See discussions, stats, and author profiles for this publication at: https://www.researchgate.net/publication/235664074 High-strength magnesium alloys for degradable implant applications. Mater Sci Eng A ARTICLE in MATERIALS SCIENCE AND ENGINEERING A · JANUARY 2011 Impact Factor: 2.57 · DOI: 10.1016/j.msea.2010.09.068 CITATIONS 41 READS 129 4 AUTHORS, INCLUDING: Alla S. Sologubenko ETH Zurich 47 PUBLICATIONS 638 CITATIONS SEE PROFILE Peter J. Uggowitzer ETH Zurich 225 PUBLICATIONS 4,051 CITATIONS SEE PROFILE Available from: Peter J. Uggowitzer Retrieved on: 04 February 2016
Materials Science and Engineering A 528 (2011) 1047–1054
Contents lists available at ScienceDirect
Materials Science and Engineering A
journa l homepage: www.e lsev ier .com/ locate /msea
High-strength magnesium alloys for degradable implant applications
P. Gundea, A.C. Hänzia, A.S. Sologubenkob, P.J. Uggowitzera,!
a Laboratory of Metal Physics and Technology, Department of Materials, ETH Zurich, CH-8093 Zurich, Switzerlandb Laboratory for Nanometallurgy, Department of Materials, ETH Zurich, CH-8093 Zurich, Switzerland
a r t i c l e i n f o
Article history:Received 24 June 2010Received in revised form19 September 2010Accepted 23 September 2010
This article describes the design principles deployed in developing high-strength and ductileMg–Zn–Zr–Ca–Mn(–Yb) alloys based on a concept, which aims to restrict grain growth considerablyduring alloy casting and forming. The efficiency of the development approach is discussed. Moreover,the microstructure and phase analysis of the alloys subjected to different thermal treatments are pre-sented and the influence of the alloy composition, particularly the addition of Yb, on the evolution ofthe microstructure is discussed in connection with the mechanical properties of the materials. The newlydeveloped alloys exhibit high strength (yield stress of up to 350 MPa) at considerable ductility (elongationto fracture of up to 19%) in the as-extruded state and reveal age hardening potential (increase in hard-ness of 10–15% compared to that in the recrystallization heat-treated state). Appropriate heat treatmentsenable tailoring of the strength–ductility relation. Thermal annealing of the material resulted in a remark-able increase in ductility (elongation to fracture of more than 20% for all heat-treated samples) while highstrength is retained (yield stress ranging from 210 to 315 MPa). We attribute the attractive mechanicalproperties of the developed alloys to their fine-grained microstructure, where the grain boundaries andlattice defects are stabilized by second phase particles formed during casting and thermal treatments.
The interest in the area of light-weight structural materials hasfor some time been directed towards high-strength Mg alloys [1].Owing to the degradation performance and biocompatibility ofmagnesium, Mg alloys are also a promising candidate for tem-porary implants [2,3], and the focus of interest in Mg alloys haswidened to include the field of degradable implant applications[2,3]. Here, high-strength Mg alloys are especially attractive inosteosynthesis because of their mechanical characteristics and inparticular because of the closeness of their Young’s modulus tothat of the human bone [4]. Because the Mg implants do not haveto be removed after the bone has healed, the burdens on boththe patient and the health care system are reduced. Demand forsuitable Mg alloys for applications in osteosynthesis is thereforeincreasing.
The history of Mg implants in osteosynthesis dates back to thebeginning of the 20th century, when the first Mg plates in combi-nation with Au-plated steel nails were used to fix a broken bone[5]. It was soon discovered, however, that for such applications thestrength of an implant is a critical issue, and it is often insufficientin the available Mg alloys. Moreover, the evolution of hydrogen
upon the degradation of the Mg in the physiological environmentis problematic [3]. During the last few decades considerable effortsto develop high-strength Mg alloys were made, but mostly in thefield of automotive applications, e.g. [6,7]. These Mg alloys featureattractive strength values at generally low ductility, but are mostlyunsuitable for implant applications because of their insufficientelectrochemical and/or biological performance. In the following wepresent the design strategy employed for the development of highstrength Mg alloys for degradable implant applications in osteosyn-thesis.
