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Progress in Natural Science Materials International Progress in Natural Science: Materials International 24 (2014) 523–530 www.elsevier.com/locate/pnsmi www.sciencedirect.com

Original Research

Preparation, microstructure and degradation performance of biomedical magnesium alloy fine wires Jing Baia,b,n, Lingling Yina,b, Ye Lua,b, Yiwei Gana,b, Feng Xuea,b, Chenglin Chua,b, Jingli Yana,b, Kai Yanc, Xiaofeng Wand, Zhejun Tange a

School of Materials Science and Engineering, Southeast University, Jiangning, Nanjing 211189, Jiangsu, China b Jiangsu Key Laboratory for Advanced Metallic Materials, Jiangning, Nanjing 211189, Jiangsu, China c College of Mechanical Engineering, Yangzhou University, Yangzhou 225127, Jiangsu, China d School of Mechanical Engineering, Nantong University, Nantong 226019, Jiangsu, China e Mechanical and Electrical College, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, Jiangsu, China Received 2 July 2014; accepted 1 September 2014 Available online 28 October 2014

Abstract With the development of new biodegradable Mg alloy implant devices, the potential applications of biomedical Mg alloy fine wires are realized and explored gradually. In this study, we prepared three kinds of Mg alloy fine wires containing 4 wt% RE(Gd/Y/Nd) and 0.4 wt% Zn with the diameter less than 0.4 μm through casting, hot extruding and multi-pass cold drawing combined with intermediated annealing process. Their microstructures, mechanical and degradation properties were investigated. In comparison with the corresponding as-extruded alloy, the final fine wire has significantly refined grain with an average size of 3–4 μm, and meanwhile shows higher yield strength but lower ductility at room temperature. The degradation tests results and surface morphologies observations indicate that Mg–4Gd–0.4Zn and Mg–4Nd–0.4Zn fine wires have similar good corrosion resistance and the uniform corrosion behavior in SBF solution. By contrast, Mg–4Y–0.4Zn fine wire shows a poor corrosion resistance and the pitting corrosion behavior. & 2014 Chinese Materials Research Society. Production and hosting by Elsevier B.V. All rights reserved.

Keywords: Biomedical magnesium alloy; Wires; Cold drawing; Microstructure; Degradation

1. Introduction Magnesium and its alloys have attracted more and more attentions in biomaterials field recent years [1] because they are of the interesting advantage of both traditional metal implant materials and biodegradable polymers, including bioabsorbability, high mechanical properties, good mechanical and biological compatibility, etc.. The huge application potential of Mg alloys in medical devices also creates the demand of Mg alloy fine wire with diameter of less than 0.5 mm [2–6]. Mg alloy fine wire, for example, can be made to prepare sutures or anastomotic nail for n Corresponding author at: School of Materials Science and Engineering, Southeast University, Jiangning, Nanjing 211189, Jiangsu, China. Tel.: þ 86 25 52090689. E-mail address: [email protected] (J. Bai). Peer review under responsibility of Chinese Materials Research Society.

anastomosis, and also can be knitted to various kinds of tubular mesh stents, or used as reinforcement to prepare the polymer based composite for orthopedic surgery, etc. However, there are only few investigations on Mg fine wires reported so far [5–10]. First, due to the poor ductility of hcp magnesium at room temperature, the preparation of Mg alloy wire is much more difficult than that of traditional fcc and bcc metal wire (i.e. steel, copper, aluminum wires, etc.) through cold drawing process, which is considered currently a most mature and effective method to produce metal fine wire. Additionally, besides few welding wires [11] (most 4 1 mm in diameter), there was a lack of application for Mg alloy wires as industrial structural materials in the past. Furthermore, the compositions of Mg alloy wires in another aspect need to be considered. For these researches [5–10] on Mg fine wires, most of the materials studied were AZ31 or AZ61 alloy based on the Mg–Al series. Although AZ alloys

http://dx.doi.org/10.1016/j.pnsc.2014.08.015 1002-0071/& 2014 Chinese Materials Research Society. Production and hosting by Elsevier B.V. All rights reserved.

