Drhomír Dvorsky ,Ji?í Kubásek ,Mihel Roudniká,b ,Filip Pr°u? ,Dvid Ne?s ,Peter Minárik ,Jitk Stráská,Dlibor Vojtěh

a University of Chemistry and Technology,Faculty of Chemical Technology,Department of Metals and Corrosion Engineering,Technická 5 166 28 Praha 6 -Dejvice,Prague,Czech Republic

b Institute of Physics,Czech Academy of Science,Na Slovance 1999/2 182 21 Praha 8,Prague,Czech Republic

cDepartment of Physics of Materials,Charles University,Ke Karlovu 5,Prague 12116,Czech Republic

Abstract Powder metallurgy is a powerful method for the preparation of materials with superior properties.This work aimed to investigate the effect of powder size on the microstructure,mechanical,and corrosion properties of advanced WE43 (Mg-4Y-3REE-Zr) alloy prepared by spark plasma sintering (SPS).At the same time,the effect of HF pre-treatment of the powder on the properties of fina compacted products is studied.Smaller powder particles yielded microstructure with more interfaces formed by Y2O3,or MgF2 and YF3.These interfaces work as barriers against corrosion,which greatly improves corrosion resistance.The suggested pre-treatment of powder in HF further reduced the corrosion rate of the compacted materials.On the contrary,fragile interfaces of YF3 decreased mechanical properties as the crack primarily propagates through these interfaces.The original powder containing the mixture of all powder fractions exerted the best combination of mechanical properties.Powder size has also shown to affect ignition temperature.The highest ignition temperature was measured for the fines powder fraction.

Keywords: Powder metallurgy;Mechanical properties;Corrosion;Interface;SPS.

1.Introduction

Magnesium is a metal with characteristics interesting for biomaterials:great biocompatibility,good mechanical properties,and Young′s modulus very close to the human bone [1].As Mg spontaneously degrades in the body,it can be used for biodegradable implants for bone fixatio or as cardiovascular stents [2].However,there is a great problem with the corrosion of magnesium.Its degradation rate is so high that it can cause medical complications by heavy hydrogen generation[3].There is also a danger of an enormous decrease in the mechanical properties if inhomogeneous corrosion of the implant occurs.Therefore,mechanical and corrosion properties have to be improved by specifi treatment which would improve the behaviour of an implant in the organism and would make it more suitable for the application in medicine [1].

The most common type of treatment that improves both the mechanical and corrosion properties of magnesium is alloying.It is important to choose the right element which would not affect the biocompatibility of the material [4].There exists a certain amount of magnesium-based alloys,but only a fraction can be used in the human body.The material should have a reasonable degradation rate so that the released corrosion products would be below the daily allowance limits for the particular element [5].Otherwise,a higher quantity can cause an allergic or toxic reaction.The most advanced magnesium alloys contain rare earth elements (REEs) which greatly improve both mechanical and corrosion properties[6-8].However,there are debates about the influenc of biomaterials with REEs on the human organism as long-term studies are still missing.It has been shown that these elements may accumulate in the livers and kidneys if they are consumed in food[9].Y and Nd seem to be non-toxic and behave as inert in the human body environment [10,11].Moreover,some of the REEs are used for cancer treatment for their anticarcinogenic effect [12-14].Generally,REEs with higher solubility in Mg matrix have lower influenc on the cytotoxicity compared to the ones with lower solubility [11].Nevertheless,cytotoxicity tests revealed no negative effects of REEs in the short-term studies.Also,many in-vivo tests were performed with positive results [10,15-17].