1.1. Design principles
When evaluating the potential of Mg alloys for applicationsin osteosynthesis the mechanical, electrochemical and biologicalproperties have to be considered. From a mechanical point of view,osteosynthesis applications require in particular a combination ofhigh strength and reasonable ductility, to provide adequate supportto the fractured bone and to allow a sufficient plastic deformationtolerance on the part of the device during implantation (e.g. whenplates must be bent to fit the bone, or in screw fixation). Froman electrochemical point of view, fairly slow and homogeneousdegradation of the alloy is desirable to assist the tissue healing pro-cess and to prevent the formation of hydrogen gas pockets uponMg degradation. Biologically, the material and substances releasedduring degradation must not be harmful to the human body.
Grain refinement of Mg alloys is a well known and efficientapproach for enhancing both the strength and the ductility ofthe material [8]. Upon grain refinement, the strength of an alloyis heightened by grain boundary hardening described by theHall–Petch relationship [9,10], and the ductility is increased due toan activation of non-basal slip deformation modes based on com-patibility stresses near grain boundaries [11,12]. Recently, designstrategies aiming at substantial grain refinement via industrialprocessing routes (also the objective of this study) have been devel-oped [13,14]. They consider the grain-growth-inhibiting potentialof solutes during solidification [8], represented by the growthrestriction factor (GRF), and the dragging influence of second phaseparticles on the grain boundaries during extrusion, represented bythe Zener pinning pressure [15,16]. Further increase in strengthmay subsequently be gained by precipitation hardening, which inturn can be triggered by specific aging heat treatments after extru-sion.
The Mg–Zn–Zr(–RE) alloying system is very attractive fordesigning material properties because it allows the formation ofgrain-refined, and thus high-strength, reasonably ductile and alsoage-hardenable Mg alloys, e.g. [7,13,17], that in addition is particu-larly tolerant of slight deviations in composition while maintainingthe overall performance. Therefore, this system was used as a basisin the present work. To improve the electrochemical and biologicalproperties, further alloying elements were considered and carefullyselected, namely Ca, Mn and Yb. The alloy compositions are indi-cated in Table 1 and the reasons for the selection of the alloyingelements are explained below.
In a previous study the Mg–Zn–Ag–Ca–Mn(–Zr) alloyingsystem was elaborated and the performance of as-extrudedMg–3Zn–0.5Ag–0.25Ca–0.15Mn(–0.5Zr) (in wt.%) alloys was inves-tigated, particularly with respect to microstructural and mechan-ical properties [13,14,18]. The Zr-free alloy was found to exhibithigh tensile elongation (27%) at reasonable strength (ultimatetensile strength (UTS) of 250 MPa), while the alloy with Zr exhib-ited high strength (UTS of 350 MPa) at a considerable tensileelongation (21%). Therefore, owing to the promising combina-tion of high strength and remarkable ductility, this Zr-containingalloy, despite its not-yet-optimal electrochemical performance,provided a basis for the alloys presented in this study andserves as reference material. To enhance the performance of thisMg–3Zn–0.5Ag–0.25Ca–0.15Mn–0.5Zr alloy, particularly in thecontext of its potential application in osteosynthesis, the composi-tion was adapted (see Table 1) and the influence of the rare earthelement Yb on the alloy’s property profile was studied. Zirconiumyielded a significant reduction in grain size (see [13]). In compar-ison with earlier studies [13,14,18] in this work the amount of Znwas increased to 5 wt.% for three reasons: (i) to improve casta-bility; (ii) to increase the GRF during solidification (see below);and (iii) to enhance the alloys age hardening response [17]. Theamount of Ca and Mn were kept at rather low levels. Calcium wasadded mainly for its grain refining effect both during solidification(considerable GRF [8,13]) and during extrusion (formation of finegrain-boundary-pinning Ca–Mg–Zn intermetallics as found in thealloys in [13] and also in Mg–Zn–Zr–RE–Ca alloys [19]). Addition-ally, the combined addition of Zr and Ca was reported to enhancethe refinement of castings [20]. Manganese was added to convertheavy metal impurities, which are detrimental to corrosion resis-tance, into compounds that sediment during alloy melting [21]. The
Fig. 1. Thermodynamic calculations after Scheil for the freezing range of the alloyZKCa50, and (manually included) the solidus temperature of ZYbK520, measured byDSC.
