Tensile creep anisotropy of a Mg-2Y alloy extruded sheet with a splitted texture

Journal Pre-proof Tensile creep anisotropy of a Mg-2Y alloy extruded sheet with a splitted texture Zhenyu Xiao, Qinghuan Huo, Yuxiu Zhang, Zhirou Zhang, Zehao Li, Aki Hashimoto, Xuyue Yang PII:

S0925-8388(20)30117-1

DOI:

https://doi.org/10.1016/j.jallcom.2020.153754

Reference:

JALCOM 153754

To appear in:

Journal of Alloys and Compounds

Received Date: 2 September 2019 Revised Date:

3 January 2020

Accepted Date: 7 January 2020

Please cite this article as: Z. Xiao, Q. Huo, Y. Zhang, Z. Zhang, Z. Li, A. Hashimoto, X. Yang, Tensile creep anisotropy of a Mg-2Y alloy extruded sheet with a splitted texture, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.153754. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.

Author contributions section Zhenyu Xiao: Methodology; Formal analysis; Writing - Original Draft; Writing Review & Editing. Qinghuan Huo: Conceptualization; Supervision; Funding acquisition. Yuxiu Zhang: Investigation. Zhirou Zhang: Investigation. Zehao Li: Investigation. Aki Hashimoto: Investigation. Xuyue Yang: Conceptualization; Supervision; Funding acquisition.

Tensile creep anisotropy of a Mg-2Y alloy extruded sheet with a splitted texture

Zhenyu Xiao a, Qinghuan Huo a, *,Yuxiu Zhang a, Zhirou Zhang a, Zehao Li b, Aki Hashimoto c, Xuyue Yang a, d, *

a

Educational Key Laboratory of Nonferrous Metal Materials Science and

Engineering, School of Materials Science and Engineering, Central South University, Changsha 410083, China b

National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, 305-0047, Japan

c

Technology Div. Nippon Paint Surf Chemicals Co. LTD., 4-1-15 Minamishinagawa,

Shinagawa-ku, Tokyo 140-8675, Japan d

Nonferrous Metal Oriented Advanced Structural Materials and Manufacturing

Cooperative Innovation Center, Changsha 410083, China Corresponding author: Tel: +86-731-88876470, Fax: +86-731-88830136 Email: [email protected] (X. Yang), [email protected] (Q. Huo)

ABSTRACT: Mg-RE based alloys usually exhibit better creep resistance than traditional RE-free Mg alloys. Creep anisotropy can be still found in Mg-RE alloys but barely investigated. Thus, in the present study, tensile creep anisotropy and creep mechanisms were studied in a dilute Mg-2Y alloy extruded sheet. It is found that tensile creep property shows not only obvious stress and temperature dependences but also strong anisotropy. Varied stress exponent n is obtained under different conditions. Under σ < 40 MPa at T = 493 K, n = 2 is shown and grain boundary sliding is the main creep mechanism. Under σ < 40 MPa at T = 523 and 553 K, n = 3-4 is obtained and dislocation motion gradually dominates the creep. Under σ > 40 MPa at all temperatures, n > 4 is seen and intense dislocation slip becomes the predominant deformation mode. Besides, steady-state strain rate

and n are always higher along

ED than TD. This creep anisotropy is caused by different dislocation slip types. Basal slip plays the most important role in the creep anisotropy under σ < 40 MPa and non-basal slip contributes to the anisotropy under σ > 40 MPa. Heavy basal slip and

pyramidal slip are the main mechanism during the creep along ED while cross-slip of dislocations plays the dominant role along TD. Furthermore, {1012} anomalous twins are more likely to be found in samples crept along ED with a smaller Schmid factor for twinning. Anomalous twinning and pyramidal slip contribute to the creep fracture of ED sample. Unlikely, cross-slip composed by basal and prismatic dislocations induces the fracture of TD sample.

Key words: Magnesium; creep; anisotropy; dislocation; twin.

1 Introduction With the superior specific properties, magnesium (Mg) alloys have drawn vast interest for decades, making them good alternatives for certain high weight materials in automobile industries [1, 2]. The major drawbacks of Mg alloys elementally lie in their insufficient cold formability and limited high-temperature strength performances [3, 4]. Addition of rare earth (RE) elements such as Y, Gd, Nd to Mg alloys becomes a very useful approach to fix these limitations [5-8]. Mg-RE based alloys usually exhibits a refined microstructure and weakened texture. RE additions also change the c/a ratio and facilitate non-basal slip [9-12]. These can help improve the ductility of Mg alloys at room temperature efficiently. In the meantime, the extraordinary solution strengthening effect of Y, Gd ensures a better performance at elevated temperatures especially the creep resistance of Mg alloys, which is the key for applications in automotive and aerospace industries [4, 13-15]. Thus, lots efforts have been put into the research and development of Mg-RE based alloys nowadays. Another limitation is the highly anisotropic properties due to the insufficient number of deformation modes in hcp metals [16-23]. Remarkable effects of orientation on creep can be found in Cd [16], Zn [17], Ti [18, 20], Zr [21] as well as Mg alloys [24-28], when dislocation slip dominates deformation. This orientation dependence is weakened when grain boundary sliding takes the dominant place [24] or become less obvious when multiple types of dislocation slip are activated at elevated temperatures [26, 27]. It relies closely on the dominant creep mechanisms