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show good comprehensive mechanical properties and are now the most common commercial Mg alloy, many Al-free Mg alloys with better biological performance, for example Mg-RE (rear earth element) and Mg–Zn based alloys, have been extensively developed and investigated recently in vitro and in vivo as new degradable biomedical Mg alloys [1,12–17]. To examine the cold drawing deformability of new biomedical Mg alloy wires, in this study, we designed three kinds of ternary Mg alloys with different RE (Gd, Y and Nd) as main alloying elements ( 4 wt%) and small amount of Zn ( 0.4 wt %) as the third element. The reason to select these alloying elements and contents is based on the following considerations: (i) the poor strength of pure Mg limits its applications in implant materials. (ii) Many evidences [12–17] support that all these alloying elements added have relatively high bio-safety at the present concentration. (iii) The addition of RE elements into Mg alloys can result in the formation RE texture and the weakening of the overall texture sharpness during deformation, contributing an improved room temperature ductility for based alloys [18– 20]. (iv) Gd, Y and Nd have relatively high maximum solid solubility in Mg matrix at eutectic temperature, about 22.5 wt%, 12.5 wt% and 3.6 wt%, respectively. These alloys, thus, can be strengthened through solution heat treatment. Simultaneously, almost single phase microstructures would help to provide a better drawing deformability and more homogeneous degradation. (v) Zn not only is an essential nutrient element in humans but also contribution effective strengthening to Mg alloy [17,21]. However, more contents of Zn will affect the solution of RE, and even form stable long-period ordered structure to deteriorate the ductility [22–24]. (vi) Due to the significant grain refinement after multiple cold drawing anneals, other microalloying elements usually added into Al-free Mg alloys to refine grain (i.e. Zr etc.), were not be selected. In this study, the microstructures and mechanical properties of the experimental alloys were examined at different states, and for the finished wires, the in vitro degradation tests were also carried out in order to investigate their degradation behavior. 2. Experimental Three kinds of ternary Mg–4%RE (Gd/Y/Nd)–0.4%Zn alloys (all compositons are griven in weight percent unless otherwise stated) were prepared. Their designed compositions and corresponding codes in this article are listed in Table 1. Commercial pure Mg, pure Zn and Mg–Gd/Y/Nd master Table 1 Chemical compositions of the experimental alloys. Alloy code Designed compositions (wt%) Gd 1# 2# 3#

Y

Analyzed compositions (wt%)

Nd

Zn

Mg

Gd

Bal. Bal. Bal.

4.26

4.0

0.4 0.4 0.4

4.0 4.0

Y

Nd

Zn

Mg

4.11

0.52 Bal. 0.46 Bal. 4.32 0.49 Bal.

alloys were used to achieve the target compositions. These alloys were melted in a mild-steel crucible under the protection of a mixed gas atmosphere of SF6 (1 at%) and CO2 (bal.), and subsequently cast into a cylindrical copper mold surrounded by water-cooling system with 60 mm in diameter and 250 mm in height. The chemical compositions of prepared alloys were analyzed by inductively coupled plasma atomic emission spectroscopy (ICP), and the results are also listed in Table 1. In order to eliminate segregation and obtain an ideal solidsolution microstructure, homogenization annealing treatment was performed for as-cast alloy billets at the temperature of 530 1C for 20 h. The annealing temperature is close to their solidus temperature measured by differential scanning calorimeter (DSC). These alloy billets then were multi-hole hot extruded to 2.7 mm diameter thick wires at an extrusion temperature of 480 1C and an extrusion ratio of 20. Starting from the original as-extruded wires, the cold drawing was performed step by step at room temperature. In this study, we used the constant strain multi-pass drawing, namely keeping the true strain as consistent as possible between adjacent drawing dies for all alloys and keeping the accumulative strain as consistent as possible between adjacent anneals for a certain alloy. Firstly, the maximum accumulative deformation (or ultimate drawing passes) of each alloy wire was respectively determined by continuous multi-pass drawing until the work hardening leaded to frequency rupture. These passes before the ultimate drawing pass were selected as the applied drawing passes. Subsequently, the annealing treatment was carried out at 400–450 1C for 5–30 min to obtain an optimum annealing system. In whole wire preparation procedure, multi-pass cold drawing and annealing treatment were performed repeatedly until the wire diameter is less than 0.4 mm. Electrochemical measurements for the final wires were performed using a three electrode system. A saturated calomel electrode (SCE) and Pt electrode were used as the reference and counter electrodes, respectively. All the measurements were carried out in SBF solution [25] at the temperature of 37 1C by a CHI760D electro-chemical workstation. The immersion tests were performed through the hydrogen evolution method developed by Song et al. [26] in SBF solution [25] according to ASTM-G31-72 [27] and the test temperature was maintained at 37 1C by water bath. During immersion test whole SBF solution was changed regularly to keep the pH value below 8. After the samples were removed from SBF, they were rinsed with distilled water and dried at room temperature. To observe the surface morphologies without corrosion products, some samples were then cleaned using a chromic acid solution. An average of three specimens was measured for each alloy. The microstructures were examined using Olympus BX60M optical metallographic microscope (OM), XL30 ESEM environmental scanning electron microscopy (SEM) with a GENESIS 60S X-ray energy-dispersive spectroscopy (EDS) equipped, and Tecnai G2 transmission electron microscope (TEM). The average grain size and grain size distribution were statistically calculated over 100 grains. The tensile strength was performed using CMT5105 electronic universal testing machine on the wires with gauge of 100 mm in length.


