Mechanical, corrosion and biological properties of advanced biodegradable Mg–MgF2 and WE43-MgF2 composite materials prepared by spark plasma sintering

Journal of Alloys and Compounds 825 (2020) 154016

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Mechanical, corrosion and biological properties of advanced biodegradable MgeMgF2 and WE43-MgF2 composite materials prepared by spark plasma sintering Drahomir Dvorsky a, *, Jiri Kubasek a, Eva Jablonska b, Jirina Kaufmanova b, Dalibor Vojtech a
5 166 28 University of Chemistry and Technology Prague, Faculty of Chemical Technology, Department of Metals and Corrosion Engineering, Technicka Praha 6 e Dejvice, Czech Republic University of Chemistry and Technology Prague, Faculty of Food and Biochemical Technology, Department of Biochemistry and Microbiology, University of
5 166 28 Praha 6 e Dejvice, Czech Republic Chemistry and Technology Prague, Technicka

a

b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 15 November 2019 Received in revised form 21 January 2020 Accepted 22 January 2020 Available online 24 January 2020

Newly developed magnesium composite materials with a continuous network of MgF2 prepared by powder metallurgy exerted enhanced corrosion resistance and seems to be suitable for application in medicine as biodegradable implants. In this work, the influence of conditions of preparation of Mg and WE43 composite materials on final mechanical and corrosion properties of spark plasma sintered samples is revealed. Immersion in HF leads to the significant improvement of corrosion properties of Mg and WE43, while mechanical properties of WE43 alloy are reduced due to the specific interface. Moreover, cytocompatibility tests revealed the nontoxic behavior of magnesium fluoride coating on both Mg and WE43 alloy. The better proliferation of cells was observed on the WE43 composite material. © 2020 Elsevier B.V. All rights reserved.

Keywords: Magnesium Fluoride coating Corrosion WE43 SPS

1. Introduction Magnesium alloys are widely used in the automotive and aviation industry due to their low density and good mechanical properties [1]. It is also considered as a material for biodegradable implants. The most common commercially used magnesium alloy is known as WE43. It is an alloy containing rare earth elements especially neodymium (2.4e4.4 wt%), yttrium (3.7e4.3 wt%), zirconium (>0.4 wt%) and other rare elements like dysprosium or gadolinium in a small amount. Those alloying elements form intermetallic phases with magnesium [1e3]. There were observed Mg41Nd5, Mg3Nd, Mg12Nd, Mg24Y5, and b Mg14Nd2Y phases in the MgeY-Nd system [4e8]. This alloy is also known for its increased corrosion resistance compared with other alloys considered as biomaterials [9]. In this field, there is a trend to increase mechanical properties and decrease the corrosion rate. * Corresponding author. E-mail addresses: [email protected] (D. Dvorsky), [email protected] (J. Kubasek), [email protected] (E. Jablonska), [email protected] (J. Kaufmanova), [email protected] (D. Vojtech). https://doi.org/10.1016/j.jallcom.2020.154016 0925-8388/© 2020 Elsevier B.V. All rights reserved.

Mechanical properties and corrosion resistance are highly affected by the preparation method [10]. WE43 alloy can be hardened by heat treatment based on solution treatment. After T4 more alloying elements are dissolved in the magnesium matrix. This leads to solid solution strengthening. Even better mechanical properties are gained after the aging of solution treated material. This increase is caused by precipitation of fine intermetallic phases [2]. The latest research is focused on the powder metallurgy route [11]. It was confirmed that this method increases mechanical properties as well as corrosion resistance [10,12]. This method is based on the preparation of compact materials out of powders. Powders are usually gained by atomization. This process provides rapid cooling of material and provides small round shaped particles with a high amount of dissolved alloying elements (Similar to the material after T4 state). Thanks to rapid cooling this material is also fine-grained. Powder metallurgy combines the advantages of solid solution strengthening and fine-grained structure. The important effect on the final properties has the method of compacting. There are few methods of compacting typical for magnesium alloys. For example, extrusion is a common method for processing of magnesium alloys [4]. Magnesium has a hexagonal structure

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D. Dvorsky et al. / Journal of Alloys and Compounds 825 (2020) 154016

which requires higher temperatures to unlock slip planes needed for deformation. Therefore the extrusion process is performed at a higher temperature and dynamic recrystallization occurs. Powders compacted by extrusion are characterized by fine-grained structure. Individual particles of the powder are disturbed and their borders are indistinguishable. Prepared ingots have good mechanical properties with zero porosity. The more common way of compacting the powders is sintering. Sintering is based on the diffusion in a solid-state at high temperature. However, this process requires long times at high temperature and therefore the final material is coarser and may be partially porous. The improved sintering method is spark plasma sintering. This method involves pressure and high current in addition to temperature. Moreover, spark plasma sintering (SPS) is a very fast process and it only takes several minutes at high temperature, therefore, the coarsening is negligible. The final material is fine-grained with minimal porosity. After sintering, the individual particles remain recognizable in the structure. The corrosion resistance of magnesium alloys can be further reinforced by protective coatings. Some researchers tried using hydroxyapatite coating to increase biocompatibility and decrease corrosion rate [13,14]. However, this layer is relatively thick and with low adhesive strengths to magnesium alloys (6e17 MPa [14,15]). Another research is focused on coatings based on magnesium fluoride. Fluoride ions are also biocompatible with organisms similar to magnesium ions. The suggested daily intake of fluorine is 2e5 mg [16]. Magnesium fluoride coatings are also combined with calcium phosphate to improve corrosion resistance, biocompatibility and bonding strength [15,17]. Fluoride layers are much thinner than hydroxyapatite layers (0.1e4 mm [17e20]). The bonding strength of this coating is ranging from 33 to 43 MPa [15e18], which is much more than hydroxyapatite coating or combined double-layer of magnesium fluoride and calcium phosphate. It was proved that the combination of coatings is the best for increasing the corrosion resistance and bioactivity. However, the corrosion resistance is mainly provided by the magnesium fluoride layer [15,17,18]. The coating increases the polarization resistance magnificently and the corrosion potential is displaced toward more noble metals [15,16,21]. The preparation of fluoride coating is much easier than the preparation of hydroxyapatite coating or their combination. Magnesium oxide and magnesium hydroxide on the surface converts in hydrofluoric acid into the passive salt layer which consists of MgF2. MgF2 is almost insoluble in the water solutions and the thickness and quality depend mainly on the type of magnesium alloy and on the immersion time and concentration of hydrofluoric acid as illustrates Table 1 [22]. In more concentrated hydrofluoric acid the less immersion time is needed to form a layer