and varies with applied stresses and temperatures [29]. With the addition of RE elements into Mg alloys, promoted non-basal slip are found and new types of texture component can be developed [30, 31]. Besides, RE addition could also obviously strengthening the grain boundaries [32, 33]. Deformation modes during the creep in Mg-RE alloys become more complicated and this would accordingly affect the creep anisotropy. But to date almost no work has reported the creep anisotropy of RE containing Mg alloys and detailed investigations are required. Heavy RE addition to Mg alloys may help improve the creep resistance of Mg alloys tremendously but also raises the cost and density [4]. Nowadays, creep resistance enhancement of the Mg alloys containing fewer RE is attracting much attentions [34-36]. Before this, deformation mechanism in dilute Mg-RE alloys needs to be clarified. In the present work, a dilute Mg-2Y (wt.%) alloy extruded sheet is used. Tensile creep properties of the Mg-2Y sheet are tested along ED and TD under various stress and temperature conditions. Creep mechanisms are then investigated carefully and creep anisotropy is analyzed in detail. This work helps provide a better understanding of the relationship between deformation mechanisms and creep properties as well as anisotropy of dilute Mg-RE alloys.

2 Experimental procedures The Mg-2Y (wt.%) binary alloy, used in the present study, was obtained through direct casting. The Mg-2Y alloy cast ingots were prepared by melting a Mg-29.99Y (wt.%) master alloy with pure Mg (99.95%, wt.%) ingots in an electric resistance furnace under a protective atmosphere of CO2 and SF6 in a ratio of 100:1 at 993 K and poured at 953 K. Then the Mg-2Y ingots were homogenization treated at 773 K for 24 h before extrusion. Hot extrusion was then carried out at a temperature of 653 K with an extrusion ratio of 30:1 and extruded sheets were obtained with a cross-section of 50 mm × 4 mm. The extruded sheets were subjected to annealing at 683 K for 1.5 h to obtain a homogenous microstructure with an average grain size of ~70 µm. Dog-bone shaped samples for uniaxial tensile tests and creep tests were fabricated with a gauge of 5 mm in length, 1.5 mm in width and 1.2 mm in thickness. These

samples were taken from the extruded sheet along extrusion direction (ED) and traverse direction (TD), using an electrical discharge machine. Creep tests were carried out at 493, 523 and 553 K under various applied stresses ranging from 10-100 MPa using the RWS-50 electronic creep testing machine. Before the creep tests, uniaxial tensile tests for each sample were performed at 493, 523 and 553 K with a strain rate of 3 × 10-3 s-1. These samples were heated and held for 20 min until the test temperature is achieved and distributed uniformly in each sample. Microstructure characterizations were conducted on the surfaces containing ED and TD, by optical microscopy (OM), scanning electron microscopy (SEM, FEI Quanta™ 650 FEG) and electron backscattering diffraction (EBSD) apparatus. The samples were first mechanically grounded with #1200, #2400, and #4000 abrasive papers. For OM observations, they were etched in a solution of 10 g picric acid, 25 ml distilled water, 25 ml acetic acid and 175 ml ethanol. For SEM and EBSD observation, they were electropolished in a solution of nitric acid (10 ml) and absolute ethanol (100 ml) after the mechanical grounding. In addition, thin foils were also made to investigate the dislocation mechanisms via a transmission electron microscopy (TEM, FEI Tecnai G2 F20). These foils were prepared firstly by punching samples into disks with a diameter of 3 mm, then mechanically grounded to a thickness of 100 µm and finally ion milling to perforation.

3 Results and discussion 3.1 Initial microstructure The initial microstructure of the extruded sheet observed by the EBSD method are presented in Fig. 1. A random distribution of color can be seen in the OIM map in Fig. 1(a)) and a splitted basal texture is found with most {0001} poles tilting away from the center towards ED in Fig. 1(b). The pole figures of {101 0} and {112 2} crystallographic planes are also given in Figs. 1(c) and 1(d) and a relatively random texture is shown for the present extruded sheet. Derived from the EBSD data, values of Schmid factor (SF) for basal slip, prismatic slip and 2nd-order pyramidal slip are investigated and their distribution charts are presented in Fig. 2. As

shown, the main discrepancy lies in the SFs of basal slip. For ED sample, most of the SFs of basal slip, mb, are higher than 0.35 and an average number of 0.37 is reached in Fig. 2(a). By contrast, the SF value of basal slip for TD sample is 0.27 on the average in Fig. 2(d), much lower than that for ED sample. The other two slip systems, prismatic and pyramidal slip, however, show smaller differences in the SF distributions between ED and TD samples and slightly lower values are found for ED sample. In this case, basal slip can be more easily activated in ED sample while a little higher initial activity of non-basal slip is found for TD sample. As the main deformation modes, their different incipient tendency could influence the deformation behaviors greatly of the present sheet.

3.2 Uniaxial tensile properties Tensile stress-strain curves for ED and TD samples tested at temperatures from 493 to 553 K are given in Fig. 3. Typical mechanical properties are recorded and presented in Table 1. As shown, higher strengths are detected along TD than along ED. This difference in strength, 0.2% proof stress (recognized as yield strength, YS) in particular, should be ascribed to the initial texture [37-39]. It is known that micro-yielding is closely related to the activation of the easiest deformation mode, i.e. basal slip in Mg alloys [40, 41]. If micro-yielding happens early, the density of basal dislocations would reach saturation rapidly and macro-yielding would also come soon. As shown, the ED-splitted texture leads to the different SF values for basal slip along ED and TD. The higher SF value for basal slip results in a favorable basal slip and thus lower proof stresses along ED. The difference is narrowed down from 42/53 (0.79) to 34/40 (0.85) when the temperature is elevated from 493 to 553 K. This can be attributed to the easier activation of non-basal slip at higher temperatures [42]. Strong anisotropy of strength is thus found in the present Mg-2Y sheet and the anisotropy will be weakened but cannot be eliminated by increasing temperature. 3.3 Creep properties Fig. 4 presents strain–time curves obtained from the creep tests carried out at 493 to