3. Results and discussion 3.1. Microstructures Microstructural observations were initially performed on ascast alloys, as seen in Fig. 1(a)–(c), showing the optical micrographs of as-cast 1#–3# alloy, respectively. The nonequilibrium solidification under the present cooling conditions results in the formation of intermetallic phases in all three alloys. With the decrease in solid solubility of RE elements in Mg matrix in the order of Gd, Y and Nd, the volume fractions of intermetallics increase, from isolated block-shaped particles in 1# alloy (Fig. 1(a)) to a nearly continuous network distribution along grain/dendrite boundaries in 3# alloy (Fig. 1(c)). To eliminate the segregation and obtain better ductility for subsequent cold drawing, we performed a 530 1C/20 h homogenization annealing treatment and then 480 1C hot extrusion to prepare wire with a diameter of 2.7 mm. Fig. 2 shows the microstructures of as-extruded alloys with their corresponding grain size distribution diagrams. After hot extrusion, most intermetellics have been dissolved into matrix and completely dynamic recrystallization occurs in all alloys. Unlike the similar dendritic size in as-cast alloys, there is an obvious difference in the average grain size between as-extruded 1#–3# alloys, 18.1 μm, 10.8 μm and 5.8 μm, respectively (Fig. 2 (a)–(c), note the scale bar in Fig. 2(c) is different from other two). Since the solid solubility of Nd in Mg matrix is much less than that of Gd and Y [28,29], it is reasonable to notice

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that there are much more intermetallics remaining in 3# alloy (Fig. 2(c)). Considering the significant effect of the secondphase particles on impeding the grain boundary migration during recrystallization, it is believed that the volume fraction of the residual intermetallics in different alloys play a major role to lead to the difference in the recrystallized grain size. The distribution and morphologies of the intermetallics in the as-extruded alloys were observed by SEM, as shown in Fig. 3. Very few intermetallic particles in 1# and 2# alloys present similar block-shaped morphology with size 1–4 μm, and their typical morphologies are displayed in the inserted higher magnification SEM images in Fig. 3(a) and (b), respectively. Obviously more intermetallic were observed in 3# alloy, as seen in Fig. 3(c), where they present two kinds of typical morphologies: relatively big isolated particles and band structure parallel to the extrusion direction with many small particles gathered. Because the volume fraction of the intermetallics in 1# and 2# alloys are so low that they can not be detected by XRD, EDS was used to measure the chemical compositions of these intermetallics and to estimate their phase structures. From 1# to 3# alloys, the compositions of the intermetallics observed are 82.3 7 5.2 at% Mg–17.27 5.0 at % Gd–0.57 0.2 at% Zn, 69.87 10.1 at% Mg–29.67 9.9 at% Y–0.8 7 0.2 at% Zn, and 93.5 7 3.6 at% Mg–6.07 3.3 at% Nd–0.5 7 0.3 at% Zn, respectively. According to the reported binary diagrams and literatures [29–33], the intermetallics in the three kinds of as-extruded alloys can be identified as Mg5Gd, Mg24Y5 and Mg12Nd phase with some Zn atoms

Fig. 1. Optical micrographs of as-cast alloys: (a) 1# Mg–4Gd–0.4Zn, (b) 2# Mg–4Y–0.4Zn and (c) 3# Mg–4Nd–0.4Zn.