of defined thickness. However, after some time, the deposited film will reach a constant value [23]. The maximum reached thickness was 4 mm for pure magnesium [20], while 1.5 mm thick coating on AZ31 was enough to increase polarization resistance significantly [21]. Moreover, it was reported that magnesium coating increases the ignition temperature of magnesium making it less dangerous for storage, especially in the form of the powder [24]. The cytotoxicity tests were successfully applied on magnesium fluoride coatings with good results [16,25e27]. Moreover, the antibacterial properties of fluoride coating were discovered [16]. Magnesium fluoride coatings were also successfully tested on animals [18,19,28e30]. However, layer protection has a disadvantage when the layer is disturbed. The corrosion process is then usually concentrated in one place with devastating consequences. This danger can be prevented using the combination of coating protection and powder metallurgy. If each particle of the powder will be immersed in the HF solution and compacted afterward then the continuous network of the fluoride layers will be created across the sample. The optimal method for gaining the best mechanical properties is compacting magnesium powders by extrusion, however, during this process, there is a huge risk of breaking the shells of the individual particles and therefore breaking the continual network of magnesium fluoride layers. So the ideal method for compacting the powder is spark plasma sintering. Sintering should preserve magnesium fluoride shells while creating compact material without pores. This material should have improved corrosion resistance even after the surface layer is disturbed because the whole material will be made of the network of magnesium fluoride layers. However, the mechanical properties might be negatively affected by the fluoride network and it may cause significant embrittlement depending on the thickness of the coating. This new method of preparation of MgeMgF2 composite material was first introduced in previous research [31] with positive results. A similar method was applied to WE43 alloy in another paper [32]. WE43 alloy exerted different behavior in HF and an inhomogeneous layer of MgF2 was created on the surface of each particle. Moreover, intermetallic phases showed a special affinity to HF and each intermetallic phase on the surface was surrounded by MgF2 and the YF3 phase was created after sintering. The principle of corrosion protection was therefore different compared with pure Mg. Nevertheless, the aim of this paper is the investigation of biological properties as well as the influence of conditions of preparation on the mechanical, and corrosion properties of prepared composites. The immersion time in HF should influence the thickness of the obtained coating and thus mechanical and corrosion properties.

Table 1 Fluoride conversion coatings on selected magnesium alloys. Material

Concentration of HF [%]

Duration [h]

Temperature [
C]

Thickness [mm]

Reference

AZ31B AZ31B AZ31 Mge3Zn-0.8Zr Mge3Zn-0.8Zr Mge3Zn-0.5Zr Mge3Zn-0.5Zr Mg Mg Mg Mge2Nd Mge2Zn-0.5Y-0.5Nd Mg-Nd-Zn-Zr WE43

50 50 48 20 20 20 20 48 48 40 40 40 40 40

48 72 24 6 12 6 2 24 24 96 96 48 12 96

30 30 Room 37 37 37 37 Room Room Room Boiled in NaOH Room Room Room

1.9 2.75 1 0.5 3 0.5 0.5 1.3 1.5 4 1.6 0.1 1.5 0.4

[16] [21] [17] [19] [33] [34] [35] [18] [36] [20] [28] [25] [27] [20]

D. Dvorsky et al. / Journal of Alloys and Compounds 825 (2020) 154016

2. Materials and methods 2.1. Powder characterization and coating Commercial atomized powders of pure magnesium and WE43 alloy (composition with dangerous trace elements, which highly increase the corrosion rate of magnesium alloys (Fe, Ni, Cu) is given in Table 2) were characterized by laser diffraction to determine the particle size. Particle size has an effect on the ratio between magnesium fluoride and magnesium matrix. Fifty grams of powder was immersed and stirred in the 300 ml of 40% HF for 1, 24 and 96 h. The HF was then poured out and the powder was rinsed with distilled water multiple times and filtered through filter paper and rinsed with ethanol multiple times. The powder was then dried out in the kiln-drying. 2.2. Compacting The coated and uncoated atomized powder was processed by the SPS method at 500
C with heat rate 100
C/min and 7 kN pressure level and with operation time 10 min. The SPS machine HP D 10 FCT system GmbH was used. The final cylindrical rods had a diameter of 20 mm and a height of approximately 10 mm.