553 K. Typical creep curves containing two stages, namely primary creep stage and steady-state creep stage, are seen in most cases. At improved stresses, obvious tertiary or accelerated creep stage is shown and followed by fracture, resulting in a creep life less than 100 h. Strong temperature dependence is found through attaining these creep curves. As temperature increases, creep strain obtained at 100 h mounts rapidly and creep life becomes shortened largely. Analogous to the tensile strength above, anisotropic creep property is shown of the present extruded sheet. Under each temperature and stress condition, a reduced strain or prolonged lifespan is found of TD sample compared to ED. As depicted in Fig. 4, strains of TD samples are only 2/3 those of ED samples in most cases. With a lower ultimate tensile strength (UTS) along ED, steady state scarcely exists in creep curve at σ = 100 MPa and T = 493 K. By contrast, creep life along TD under the same condition can achieve 7.6 times those along ED. Difference in creep strain would become less obvious at lower stresses. At σ = 20 MPa and T = 493 K, strains obtained at 100 h are 0.34% and 0.22% along ED and TD, respectively. When the stress is decreased to 15 MPa, the strains between ED and TD are narrowed down to 0.18% and 0.14%. Strain rate of steady-state stage of each creep curve is calculated and its dependence on applied stress and temperature are investigated. The dependence of the steady-state strain rate ( ) on the applied stress (σ) can be characterized by the stress exponent, n =

|

(1)

Fig. 5(a) presents the log-log plot of the steady-state creep rate against applied stress along ED and TD for Mg-2Y sheet at various temperatures. As shown, n values for ED and TD samples range from 2 to 13 in Fig. 5(a). Different n values are obtained with different stress ranges at the same temperature while the n value changes are less conspicuous at different temperatures under the same stress. On the other hand, the temperature (T) dependence of the steady-state strain rate ( ) can be characterized by the activation energy, Q =



⁄

|

(2)

where R is the universal gas constant. Fig. 5(b) shows the Arrhenius plot of the logarithmic strain rate versus reciprocal temperature for Mg-2Y sheet under various applied stresses along ED and TD. In Fig. 5(b), the values of activation energy are higher than 135 kJ/mol for all the given samples. Like the n value, the Q value changes with applied stress and temperature as well. The n and Q values not just reflect the stress and temperature dependences but also denote the dominant creep deformation mechanisms for Mg alloys. Based on the large amounts of research works, diffusional creep is the dominant mechanism when n = ~1; grain boundary sliding (GBS) takes the dominant role of rate-controlling when n = ~2. And when n = 3-7, dislocation motions become the main deformation mechanism during creep (n = 3 for solute drag / viscous gliding and n = 4-7 for gliding and climbing). With n above 7 or 8, power-law breakdown happens, as a result of cross-slip in precipitation-hardened alloys or long-range plastic strain [4, 29]. Besides, Q = 80-92 kJ/mol indicates grain-boundary/pipe diffusion-controlled creep while 135 kJ/mol (which equals the value of self-diffusion activation energy in Mg, Qs) denotes the lattice diffusion-controlled creep in Mg alloys [43, 44]. The variation of n and Q values with stress and temperature in Fig. 5 infers that different mechanisms are dominating the creep in different stress and temperature ranges. For a low applied stress σ = 20 MPa, the values of n lie in between ~2 to ~4 from T = 493-553 K, suggesting a transition of prevailing creep mechanism from GBS to dislocation motions from low to high temperatures. At the meantime, the values of Q at σ = 20 MPa for all testing temperatures are close to the value of self-diffusion activation energy of Mg, 135 kJ/mol, indicating lattice diffusion assists the GBS or dislocation motion under low-stress creep processes. For an intermediate stress σ = 60 MPa, the n values are much enhanced to 4-7 and dislocation motions should be the dominant creep mechanism. Simultaneously, the Q values are increased to 190-270 kJ/mol, much larger than that of Qs. This higher activation energy in Mg-RE alloys has been reported in much previous research works and recognized as a result of dislocations gliding on non-basal planes or cross-slip [4, 29, 43, 45-47]. When temperature increases, both the n and Q values increase remarkably at the same stress of 60 MPa.

With n = 7 and Q > 240 kJ/mol, dislocation gliding becomes thermally promoted when stresses larger than 60 MPa become unbearable for the samples tested at T = 553 K. The n values are further improved to 11.3-13.1 under σ = 100 MPa at T = 493 K. It proposes that power-law breakdown occurs at the stress close to the UTS. Under such a high stress, heavy dislocation slip usually becomes the dominant creep mechanism [4]. In addition, orientation dependence of steady-state strain rate ( ) is seen in the sheet. Higher values of

are generally found in the creep tests of ED samples,

indicating that creep resistance is poorer along ED. Despite the different

values,

the obtained n and Q values of the two types of samples always fall into same categories in Fig. 5. Thus, the dominant type of creep mechanism of ED and TD samples maintains the same within each stress and temperature ranges and changes synchronously.

