philips XL30 ESEM-FEG/EDAX) and transmission electron microscope (TEM, JEOL 2010). The OM samples were etched in a solution of 4vol% HNO3after mechanical polishing, with no etching for the SEM samples. The thin foil specimens for TEM were prepared by a precision ion polishing system(Leica EM RES101) operated at 4.5kV accelerating voltage and 15° incident angle. A 2kg load and holding time of 25s was applied for Vickers hardness (Hv) measurement. Ten Hv hardness measurements were conducted to obtain the average value and standard deviation for each condition. For tensile testing, samples have a gage size of 10×3.5×2.5mm(length × width × thickness), and were tested at room temperature with a strain rate of 1×10−3s−1.For each condition,three tests were performed to obtain the average mechanical properties.

3. Results and discussion

3.1. Age-hardening behavior

Fig. 2. Age-hardening curves of the pre-deformed WE43 alloy compared with the undeformed alloy at 250 °C and 300 °C: T3: pre-deformed alloy;T82 and T83: peak-aged of the pre-deformed T3 alloy at 250 °C (2h) and 300 °C (2h), respectively; T4: un-deformed alloy; T62, T6 and T63: peakaged of the un-deformed T4 alloy at 250 °C (4h, 16h) and 300 °C (2h),respectively.

The age-hardening curves of the pre-deformed T3 WE43 alloy, in comparison with the un-deformed T4 alloy, artificially aged at 250 °C and 300 °C are shown in Fig. 2. The initial hardness value of the T4 alloy is 76 Hv. In comparison, the pre-deformation process largely raises the hardness of the T3 alloy up to 89 Hv, an increment of 17%. During isothermal aging at 250 °C, the hardness of the un-deformed T4 alloy increases gradually to a peak hardness of 97 Hv after 4h (the T62 condition). In contrast, the pre-deformed T3 alloy reaches a maximum hardness of 95 Hv in 2h (the T82 condition). This indicates although the age-hardening response of the WE43 alloy is accelerated, the maximum hardness achievable at 250 °C is not significant affected by the pre-deformation. However, it is not consistent with the result reported by Nie and Muddle [7] that the maximum hardness of WE43 alloy has been largely increased by a 6% prestretch strain prior to aging at 250 °C. They further obtained that the maximum hardness was not increased with the prestretch strain increase to 12%, although the age-hardening response was further accelerated. In present work, a different pre-deformation process of multi-directional cold free forging and a lager pre-strain of 15% were performed for the WE43 alloy. This could lead to a higher density of DTs and dislocations. Especially, a partial decomposition of the supersaturated magnesium solid solution has occurred within the interior grains and relatively coarse second particles formed at the twin boundaries (TBs) as seen from our previous work of [38]. This explains the large hardness increase by 13 Hv in the T3 alloy. However, the formation of coarse second particles at the TBs could consume a large of solute due to the abundant TBs. As a result, only a slight hardness growth was obtained in the T3 alloy when aged at 250 °C, and finally no hardness increases in the pre-deformed T82 alloy when compared to the un-deformed T62 alloy. With prolonged aging at 250 °C, the hardness of the un-deformed T4 alloy begins to pronouncedly decrease, and the conventional standard T6 treatment (aging for 16h) becomes over-aged with a hardness of 90 Hv in present work [6]. In contrast, the hardness of the pre-deformed T3 alloy has only a little decrease and retains stable up to 144h before the occurring of a significant hardness decrease.

During isothermal aging at 300 °C, both the un-deformed T4 and pre-deformed T3 alloys only exhibit a little hardness increase after 2h. The conditions are assigned as the T63 and T83, respectively. The T83 alloy has a hardness of 90 Hv,which is much larger than that of the T63 alloy (78 Hv).However, during the over-aging time the hardness of the predeformed T3 alloy begins to largely decrease. Nevertheless,the pre-deformed T3 alloy achieves a much higher hardness than the un-deformed T4 alloy during the aging.

3.2. Microstructural evolution

During the artificial aging at 250 °C, the microstructural evolution of the pre-deformed alloy from aging time of 2–258h, revealed by OM and SEM, is presented in Fig. 3. After solution treatment, the T4 alloy obtains a single solute supersaturated Mg matrix with an equiaxed grain structure of average grain size of 184 μm. In addition to a small amount of Zr particles for grain refinement, a low number of Y-rich cuboidal particles within the Mg matrix are not dissolved, as shown in Fig. 2b and the inset of a high magnification view.