reasons for the ytterbium addition were the following. Ytterbiumis soluble in Mg [22] and yields considerable growth restrictionduring solidification [8]. It was also observed that the solidus tem-perature (TS) of the castings increased upon the addition of Yb. Thisincrease in TS resulted in a smaller terminal freezing range (TFR)than that of the Yb-free alloy and, because the TFR correlates withthe hot tearing susceptibility [23], also caused a reduction in the(generally high) hot tearing susceptibility of Mg–Zn alloys. Fig. 1shows Scheil calculations of the TFR of the Yb-free alloy, which wereperformed using the Pandat software (PanMg8 database) [24]. Thesolidus temperature derived from differential scanning calorime-try (DSC) measurements of the Yb-containing alloy was includedmanually, because the element Yb is not available in the database.Other reasons for choosing Yb were its beneficial influence on thecorrosion resistance of Mg–Zn alloys [25] and the age harden-ing potential of Mg–Yb alloys [22]. Although a higher RE contentwas shown to increase the strength and hardness of Mg–Zn–Zr–REalloys [6], Yb addition was kept at intermediate values because of itspotential influence on inflammatory genes in human cells [26]. Tostudy the influence of Yb on the properties of the Mg–Zn–Zr–Ca–Mnsystem, both Yb-free and Yb-containing alloys were produced.
2. Experimental
2.1. Sample preparation
The alloys with the compositions given in Table 1, denoted asZKCa50 (Yb-free alloy) and ZYbK520 (Yb-containing alloy), werecast by vertical direct chill casting on an industrial scale (billetdiameter: 185 mm). Magnesium and the alloying elements werefirst melted in an electric furnace at a temperature of approximately690 "C and then cast at a speed of approximately 1.6 mm s#1 withcontinuous water-cooling. Afterwards, the castings were formed bydirect extrusion: the extrusion ratio was approximately 30:1, thespeed 0.3–0.6 mm s#1 and the temperature of the recipient 300 "Cfor both alloys. To avoid dissolution of the grain-boundary-pinningparticles, which formed during solidification (Zn–Zr or Ca–Mg–Znintermetallic phases (IMP)), the extrusion temperature was chosencarefully based on thermodynamic calculations using the Pandatsoftware and DSC analysis.
In order to modify the alloys mechanical properties, heat treat-ments at different temperatures and for various time periods wereperformed and their influence on microstructural and mechani-cal properties were studied. The temperatures and durations werechosen so as to avoid grain growth. In order to enhance the degreeof recrystallization due to incomplete recrystallization after the
P. Gunde et al. / Materials Science and Engineering A 528 (2011) 1047–1054 1049
Fig. 2. Optical micrographs of (a and b) the as-cast and (c and d) the as-extruded material of (a and c) the Yb-free ZKCa50 and (b and d) the Yb-containing ZYbK520 alloys.
extrusion process, a recrystallization heat treatment (RHT) was car-ried out for both alloys at 360 "C (well below the respective TS ofthe extrusions) for various time periods in the range of 0.5–168 h.Afterwards, the 4 h/RHT and 48 h/RHT samples, differing by theirdegree of recrystallization, were subjected to a brief solution heattreatment (SHT) at 379 "C and 387 "C for ZKCa50 and ZYbK520,respectively, which is approximately 15 K below the correspondingTS. Due to the relatively high SHT temperatures, the solution heattreatments were performed for only 0.25, 0.5 and 1 h to avoid graingrowth. Both RHT and SHT were carried out in an air-circulatingoven. Additionally, an aging heat treatment was performed on4 h/RHT and 48 h/RHT samples, and on RHT + 0.5 h/SHT samples.The aging was performed at 160 "C in an oil bath for various peri-ods of time (up to 72 h) and complemented by the Vickers hardnessmeasurements carried out on a regular basis. Altogether, four heattreatment states of Yb-free and Yb-containing alloys were studied:as-extruded, RHT, RHT + SHT, and RHT + SHT + aging.
The mechanical properties of the as-extruded and the vari-ously heat-treated material were evaluated by standard tensileand compression tests at room temperature. The cylindrical tensilespecimens had a gauge length of 18 mm and a diameter of 3 mm,and the compression samples measured 5 mm in diameter and10 mm in length. The crosshead speeds were 1.08 mm min#1 and0.6 mm min#1 for the tension and compression tests, respectively,corresponding to an initial strain rate of 10#3 s#1. Three sampleswere tested per material state.