3.4 Creep mechanism 3.4.1 Microstructure developed The n and Q values, which are attributed to the subtleties of deformation mechanism during the creep tests, will be discussed hereafter. Microstructures evolved after creep tests at different stresses and temperatures are shown by optical microscopies (OM) in Figs. 6 and 7. No obvious GBS or formation of new fine grains can be observed under the given stress and temperature ranges. Instead, deformation twins are observed in some of these samples and the twin amounts increase with increasing stress and strain in Fig. 6. The twins are more likely to develop during the creep tests along ED rather than TD. Nearly no twins are detected in both ED and TD samples crept at σ = 40 MPa. As the stress as well as strain increases, the number of twins increases more quickly in ED samples than that in TD samples. When fracture occurs at σ = 100 MPa, twins can be largely observed in ED sample but still seldom in TD sample. Somehow, the twin amounts drop quickly as the creep temperature mounts. As shown in Fig. 7, large amounts of twins are observed in the samples tested at lower temperatures. By contrast, few twins can be found in the samples tested at T

= 523 and 553 K even after larger creep strains. With the help of EBSD technology, Figs. 8(a) and 8(b) present the typical OIM maps and their corresponding {0001} pole figures for ED and TD samples after creep tested for 50 h at T = 493 K and σ = 60 MPa. Their OIM maps with colors and pole figures on the left side show little difference compared with the initial ones in Fig. 1. This indicates that the majority of grains have their orientations scarcely changed (corresponding to the creep strain ε = ~2%). OIM maps without colors are presented on the right side in Fig. 8 to identify boundaries’ characteristics. In such maps, grain boundaries (GBs) with a misorientation angle above 15 deg are delineated by black lines while twin boundaries (TBs) of {1012} tensile twins are delineated by blue and {1011} compressive twins by red, respectively. As shown, only blue lines denoting {1012} TBs are found and no {1011} TB is formed during creep. The fraction of {1012} TBs amounts to 20.6% in ED sample but only to 6.9% in TD sample under the same creep condition. Typical OIM maps and {0001} pole figure for ED sample crept at a higher temperature of T = 523 K is also given in Fig. 8(c). To achieve a similar small strain, the lower stress of σ = 40 MPa is selected. The fraction of {1012} TBs decreases rapidly from 20.6 to 2.1% though the higher strain of 4.04% is finally achieved after test at T = 523 K (Fig. 8(c)). Hence, all the variations of twin fractions with tensile direction and temperature in the EBSD data is in accordance with the results observed in Figs. 6 and 7. In general, {1012} twinning is one of the most common deformation modes in Mg alloys but its activation shows great orientation dependence. Combined with the Schmid law, grains are mostly prone to triggering {1012} twinning when tension is applied parallel to the basal poles [4]. According to the initial {0001} pole figure in Fig. 1(a), most of the basal poles are at large angles with respect to ED and TD. More information details are given in Fig. 9 pertaining to the matrices’ orientation of these twinned grains. As shown in Fig. 9(a), the twinned grains in ED sample take up 28.1% in number fraction. The majority of the twinned grains have their matrices’ basal poles inclined 45 deg to ED (see Fig. 9(b)) along which tensile stress is applied. By contrast, only 6.1% of grains in TD sample contain twins with most of their basal

poles at an angle of 10 deg to TD. The SFs are calculated for {1012} twinning as well as basal slip in the twinned grains of both samples. The average value of SFs of {1012} twinning is 0.21 in ED sample while it becomes 0.32 in TD sample. This result is in contradiction to the present fact that a higher twinning activity is observed in ED sample instead of TD. Additionally, only 3% of these SFs exceed 0.35 and some of the SFs are even below 0 of the twinned grains in ED sample, in which twinning is expected to be very difficult to initiate according to the Schmid law. In other words, {1012} twinning in ED sample doesn’t conform to the Schmid law. The unconformable twins are referred to as “anomalous twins” and have been reported by lots of research works in kinds of Mg alloys [49-52]. Their activation derives from the complex internal stress state, to compensate strain incompatibility induced by other early deformation modes [52, 53]. In our previous work [27, 55], anomalous {1012} twins were largely detected in cyclic tensile tests along ED or pre-compression treatment along the ND of the Mg-Y sheets. The anomalous twins were largely formed in the grains with high basal slip tendency. In the present study, the same phenomenon is observed and the twinned grains in ED sample also possess a very high average SF value for basal slip. The average value of 0.45 in Fig. 9(c) is much higher than that of the total grains along ED (i.e. 0.37 in Fig. 2(a)). 98% of the SFs are above 0.35, indicating that almost all the twinned grains possess facilitating orientations for basal slip. Positive relationship is then manifested between the activity of basal slip and anomalous twins in ED sample. Inversely, a lower average SF value of basal slip is found of the total grains (i.e. 0.27 in Fig. 2(d)) along TD and consequently twins are seldom found in TD sample. With their matrices’ c-axis aligned to the tensile direction, the majority of the twins in TD sample are clarified as normal twins.

3.4.2 Dislocation mechanism during the creep along ED Furthermore, dislocation motions under creep are investigated with the help of TEM. TEM images of samples creep tested under different applied stresses along ED at T = 493 K are presented in Fig. 10 (incident beam direction B = [2110]). In Fig.