Fig. 2. Optical micrographs of as-extruded alloys with corresponding grain size distribution diagram inserted: (a) 1# Mg–4Gd–0.4Zn, (b) 2# Mg–4Y–0.4Zn and (c) 3# Mg–4Nd–0.4Zn (note that scale bar in (c) is different from (a) and (b)).

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dissolved into crystalline lattices, respectively, although the recent work of S. Gorsse et al. [34] confirmed that Mg12Nd phase is metastable in Mg–Nd system. As-extruded wires were used to fabricate fine wires through the constant strain multi-pass cold drawing combined with intermediate annealing process. More details were described in Section 2. Three experimental alloys show different drawing formabilities, the maximum accumulative true strain before annealing of 1#–3# is 110%, 96% and 68%, respectively, and their corresponding applied accumulative true strain in this study is 96%, 82% and 55%, respectively. Although asextruded 1# alloy has relatively coarser grain, it conversely shows the best drawing formability. In general, these alloys show good cold drawing performance in contrast to AZ31, which has a similar formability with 3# alloy under the same conditions in terms of our previous works. The difference of drawing formability between alloys may be attributed to the different amounts of intermetallics and the effect of solute atoms on the deformation texture of Mg matrix. The micrographs of as-drawn 1# alloy under the conditions of one-pass drawing with true strain 14% and ultimate-pass drawing with accumulative true strain 110% are shown in Fig. 4(a) and (b), respectively, where the drawing direction is parallel to the vertical direction. In the initial stage of cold drawing, many deformation twins appear within the intragra-

nular fields, while the original equiaxed grain boundaries are still visible in Fig. 4(a). Larger strain imposed by cold drawing results in an apparent transformation in the microstructures. When the accumulative true strain reaches to 110%, the early morphological features, including original grain boundaries, twins etc., cannot be distinguished in the present resolution of optical microscope in Fig. 4(b). TEM was used to conduct more detailed observation, as seen in Fig. 4(c), showing the bright-filed image and corresponding selected area electron diffraction (SAED) patterns with an aperture size of 4 μm. TEM image shows that with a mass of dislocation propagation, very high dislocation density resulted in the formation of dislocation cell. Meanwhile, many small subgrians with a size 100–150 nm can be also found in the dislocation high-density areas, although these subgrain boundaries are generally ill-defined and irregular. The observations also are supported by the SAED pattern, where several relatively clustered diffraction spots regularly distribute along around circles. Severe plastic deformation (SPD), like high pressure torsion (HPT), usually can achieve significant grain refinement effect to metal materials [35]. During room temperature HPT processing, intense deformation even can drive the subgrain boundaries to transform into high-angle grain boundaries through an increase in their misorientation by continuously absorbing dislocations or by

Fig. 3. SEM micrographs of as-extruded alloys with inserted figures showing typical intermetallic particles and corresponding compositions measured by EDS: (a) 1# Mg–4Gd–0.4Zn, (b) 2# Mg–4Y–0.4Zn and (c) 3# Mg–4Nd–0.4Zn.

Fig. 4. Optical micrographs of cold drawn 3# Mg–4Gd–0.4Zn alloys with cumulative true stain (a) 18% and (b) 110%; (c) TEM bright field image of cold drawn 3# Mg–4Gd–0.4Zn alloys with cumulative true stain 110% and corresponding SAED pattern.