3

Three-point bending properties were carried out on the same machine with a constant deformation speed of 0.001 s1 on rectangular specimens with dimensions of 3  3  15 mm. Bending yield strength (BYS), ultimate bending strength (UBS) and total deformation (B) were determined. 2.5. Corrosion test Immersion tests were performed in simulated body fluid (SBF) (Table 3) at 37
C for 14 days. The ratio of solution volume to the surface area was 100 ml/cm2. Samples were removed from the immersion solution and were rinsed in distilled water and dried. The corrosion products were removed by a solution of 200 g/l CrO3, 10 g/l AgNO3, 20 g/l Ba(NO3)2. Samples were then dried and weighed. The corrosion rate was calculated from weight changes after immersion and removing corrosion products. Electrochemical measurements were performed in the same medium with a volume of 100 ml, where 1 cm2 of the sample was exposed to the solution. The measurement of polarization resistance and potentiodynamic curves were performed after 1 h of stabilization of open circuit potential. The potentiodynamic curves were measured within the range of potentials from 0.02 V vs OCP to 0.6 V vs OCP. 2.6. Cytocompatibility e test with extracts

2.3. Surface and microstructure characterization The input powders and microstructure of the compact materials were characterized by electron scanning microscope SEM TescanVEGA3 with energy dispersion spectrometry (EDS e Energydispersive X-ray spectroscopy, AZtec). Samples were mounted into an epoxy resin. Subsequently, they were ground on SiC grinding papers (P80eP2500) and polished on diamond paste D3, D2, D0.7. The final polishing was done on Eposil F. Porosity of samples was evaluated by ImageJ software by evaluation of 10 images. The grain size was measured by the same software after etching the sample in the solution of 10 ml HNO3, 30 ml of acetic acid, 120 ml of ethanol and 40 ml of H2O. The phase composition of the alloys was determined by X-ray diffraction (X’Pert Philips, 30 mA, 40 kV, X-ray radiation Cu Ka). Transmission electron microscopy (TEM, Jeol 3010) operated at an accelerating voltage of 300 kV was used for TEM examination. TEM foils were first prepared from the samples by mechanical grinding, followed by dimpling (Gatan Dimple Grinder 656) to a thickness of 3 mm. Final TEM foils were obtained by subsequent precision ion polishing (Gatan PIPS 691). 2.4. Mechanical properties Compressive testing was carried out on rectangular samples with dimensions of 5  5  7 mm using LabTest 5.250SP1-VM at room temperature. A constant strain rate of 0.001 s1 was used. Compressive yield strength (CYS), ultimate compressive strength (UCS) and total deformation to fracture were determined (D). Tensile properties were measured using the same machine at room temperature and on a bone-shaped sample that was 2  4 mm thick and 10 mm long in the narrowed area. A constant deformation speed of 0.001 s1 was used. Tensile yield strength (TYS), ultimate tensile strength (UTS) and total elongation (A) were determined. Table 2 Composition of input powders with dangerous impurities [ppm].

Mg WE43

Mg

Y

Nd

Gd

Dy

Zr

Fe

Ni

Cu

Bal. Bal.

7.0 41317.4

4.2 23652.7

1.1 5229.5

0.3 4331.3

3.6 4041.9

89.9 19.0

20.0 21.0

10.0 32.9

Samples of WE43 and Mg immersed for 96 h in HF were tested for cytotoxicity according to the ISO 10993-5 (resazurin test with extracts for metabolic activity). These samples were compared to the materials before the immersion. Three samples of each type were used. They were sonicated in ethanol and then sterilized in 70% ethanol for 2 h and under UV (ultraviolet) for another 2 h. Samples were then shaken (125 RPM) in MEM (Minimum Essential Medium) þ 5% FBS (Fetal Bovine Serum) at 37
C in closed vessels for 24 h. 1 ml of medium was used for 87.5 mm2 of the sample surface. Murine fibroblasts L929 (ATCC® CCL-1 ™) were cultured under standard conditions in MEM þ 10% FBS. For the test, the cells were seeded in 96-well plates at a density of 104 cells per well in MEM þ10% FBS. After 24 h, the medium was replaced by the extracts of the tested samples. MEM with 5% FBS was used as a negative control. After 24 h of incubation with the extracts, cell metabolic activity was determined using resazurin assay. The cells were first washed with PBS (phosphate-buffered saline) and then, 100 ml of resazurin (final concentration 25 mg/ml) in MEM with 10% FBS was added to each well. Fluorescence was measured after 1 h of incubation (excitation/emission wavelength being 560/590 nm). The percentage of metabolic activity of the cells exposed to extracts relative to the negative (untreated) control was evaluated. The test was performed in hexaplicates for each extract. The inductively coupled plasma mass spectrometry (ICP-MS, Elan DRC-e) was used for the determination of the concentration of released ions in extracts. Two sample t-test was used to evaluate the statistical difference between the coated and uncoated state. 2.7. Cytocompatibility e direct contact test Samples of WE43 before and after the immersion were also used for direct contact test. The samples after the test with the extracts were used, i.e., the samples were preincubated for 24 h in the culture medium. Murine fibroblasts L929 in MEM þ10% FBS were Table 3 Composition of SBF [mmol/l]. Ions

Naþ

Kþ

Mg2þ

Ca2þ

Cl

HCO 3

HPO24

SO24

SBF

142.0

5.0

1.5

2.5

147.8

4.2

1.0

0.5

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occasionally small satellites adherent to the greater particle. The size of the Mg powder ranged between 50 and 200 mm, while in the case of WE43 alloy it was between 10 and 200 mm. Even though the range of particle size is wider in the case of WE43 (Fig. 2) the average particle size is much smaller. Nevertheless, the size distribution of Mg is more homogeneous. The grain size of the magnesium powder ranged between 10 mm and 40 mm. The grain size of the WE43 powder cannot be obtained as the structure was characterized by fine dendritic structure. The dendrites consist of magnesium solid solution (3 wt% of Y, 1.7 wt% of Nd, 0.3 wt% of Dy, and 0.3 wt% of Gd) and the interdendritic space was enriched on alloying elements (6.6 wt% of Y, 4.7 wt% of Nd, 0.7 wt% of Dy and 0.4 wt% of Gd). By XRD analysis the metastable b0 -phase, with a composition close to b-phase (Mg14Nd2Y) was identified. Each material behaved differently in HF. Pure magnesium created a homogeneous continuous layer, while the coating on WE43 was very irregular (Fig. 1C and D). The reason behind this irregularity is discussed later. 3.2. Microstructure