10(a), few long basal dislocations are observed adjacent to grain boundary in the sample crept at σ = 20 MPa. As the applied stress increases to 60 MPa, {1012} twin can be detected and short non-basal dislocations lie in the spaces between the long basal dislocations. Certain dislocation trace lines are stopped at the left side of the twin and almost no dislocation line along the same direction is shown at the other side. Fig. 11 gives the TEM bright fields under two-beam mode of Mg-2Y alloy sheets crept along ED for 50 h under different conditions: (a), (b) T = 493 K and σ = 60 MPa; (c), (d) T = 523 K and σ = 40 MPa; (e), (f) T = 553 K and σ = 20 MPa. Such conditions are selected to keep the creep strain at the same low level (2 to 4%) where creep has come into the steady-state stage. The incident beam direction for these images is B = [2110] and two perpendicular diffraction vectors g = [0110] (Figs. 11(a), 11(c) and 11(e)) and g = [0002] (Figs. 11(b), 11(d) and 11(f)) are selected, where c- or a- dislocation components disappear for the g·b = 0 invisibility criterion. In Figs. 11(a), 11(c) and 11(e) with g = [0110], dislocations can be obviously seen in these ED samples after creep. In ED sample crept at T = 493 and 523 K (Figs. 11(a) and 11(c)), most dislocations are long and straight. These long straight dislocations are parallel to the trace of (0002) plane and recognized as basal dislocations. Meantime, several short dislocations (marked by yellow arrows) are also found, perpendicular to the long basal dislocations. These are recognized as prismatic dislocations. As temperature increases to 553 K, the prismatic dislocations can be also found at the low stress of 20 MPa (Fig. 11(e)). In addition to the long basal dislocations, a few short dislocations are arranged in arrays parallel to the basal plane trace in Fig. 11(e), which are probably resulted from climbing. On the other hand, with g = [0002], only dislocations containing a c-component can be seen in Fig. 11(b), (d) and (f). As dislocation is sessile and cannot be activated alone, it must be accompanied by glissile dislocation [56, 57]. Thus, these dislocations having c-component are resulted from pyramidal slip during creep and are usually found long but curly. More pyramidal slip can be found

under σ = 60 MPa/T = 493 and 40 MPa/523 K but seldom detected under σ = 20 MPa at T = 553 K. It is suggested that pyramidal slip is more likely to occur under high applied stresses regardless of the creep temperature. Combined with the n values given in Section 3.3, the creep mechanism is closely dependent on the applied stress σ. At σ > 40 MPa (or 30 MPa at 553 K), dislocation motion is the dominant creep mechanism with n > 4. Basal slip is the main dislocation mechanism in ED samples, inferred from the long basal dislocations in Fig. 11. Non-basal slip, including evident pyramidal slip and minor prismatic slip, has also occurred in addition to basal slip under higher stresses along ED. At σ < 40 MPa, a transition of n = 2 to 4 is found from T = 493 to 553 K. With n = ~2, GBS is suggested as the dominant creep mechanism at 493 K. Furthermore, basal slip, as the easiest slip mode, can be found assisting the creep at σ = 20 MPa. When temperature elevates, dislocation motions become profound and contribute to the creep more than GBS. In addition to the heavy basal slip detected, both prismatic slip and dislocation climb has been gradually activated at T = 553 K. As a result, the n value increases to above 3 and even close to 4 at T = 523 and 553 K. On the whole, GBS dominates the creep process under applied stress σ < 20 MPa. Motions of basal dislocations are the main dislocation mechanism during creep σ > 20 and simultaneously contribute to the creep. Prismatic slip can be detected at σ > 20 MPa, and furthermore, pyramidal slip is frequently found at σ > 40 MPa.

3.4.3 Comparison of creep mechanisms of ED and TD samples Low applied stresses According to previous research works, the anisotropy of creep properties shown in Section 3.3, can be closely related to the activation of dislocation slip [24, 25, 27, 28]. Thus, when GBS is considered as the main creep mechanism under low applied stresses at T = 493 K (n = 2), isotropic creep property is expected under such conditions. But the differences in values of

and n cannot be ignored between ED

and TD samples in the present study. To achieve a fixed creep strain of 0.2% at T =

493 K, a low stress of ~20 MPa needs to be applied along TD while only ~15 MPa along ED (Fig. 4(a)), which can be treated as the creep strengths at 493 K [4]. The ratio of creep strength in two directions can be ~0.75 (15 : 20). At the meantime, the average values of SF for basal slip along ED and TD are 0.37 and 0.27, which also have the ratio of 0.73 (τc/mb,ED : τc/mb,TD). The value of the two ratios is very close. Therefore, the anisotropy under low stresses is mainly caused by the different activities of basal slip along two loading directions. When the higher activity of basal slip is shown along ED, the n value is enhanced to 2.4 and becomes closer to 3, which means basal slip contributes a larger part to the creep along ED than along TD. It is noted that during the uniaxial tensile tests at 493 K, the stress corresponding to the same strain of 0.2%, namely the proof stress, is 42 and 53 MPa along each direction. This is much higher than the “creep strength” at T = 493 K and the ratio of proof stresses is also larger (0.79), indicating that more frequent non-basal slip happens during the uniaxial tensile tests rather than the creep tests at T = 493 K. As temperature increases to 553 K, dislocation motions take the dominant place of GBS, resulting in n = 3.5. Non-basal slip, especially prismatic slip, can be detected at σ = 20 MPa (lower than the 0.2% proof stress). The activation of the non-basal slip somehow helps weakening the anisotropy at 553 K. As shown, n values along the two directions become the same in Fig. 5(a) and strain rate difference becomes narrowed down in Fig. 5(b).