2.2. Characterization of the microstructure
Microstructure analysis of the castings and the heat-treatedmaterial was performed by optical microscopy (OM), scanningelectron microscopy (SEM) and transmission electron microscopy(TEM). The samples for OM and SEM studies were ground, pol-ished and then etched with acetic–picric acid (1 g picric acid, 5 mlacetic acid, 10 ml deionized water, 100 ml ethanol). The averagegrain size was determined using the linear intercept method. SEM
analysis was performed on a Hitachi SU-70 microscope, operatedat 10 kV in the secondary electron (SE) and back-scattered elec-tron (BSE) imaging modes. Conventional TEM (CTEM) and scanningtransmission electron microscopy (STEM) studies complementedby energy dispersive X-ray (EDX) analysis in the STEM mode, wereemployed for the morphological and phase analysis of the alloysof interest. STEM was performed in a high-angle annular dark-field (HAADF) operation mode using a 0.6–0.8 nm probe size. TEMstudies were performed on a FEI Tecnai F30 machine (operated at300 kV). For TEM, thin disks of 3 mm in diameter were first mechan-ically ground to a thickness of about 100 !m, and further thinneddown to approximately 20–30 !m by dimple grinding (Gatan dim-ple grinder) on both foil sides. Electron transparency was achievedby chemical etching of the dimpled foils for a few tens of seconds ina solution of 90 vol.% methanol and 10 vol.% nitric acid performedat room temperature. The chemical etching of the foils was carriedout immediately before the TEM sessions, to minimize the imagingartifacts due to the oxidation and contamination of the foils.
3. Results
3.1. Metallographic characterization
The cross-sections of the castings of both alloys featured ahomogeneous and very fine-grained microstructure across the bil-let diameter, as shown in Fig. 2a and b. The grain-refining effect ofYb is clearly visible. Upon its addition, the grain size decreased from100 ± 15 !m in ZKCa50 (Fig. 2a) to 70 ± 10 !m in ZYbK520 (Fig. 2b).In both alloys the shape of the grains was globular, and intermetallic(eutectic) phases were found at the grain boundaries. After extru-sion both alloys exhibited a fine-grained microstructure with grainsof <2 !m in size (Fig. 2c and d). The grain size of the Yb-containingalloy (Fig. 2d) was slightly smaller than that of the Yb-free alloy(Fig. 2c) at the same extrusion conditions. However, in the centre ofall extruded bars, “worm-like” sections of the microstructure wereobserved by optical microscopy (Fig. 3a) evidencing the incomplete
1050 P. Gunde et al. / Materials Science and Engineering A 528 (2011) 1047–1054
Fig. 3. Optical micrographs of the ZYbK520 alloy (a) in the as-extruded and (b and c) in the 48 h/RHT states. (a) Shows the centre of the extruded bar. (c) Acquired at highermagnification from the area depicted in (b).
dynamic recrystallization, which is typical in many Zr-containingMg alloys [27]. After annealing at 360 "C (48 h/RHT) recrystal-lization was observed, accompanied by the disappearance of the“worm-like” sections (Fig. 3b and c) and slight growth of the grainsize and the size of the IMP particles (Fig. 3c). Generally, in bothalloys the grain size remained small (average grain size reached$2.5 !m for 48 h/RHT, with single grains growing to approximately5–10 !m). Simultaneously, the hardness of ZYbK520 alloys con-tinuously decreased from 80 to 85 HV in the as-extruded state toapproximately 70 HV after 48 h/RHT. The hardness values of theYb-free alloy were generally a bit lower. The following SHT of therecrystallized material led to an additional marginal reduction inhardness and to little grain growth in both alloys; the averagegrain size reached around 4 !m after 48 h/RHT + 1 h/SHT. The agingheat treatment at 160 "C produced an improvement in hardness. Inboth alloys hardness peak values of 75–80 HV were reached afterapproximately 24 h of aging, representing an increase of 10–15%.Upon further aging the hardness decreased only marginally.
3.2. Mechanical performance
Tensile tests on the as-extruded material revealed high yieldand tensile stress values at a high ductility level for both alloys(Table 2). In general, the Yb-free alloy exhibited somewhat loweryield stress at higher ductility compared to ZYbK520. For ZKCa50,yield and tensile stresses of 330 and 360 MPa were achieved at anelongation to fracture value of 19%. For ZYbK520, even a higheryield stress value of 350 MPa was obtained at an elongation to frac-ture of 15%. In both alloys the uniform elongation was around 10%.The RHT had a positive influence on ductility of both alloys evenif entire recrystallization could not be achieved, as indicated bythe strain hardening behavior of the Yb-containing alloy presented
Fig. 4. Stress–strain diagram of as-extruded ZKCa50 and ZYbK520, and ZYbK520 invarious heat treatment states.
in Fig. 4: uniform elongations over 10% and elongations to frac-ture of more than 20% were reached. With increasing RHT time,the yield and tensile stress decreased and the ductility increased;after 48 h/RHT minimum yield stresses of 210 and 230 MPa at anelongation to fracture of 26% and 23.5% for ZKCa50 and ZYbCa520,respectively, were measured. Upon the subsequent aging at 160 "Cthe yield stress increased again by 15–20%, accompanied by an onlymarginal loss of ductility. The brief SHT prior to the aging resultedin a slight rise in the yield and tensile stress values compared to thenon-SHT state for both alloys.