The improvement of mechanical properties is associated with the solid solution strengthening and strengthening by intermetallic phases.Precipitation hardening is also possible due to the partial solubility of REEs in magnesium solid solution[18].Corrosion resistance improves with a higher concentration of alloying elements in the solid solution [19].If intermetallic phases are present,these may work as cathodic places due to the different potential between them and the magnesium matrix.Nevertheless,most RE intermetallic phases have quite similar corrosion potential to the magnesium matrix,so the danger of galvanic corrosion is reduced compared to for example AZ (Mg-Al-Zn) alloys,where intermetallic phases increase the corrosion rate [20].The element which greatly improves the corrosion resistance of magnesium is Y.The most advanced alloy WE43 (Mg-4Y-3REE-Zr) is known for its good corrosion resistance which is associated with the enrichment of corrosion products by Y.Yttrium is incorporated in the corrosion products in the form of Y2O3and Y(OH)3which slows down the diffusion between the material and the corrosion medium [21].This layer is easily created if the amount of Y in the solid solution is high enough [22].Otherwise,an inhomogeneous structure with large intermetallic phases may lead to localized forms of corrosion as intermetallic phases usually stay incorporated in the corrosion products without any change.This greatly increases the corrosion rate of WZ(Mg-Y-Zn)alloys,where there are LPSO phases with a very large surface.The negative effect of intermetallic phases might be removed by proper heat treatment [23-26].Generally,magnesium alloys are very sensitive to impurities like Fe,Ni,and Co,which greatly deteriorate corrosion resistance due to the galvanic corrosion as those impurities are not soluble in the Mg matrix.Therefore,those elements have to be held below the allowance limits (Fe:35-50ppm,Ni:20-50ppm,Cu:100-300ppm) [23].Nevertheless,some alloying elements can absorb those impurities and therefore reduce their negative effect [27].

Material properties also strongly depend on the preparation method.It is known that the progressive method of powder metallurgy improves mechanical and corrosion properties[21].The improvement is associated with small grain size and a high amount of alloying elements in the solid solution due to the rapid cooling.The fina material is usually characterised by a fine-graine homogeneous structure with a high amount of alloying elements in the solid solution and very fin intermetallic phases.Properties depend highly on the powder processing and method of compaction [28].The advanced method of sintering with the ability of fast production is spark plasma sintering (SPS).This method consists of powder pressing,heating,and high current which fl ws through the powder.Places between particles have high resistance which produces Joule’s heat.This process takes only several minutes and therefore,the structure of the powder may stay almost unchanged after sintering [29].The fina material is therefore fine-graine and with a high amount of alloying elements in the solid solution [30].

The Joule’s heat during sintering is associated with the resistance of the materials and the interface between them[31-33].Therefore,more interfaces with high resistivity might improve the effectiveness of sintering.This can be done by powder separation by size through sieving or altering the surface of particles.

Another effective method for the reduction of corrosion rate is the application of coatings.Conversion coatings usually have a better adhesive strength.MgF2conversion coating is relatively thin (up to 4μm),highly effective (reduction of corrosion rate by up to ten times),and easy to prepare by immersion in HF [7,24,34-37].Some danger exists if coating is disrupted,which leads to localized forms of corrosion.

Powder treatment might be also interesting for the additive manufacturing of magnesium alloys.In this process,there is a great risk of possible ignition of the magnesium-based powder inside the chamber of the 3D printer [38].Therefore,most researchers are focused on the additive manufacturing of WE43 alloy which has a relatively high ignition temperature especially due to the high amount of Y in the solid solution[39-42].A recent paper reveals that it is now possible to prepare the material with a high density and reasonable mechanical properties by 3D printing [42].Possibly,pure magnesium could by processed too with an appropriate powder treatment as it was reported that the ignition temperature of the pure magnesium powder might be increased up to 750°C by HF treatment [43].

In this work,the effect of particle size and HF treatment on mechanical,corrosion and ignition properties of the powder and SPS products of WE43 alloy is investigated.

2.Materials and methods

2.1.Powder treatment

Materials were prepared from the commercially available atomized powder of WE43 alloy (LUXFER MEL TECHNOLOGIES).Its composition,including impurities (Fe,Ni,Cu),is listed in Table 1.The powder of WE43 alloy consisted of particles with sizes ranging from 5 to 200μm.Powder fractions were separated by a sifter with the holes sizes of 25,36,45,63,100,125,and 180μm.As a result,fractions with the size of particles between 25 and 36μm,36-45μm,45-63μm,63-100μm,100-125μm,and 125-180μm were separated.For specifi surface treatment,10g of each powder fraction were stirred in a plastic container with 38-40% HF.The container with the powder and HF was left without stirring for a while to settle the powder and then HF was poured out.The powder was subsequently washed with distilled water and ethanol several times.Subsequently,the powder was put in a dryer at 60°C for 24h.

Table 1 The composition of the WE43 atomised powder.

Table 2 Fraction of interfaces between particles [%] and the amount of individual elements in the structure [wt.%].