High applied stresses Under stresses above 40 MPa, non-basal slip can be easily activated, which makes the anisotropy situation more complicated. Similarly, the dislocation mechanism in TD samples crept under high stresses is investigated by using TEM under two-beam mode. Fig. 12 presents TEM bright fields of Mg-2Y alloy sheets crept along TD for 50 h under different conditions: (a), (b) T = 493 K and σ = 60 MPa; (c), (d) T = 523 K and σ = 40 MPa. The incident beam direction is B = [2110] and two perpendicular diffraction vectors g = [0110] (Figs. 12(a), 12(c)) and g = [0002] (Figs. 12(b), 12(d))

are selected. Apart from basal dislocations, other non-basal dislocations, mostly prismatic dislocations, are largely observed in the crept TD samples. At T = 493 K, these short prismatic dislocations are perpendicular or at large angles to the trace of (0002) plane (marked by yellow arrows in Fig. 12(a)). They are found intersected and connected with the sparse basal dislocations in TD sample, at a higher frequency compared to ED sample in Fig. 11(a). It suggests that basal slip is difficult to occur while prismatic slip is promoted in TD sample. Since prismatic slip shares a same component with basal slip in Mg alloys; when basal slip is less likely to continue, dislocations can keep on gliding on prismatic planes. As a result, cross-slip composed by dislocations between basal and prismatic planes becomes the main deformation mode in TD sample. As temperature increases to 523 K, cross-slip of dislocation becomes more noted (see Fig. 12(c)) due to the elevated temperature [58, 59]. ery few dislocations having c-component are shown in TD samples crept at 493 K in Fig. 12(b) with g = [0002]. As temperature increases to 523 K, number of these dislocations increases in TD sample but still much less than those in ED sample under the same temperature and stress condition. Thus, under higher stresses, cross-slip of dislocation becomes the dominant dislocation mechanism in TD samples. It is found that the creep anisotropy under higher stresses can be mainly attributed to the different type of dislocation slip along different direction. Basal slip and pyramidal slip dominate the creep process along ED while suppressed basal but promoted prismatic slip plays the main role along TD. Different amounts and types of dislocation slip take part in the creep process along the two directions. Among these slip modes, the intense activation of basal slip is the most likely to accelerate the creep and deteriorate the creep resistance. The creep strain and corresponding steady-state strain rate are accordingly higher along ED with a higher SF value for basal slip. At the same time, prismatic slip can be easily activated under the low stress of 20 MPa when the temperature is higher at 553 K. The activation of prismatic slip could also undermine creep resistance but it causes less harm compared to basal slip, inferred from the better creep resistance of TD

samples. In these TD samples, promoted prismatic slip instead of basal slip is found. It should be noted that a lower value of SF for pyramidal slip is shown along ED than TD. But after creep tests under higher stresses, higher proportions of dislocations with c-component are found in ED sample. The same situation can be seen for the anomalous {1012} twins, which deviate from the Schmid law as well, as presented in Section 3.4.1 above. Both pyramidal slip and anomalous twinning can accommodate strains induced by heavy basal slip during creep under stresses above 40 MPa. As temperature increases, more pyramidal slip is shown in the crept samples but less twinning. Namely, pyramidal slip is activated in replacement of {1012} twinning in ED samples. The reason can be explained using some previous researches [52, 59]. Anomalous twinning usually cause latent strain [54] while pyramidal slip cause longitudinal strain [54]. When the creep temperature increases, pyramidal slip becomes more active to compensate the strain incompatibility along the loading direction. Additionally, the effect of anomalous twins on creep behaviors should be clarified here. It is unexpected that the largest difference in n values between ED and TD samples is obtained at T = 523 K rather than 493 K. Based on the observation results in the present work and the authors’ previous works [27, 28, 61], the anomalous twins developed in dilute Mg-RE alloys can inhibit dislocation motions and thus affect the steady-state strain rate. Considerable anomalous twins are shown in ED sample (Figs. 7 and 8) crept at T = 493 K whereas very few twins are seen at T = 523 K. This suggests that most basal dislocations are hindered by these anomalous twins at T = 493 K while basal dislocations can still move smoothly at T = 523 K. For comparison, much fewer anomalous twins have been developed and thus shows a minor effect on

and n values of TD samples crept at T = 493 and 523 K. Hence,

the most obvious creep anisotropy between ED and TD samples exists at T = 523 K instead of 493 K. In addition, the stress exponent n also reflects the undermining effect of increasing stress on creep resistance. As shown in Fig. 6, more twins are developed with increasing stresses at T = 493 K, which postponed the deterioration in ED samples.

3.4.4 Fractography At 493 K, n > 8 is shown in Fig. 5(a) under σ > 80 MPa for ED sample and σ > 90 MPa for TD sample, where power law breakdown occurs. Large creep strain is observed for each sample at such high stresses close to the UTS. Strain rate difference become amplified again, which can be attributed to the lower UTS along ED than TD. Fig. 13 presents the SEM images of typical surfaces and fracture morphologies of samples crept (a), (c), (e) along ED under σ = 90 MPa and (b), (d), (f) along TD under σ = 100 MPa at T = 493 K. The surfaces of both samples were electropolished only once before the creep tests and any characteristic observed from the surfaces is resulted from the creep. Obvious slip lines are observed on the surfaces of both ED and TD samples. In ED sample, mainly one set of parallel slip lines is shown in the majority of grains in Fig. 13(a). The magnified area given in Fig. 13(c) shows that some of the slip lines in ED sample are wavy. Crack can also be observed along the wavy slip line. The wavy slip lines are usually reported as the pyramidal slip traces [55, 61]. By contrast, in TD samples, two sets of the intersected lines are frequently found and they cross each other at an angle of ~120° (marked by yellow circles) in Fig. 13(b). In the magnified area, it is seen that meshes has been formed (see Fig. 13(d)), induced by the intersected cross-slip lines [55, 62]. In addition, the crack of TD sample propagates along two sets of slip lines and the crack edge become twisted. The corresponding fracture morphologies are also presented for ED and TD samples in Figs. 13(e) and 13(f), respectively. Shallow dimples are largely found in ED sample, indicating the high creep strain. Sharp cleavages are observed near these dimples in Fig. 13(e), which means twinning contribute a lot to the fracture of ED sample. In TD sample, very few dimples are found. Besides, striation regions developed by dislocation pile-ups are observed in the fracture morphology of TD sample (Fig. 13(f)), indicating that the fracture of TD sample is induced by dislocation slip. Although cleavages are also observed in TD sample, the twins should not be regarded as the main fracture origin due to its small amount (see Fig. 6(h)).