In both alloys, the RHT also had a positive influence on thetension-compression yield stress asymmetry, which was generallymoderate. In compression the yield stress was lower than in ten-sion for all heat treatment states. In all the as-extruded material the
Table 2Mechanical properties in compression and tension of the as-extruded and heat-treated alloys ZKCa50 and ZYbCa520 at room temperature (standard deviation of stress data<10%; standard deviation of strain data <20%).
a !0.2,t , yield stress in tension; !UTS, ultimate tensile stress; "u, uniform elongation; "f , elongation to fracture; !0.2,c , yield stress in compression. RHT was performed at360 "C, SHT at 379 "C for ZKCa50 and at 387 "C for ZYbK520, and aging at 160 "C.
P. Gunde et al. / Materials Science and Engineering A 528 (2011) 1047–1054 1051
Fig. 5. BF TEM micrograph of a polycrystalline region in as-extruded ZYbK520. Thegrain morphology and the high dislocation density within the grains evidence theincomplete recrystallization. The large black areas in the image, labeled “I”, corre-spond to the (Yb,Ca)–Mg–Zn intermetallic phase particles.
tension-compression yield stress asymmetry was approximately1.25 and generally decreased with RHT time, subsequently slightlyincreasing again due to aging.
3.3. SEM and TEM studies
SEM and TEM analysis was performed on both alloys. As themicrostructural features of the Yb-containing and Yb-free alloyswere similar and as the main interest of the work laid on theinfluence of Yb on the alloy properties, only the results of theYb-containing alloy ZYbK520 are presented here. The TEM morpho-logical analysis of the as-extruded ZYbK520 alloy (Fig. 5) revealeda not completely recrystallized grain morphology and a high den-sity of lattice defects in the material. The OM observations of theas-extruded and 48 h/RHT ZYbK520 specimens (Fig. 3a–c), com-plement these TEM results showing that hot extrusion and thesubsequent heat treatments did not yield complete dynamic andstatic recrystallization. In addition, the SEM micrographs (Fig. 6aand b), taken in the back-scattered electron (BSE) imaging modeand thus featuring the chemical contrast, show the presence ofheavy-atom-enriched IMP particles decorating grain boundariesand lattice defects (small bright spots in the images marked with
black arrows for the grain boundary and with white arrows forthe lattice defect decorations). The comparison of the images inFig. 6 demonstrates that the IMP particles grow upon the heattreatment.
Microstructural analysis of the Yb-containing alloys in all heattreatment states by CTEM, SEM and HAADF-STEM (the latter twoare atomic number sensitive) revealed that a few distinctly differ-ent types of IMP particles occur in the material (Figs. 6 and 7). Uponheat treatments all IMP particles grow and evolve (compare Fig. 6awith b and the vertical rows of images in Fig. 7). The comparison ofthe particle morphologies allows to suggest the occurrence of threedifferent intermetallic phases in the Yb-containing alloys.
(i) Large globular particles of about 0.5–2 !m in size are the mostprominent and of the largest volume fraction. They were foundpredominantly at the matrix grain boundaries and were eval-uated by OM, SEM and TEM. They are seen as large black grainsin the TEM micrograph in Fig. 5 (labeled “I”) and they are thelargest and brightest particle in the SEM image in Fig. 6a.
(ii) Roughly spherical, a few nanometer large (depending on theheat treatment state the diameter ranges within 3–50 nm) par-ticles were observed at the (small-angle) grain boundaries anddislocations (short white arrows in Figs. 6 and 7). Larger parti-cles tend to be located at grain boundaries, as seen in Fig. 7a–c.In the as-extruded state (Figs. 6a and 7a), the diameter of theseIMP particles ranged within 10–15 nm for particles pinning thegrain boundaries and within 3–10 nm for those decorating thesmall-angle grain boundaries and dislocations. Upon 48 h/RHT,these particles grew to approximately 15–50 nm (see Fig. 7b).No significant particle growth was observed upon the subse-quent SHT + aging heat treatments (see Fig. 7b and c).