Fig. 1. The cut of the A) Mg powder, B) WE43 powder, C) Mg immersed in HF (24h) with EDS map of F, D) WE43 immersed in HF (24h) with EDS map of F.

then seeded directly on to the samples. The test was performed in 6-well plates in 5 ml of the media per well. The initial cell density was 24 000 cells per cm2. After 24 h of incubation at 37
C in 5% CO2, the cells growing on the samples were rinsed with PBS, fixed with Karnovsky solution for 1.5 h and dehydrated in series of ethanol (50, 70, 90, 100%). They were then dried using CPD (Leica EM CPD300) followed by sputter coating (10 nm of gold) and scanning electron microscopy. 3. Results and discussion 3.1. Powder characterization Both input powders were characterized with round-shaped particles (Fig. 1). In the case of WE43 powder, there were also

Samples were successfully prepared by spark plasma sintering. The microstructures of prepared compact samples are summarized in Fig. 3. Individual particles are preserved and distinguishable in all specimens, which is a presumption for the creation of a continuous network of MgF2. All materials were well sintered with minimal porosity, which was beneath 0.3%. The presence of F leads to a decrease of porosity to less than 0.1%. The porosity of obtained material is similar to the results published by Minarik et al. [8] who prepared MgeY-Nd alloy by SPS. The grain size of both materials was evaluated by image analysis of 10 images of etched samples. The fluoride treatment had no effect on the final grain size. The average grain size of pure Mg was approximately 20 mm after sintering, which stayed almost unchanged compared with the unsintered powder. The structure of WE43 alloy changed after sintering significantly. The dendritic structure was no longer visible in the sintered material and the metastable b0 -phase transformed into more thermally stable phases like Mg24Y5 which was detected by XRD analysis (Fig. 4). Nevertheless, other phases rich on Nd like Mg14Nd2Y and Mg41Nd5 were also observed and identified only by the EDS analysis as their amount was too low, and therefore they were beneath the detection limit of XRD analysis. Other phases like Mg3Nd and Mg12Nd were observed in the WE43 alloy prepared by

Fig. 2. The size distribution of particles: a) Mg, b) WE43.

D. Dvorsky et al. / Journal of Alloys and Compounds 825 (2020) 154016

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Fig. 3. The microstructure of A) Mg, B) Mg 1h HF (With EDS analysis of F), C) Mg 24h HF (With EDS analysis of F), D) Mg 96h HF (With EDS analysis of F), E) WE43 (With EDS analysis of O and Y), F) WE43 1h HF (With EDS analysis of Y and F), WE43 24h HF (With EDS analysis of Y and F), WE43 96h HF (With EDS analysis of Y and F).

SPS method in a low amount by other authors [8,37]. Alloying elements segregated at the particle and grain boundaries (Figs. 3E and 5A, B). Individual particles were surrounded by the yttrium and oxide rich layer (Fig. 3E). The area just beneath the borders made of yttrium rich layer is depleted on alloying elements and intermetallic phases. Yttrium needed for the creation of the Y2O3 on the boundary between particles was consumed from the nearby solid solution. The oxygen needed for this reaction was probably acquired from the residual air between particles. This transformation created a gradient of Y content in each individual particle (Y poor near the surface e 0.8 wt%, Y rich in the center of the particle e 3.7 wt%). Grains could be better distinguished after etching and according to the image analysis the grain size of WE43 ranged between 1 and 10 mm after sintering. In the case of pure Mg, the borders between particles are less evident than in other cases as the boundary is made of a very thin oxide layer. Otherwise, the longer time of immersion in HF leads to

Fig. 4. XRD patterns of composite materials.

more distinguishable boundaries, and therefore, a thicker layer of magnesium fluoride. The thickness of the MgF2 layer was identified by line scan analysis, and it ranged between 3.4 and 3.8 mm. The difference between 24 h and 96 h was, however, very low which is in accordance with conclusions for bulk materials immersed in HF, where the maximum thickness was achieved after 8 h and was constant after 24 h (for AZ91D alloy) [23]. The coating is homogeneously distributed around each particle for powders immersed for 24 and 96 h (Fig. 3C and D). Powder immersed for just 1 h is uneven and poor on F (Fig. 3B). The amount of F detected by EDS analysis of the surface increased from 0.4 wt% after 1 h to 2.0 wt% after 96 h of immersion in HF. Individual phase transformation during the processing is summarized in Table 4. Sintering of powder of WE43 alloy immersed in HF resulted in the structures depicted in Fig. 3FeH. The alloy WE43 was much more resistant to HF then pure Mg, which was confirmed by previous research on bulk material [20]. The fluoride coatings were uneven on powders as well as on the bulk material. There were increased concentrations of F around some places and contrary areas poor on F. Such behavior could be contributed to the intermetallic phases on the surface of particles. As was observed by € hlinger et al. [38] intermetallic phases tend to lead to enhanced Ho layer growth above them and it makes the layer inhomogeneous. Also, Trinidad et al. [39] observed the creation of MgF2 specifically around the eutectic phase in the as-cast ingot. However, in our case, the increased temperature during spark plasma sintering allows reaction between MgF2 and yttrium dissolved in the solid solution creating YF3 phase (Table 4) due to the lower standard molar Gibbs energy of formation of YF3 (1745 kJ/mol and 1871 kJ/mol at room temperature and 527
C, respectively) compared to that of MgF2 (1141 kJ/mol and 1190 kJ/mol at room temperature and 527
C, respectively) [40,41] (Fig. 3FeH, Fig. 5C and D). The presence of YF3 phases was confirmed by XRD analysis (Fig. 4) and by TEM analysis of the interface (Fig. 5D). These phases were accompanied by the MgF2 layer. In this case, the solid solution is also slightly depleted on Y near the borders between particles due to the creation of the YF3 phase. Nevertheless, the amount of yttrium consumed for those phases is much lower compared to the Y2O3 due to the stoichiometry. Longer times of immersion lead to slightly more even fluoride coating according to EDS analysis (Fig. 3FeH). The amount of F measured by the EDS analysis of the material