By the pre-deformation process, a large number of{10¯12}<10¯1¯1> extension twins and basal 〈a〉dislocations are introduced into the T3 alloy. Furthermore, a small number of spheroid-shaped particles discontinuously formed at the twin boundaries (TBs). In addition, a portion of fine metastable β′′and β′precipitates dynamically formed within the matrix [14,24]. It is worth noting that in pervious works the metastable β′phase formed in the Mg–Y–Nd system alloys is assigned as β′′′ phase by Solomon [13] in his recent work. The microstructure evolution during the predeformation process has been presented in detail in our previous paper of [38]. It is observed in Fig. 3 (h–j) that the DTs are retained within the alloy for prolonged aging time of 258h, the longest preformed in this work. Particularly, with the prolonging of aging time the number density of particles formed at TBs increases and meanwhile they have a growth especially during the over aged stage.

The precipitates characterized by TEM under incident beam near [0001]Mgor [0001]Twinwithin the pre-deformed T82 (peak aged) alloy compared with the un-deformed T62(peak aged) and T6 alloy are presented in Fig. 4. In the undeformed T62 alloy, a uniform dispersion of fine plate-shaped and globular precipitates are observed, and more than two of plate-shaped precipitates are always associated with a globular precipitate(Fig.4a).According to previous detailed investigations on the precipitation behavior of Mg–Y–Nd system alloys at 250 °C, these two can be indicated as metastable plateshaped β′′′and globular β′phase, respectively [5,6,12,24].

Fig. 3. OM and SEM of the pre-deformed (T3) WE43 alloy aged at 250 °C for (a–c) 2h, (d–f) 16h, and (h–j) 258h. The inset in (b) indicates a high magnification view of a Y-rich cuboidal particle.

In the conventional standard T6 alloy, a uniform dispersion of plate-shaped β′′′ and globular β′ precipitates are also observed, but with a larger size than that in the T62 alloy (Fig. 4b). Additionally, a small number of relatively coarse plate-shaped precipitates (indicated by a white arrow in Fig. 4b) are observed, and its two ends attach to the globular β′precipitates. It can be characterized as metastable β1phase based on the investigation of Nie and Buddle [5]. The microstructure of the T6 alloy is consistent with previous reports on WE43 (T6) alloy [6,24]. The growth of β′′′and β′strengthening precipitates and especially the formation of relatively coarse β1precipitates occur during the overaging of the T6 alloy. In contrast, within the pre-deformed T82 alloy the precipitates are mainly in plate-shape (Fig. 4c). A large view of such precipitates within the DTs and original Mg matrix, together with their corresponding selected area electron diffraction (SAED) patterns, are presented in the Fig. 4d, f and Fig. 4e, g, respectively. The SAED patterns indicate the plate-shaped precipitates are β1and β phases [47,48]. We have observed that the β1precipitates form heterogeneously on dislocations introduced by the pre-deformation when the WE43 alloy is aged at 200 °C [38]. Nie [7] has also observed that the dislocations introduced by a cold work would promote the nucleation and growth of β1precipitate plates at the expense of β′phase in WE54 while aged at 250 °C. They reported that the β1phase in the WE54 alloy is similar to the intermediate precipitate phase Mg3X (X=Nd or Ce) in Mg–Nd and Mg–Ce alloys, both nucleating preferentially on dislocations. However, in present work, the pre-deformation has further promoted the formation of equilibrium β phase. This is due to a larger deformation strain performed by the predeformation process in this work, and it leads to the in-situ transformation of metastable β1phase into the equilibrium β phase in the pre-deformed T82 alloy.

Fig. 5 shows the pre-deformed alloy microstructures aged at 300 °C for 0.5 and 96h by OM and SEM. It is found that the DTs are thermal stable up to the aging time of 96h,which is the longest aging time preformed in this work at 300 °C. It is worth noting that the precipitates formed at twin boundaries become coarsened rapidly at the higher aging temperature.Free precipitate zones around the TBs are observed. It has been reported that, while aging at 300 °C, the equilibrium β precipitates would immediately form in the Mg–Y–Nd system alloys [49]. The TEM micrographs under the incident beam near [0001]Mgwith SAED pattern of the pre-deformed T83(peak aged) alloy, compared with that of the un-deformed T63 (peak aged), and their SEM micrographs of the overaged alloy (96h), are shown in Fig. 6. The SAED patterns indicate the plate-shaped precipitates both in the T63 and T83 alloys are equilibrium β phase. It is worth noting that the pre-deformation has obviously refined the β precipitates. A summary of precipitate with its width and thickness presented in the WE43 alloy under various conditions are indicated in Table 1.