subgrain boundary migration and coalescence [36]. However, for cold drawing, the strain to fracture subjected by two compressive stresses and one tensile stress is still not large enough to induce the high-angle boundary transformation, as seen in Fig. 4(c). Besides drawing, intermediate annealing is equally important for fine wire preparation. Fine and uniform recrystallized grains are necessary for followed drawing. In order to obtain the optimum annealing process, for each alloy and deformation state, we respectively examined various annealing temperatures and time in the range of 400–450 1C and 5–30 min in according to the pre-experimental results. The annealing temperatures to achieve completely static recrystallization of Mg–RE alloys are much higher than that of cold drawn AZ31 alloy ( 200–250 1C) [7,9], suggesting that RE solute can intensively retard the nucleation and growth of recrystallization. More detailed results about the microstructural evolution during drawing and annealing will be published elsewhere. After drawing and annealing repeatedly for several times, the diameters of the final fine wires were reduced to less than 0.4 mm, corresponding to a cumulative true strain of 385%, 410% and 440%, respectively. Meanwhile, their grains were significantly refined to a similar level with average gran size of 3.8 μm, 3.8 μm and 3.6 μm, respectively, as shown in Fig. 5 (a)–(c). The average grain sizes of the three alloys after every annealing processes are plotted in Fig. 6 as a function of the cumulative true strain. The average grain size of original asextruded wires are also delineated as strain ¼ 0. Three curves display similar changing trend: rapidly decreasing in the initials stage and gradually reaching an approximate saturation level 3–4 μm after 2–3 times annealing. These results suggest that there is a similar threshold value of the grain refinement for the alloys studied through the present processing method in spite of the obvious difference in their original grain size and the amount of the residual intermetallics. Moreover, by comparison with the microstructures of the corresponding as-extruded alloys in Fig. 2, there were no obvious difference of intermetallics in the volume fraction and morphologies in the finished wires, suggesting these second phase particles are stable during the present drawing and annealing processes. The macro photos of the finished wire taken by digital camera and stereo microscope are displayed in Fig. 7(a) and (b), respectively.

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3.2. Mechanical properties Room temperature tensile tests were carried out for the asextruded thick wires and the final fine wires. The testing results are illustrated in Fig. 8. Industrial as-extruded AZ31 alloy commonly has an approximate 300 MPa ultimate strength and 10–15% elongation. By contrast, all studied as-extruded Mg–RE based alloys show moderate strength and excellent ductility above 20% on average. With grain refined after repeated cold drawing and annealing, the tensile properties of the finished wires present a notable change. For the same alloy, the yield strength of fine wire increases, but accompanying with the decline of elongation simultaneously. It implies that there is another determinant instead of grain refinement influencing the plastic performance. Al is one of the most effective alloying element to strengthen Mg alloys [37,38]. From the present results, the strengthen effect of RE is less than that of Al in similar solute concentration level. However, considering the drawing and tensile test results, it is concluded that the additions of RE elements can remarkably improve the ductility and cold drawing formability of Mg alloys. It has been well known that the texture plays an important role in the mechanical properties of Mg alloys due to the intensive {0001} basal slip and f1012g twining. Many investigations [18–20,39–41] have proved the improvement effect of RE on ductility of Mg alloys. This is attributed to the weakening (or randomizing) of texture by the new 〈1121〉 RE texture by Santford

Fig. 6. The change of the average grain size after every time drawing anneals with the increase of cumulative true strain.

Fig. 5. Optical micrographs of the final fine wire with corresponding grain size distribution diagram inserted: (a) 1# Mg–4Gd–0.4Zn, (b) 2# Mg–4Y–0.4Zn and (c) 3# Mg–4Nd–0.4Zn.

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Fig. 7. Macro photos of ∅ 0.3 mm wire: (a) whole roll taken by digital camera and (b) partial taken by stereo microscope.

Fig. 8. The tensile test results of the experimental as-extruded alloys and corresponding fine wires.