Fig.1.Time profile of the SPS process conditions.

2.2.Powder compaction

Each fraction of the treated and non-treated powder was processed by the SPS method using an HP D 10 FCT machine with a GmbH system.The powder was fille into a graphite die,pre-pressed at 9MPa and subsequently,compressed at 32MPa at a temperature of 500°C for 10min.The heating rate was set to 100°C·min-1.The pressure inside the chamber was 3kPa.The fina sample in the shape of a cylinder had a diameter of 20mm and a height of 15mm.The diagram showing the time profile of the process conditions during sintering is presented in Fig.1.

2.3.Microstructure

Samples were ground on SiC grinding papers,polished on a diamond paste D2 and a SiO2suspension Etosil E.Microstructure was studied using an electron microscope (SEMTESCAN VEGA3)with an EDS analyser(AZtec-Oxford Instruments).Porosity and the area of particle boundaries were evaluated by an image analysis in ImageJ software by processing 10 images for each material.

2.4.Mechanical properties

Mechanical properties in compression and tension were measured by a universal LabTEST 250SP1-VM machine on samples with the dimensions of 5mm x 5mm x 7mm for the compressive tests and bone-shaped samples with 10mm in height and 2mm x 4mm in the cross-section for the tensile tests.The average mechanical properties were evaluated from three measurements.

2.5.Corrosion properties

Corrosion behaviour was studied in the simulated body flui (SBF) prepared according to Müller [44].The starting pH was set to 7.4 at 37°C.An SBF volume of 100ml per 1 cm2of sample surface was used.Samples were fi ed in plastic holders and immersed in SBF for 14 days at 37°C.The corrosion rate was evaluated based on the weight changes after the removal of corrosion products and from the amount of magnesium ions released to the solution,measured by ICPMS (Elan DRC-e).Three measurements were performed for all materials.

2.6.Ignition temperature

Approximately 2g of a tested powder were put into the Al2O3crucible which was inserted into the chamber of a resistance furnace.One thermocouple was in the direct contact with the powder and the other one was situated in the middle of the crucible.The tests were performed at a stable fl w of technical air (100 l·h-1).The heat rate corresponded to the 25°C·min-1.The temperature was regularly increased with time during the test until the rapid increase due to the ignition of the sample.Each sample was measured three times.

3.Results

3.1.Microstructure

Particles of the used WE43 powder were characterised by a regular spherical shape typical for gas-atomised powders.Nevertheless,some particles were surrounded by small satellite particles.Fig.2 shows the inner structure of a particle.In a typically dendritic structure,dendrites consist of anα-Mg solid solution which contained about 3.0wt.% of Y,1.7wt.%of Nd and 0.3wt.%of Gd and Dy.The interdendritic area was enriched in alloying elements,thus containing approximately 6.6wt.% of Y,4.7wt.% of Nd,0.7wt.% of Dy and 0.4wt.%of Gd.X-ray diffraction revealed the existence of metastableβ1-phase,which is compositionally very close to theβ-phase,Mg14Nd2Y (Fig.3).

Fig.2.Microstructure of powder particles:A) original powder of the WE43 alloy,B) HF-treated powder.

Fig.3.Diffraction of powders and compacted products.

The powder immersion in HF resulted in an inhomogeneous coating (Fig.2B),which is in accordance with the finding of other authors[45,46]for bulk WE43 alloy.A coating of MgF2was formed predominantly near the intermetallic phases and the areas without intermetallic phases were covered just by a very thin coating.This is related to the corrosion resistance of WE43 alloy and the difference between the potentials of the matrix and the intermetallic phases [45,46].

After powder fractionation,each fraction was processed by SPS.Final microstructures are shown in Fig.4.Individual particles are very well recognizable in all microstructures.As can be noticed,the microstructures of materials prepared from larger particles contain smaller particles too.That is because these small particles stacked on the bigger ones during atomization and were not separated during sieving.Moreover,particles with a size lower than 100μm were characterised by a relatively round shape,while larger ones became more hexagonal in order to fil large gaps in between powder particles during sintering (Fig.4).The microstructure of WE43 after SPS (Figs.3,4) consisted ofα-Mg solid solution and phases identifie as Mg24Y5andβ-phase (Mg14Nd2Y).βphase was formed from metastableβ1-phase during the compaction by SPS.Samples prepared from the pre-treated powder were macroscopically very similar,thus are not presented here.The porosity of all samples was less than 0.3%.The addition of the HF pre-treatment step led to a decreased porosity (less than 0.1%).This may be due to the higher resistance of MgF2compared to the metal matrix which increased the Joule’s heat during SPS and improved the compaction.