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Fig. 5. Microstructure - SEM and TEM with EDS analysis A, B) WE43 and C), D) WE43 24h HF.

Table 4 Compounds on the surface of the particles and their transformation at each stage of processing (MgRE ¼ intermetallic phases with rare elements, *without HF treatment).

Mg WE43

Powder

Immersion in HF

SPS

Mg, MgO Mg, MgO, MgRE

MgF2 MgRE, MgF2

MgF2 (MgO*) YF3, MgF2 (Y2O3*)

increased slightly from 1.1 wt% after 1 h to 1.5 wt% after 96 h of immersion in HF. Nevertheless, the coating is still inhomogeneous after 96 h compared with the pure Mg. The inhomogeneity is associated with the creation of very fine YF3 phases on the interface between particles. 3.3. Mechanical properties The compressive, tensile and bending properties were measured and the results are summarized in Fig. 6. Tensile and bending properties could provide us information about the adhesive properties between particles as previous research revealed the fracture happens mainly between particles and the fracture is predominantly covered by oxide or fluoride layer [31]. Compressive properties should reveal information about particle strengthening or weakening. The mechanical properties of the materials prepared by spark plasma sintering were comparable or even better than the properties of the casted ingots [42,43]. Such behavior could be attributed to the minimal porosity, good sintering and good adhesion of fluoride layer [16,17]. In the case of magnesium, the compressive and tensile properties of the composite materials were better than of the pure magnesium for all times of immersion while the best mechanical properties were achieved after 24 h of

immersion. After 96 h of immersion, the layer was probably thicker and therefore more brittle while 1 h leads only to the creation of a very thin layer which resulted in only slight improvement of mechanical properties. Similar results were obtained by bending tests. Again 24 h of immersion in HF resulted in the increase of ultimate strength due to the improved adhesion between particles, which is greater for the MgeMgF2 bond than for the MgeMgO bond [16,17,31]. This can be deduced from the fractography with the EDS analysis in Fig.7A, C. The fracture of pure Mg and MgeMgF2 composite proceeds between particles. The surface of pure Mg contains an increased amount of O (Fig. 7A), while in the case of MgeMgF2 composite there is a high amount of F (Fig. 7C). Totally different behavior was observed for the WE43 composite material. Plasticity and total deformation were decreasing with the increasing time of immersion in HF. The dramatic decrease of mechanical properties was noticeable even after 1 h of immersion. This could be associated with the creation of a specific interface that consists of the YF3 phase and the inhomogeneous fluoride layer. This interface has probably much lower adhesion strength compared to the yttrium oxide bonding strength. Even though WE43 alloy exerted specific resistance against HF and mostly only intermetallic phases and nearby surroundings were turned into MgF2 it seems it had a serious impact on the mechanical properties. Very fine intermediate YF3 phases that were densely distributed around each particle deteriorated mechanical properties. Longer immersion times in HF lead to an increased amount of YF3 phases. Therefore, shorter times of immersion in HF (up to 24 h) are suggested in order to sustain at least half the values of mechanical properties. The impact of YF3 phases on the plasticity is illustrated on the fracture surfaces with EDS analysis in Fig. 7B, D. The fracture of bare WE43 alloy was characterized with the partial plasticity (Fig. 7B) as there are no sharp edges made of particle boundaries.

D. Dvorsky et al. / Journal of Alloys and Compounds 825 (2020) 154016

7

Fig. 6. Mechanical properties of composite materials (CYS ¼ compressive yield strength, UCS ¼ ultimate compressive strength, TYS ¼ tensile yield strength, UTS ¼ ultimate tensile strength, BYS ¼ bending yield strength, UBS ¼ ultimate bending strength): A) Compressive, B) Tensile, C) Bending.

Fig. 7. Fractography with EDS analysis of A) bare Mg, B) bare WE43, C) Mg-24h HF and D) WE43 24h HF.

On the other hand, WE43-MgF2 composite material exerted fragile behavior as the borders between particles are distinguishable similarly as in the case of Mg (Fig. 7D). The fracture surface consists of the specific interface of very fine YF3 phases and a layer of MgF2. Out of these results, it is obvious that the change of mechanical properties after immersion in HF depends strongly on the magnesium alloy and the specific reactions between MgF2 and alloying elements. Longer times of immersion in HF than 24 h resulted in a decrease of mechanical properties for both materials. Magnesium composite materials prepared in this work resulted in similar mechanical properties as materials prepared by spark plasma sintering in the literature sources [37,44e48] (Table 5). Compressive properties of the WE43 composite material are much better than the AZ91 alloy prepared by spark plasma sintering [49]. Also, compressive properties of WE43 alloy are similar to the material produced by additive manufacturing and it exceeds the mechanical properties of bone tissue [50,51]. Nevertheless, the application of such materials would rather be for the non-bearing

applications such as maxillofacial surgery as the products made by extrusion are better for bearing application due to the better mechanical properties [12,52].