et al. [39–41]. As a result, the alignment of grains for the basal slip and the f1012g twining is better, causing the yield strength is reduced and the ductility increases. It is in good agreement with the present experimental phenomena with better cold drawing formability and tensile ductility of Mg–RE alloy in contrast to that of AZ31. However, with drawing strain imposed continuously, the texture intensity peaks probably rise gradually, and eventually reduce the ductility of fine wire, as seen in Fig. 8. Additionally, it seems that the annealing cannot completely eliminate the influence of texture introduced by cold drawing. Although Mg alloy textures have been widely investigated, there is still a lack of understanding on the texture characteristics and the effect of solute atoms on texture evolution in cold drawing processing. These researches will be carried out in future works. 3.3. In vitro degradation The in vitro degradation performance of the finished fine wire was evaluated by electrochemical and immersion tests in SBF solution at 37 1C. Fig. 9 illustrates the electrochemical polarization curves of three alloy wires. Three curves shows very similar tendency with an obvious current plateaus in their anodic curves, indicating the existence of a temporary protective film on the surface. It is also apparent from Fig. 9 that there is a remarkable difference in the corrosion potential (Ecorr) between these curves. The Ecorr value of 1# ( 1.40 V) and 3# ( 1.43 V) are much more positive than that of 2# ( 1.62 V), indicating that the 2# wire is more susceptible to corrosion in SBF. Simultaneously, the slightly lower breakdown potential (EPt) of 2# ( 1.25 V) than that of 1#

Fig. 9. Polarization curves of the three kinds of experimental alloy wires in SBF solution.

Fig. 10. Hydrogen evolution curves of the experimental alloy wires in SBF solution with an inserted figure showing the detailed information of 1# and 3# curves.

( 1.22 V) and 2# ( 1.20 V) suggests that there was a relatively less protective corrosion layer on the 2# wire surface. Similarly, the hydrogen evolution curves measured by immersion tests also show an clearly difference in the corrosion rate between these alloy wires, as plotted in Fig. 10 with a small figure inserted to offer detailed information of 1# and 3# curves. In comparison with 1# and 3#, 2# showed the fastest corrosion rate, and released approximately one order of magnitude more hydrogen within a quite long immersion period. On the other hand, the 1# and 3# present very similar hydrogen evolution


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Fig. 11. SEM micrographs of the three kinds of wires after immersion in SBF solution for 5 days: (a) and (d) 1# before and after removing the corrosion products; (b) and (f) 2# before and after removing the corrosion products; (c) and (f) 3# before and after removing the corrosion products, respectively.

process with a low corrosion rate except that there is a visible transition at 350 h in 1# curve. Similar transition also can be found in 2# curve when the immersion time is more than 200 h, indicating there is an accelerate corrosion after a parabolic corrosion behavior. The accelerate corrosion can be attributed to the occurrence of the localized corrosion over prolonging the immersion time and especially a subsequent rapid increase in corroded areas resulted from the corrosion fractures of the testing wire into a lot of short segments. To investigate the degradation behavior, surface morphologies of the specimens immersed for 5 days in SBF solution were observed by SEM, as shown in Fig. 11, where three columns from left to right are 1#, 2# and 3#, respectively, and the upper and lower rows relate to the morphologies with and without corrosion products, respectively. After 5 days immersion, a lot of corrosion products presenting granule and cluster morphologies with bright contrast can be observed on the surfaces of all specimens. In addition, from Fig. 11(c), it is clearly identified that the corrosion surface is a thin shrinkage cracking layer with a thickness less than 10 μm, and with cracking layer partially falling off, the relatively smooth substrate appears. By contrast, besides drying cracks, many deep corrosion cracks can be found in 2# specimens, as shown in Fig. 11(b). After removing the corrosion products, the differences in surface morphologies of naked wires are visible in Fig. 11(d)– (f). Many big corrosion holes, as shown in Fig. 11(e), were found in the 2# wire. On the contrary, 1# and 3# show relatively flat surface with a homogeneous distribution of small and shallow corrosion pits (Fig. 11(d) and (f)). The observations are consistent with the above hydrogen evolution test results, suggesting the uniform corrosion behavior for 1# and 3# wires but the pitting corrosion behavior for 2# wire. According to our researches on the present alloys, their corrosion performance can be influenced by many factors including materials state, grain boundary, intermetallic, etc., and these results will published later. For the present work, the