A more detailed view (Fig.5B) reveals that new phase(β-phase (Mg14Nd2Y)) precipitated predominantly at grain boundaries.Grain size after sintering ranged from 1 to 10μm.Although the materials prepared from larger powder particles had slightly larger grains,the difference was almost neglectable (7±3μm and 6±2μm for the 125-180μm fraction and 25-36μm fraction,respectively).The main difference between the materials sintered from the pre-treated and untreated powders is in the interface between particles(Fig.5).

Fig.4.Microstructures of the materials prepared from the specifi fractions of the original untreated powder -SEM images.

Fig.5.Details of the interface between particles in sintered products:A) WE43,B) WE43 -HF.

The interface amongst the particles of the untreated powder consisted of continuous layers of Y2O3(Fig.5A).Even though the SPS process was performed under near vacuum conditions,residual air was trapped in between the powder particles,which allowed alloy oxidation.Yu et al.[47] studied the annealing of WE43 alloy and discovered that if the alloy is exposed to a higher temperature the layer of Y2O3is created on its surface from the Y in the solid solution and this layer protects the material against further oxidation.This barrier should also theoretically decrease the corrosion rate as Y2O3is hardly soluble.Normally,Y2O3is formed on the surface of the WE43 alloy during corrosion together with Y(OH)3,which slows down the diffusion of corrosion media to the surface of the sample [22].

In the case of coarse powder fraction (125 -180μm),the effect of continuous Y2O3layer at particles interface on the corrosion resistance might be,however,adverse.The formation of the Y2O3layer on the surface of powder particles during rapid SPS compaction caused the depletion of the solid solution by Y.The area below the Y2O3layer thus containedonly 0.8±0.2wt.% of Y,while the centre of powder particles contained 3.8±0.2wt.% of Y.This gradient of depleted zone reached up to 15μm from the interface to the centre.The levels of other elements were the same.Such inhomogeneity is unfortunately undesirable from the corrosion point of view;after the disruption of the surface layer of Y2O3,corrosion propagates easily through these continuous Y-depleted areas.

In the case of fine powder fractions (25 -36μm),the difference in Y concentration between the centre of the particles and the area close to their interface was much lower.There was about 2.8±0.1wt.% of Y in the centre of the particles versus 1.9±0.1wt.% near the interface.Although the total amount of alloying elements (including Y) remained almost the same in all materials (Table 2),lower Y amount was detected in the solid solution in fine particles.(The different amount of Y in the fin particles can be associated with a shorter diffusion route from the centre to the surface of these particles.Also,the cooling rate during atomization has an important effect.As a consequence of higher cooling rates for small particles during atomization,the solid solution is more oversaturated by Y than in the case of coarse particles.Higher amount of Y in the solid solution then favours the formation of Y2O3on the interface,and so lower concentration of Y in the solid solution and higher content of Y2O3at the particle interface is observed.

Powder particles that were pre-treated in hydrofluori acid were covered by an inhomogeneous coating of MgF2(Fig.2B) that protected the material against oxidation during sintering,therefore no additional Y2O3was formed at particles interface.On the contrary,MgF2reacted with the Y in the solid solution and formed very fin YF3phases(Fig.5B).Previously,it has been observed that MgF2on WE43 alloy is predominantly created near intermetallic phases due to the differences in potentials between intermetallic phases and Mg matrix [45,46].At room temperature,the reaction between Y and MgF2is not possible [45,46],however,increased temperature during sintering allows it due to the diffusion and lower standard molar Gibbs energy of YF3formation (-1871kJ/mol at 527°C) compared to that of MgF2(-1190kJ/mol at 527°C) [36,48,49].The fina SPS compacts,therefore,consisted of residual MgF2and very fin(<1μm) YF3phases (Fig.5B).Presented reaction consumed Y from the solid solution ofα-Mg similarly as Y2O3,however,the observed depletion was not so significan due to the different stoichiometry and low amount of MgF2.In the coarse powder (125 -180μm),the solid solution near the interface contained about 2.8±0.2wt.% of Y compared to the 4.0wt.% of Y in the centre.The fin powder (25 -36μm)contained about 2.1±0.1wt.% of Y near the interface and 2.8±0.1wt.% in the centre.