(iii) The third type of second phase particles are rod- and plate-like inclusions found at dislocations and also in defect-freeregions of the matrix grains (Fig. 7d–f). The rods are alignedperpendicular to the basal planes of "-Mg (elongated in the cdirection of the "-Mg lattice), whereas the plates are parallel tothe basal planes of "-Mg. Among all second phases, these rod-and plate-like particles grew most prominently upon each heattreatment step (Fig. 7d–f). Additionally, aging induced a sig-nificant increase in the volume fraction of the rods and plates(compare Fig. 7e and f and see schematic in Fig. 7g–i).
The TEM micrograph of a 48 h/RHT ZYbK520 specimen (Fig. 8)shows the interactions of roughly spherical IMP particles (arrowed)with lattice defects. These IMP particles pin the dislocations andmost predominantly the small-angle grain boundaries (arrows inFig. 8 point to some particles on small-angle grain boundaries). Inthis image, the rod- and plate-like precipitates are barely visible due
Fig. 6. BSE SEM images of the ZYbK520 alloy acquired from specimens in (a) the as-extruded and (b) the 48 h/RHT state. Very dark areas in the image are artifacts dueto preferential etching at the large IMP particle and grain boundaries. Black arrows point to grain boundaries decorated with second phase particles. White arrows markdislocations or small-angle grain boundaries pinned by IMP particles.
1052 P. Gunde et al. / Materials Science and Engineering A 528 (2011) 1047–1054
Fig. 7. HAADF-STEM micrographs of ZYbK520 in (a and d) the as-extruded, (b, and e) the 48 h/RHT, and (c and f) the 48 h/RHT + 0.5 h/SHT + peak-aged states. The (a, b, andc) images are taken at lower and the (d, e and f) images at higher magnifications. The latter were acquired from singular grains in high symmetry zone axis orientations (thelarge white arrows indicate the c direction of the matrix crystal) and illustrate the intra-granular defect structure and specific crystallographic correspondence of the matrixand the IMP particles. Black arrows mark grain boundaries decorated with IMP particles. The small white arrows indicate dislocations or small-angle grain boundaries pinnedby IMP particles. The (g, h, and i) images are schematics illustrating the crystallographic correspondence of the rod- and plate-like IMP particles and the matrix of ZYbK520in (g) the as-extruded, (h) the 48 h/RHT, and (i) the 48 h/RHT + 0.5 h/SHT + peak-aged states. The rod-like particles are parallel to the c direction of the "-Mg whereas theplate-like particles lie on the matrix basal planes. The increase in number density and size of the IMP particles upon heat treatments can be seen from the image comparison.
Fig. 8. BF TEM micrograph acquired in the [0 1 1̄ 0] zone axis orientation of a48 h/RHT ZYbK520 grain. The interactions of roughly spherical IMP particles withsmall-angle grain boundaries can be seen (arrowed). The rod- and plate-like particlesare also present, but hardly visible in the image.
to their small size and thus low contrast, but in higher magnificationCTEM images the rods and plates were discerned.
EDX analysis (not shown) indicated the presence of Mg and Znin the "-matrix. In the case of the largest, micrometer-sized IMPparticles, Mg, Zn, Yb and Ca signals were detected. According toEDX-data, the roughly spherical IMP particles contain Mg, Zn, Zrand Mn, while only Mg and Zn were found in the rod- and plate-likeparticles.
4. Discussion
4.1. Design principles
As a result of our design strategy efficient grain refinementduring solidification was achieved in both the Yb-free and Yb-containing alloys. Considering the identical casting parametersused, the smaller grain size in ZYbK520 is ascribed to the sup-plemental grain-growth-restricting influence of ytterbium. Thesomewhat larger grain size of our castings compared to simi-lar Mg–Zn–Zr–Ca alloys in [13,17] is accredited mainly to thedifferences in casting conditions, in particular the lower coolingrates. After the thermo-mechanical processing the grain size ofthe ZKCa50 and ZYbK520 alloys was very small and similar tothat reported in [13,17]. We attribute this fact to the fine-grainedmicrostructures of the castings, the appropriately chosen extru-sion parameters and the presence of grain-boundary-pinning IMPparticles, which hinder grain growth during thermo-mechanical
P. Gunde et al. / Materials Science and Engineering A 528 (2011) 1047–1054 1053
forming. Our analyses show that the extend of recrystallizationwas moderate in as-extruded alloys, but noticeably enhanced inRHT specimens. Complete recrystallization, however, could not beachieved in our experiments. We explain this by the effect of IMPparticles on the dislocation mobility.