Combined with the surface information above, it is suggested that anomalous twins developed during the creep lead to the failure of ED sample while intense cross-slip are responsible for the fracture of TD sample.

4 Conclusions In this study, the tensile creep property and its anisotropy of a Mg-2Y alloy extruded sheet, which has an ED-splitted basal texture, are investigated under various stress and temperature conditions. The main conclusions can be drawn as follows: (1) Tensile creep property of the sheet shows not only obvious stress and temperature dependences but also strong anisotropy at all stresses and temperatures. The creep strain and steady-state strain rate are always higher when stresses are applied along ED than along TD. (2) Different n values are obtained under different stresses and temperatures. Under σ < 40 MPa at T = 493 K, n = 2 is shown and grain boundary sliding is the dominant creep mechanism. Under σ < 40 MPa at T = 523 and 553 K, n = 3 - 4 is obtained and dislocation motion gradually plays the dominant role. Under σ > 40 MPa at all temperatures, n > 4 is seen and intense dislocation slip becomes predominant. (3) The creep anisotropy is caused by different dislocation slip types at different applied stresses. Basal slip plays the most important role in the creep anisotropy under σ < 40 MPa and non-basal slip contributes to the anisotropy under σ > 40 MPa. Heavy basal slip and pyramidal slip are the main mechanism during the creep along ED while cross-slip of dislocations plays the dominant role along TD. (4) Under high stress ranges, {1012} anomalous twins are more likely to be found in samples crept along ED with a smaller Schmid factor for twinning. Anomalous twinning and pyramidal slip contribute to the fracture of ED sample. Unlikely, cross-slip composed by basal and prismatic dislocations induces the fracture of TD sample.

Data Availability Statement The raw and processed data required to reproduce these findings cannot be shared

at this moment due to technical and time limitations.

Acknowledgements The authors gratefully acknowledge the financial supports received from the National Natural Science Foundation of China (Grants No. 51974376, 51771230 and U1864209), the Natural Science Foundation of Hunan Province (Grant No. 2019JJ40376) and Key Research Program of Hunan Province (Grant No. 2019WK2062). Q.H. Huo acknowledges the Distinguished Professor Project of Central South University (Grant No. 202045009).

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Table 1 Values of 0.2% proof stress (YS), ultimate tensile strength (UTS) and fracture elongation (FE) obtained from uniaxial tensile tests at temperatures from 493 to 553 K.

YS / MPa

UTS / MPa

FE / %

T/K ED

TD

ED

TD

ED

TD

493

42

53

104

116

38.0

29.4

523

40

50

91

97

43.6

32.6

553

34

40

76

83

44.4

42.6

(a)

(b)

ED TD (c)

200 µm

(d)

Fig. 1. (a) Typical OIM map and corresponding (b) {0001}, (c) {1010}, (d) {1122} pole figures of Mg-2Y alloy extruded sheet.

Fig. 10. TEM bright fields of Mg-2Y alloy sheets crept along ED under (a) σ = 20 MPa and (b) σ = 60 MPa at T = 493 K for 50 h with incident beam direction B = [2110].

Fig. 11. TEM bright fields under two-beam mode (B = [2110]) of Mg-2Y alloy sheets crept along ED for 50 h under different conditions: (a), (b) T = 493 K and σ = 60 MPa; (c), (d) T = 523 K and σ = 40 MPa; (e), (f) T = 553 K and σ = 20 MPa. Such conditions are selected to keep creep strain at the same low level (2 to 4%).

Fig. 12. TEM bright fields under two-beam mode (B = [2110]) of Mg-2Y alloy sheets crept along TD for 50 h under different conditions: (a), (b) T = 493 K and σ = 60 MPa; (c), (d) T = 523 K and σ = 40 MPa.

Tensile direction

Tensile direction

120°

d

c

120° Tensile direction

Tensile direction

dimple cleavage

cleavage

striation

Fig. 13. SEM images of (a), (b) typical surfaces, (c), (d) local magnified areas and (e), (f) fracture morphologies of samples crept (a), (c), (e) along ED under σ = 90 MPa and (b), (d), (f) along TD under σ = 100 MPa at T = 493 K.

10 5

(d)30

Basal slip 25 SF : 0.27 avg

TD

20 15 10 5 0 0.0

0.1 0.2 0.3 0.4 Schmid factor, m

0.5

20 15 10 5 0 0.0

0.5

ED

Prismatic slip 25 SF : 0.31 avg

0.1 0.2 0.3 0.4 Schmid factor, m

(e) 30

Prismatic slip 25 SF : 0.37 avg

TD

15 10 5 0.1 0.2 0.3 0.4 Schmid factor, m

0.5

Pyramidal slip

ED

25 SF : 0.36 avg 20 15 10 5 0 0.0

0.5

20

0 0.0

(c) 30 Frequency, f / %

15

0.1 0.2 0.3 0.4 Schmid factor, m

(b) 30

(f) 30 Frequency, f / %

20

0 0.0

Frequency, f / %

ED

Frequency, f / %

Basal slip 25 SF : 0.37 avg

Frequency, f / %

Frequency, f / %

(a)30

0.1 0.2 0.3 0.4 Schmid factor, m

Pyramidal slip

25 SF : 0.41 avg

0.5

TD

20 15 10 5 0 0.0

0.1 0.2 0.3 0.4 Schmid factor, m

0.5

Fig. 2. Distribution charts of Schmid factor for (a), (d) basal slip, (b), (e) prismatic slip and (c), (f) pyramidal slip along ED and TD of Mg-2Y alloy extruded sheet.