As a consequence of the different size of powder particles,there exist variations in the volume fraction of particle interfaces (consisting of oxides or fluorides) as shown in Table 2.These values were calculated by image and EDS analysis and correspond to the percentual area fraction of particle interfaces.The materials prepared from powder fractions with a lower particle size were characterised with a higher volume fraction of oxides or fluorides The interface fraction affects the total amount of F in the material as well as fina mechanical and corrosion properties (Table 2).

The interface fraction increased with the lower particle size(Table 2).This dependency can be demonstrated using a simplifie model of an area fille with circles of a certain diameter in the tightest arrangement:

The perimeter of an individual particle is proportional to the average particle diameterdaccording to the Eq.(1).The sum perimeter of all particles is gained by multiplying the perimeter of one particle by the number of particles in a certain area.The surface area of the interface is then gained by multiplying half of the measured thickness of the Y2O3/MgF2surface layer by the sum perimeter of all particles in that area.Finally,the interface fraction is gained by dividing the total perimeter of all particles by the area which gives us the Eq.(2),whereAis the percentual fraction of interfaces,xis the size of the examined area,ais the thickness of the layer anddis the average particle size.After simplification we get the Eq.(3).This equation is represented by the black line in Fig.6 and is compared with the measured results of the fraction of interfaces obtained by image analysis.Those results fi well with the proposed equation,therefore the fraction of interfaces in the materials could be estimated by this equation.

Fig.6.Dependence of interface fraction on the particle size.

3.2.Mechanical properties

Compressive and tensile mechanical properties were measured.The results of these tests are summarized in Fig.7.Compressive yield strength (CYS),ultimate compressive strength (UCS),relative deformation (D),tensile yield strength (TYS),ultimate tensile strength (UTS),and relative elongation (A) were evaluated.The measurements were performed without an extensometer due to the small size of the specimen.Therefore,the values of deformation (D) and elongation (A) should be considered only for the comparison amongst presented materials.

Fig.7.Compressive (A) and tensile (B) properties of the prepared materials.Powder fractions are designated on the x-axes.

There is a trend of compressive properties slightly improving with increasing particle size for materials prepared from non-treated powders (Fig.7A).As compressive properties are less sensitive to defects and inhomogeneities,this trend is more visible in case of tensile properties (Fig.7B).This trend might be partly associated with the amount of Y in the solid solution,which is increased for larger particles.The Y from the solid solution in the small particles was consumed for the Y2O3interface,while centres of the large particles stayed saturated by Y.Nevertheless,the interface fraction is considered as the main factor affecting the observed differences in mechanical properties.Materials prepared from fine powder fractions contain more interfaces amongst powder particles.These interfaces are formed by oxides or flu orides which cause fragility.The comparison between materials prepared from smaller and larger particles is visible on fracture surfaces in Fig.8.In the material with smaller particles and high interface fraction,the crack propagated primarily through these interfaces (Fig.8D).Compared to that,the combination of ductile fracture (dimple like morphology)with an occasional brittle fracture through the particle interface occurred in the material with large particles (Fig.8C).Evidently,a higher amount of Y2O3at particle interfaces increases brittleness.The material prepared from the powder fraction of 125-180μm exerted the best mechanical properties amongst the products of other powder fractions.Yet,in terms of mechanical properties,especially elongation,a significantl superior material was prepared from the original powder mixture.This relates to the fact that in the original powder,the mixture of various-sized particles has a better space fillin capacity,moreover it makes the way of crack propagation through interfaces longer and more complicated.