4.2. Microstructural and mechanical performance
Evaluation of the mechanical properties by tensile and compres-sive testing shows a good combination of strength and ductilityfor both alloys in all states (as-extruded and heat-treated). More-over, the tests reveal that the appropriate choice of heat treatmentparameters enables controlling the alloys stress–strain perfor-mance. The attractive mechanical properties of the alloys can beprincipally attributed to the appropriately designed fine-grainedalloy microstructures that result in reduced grain boundary mobil-ity, which strengthens the alloy (as described by Hall–Petch).Moreover, the small grain size promotes the activation of non-basal dislocation modes enhancing the ductility of the alloys, asseen also in [11,12]. In addition, the presence and formation of sec-ond phase particles of different morphologies on the lattice defectsleads to effective pinning of the dislocations and the very abundantsmall-angle boundaries. As a consequence the dislocation mobilityis lowered and the strength is enhanced.
The high yield strength in the as-extruded state is attributedmainly to the fine-grained microstructure and the reduced dis-location mobility due to the obstacle action of the IMP particlesand the dislocation network. Upon the RHT, slight grain growthand recrystallization occur, facilitating dislocation movement andcausing a decrease in tensile and compressive yield stress. Agehardening, on the other hand, results in an increase in the vol-ume fraction and size of the rod- and plate-like IMP particles inthe matrix grains, which contributes to the yield stress increase.From the morphological and chemical analysis of these particles,it is assumed that the small rod- and plate-like IMP particles, thatwere also found in the Yb-free alloy, represent the same Mg–Znhardening phase (MgZn2) as reported by Mendis et al. [17]. How-ever, for a confident identification of this phase, further studies arenecessary.
Besides the attractive combination of strength and ductility,another advantage of our alloys is their moderate tension-compression yield stress asymmetry, which is mainly ascribedto the small grain size [12]. It is known for Mg alloys that withdecreasing grain size an activation of the twinning deformationmode becomes increasingly difficult. The energy caused by grainboundary incompatibility stresses is assumed to be consumed firstby non-basal slip than twin formation [12]. Upon the RHT, thetension-compression yield stress asymmetry decreases in all alloyscompared to the as-extruded state. Reduced tension-compressionyield stress asymmetry can be attributed to a weakening of thecrystallographic texture as well as a decreased dislocation densityand the associated reduced stress concentrations at which twinningwould preferentially be initiated. The latter effect was observed inthis study, while detailed analysis of the crystallographic texture ispart of an up-coming work. Weak texture, however, was observedin alloys similar in composition and fabrication as those presentedhere [28]. Upon aging the tension-compression yield stress asym-metry increases somewhat. We assume that this can be attributedto the prompt and extensive increase of the volume fraction ofcoherent rod-like IMP particles (detected by TEM, Fig. 7f). On theone hand, they decorate dislocations and impede their motion. Onthe other hand, the preferential growth of these rod-like IMP par-ticles causes strain fields, which may promote twinning becauseof stress concentrations related to them. A similar effect is sug-gested to take place also in the material after the RHT + SHT + agingtreatment.
4.3. Influence of ytterbium
Ytterbium is not a widely employed alloying element for Mgalloys and has rarely been used in Mg technology. Therefore onlylittle literature data exists about the effect of Yb on the mechanicalbehavior of Mg alloys. Thus we present its role in more detail in thefollowing.
Due to the similar mechanical performance of ZKCa50 andZYbK520 and the experimentally confirmed presence of Yb in themicrometer-sized IMP particles only, it is assumed that the alloyingelement Yb exerts no significant influence on the active deforma-tion and strengthening processes. However, the following effectscan be attributed to ytterbium: (i) Yb reduces the alloy’s grainsize due to its contribution to the grain growth restriction factor[8]; (ii) Yb reduces the hot tearing susceptibility during casting byincreasing the solidus temperature and consequently decreasingthe terminal freezing range [23]; and (iii) Yb is present in the large0.5–2 !m IMP particles on the grain boundaries and thus affectsthe grain-boundary pinning. It is noteworthy that Yb additions donot considerably increase the cost of the alloy (price of Yb: circa120 Dkg#1). From the results of the EDX TEM measurements thatshow the presence of Yb only in the large, grain-boundary-pinningparticles, from Pandat simulations of the phase configurations ofZKCa30 [13], which predict the occurrence of Ca–Mg–Zn IMP parti-cles (Ca2Mg6Zn3), and taking into account the complete miscibilityof Ca and Yb, it is suggested that the ternary (Ca,Yb)2Mg6Zn3 phasehas formed as large globular particles in the alloy ZYbK520. How-ever, the precision of the EDX elemental analysis in TEM is notsufficient to make reliable conclusions about the exact composi-tion of these IMPs. These globular particles are likely to have formedduring solidification of the remaining melt (eutectic phase).