3.4. Corrosion properties Prepared samples were immersed in simulated body fluid at 37
C for 14 days in order to investigate the corrosion behavior. The comparison of corrosion rates of the similar materials and the corrosion rates calculated out of the weight changes together with the results from the electrochemical measurements is summarized in Table 6. Magnesium composite materials resulted in more than three times lower corrosion rate compared with the pure Mg. Otherwise, there were no great differences in the corrosion rates depending on the time of immersion in HF. On the other hand, samples were immersed in SBF 1 h before measurement and therefore, the corrosion front was probably already slowed down on the MgF2 interface, so electrochemical measurements revealed a

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Table 5 Mechanical properties of prepared alloys compared with literature sources.

Mg Mg1HF Mg24HF Mg96HF WE WE1HF WE24HF WE96HF Mg [44] Mg [45] Mg [46] Mg [47] Mg [48] WE43 [37] WE43 [50] AZ91 [49] Compact Bone [51] Trabecular bone [51]

CYS [MPa]

UCS [MPa]

D [%]

62 ± 3 82 ± 4 93 ± 4 88 ± 4 221 ± 11 194 ± 9 209 ± 10 217 ± 10

178 201 211 194 415 388 400 359 156

± ± ± ± ± ± ± ±

15.3 17.8 18.1 15.1 22.1 17.4 18.7 14.9 10.2

112 58

170 102 221 367 395 227 205 6.8

8 10 11 10 21 19 20 18

± ± ± ± ± ± ± ±

0.8 0.9 0.9 0.8 1.1 0.9 0.9 0.7

TYS [MPa]

UTS [MPa]

E [%]

BYS [MPa]

UBS [MPa]

B [%]

30 ± 2 31 ± 2 37 ± 2 36 ± 2 178 ± 9 98 ± 5 110 ± 6 25 ± 1

39 ± 2 42 ± 2 59 ± 3 48 ± 2 228 ± 11 142 ± 7 152 ± 8 36 ± 2

3.5 ± 0.2 4.9 ± 0.2 6.6 ± 0.3 4.3 ± 0.2 22.8 ± 1.1 8.7 ± 0.4 10.3 ± 0.5 0.7 ± 0.1

77 ± 4 66 ± 3 92 ± 5 65 ± 3 266 ± 13 249 ± 12 222 ± 11 106 ± 5

81 ± 4 75 ± 4 101 ± 5 72 ± 4 420 ± 21 265 ± 13 241 ± 12 106 ± 5

2.1 ± 0.1 1.8 ± 0.1 2.9 ± 0.1 1.5 ± 0.1 27.5 ± 1.4 4.8 ± 0.2 4.3 ± 0.2 1.5 ± 0.1

70

200 208 148

12.4 135 5.3

though WE43 alloy is known for its good corrosion resistance. The corrosion resistance of WE43 alloy is associated with the creation of stable corrosion products of Y2O3 and Y(OH)3 [12]. The corrosion resistance, therefore, depends on the amount of Y in the solid solution. Otherwise, intermetallic phases work as slightly cathodic places and to some extent increases corrosion rate due to the galvanic corrosion. In this work, the high corrosion rate of bare WE43 is associated with the different yttrium content in the solid solution within the particle [52]. Even though the created Y2O3 barrier around each particle worked similarly to the MgF2 coating and slowed down the corrosion front (Fig. 8F). The creation of the Y2O3 layer consumed Y from the nearby solid solution. Therefore, the depleted solid solution is after breaching the Y2O3 barrier vulnerable to the corrosion due to the galvanic cell between Y rich and Y poor areas. The high corrosion rate is, therefore, associated with the inhomogeneous distribution of Y inside each particle and the cathodic effect of intermetallic phases. The fluoride coating on the WE43 alloy worked differently as well (Fig. 8G). One can see, that WE43-MgF2 composites exerted three times lower corrosion rate compared with the uncoated one (Table 6). The improvement of the corrosion resistance was associated with the creation of a specific interface made of chains of YF3 phases together with the shell of MgF2. The corrosion front is slowed down on the border between particles similarly as in the previous case (Fig. 8G). However, the reaction between MgF2 and Y in the solid solution consumes less Y than reaction with O2 due to the stoichiometry of the products. The gradient between Y poor and rich areas is not that significant. Moreover, the MgF2 present on the surface of the particles prevents the creation of Y2O3. Longer times

great increase of polarization resistance as well as significant decrease of current density. The corrosion potential is slightly shifted to the more negative potential. It seems even a very thin layer of MgF2 is enough to slow down the corrosion magnificently. The different impacts of MgF2 interfaces between particles on the progress of the corrosion fronts are illustrated in Fig. 8. In Fig. 8A one can see that the corrosion of pure Mg proceeds through the vulnerable boundaries between particles. The corrosion front may cause the whole layer of particles to peel away as the corrosion progresses through the material. This corrosion is very similar to the intercrystalline corrosion. On the other hand, the MgeMgF2 composite degraded uniformly as the corrosion front was effectively slowed down on each boundary between particles (Fig. 8B). This interface is made of the hardly soluble MgF2 layer. This barrier effect is responsible for the reduction of the degradation rate of the composite material. Even thin MgF2 coating created after 1 h of immersion in HF was enough to trigger this effect. The effectiveness of the coating can be also observed on the surface of the sample after the removal of the corrosion products (Fig. 8CeE). The surface of MgeMgF2 composite material is covered with the hardly soluble MgF2 shells which slowed down the corrosion front (Fig. 8D and E). After breaching the barrier, exposed insides of particles were consumed by corrosion, but the remains of the MgF2 shells stayed preserved as can be seen in Fig. 8E on the surface after the removal of the corrosion products. The total amount of fluorine on the surface according to the EDS analysis increased from 2 wt% to 16 wt% (Table 7). The corrosion mechanism of the WE43 alloy was different. The corrosion rate of bare WE43 alloy (Table 6) was very high even