In the crack propagation,the influenc of grain size and grain boundaries are considered as one of the main factors affecting the fracture mechanism [50-53].Even though the grain boundaries and particle interfaces are different subjects,some parallels might be found.Ma et al.[54] and Qui et al.[55] found out that crack growth is decelerated with an indirect path of crack propagation.Therefore,crack propagates easily if it goes straight and mostly through particle interfaces as is proposed in Fig.9A and supported by a cross-section in Fig.9E (red lines).It is evident that there exist just a few ductile transparticle cracks (blue lines).Moreover,He et al.[56] observed that even though crack propagation at the grain boundary require less energy,the crack might spread through the grain along slip bands if it makes the crack path more straight.This might explain the more ductile behaviour of the materials with larger particles,as it is probably easier for the crack to spread through the particle rather than change the path as is illustrated in Fig.9B (blue line) and supported on the cross-section in Fig.9D.This is also evident from Fig.8C,where a lower amount of surface of original particles is clearly evident on the surface of the fracture of samples prepared from larger powder particles compared to the small particles (Fig.8D).

If the particles are too small,the amount of brittle interfaces is huge,the crack propagates more directly through this interface,and therefore,plasticity,TYS and UTS are reduced.On the contrary,if the particles are large enough,the particle interfaces are relatively long and linear.If the interface is favourably orientated,then the crack might propagate fast in this direction for a long distance,which would result in a lower strength.

Samples prepared from the powders pre-treated in HF exerted almost similar behaviour under compression tests like materials from untreated powders (Fig.7A),although they were characterised by worse mechanical properties in tension (Fig.7B),especially much lower ductility.This can be attributed to the presence of the brittle YF3phases at the particle interfaces (Fig.5B).The interface containing small particles of YF3is much more fragile than Y2O3forming a continuous film Compared to the untreated samples,the highest TYS and UTS were measured for the materials with midsized particles(45-63μm and 63-100μm).This phenomenon can be again explained by multiple contradictory factors.

The amount of Y in the solid solution decreases with lower particle size due to the intensifie depletion of solid solution by Y consumed for the formation of the YF3layer(Fig.5B).Also,the amount of brittle YF3interfaces increases with lower particle size.Smaller particles then allow macroscopically direct crack propagation through the brittle interfaces (Fig.9E).Those factors should lead to the improvement of mechanical properties with increasing particle size.On the other hand,microstructures formed by smaller particles should improve the quality of particle bonding due to the increased number of contact points between powder particles improving heat transfer during sintering.Smaller particles also complicate crack propagation;the crack has to change direction frequently (Fig.9A).Moreover,particles larger than 100μm characterised with a hexagonal shape (Fig.4) create almost linear interfaces,which allows direct crack propagation through the interface for a long distance (Fig.9B -red lines).Conversely,these factors should lead to the decrease of mechanical properties with increasing particle size.In the global point of view,the contradictory factors are thus compensated for the mid-size chemically treated particles,providing these materials with the best mechanical performance.The observed peak in mechanical properties (Fig.7B) can be expected for the untreated material too.However,due to the better ductility of the Y2O3interface,it is probably wider and shifted to a larger particle size.In an extreme condition where there would be just two large particles,the crack would propagate through the interface resulting in a low strength equal to the adhesive strength of Y2O3.Similarly,in case of the smallest and largest particle sizes of chemically treated powder,the crack propagated predominantly through the particle interface containing MgF2/YF3(Fig.9E,F),which means that the measured UTS is close to the adhesion strength of the YF3to the matrix (～47MPa) or to the adhesion strength of MgF2to the matrix (33-43MPa) [57-59].

Fig.8.Fracture surfaces of A) WE43 mix,B) WE43 HF mix,C) WE43 125-180μm,D) WE43 25-36μm.

As was discussed above,there were some analogies between grain size and particle size.It is thus interesting to check the similarity for the Hall-Petch relation as well.In Fig.10,there is a dependency of TYS on particle size.The materials prepared from the untreated powder meets pretty well the inverse Hall-Petch relation according to Eq.(4),wheredis the average particle size:

Interestingly,materials prepared from HF pre-treated powders almost follow this equation for three distinct average particle sizes.In these particular materials,the contradictory factors mentioned above led to the improved properties.On the contrary,the remaining materials showed much lower strength as the crack propagated predominantly through the direct brittle interface.Eq.(4) seems reasonable also due to the subsequent consideration:If the particle size would be large enough,then the second member of Eq.(4)would be neglectable and the value ofσwould represent the yield strength of the material without particle interfaces.In such case,the value of 155MPa is in accordance with the measured values of TYS for as-cast WE43 alloy [60].