3.3. Mechanical properties

Fig. 4. TEM bright-field images with corresponding selected area electron diffraction (SAED) pattern under the incident beam near [0001] of the Mg matrix(a, b) and (c) twin showing the effect of pre-deformation on the precipitates of WE43 alloy at 250 °C: (a, b) un-deformed alloy aged for 4h (T62) and 16h(T6), respectively; (c) pre-deformed alloy aged for 2h (T82), (d) and (e) a large view of precipitates within the square A (twin) and B matrix) in (c) for T82,respectively; (f) and (g) a corresponding SAED pattern of the twin and matrix, respectively.

Fig. 5. OM and SEM of the pre−deformed (T3) WE43 alloy aged at 300 °C for (a, b) 0.5h and 96h (c, d).

Fig. 6. (a, c) TEM micrographs under the incident beam near [0001]Mg with corresponding selected area electron diffraction (SAED) pattern (inset), and (b,d) SEM images showing the effect of pre-deformation on the precipitates of the WE43 alloy at 300 °C: (a, b) un-deformed alloy aged for 2h (T63) and 96h,respectively; (c, d) pre-deformed alloy aged for 2h (T83) and 96h, respectively.

Table 1 Summary of precipitates present in the WE43 alloy under various conditions. The width and thickness of precipitate (in unit of nm) are indicated in this Table.

Table 2 Tensile properties of WE43 alloy in different conditions at room temperature (YS: yield stress (MPa); UTS: ultimate tensile stress (MPa); El: elongation).

Typical engineering stress−strain curves and mechanical properties of the pre-deformed WE43 alloy, compared with the un-deformed alloy, under different conditions at ambient temperature are shown in Fig. 7 and Table 2. The undeformed T4 alloy has a tensile yield strength (YS) and ultimate tensile strength (UTS) of 145MPa and 204MPa, respectively. In contrast, for the pre-deformed T3 alloy, a pronounced increment in YS and UTS of 63MPa and 55MPa,reaching 208 and 259MPa has been achieved by the predeformation. However, the elongation has been degraded by approximate 50%. Based on the microstructural evolution, the strength enhancement by the pre-deformation is attributed to the generation of many dislocations and TBs, plus a small amount of precipitates. These dislocations will strengthen the alloy by interacting with themselves and impeding their own movements (i.e. dislocation strengthening) [50]. For the effect of TBs on the material strength, three kinds of effects could be performed based on the properties of TBs analyzed by Konopka and Wyrzykowski [51]. The first case is that TBs would contribute slightly to the strength because the TBs constitute a weak barrier to the dislocation movement. In the second case, TBs can largely strengthen the material due to the strong oppose against dislocation movement. Oppositely,the third case is that TBs would decrease the strength for that they become effective dislocation sources to facilitate the dislocation slip. In pure Mg contained a large number of DTs,there is no apparent hardening after annealing treatment [52].However, a strong pinning effect on the further migration of TBs and dislocation slip was observed in a pre-deformed ductile Mg–Gd alloy by Nie et al. [41] due to the segregation of solute Gd atoms to these TBs. For Mg alloys, Serra et al. [53] suggested TBs themselves are barriers to basal slip.Therefore, it is reasonable to suggest that the TBs decorated with solutes and precipitates in the pre-deformed WE43 alloy could lead to a TBs strengthening. In our previous work of[38], we have quantitatively calculated the strength contribution from each strengthening mechanism to the YS for the pre-deformed T3 alloy. A large dislocation strengthening of 38MPa, precipitation hardening of 31MPa and a moderate TBs strengthening of 15MPa is estimated by using the strengthening analysis method presented in [38].

Fig. 7. Typical engineering stress−strain curves of WE43 alloy in different conditions: T4: solution treated condition (525 °C × 8h); T3: pre-deformed condition; T62: peak-aged of the un-deformed T4 alloy at 250 °C (4h); T6:conventional standard peak-aged condition (250 °C × 16h); T82: peak-aged of the pre-deformed T3 alloy at 250 °C (2h); T63: peak-aged of the undeformed T4 alloy at 300 °C (2h); T83: peak-aged of the pre-deformed T3 alloy at 300 °C (2h).