140 493 K Engineering stress / MPa

120 100 80 60 523 K

40 20 0 0

ED TD

553 K

10 20 30 40 Engineering strain / %

50

Fig. 3. Tensile engineering stress – engineering strain curves along ED and TD of Mg-2Y alloy extruded sheet under various temperatures.

Creep strain, ε / %

(a) 30

100 MPa

ED TD T = 493 K

25 Fracture

20

90 MPa

15

80 MPa

10 60 MPa

5 0 0

20

40 60 80 Time, t / h

Creep strain, ε / %

(c) 50 40

80 MPa

40 MPa 20 MPa 15 MPa 0 10 20 30 40 50 60 70 80 90 100

120

60 MPa

(d)

20 40 MPa

10 0 0

(e) 50 Creep strain, ε / %

1.2 0.9 0.6 0.3 0.0

ED TD T = 523 K

Frature

30

100

(b)

20 60 MPa

40

40 60 80 Time, t / h 40 MPa Frature

100

1.6 1.2 0.8 0.4 0.0

20 MPa 10 MPa

0 10 20 30 40 50 60 70 80 90 100

120

ED TD T = 553 K

30 20

(f)

10

1.2 0.9 0.6 0.3 0.0

30 MPa 20 MPa

0 0

20

40 60 80 Time, t / h

100

10 MPa

0 10 20 30 40 50 60 70 80 90 100

120

Fig. 4. Creep strain – time curves of Mg-2Y alloy extruded sheet along ED and TD at different temperatures: (a) T = 493 K; (b) T = 523 K; (c) T = 553 K.

(a)

(b)

Fig. 5. (a) Variation of steady creep rate with applied stress at various temperatures in the log-log form and (b) Arrhenius plot of the logarithmic creep rate versus reciprocal temperature under various applied stresses along the ED and TD.

TD 200 µm

20 µm

ED

ED 50 µm

TD

Tensile direction

Fig. 6. Microstructures developed after creep tests along (a) - (d) ED and (e) - (h) TD for 50 h under various applied stresses at T = 493 K: (a), (e) σ = 40 MPa; (b), (f) σ = 60 MPa; (c), (g) σ = 80 MPa; (d), (h) σ =100 MPa (fracture occurs within 50 h).

TD 200 µm

ED

ED

TD

Tensile direction

Fig. 7. Microstructures developed after creep tests along (a) - (d) ED and (e) - (h) TD for 50 h under various temperatures and stresses: (a), (e) T = 473 K, σ = 80 MPa; (b), (f) T = 493 K, σ = 60 MPa; (c), (g) T = 523 K, σ = 40 MPa; (d), (h) T = 553 K, σ = 20 MPa. Such conditions are selected to keep the creep strain at a same low level (2 to 4%).

(a)

GB TB(20.6%)

493 K / 60 MPa

ED 200 µm

TD (b)

GB TB(6.9%)

493 K / 60 MPa

TD 200 µm

ED (c)

523 K / 40 MPa

GB TB(2.1%)

ED TD

200 µm

Fig. 8. Typical OIM maps and corresponding {0001} pole figures of Mg-2Y alloy sheet after creep tests for 50 h under different conditions: (a) along ED and (b) along TD at 493 K and 60 MPa; (c) along ED at 523 K and 40 MPa.

Intensity, I / m.r.d

(b) 16

E (c) 60

Basal slip 50 SF : 0.45 avg

ED f0.35: 98%

40 30 20 10 0 0.00

0.10

0.20 0.30 0.40 Schmid factor, m

45°

12 8 4

0 ∥ED 0 10 20 30 40 50 60 70 80⊥90 ED Angle, θ / ° (d) 30 ED {1012} twinning 25 SF : 0.21 : 3% f avg 0.35 Frequency, f / %

Frequency, f / %

D

ED

20 15 10 5 0 -0.50

0.50

-0.25 0.00 0.25 Schmid factor, m

Intensity, I / m.r.d

(f) 48 36

0.50

TD 10°

24 12 0

T 30 25

Basal slip SFavg: 0.39

TD f0.35: 66%

20 15 10 5 0 0.00

0.10

0.20 0.30 0.40 Schmid factor, m

0.50

Angle, θ / ° TD {1012} twinning f : 33% SF : 0.32 20 0.35 avg

(h) 25 Frequency, f / %

(g) 35 Frequency, f / %

D

∥0 TD 10 20 30 40 50 60 70 80 ⊥90 TD

15 10 5 0 -0.50

-0.25 0.00 0.25 Schmid factor, m

0.50

Fig. 9. Twinned grains and their orientation information in (a) - (d) ED and (e) – (h) TD samples after creep tests at 493 K and 60 MPa for 50 h: (a), (e) typical OIM maps; (b), (f) angle distribution between {0001} poles and the loading directions; Schmid factor distribution charts for (c), (g) basal slip and (d), (h) {1012} twinning of the twinned grains.

>Creep property shows obvious stress, temperature and orientation dependences. >Basal slip induces creep anisotropy and non-basal slip aggravates it. >Heavy basal and pyramidal slip are the main creep mechanism along ED. >Suppressed basal slip results in better creep resistance along TD. >Cross-slip dominates creep and induces fracture along TD.

Declaration of Interest Statement

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.