Muhammad et al.[61] studied alloy AZ31 and pure Mg prepared by SPS at different temperatures.Although the hardness of AZ31 alloy was the same regardless of sintering temperature,there was a difference in the tensile properties.They found out that UTS is improved with increasing sintering temperature because oxides on the surface of particles are removed by the SPS process at higher temperatures.The maximum value of observed UTS for AZ31 alloy was 160MPa.Opposite results were observed by Minarik et al.[62] during tests on alloy AE42,where the compressive properties were reduced with increasing sintering temperature from 440°C to 550°C.This behaviour was justifie by the formation of coarser Al11RE3incoherent phases in the matrix.Also,there was an increase in the amount of MgO caused by oxidation of the surface.Despite these difficulties the material exerted much better results than the extruded one.Mondet et al.[63] confirme that AZ91 alloy prepared by SPS can achieve better mechanical properties than a cast ingot.There was an improvement by 18%in hardness and an improvement in CYS by 68%.Spark plasma sintering was designated as a very promising method of material preparation as mechanical properties were superior compared to the cast ingot.

Fig.9.Suggested propagation of the crack through materials prepared from small and large powder fractions (A,B) and cross-sections of fractures of those materials prepared from untreated and pre-treated powder (C,D,E,F).

3.3.Corrosion properties

Corrosion tests were performed in SBF for 14 days at 37°C.The fina corrosion rate was calculated from weight changes after removing the corrosion products and from the amount of released Mg2+ions into the solution.The results are presented in Fig.11.Values calculated according to the measured amount of released ions are slightly lower as the corrosion products originally formed on sample surface are not included.

Fig.10.Dependency of tensile yield strength on the particle size.

The high corrosion rate of WE43 prepared by SPS is associated with the inhomogeneous structure and specifi Y2O3interface between original powder particles which protects the material by the barrier effect.The corrosion front is slowed down on such an interface as can be seen in Fig.12.However,after overcoming the barrier,the corrosion front progresses very fast as the solid solution near the interface is poor in Y which was consumed for the formation of Y2O3.Generally,the good corrosion resistance of the WE43 alloy is associated with the high amount of Y in the solid solution,which can easily incorporate into the corrosion products in the form of Y2O3or Y(OH)3[21].Therefore,the corrosion resistance is weak in the areas which are poor in Y.This effect is amplifie by the formation of galvanic corrosion between the Mg matrix poor in Y and intermetallic phases working as cathodic places.Despite lower concentration of Y in the solid solution of fine particles,better homogeneity and higher amount of particle interfaces led to an improved corrosion resistance.As a consequence,the corrosion rate was reduced from 5.6 mm·y-1for materials prepared from coarse powder particles to only 1.3 mm·y-1for materials prepared from fin powder particles.

Materials prepared from HF pre-treated powders corroded similarly to the untreated ones,see Fig.12.Here,YF3with the residues of MgF2exerted the barrier effect.As the concentration of Y in the solid solution near the particle interface was not reduced as much as in the previous case,the corrosion resistance was partially preserved,and improved by additional fluoride on particle interface.Chiu et al.[64] observed the corrosion rate almost three times reduced due to the 1.5μm thick MgF2layer.In the work of T.Yan et al.[65],where the effect of hydrofluori acid on alloy AZ31B was studied,the formed layer of MgF2and MgO reduced the corrosion rate by up to four times.The reduction of the corrosion rate of SPS products of chemically treated powders was also observed in previous works [34,36].The influenc of the volume fraction of the particle interface is illustrated in Fig.13.With higher interface content,corrosion rate decreases.

Fig.11.Corrosion rates based on weight changes and the amount of released Mg2+ ions.

Fig.12.Progress of corrosion front A) WE43 25-36μm,B) WE43 125-180μm,C) WE43 HF 25-36μm D) WE43 HF 125-180μm.

3.4.Ignition temperature

Ignition temperatures were measured on the individual powder fractions before sintering.One can see,that the ignition temperature paradoxically increases with lower particle size (Fig.14A,Table 3).This is the opposite trend than observed for pure Mg powder previously [43].In the case of pure Mg,lower particle size means greater total surface which can be oxidized.Pure Mg is unable to create a protective oxide layer and,therefore,burns easily.On the contrary,WE43 alloy is characterised by the formation of surface Y2O3which protects the material against further oxidation.