After peak aged at 250 °C, the un-deformed T62 alloy achieves a significant of precipitation strengthening, in which YS has an increase of 77MPa (reaching 222MPa) compared to that of the T4 alloy. It is worth noting that the standard T6 treatment was over-aged in present work, and leaded to a lower YS(187MPa)than that of the T62 alloy.This is mainly due to the formation of coarse β1precipitates(Fig.4b),which leads to a decrease of precipitation hardening. In contrast,only a moderate YS increase of 33MPa has been gained from T3 after aging at 250 °C for the pre-deformed T82 alloy,reaching 241MPa. As indicated above, there is a partial decomposition of the supersaturated magnesium solid solution into fine metastable β′′′and β′precipitates within the interior grains and relatively coarse second particles at the abundant TBs in the T3 alloy. This contributes to a precipitation hardening of 31MPa and consumes a large of solute. Thus, the subsequent precipitation strengthening during aging could reduce due to a decrease of volume fraction of precipitate in the T82 alloy.Especially,the pre-deformation has transformed the fine metastable β′′′and β′precipitates in the un-deformed T62 alloy into relatively coarse β1and β precipitates in the pre-deformed T82 alloy. This also results in a weaker precipitation strengthening. Besides, the dislocation strengthening could be partially diminished due to its recovery at 250 °C.Therefore, the pre-deformed T82 alloy only achieved a moderate YS increase. However, the pre-deformation has still enhanced the T82 alloy by 19MPa when compared to the undeformed T62 alloy. Particularly, the tensile elongation is not degraded much (Table 2). The strength improvement could be attributed to the remnant of dislocation strengthening and TBs strengthening.The DTs are thermally stable,which could retain the TBs strengthening in the T82 alloy (Fig. 3a).

After peak aged at 300 °C, only a slight YS increase of 27MPa was obtained in the un-deformed T63 alloy (Table 2).This is due to the formation of relatively coarse β precipitates (Fig. 6a), which leads to a weak precipitation hardening. In contrast, the pre-deformed T83 alloy has an even lower YS increase of 18MPa than that of the un-deformed T63 alloy, when compared to that of the T3 alloy (Table 2).The main reason is also that the T3 alloy has an initial precipitation hardening of 31MPa, which offsets the subsequent age hardening. Especially, the dynamically precipitated fine metastable β′′′and β′in the T3 alloy has transformed into the relatively coarse β precipitates in the T83 alloy, which results in a weaker precipitation strengthening. In addition, the dislocation recovery at 300 °C would also decrease the dislocation strengthening. These factors can explain the lower YS increase in the pre-deformed T83 alloy than that of the undeformed T63 alloy.Nevertheless,there is a large YS increase of 54MPa obtained in the pre-deformed T83 alloy (226MPa),as compared to that of un-deformed T63 alloy (172MPa in Table 2). This could be attributed to the improvement of precipitation strengthening due to the refinement of β precipitates and an addition of TBs strengthening for their thermal stability at 300 °C (Fig. 6). Furthermore, the tensile elongation is not much degraded.

4. Conclusions

In present work, a multi-directional cold impact forging(i.e. pre-deformation) prior to aging at 250 °C and 300 °C was applied to the WE43 Mg alloy. The following results are obtained:

1. The pre-deformation would completely transform the metastable β′′′and β′phases in the un-deformed alloy into β1and β phases in the peak-aged alloy at 250 °C.The width and thickness of the β1and β precipitates are about 100 and 10nm, respectively.

2.The pre-deformation obviously refined the β precipitates when aging at 300°C.The β precipitate has a width and thickness of 210 and 30nm, respectively. Besides, the deformation twins introduced by the pre-deformation were thermal stable up to 300 °C.

3. The pre-deformation can increase yield strength from the T62 un-deformed alloy (222MPa) by 19MPa, to 241MPa in the peak-aged T82 alloy at 250 °C. The same can increase yield strength from the un-deformed T63 alloy (172MPa) by 54MPa, to 226MPa in the peak-aged T83 alloy at 300 °C. In addition, the tensile elongation is not obviously deteriorated.

Declaration of Competing Interest

The authors declare no conflicts of interest.

Acknowledgments

We thank the GDAS’ Project of Science and Technology Development (Grants No. 2018GDASCX-0966, 2019GDASYL-0203002, 2018GDASCX-0117) and Guangzhou Science and Technology Planning Project (Grant No. 201904010309) for the financial support.