major difference in microstructures between the three kinds of alloy wires is the amount and distribution of the intermetallic particles (Fig. 5). It was reported [42,43] that the different intermetallics may influence the corrosion behavior from negative and positive aspects: fine and continuous distribution of intermetallics could retard the corrosion, but irregular distribution would conversely accelerate the corrosion. In this study, although there were a lot of intermetallics remaining in the microstructures of 3# wire, the 3# wire still shows a good corrosion resistance and is very similar with the 1# which barely contains intermetallic, indicating that there is no significant effect of the intermetallics in 3# alloy wire on its corrosion processing. Besides, it is estimated that for the present drawing wires, the alloy compositions, namely the interaction between Mg and solute atoms, may also play an important role to influence the corrosion behavior. For further understanding on the corrosion behavior of Mg–RE alloys in different states, more works need to be done in future. Furthermore, in order to adjust the degradation rate, surface modification investigations, including micro-arc oxidation, chemical conversion, poly-L-lactic acid (PLLA) coating, etc. are being carried out by our group. 4. Conclusions In this study, three kinds of Mg–4RE (Gd/Y/Nd)–0.4Zn alloy fine wires with the diameter less than 0.4 μm were fabricated through casting, hot extruding and multi-pass cold drawing combined with intermediate annealing process. Although there is an obvious difference in original grain size, the grains of the finished wires of all alloys are significantly refined to a very approximate level with an average grain size of 3–4 μm. In comparison with corresponding as-extruded alloy, the fine wire shows higher yield strength but lower ductility at room temperature. The degradation tests results and

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surface morphologies observations indicate that Mg–4Gd– 0.4Zn and Mg–4Nd–0.4Zn fine wires show similar good corrosion resistance and the uniform corrosion behavior in SBF solution. By contrast, Mg–4Y–0.4Zn fine wire shows a relatively poor corrosion resistance and the pitting corrosion behavior. Acknowledgments This research was supported by the Project supported by the National Natural Science Foundation of China (No. 51105203), Natural Science Foundation of Jiangsu Province of China (No. BK20130319), the Ph.D. Programs Foundation of Ministry of Education of China (Nos. 20120092120048 and 20120092120050) and Colleges and Universities in Jiangsu Province Natural Science Fund (13KJB430026). References [1] Y.F. Zheng, X.N. Gu, F. Witte, Mater. Sci. Eng. R. 77 (2014) 1–34. [2] Y.H. Wu, N. Li, Y. Cheng, Y.F. Zheng, Y. Han, J. Mater. Sci. Technol. 29 (2013) 545–550. [3] C.L. Chu, X. Han, F. Xue, J. Bai, P.K. Chu, Appl. Surf. Sci. 271 (2013) 271–275. [4] C.L. Chu, X. Han, J. Bai, F. Xue, P.K. Chu, Surf. Coat. Technol. 213 (2012) 307–312. [5] J.M. Seitz, D. Utermohlen, E. Wulf, C. Klose, F.W. Bach, Adv. Eng. Mater. 13 (2011) 1087–1095. [6] J.M. Seitz, E. Wulf, P. Freytag, D. Bormann, F.W. Bach, Adv. Eng. Mater. 12 (2010) 1099–1105. [7] H.F. Sun, H.Y. Chao, E.D. Wang, Trans. Nonferrous Met. Soc. China (Eng. Ed.) 21 (2011) 215–221. [8] Y. Oishi, N. Kawabe, A. Hoshima, Y. Okazaki, A. Kishimoto, SEI Tech. Rev. 56 (2003) 54–58. [9] H.Y. Chao, H.F. Sun, E.D. Wang, Trans. Nonferrous Met. Soc. China (Eng. Ed.) 21 (2011) 235–241. [10] S.A. Farzadfar, M. Sanjari, I.H. Jung, E. Essadiqi, S. Yue, Mater. Sci. Eng. A 528 (2011) 6742–6753. [11] W.Z. Jin, S.H. Liu, Mater. Sci. Technol. 13 (2005) 466–469. [12] N. Hort, Y. Huang, D. Fechner, M. Stormer, Acta Biomater. 6 (2010) 1714–1725. [13] R. Waksman, R. Pakala, P.K. Kuchulakanti, R. Baffour, D. Hellinga, R. Seabron, F.O. Tio, E. Wittchow, S. Hartwig, C. Harder, R. Rohde, B. Heublein, A. Andreae, K.H. Waldmann, A. Haverich, Catheter. Cardiovasc. Interv. 68 (2006) 607–617.

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