Table 6 Corrosion rates calculated from weight changes and electrochemical measurements with the comparison with the literature sources. Sample Mg Mg 1h HF Mg 24h HF Mg 96h HF WE43 WE43 1h HF WE43 24h HF WE43 96h HF Mg/MgeZnO [53] Mg-HA [54] WE43 ex [12] Mg ex [20] AZ31-CaP-MgF2 [17]

vcor [mg/cm2/day] 1.101 ± 0.042 0.351 ± 0.059 0.348 ± 0.043 0.346 ± 0.047 1.553 ± 0.135 0.417 ± 0.037 0.623 ± 0.056 0.875 ± 0.062 22.800/6.000 0.810 0.605 1.762 1.340

E [V] 1.54 1.55 1.61 1.67 1.65 1.63 1.61 1.50 1.53/-1.47

Log I [A/cm2] 3

2.134  10 2.057  103 3.428  104 3.555  104 5.990  105 4.370  105 4.100  105 2.810  105 5.5  104/9.5  105 5.9  105

1.56

2.090  106

Rp [Ucm2] 467 967 4266 4859 9096 11512 10950 10538

D. Dvorsky et al. / Journal of Alloys and Compounds 825 (2020) 154016

9

Fig. 8. Structure of the material after corrosion A) Mg-Cut with EDS analysis of O and Mg, B) Mg 96h HF-Cut with EDS analysis of Mg and F, C) Mg surface after removing of the corrosion products, D) Mg 96h HF surface after removing of the corrosion products with EDS analysis of F and Mg, E) Mg 96h HF surface after removing of the corrosion products detail, F) WE43-Cut with EDS analysis of Y, G) WE43 96h HF-Cut with EDS analysis of Y and F, H) WE43 surface after removing of the corrosion products with EDS analysis of Y, I) WE43 96h HF surface after removing of the corrosion products with EDS analysis of Y and F, J) WE43 96h HF surface after removing of the corrosion products - detail.

Table 7 The amount of F and Y on the surface by EDS analysis before and after corrosion and removal of corrosion products (wt.%).

Before After

Element

Mg 1h HF

Mg 24h HF

Mg 96h HF

WE43 1h HF

WE43 24h HF

WE43 96h HF

F Y F Y

0.4 e 0.8 e

1.4 e 12.0 e

2.0 e 16.0 e

1.1 4.0 3.4 7.1

1.4 3.8 3.3 5.9

1.5 3.9 4.7 8.7

of immersion in HF increased the volume fraction of YF3 phases. It also increased the thickness of the MgF2 layer especially in the areas near the intermetallic phases on the powder surface before sintering. Therefore, with longer immersion time in HF, there were increasing differences between regions rich and poor on F. Also, a higher amount of Y is consumed from the solid solution for the creation of YF3 phases due to the presence of a higher amount of F. This inhomogeneity might be the reason behind the increasing corrosion rate with longer times of immersion in HF (Table 6). The greatest improvement of corrosion resistance was observed for the sample immersed for 1 h in HF (Table 6) as a low amount of F was enough to prevent the creation of Y2O3, while the amount of YF3 phases remained low. This prevented the consumption of Y from the solid solution as a lower amount of Y is needed for the creation of YF3 phases than Y2O3 due to the stoichiometry. Regions poor on Y could be easily attacked and such irregular corrosion may lead to pitting corrosion, which may go deep in the material and beneath the surface layer. Corrosion products which are more space demanding then cause pressure from inside of the material similarly as in the case of pure Mg. This action results in the peeling of the uncorroded particles on the surface due to the bad mechanical properties of the interface between particles. Increasing the corrosion rate with longer immersion time in HF is, therefore, also attributed to the poor mechanical properties. The effectiveness of the coating is illustrated on the surface after the removal of the corrosion products (Fig. 8HeJ). The surface of the composite materials is covered with very fine YF3 phases (Fig. 8J). The amount of F increased from 1.5 wt% to 4.7 wt% after corrosion (Table 7). The increase of the amount of F is not that evident that in the case of MgeMgF2 composite which is associated with the lesser effectiveness of the coating. One can see, that there was also an increased amount of Y on the surface of the composite due to the

presence of YF3 phases. On the other hand also bare WE43 alloy was characterized with an increased amount of Y which indicates that also yttrium oxide worked as a barrier in this alloy (Fig. 8H). Results of corrosion measurements revealed great differences between the impact of fluoride protection on pure Mg and WE43 alloy. Fluoride treatment improved the corrosion resistance for both alloys, however, in the case of Mg longer time of immersion in HF resulted in better corrosion resistance while the opposite was true for WE43 alloy. In comparison with other composite materials prepared by SPS method [53,54] products prepared in this paper were characterized with better corrosion resistance. Composite materials exerted even better corrosion resistance than the same extruded materials [12,20] or AZ31 alloy with CaPeMgF2 coating [17]. Out of the obtained results we can assume, that fluoride network has a positive effect on the corrosion properties. The corrosion rate is three times lower compared with uncoated samples and the corrosion rate is even very low compared with other researches. Moreover, there is even hidden potential to further increase the corrosion resistance by the creation of the coating on the surface of the sintered material. 3.5. Cytocompatibility e test with the extracts The cytotoxicity tests were performed only on the samples immersed for 96 h in order to determine the effect of fluoride treatment and YF3 phases. The concentration of Mg and other elements is shown in Table 8. The amount of released Mg into the media was much higher in the case of Mg compared to WE43. On the other hand, bare WE43 alloy compared to pure Mg was characterized by a higher corrosion rate in SBF. It could be due to the occurrence of Y in the