Table 3 Ignition temperature of individual powder fractions.

Fig.13.Influenc of particle interface area on the corrosion rate.

Fig.14.Representative curves of ignition properties of powders:A) Fractions of atomised powder,B) Fractions of HF treated powder.

Nevertheless,the original powder mixture of WE43 alloy ignites approximately at the same temperature as an ingot of pure Mg (655°C) [66].That can be explained by the effect of particle size.The ignition temperature depends on the quality and protective effect of surface layer of Y2O3formed during powder oxidation.As smaller powder particles are more supersaturated with Y and require shorter diffusion path for Y to reach the surface,better quality of Y2O3surface layer is achieved.That explains why lower powder fractions exerted higher ignition temperatures (Fig.14A,Table 3).However,in case of the original powder mixture which contained both small and large particles(even larger than 180μm),larger particles decide about the ignition temperature.A pile of larger particles has a smaller total surface area than a pile of smaller particles.This should mean that a lower amount of Y is necessary to form a protective layer of Y2O3,and therefore,higher ignition temperature should be achieved for larger particles.However,the results show the opposite trend.It is theoretically a single particle which ignites as firs and causes a chain reaction so that the whole pile starts to burn.Therefore,such behaviour depends rather on the individual powder particles than a whole pile of particles.A single large particle has a larger surface area than a smaller one.Besides,there exists shorter distances for the diffusion of Y to the surface in the small particles and due to the higher cooling rate,there is higher concentration of alloying elements in the solid solution which helps with the creation of Y2O3on the surface.Therefore,it is easier to form a Y2O3protective layer on the surface of the small particles than on the large particle,which cause the increase of ignition temperature.

Powders treated in HF were characterised with the ignition temperature ranging between 700°C and 750°C with an average value of 723±15°C regardless of particle size(Fig.14B).Generally higher ignition temperature compared to the untreated powder may be associated with the protective effect of MgF2surface layer.The protective effect of MgF2layer is also used in melting Mg under SF6atmosphere to protect it against oxidation [67].The decomposition of MgF2occurs at about 750°C [68],which corresponds with the measured values of ignition temperature.Some inhomogeneities of the coating and possible reactions with Y might be responsible for the observed deviations.From the practical point of view,it is thus recommended to store fine powders of WE43 alloy or HF-treated powders for safety reasons.Moreover,using smaller particles of WE43 powder or HF-treated powder might be advantageous for additive manufacturing as it might reduce the probability of powder burning in the 3D printing chamber.

4.Conclusion

This study revealed the effect of powder size on mechanical,corrosion,and ignition properties of WE43 alloy prepared by SPS method.The effect of surface conditions of the processed powder was demonstrated by the HF pre-treatment of the powder.The specifi interface consisting of Y2O3or YF3was formed between particles during sintering.The materials prepared from lower powder fractions contained a larger fraction of these interfaces which has a significan effect on resulting properties.It was observed that the presence of YF3at the interface of particles causes high brittleness,which causes deterioration of elongation and also TYS and UTS.Besides,there was the effect of the specifi crack propagation through the sample,which depended on the microstructure.The larger volume fraction of interfaces dramatically improved the corrosion resistance of the prepared materials as the corrosion front slows down on each barrier represented by the interface.Nevertheless,the solid solution near the interface is depleted in Y,which was consumed for the formation of the Y2O3or YF3interface.This inhomogeneity then causes the rapid corrosion rate of larger particles.The ignition temperature of the powders increased with lowering the particle size,which was associated with a better protective effect of as-formed Y2O3surface layer due to the higher Y supersaturation of the solid solution and shorter distance for its diffusion towards the surface in small particles.Chemical treatment of the powder led to the improvement of the ignition temperature to 700°C.

With respect to mechanical properties,it seems better to use the powder mixture.However,corrosion properties may be severely improved if the small particles or chemically treated powder are selected.Such powders are also safer to store as well as to use for additive manufacturing due to their higher ignition temperature.

Acknowledgement

The authors wish to thank the Czech Science Foundation(Project No.GA19-08937S) and specifi university research(A2_FCHT_2020_027 and A1_FCHT_2020_003) for the fi nancial support of this research.Dedicated to the memory of my grandfather Drahomír Dvorsky.