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D. Dvorsky et al. / Journal of Alloys and Compounds 825 (2020) 154016

Table 8 Concentration of the elements in the extracts of the samples (ng/ml). Sample WE43 WE43 96h HF Mg Mg 96h HF

Concentration (ng/ml) Y 28.9 ± 11.6 92.6 ± 14.2

Nd 20.5 ± 8.5 64.5 ± 10.1

Gd 4.5 ± 1.9 14.4 ± 2.2

Dy 3.5 ± 1.4 11.6 ± 1.7

Mg 64833 ± 3253 81500 ± 6000 267 000 ± 2646 219 000 ± 13528

Fig. 10. SEM images of L929 cells growing on the surface of A) WE43, B) WE43 24h HF.

4. Conclusion Fig. 9. The relative metabolic activity of L929 cells exposed to extracts from Mg and WE43 alloy (cells treated with sole MEMþ5% FBS used as a negative control had 100% metabolic activity). The line stands for normative cut-off for cytotoxicity. Error bars represent SSD from hexaplicates.

corrosion products, which probably slows down the dissolution of corrosion products into the medium. Otherwise, the corrosion products are removed when it is calculated out of the weight changes. In the case of pure Mg, the release of Mg was (p < 0.01) decreased after the treatment with HF. On the other hand, HF treatment of WE43 significantly (p < 0.05) increased the amount of Mg and RE released into the medium. It was the opposite effect of the coating compared to the test in SBF. It could be caused by rapid initial corrosion of the surface, which is significantly decreased over time as the corrosion front slows down on the fluoride interface in case of immersion tests in SBF. On the other hand, the extracts were taken after 24 h of immersion, where the positive effect of fluoride barrier is less evident. The metabolic activity of L929 cells incubated with the extracts of the samples is summarized in Fig. 9. One can see decreased metabolic activity (around 70%) in the case of bare Mg, which is probably associated with its high corrosion rate in the medium. Neither of the other tested materials shows cytotoxic effect towards L929 cells (metabolic activity decrease below 70%). Fluoride coating significantly improved cytocompatibility for pure magnesium, but slightly decrease it for WE43 alloy. Nevertheless, even coated WE43 alloy exerted metabolic activity above the normative limit of 70%. Therefore, out of the obtained cytotoxicity tests, fluoride coating is fully cytocompatible. Moreover, direct contact tests showed that more cells with healthier morphology were present on the surface of the fluoride coated WE43 samples compared to the uncoated sample (Fig. 10). It seems that after one day of preincubation, the coated surface is more acceptable for the cells as they can adhere and proliferate better. It could be due to the lower corrosion rate after the preincubation when the cells are not disturbed with the released hydrogen.

Powders of pure Mg and WE43 alloy were successfully immersed in HF for 1, 24 and 96 h to form protective MgF2 coating around each particle. Such treated powders were compacted by spark plasma sintering to form a continuous network of MgF2 coating. Continuous network with a thickness of about 3.5 mm was successfully created on the pure Mg after 24 h of immersion. Alloy WE43 exerts resistance against hydrofluoric acid and the creation of the continuous network was much more difficult. Intermetallic phases and nearby surroundings were enveloped by the MgF2 layer. The MgF2 coating partially converted into YF3 phases after sintering. So the final shell of each particle was inhomogeneous. Mechanical properties of pure Mg improved after immersion in HF, while the best mechanical properties were achieved after 24 h of immersion. This improvement was due to the higher adhesion strength of the MgF2 than MgO to the Mg matrix. WE43 alloy exerted loss of mechanical properties after immersion in HF due to the creation of YF3 and MgF2 around each particle. Immersion tests in SBF revealed that protective coating reduced corrosion rate on one third and one fourth for magnesium and WE43 alloy respectively. Specific corrosion behavior was observed for uncoated and coated samples. Different kind of protection was also observed for Mg and WE43 coated samples. The improved corrosion resistance of MgeMgF2 composite was due to the homogeneous layer of MgF2 between each particle, while WE43 alloy was characterized with YF3 phases surrounded by the MgF2 layer. According to the results, the MgeMgF2 composites prepared by spark plasma sintering have the potential to be used as biodegradable materials. The final mechanical and corrosion tests depend on the time of immersion in HF. Shorter immersion times seems to significantly improve corrosion properties, while the mechanical properties are the least affected. Nevertheless, the result strongly depends on the alloy which is applied for this purpose. Moreover, the cytocompatibility tests revealed no negative effect of fluoride treatment on the cells in the case of WE43 alloy. It can be therefore assumed that the YF3 phases have no negative effects on the cells.

D. Dvorsky et al. / Journal of Alloys and Compounds 825 (2020) 154016

Declaration of competing interests 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. CRediT authorship contribution statement Drahomir Dvorsky: Visualization, Writing - original draft, Formal analysis, Funding acquisition. Jiri Kubasek: Conceptualization, Resources. Eva Jablonska: Formal analysis, Validation. Jirina Kaufmanova: Investigation, Data curation. Dalibor Vojtech: Supervision, Writing - review & editing.

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[23]

[24]

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[26]

Acknowledgment

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Authors wish to thank specific university research (MSMT No 21-SVV/2019) for the financial support of this research.

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