When considering Yb as alloying element in our design strategy,it was expected that addition of Yb would result in an improvementin the degradation resistance by incorporating ytterbium into theMg matrix and in an increase in strength by the formation of Mg–Ybstrengthening phases. EDX measurements, however, indicated noYb signals from the matrix or the small IMP particles. It is likely thatall Yb added to the melt becomes consumed by the (Ca,Yb)–Mg–Znphases. This fact that no Yb was found solved in the "-Mg matriximplies two conclusions. At variance to observations reported in[22] or [29], no Yb-containing strengthening phases form, and thecorrosion behavior is not particularly improved, as the matrix is notrendered nobler.
4.4. Summarizing comments
Fig. 9 presents a review of the yield stress versus elongation tofracture performance of Mg–Zn based alloys potentially suitable formedical applications [13,14,17,29–31]. The data also give a compar-ison to the performance of the newly developed alloys. All alloyspresented in the figure were produced similarly to ours: they werecast and then hot extruded, with extrusion ratios between 10:1and 30:1. It can be observed that our extruded alloys are situatedin the upper central area of the diagram, i.e. they exhibit a highyield stress at moderate fracture strain. In the as-extruded state,they feature significantly higher yield stress at improved ductilityin comparison to the commercial alloys ZK31 or ZK60. In addition,high-strength alloys presented in literature which show tensileyield stresses similar to our alloys predominantly feature a signif-icantly lower fracture strain, e.g. [31]. The RHT (with and withoutsubsequent aging) results in a shift of our alloys’ mechanical perfor-mance towards an improvement of fracture strain and a decreasein yield stress. Only alloys ZQCa30, WZ21 and ZW21, which weredeveloped for applications in vascular intervention (where a highuniform elongation at moderate strength is important), and theMg–6Zn–0.4Ag–0.2Ca alloy, which is similar to ZQCa30, feature
1054 P. Gunde et al. / Materials Science and Engineering A 528 (2011) 1047–1054
Fig. 9. Schematic presenting a review of literature data on the dependence of yieldstress in tension versus fracture strain of various commercial Mg–Zn based alloys[13,14,17,29–31]. All alloys were produced in a similar way to the alloys in thisstudy. The alloys developed in our group are encircled.
higher ductility than our heat-treated material, however also atmuch lower strength. In summary, high-strength Mg alloys withtailorable strength–ductility relation were developed.
5. Conclusions
In this study Mg–Zn–Zr–Ca–Mn(–Yb) alloys were devel-oped which exhibit high strength (yield stress up to 350 MPa)at considerable ductility (fracture strain >15%) and relativelylow tension-compression yield stress asymmetries. The designapproach presented in this work involved grain growth restric-tion mechanisms which act during solidification (solute drag)and thermo-mechanical processing (grain-boundary-pinning). Theextraordinary combination of high yield strength and high ductilityobtained upon thermo-mechanical processing and subsequent heattreatments is attributed to the fine-grained microstructure, withproperly and optimally dispersed lattice-defect-pinning IMP par-ticles (i.e. (small-angle) grain-boundary- and dislocation-pinning).The high ductility of the material is mainly ascribed to the smallgrain size responsible for the activation of additional deformationmodes (non-basal slip).
Moreover, the mechanical properties can be tailored by appro-priate thermal treatments at different temperatures and for variousperiods. From high strength in the as-extruded state a good combi-nation of strength and ductility can be achieved upon subsequentannealing. This change in mechanical properties is explained bya slight increase in grain size, a reduction in lattice defect den-sity during the recrystallization heat treatments, and precipitationhardening.
The influence of Yb on the alloy performance was studied. Themain function of Yb turned out to be the reduction of the hot tearingsusceptibility during direct chill casting by decreasing the terminalfreezing range, and the refinement of the microstructure of the castand extruded material. Adding Yb to ZKCa50 significantly improvedthe alloy’s castability and slightly improved its mechanical per-formance. Yb is found in grain-boundary-pinning particles, which
hinder grain growth during hot extrusion and during subsequentheat treatments. In summary, the mechanical performance of thealloys is considered promising for temporary implant applicationsin osteosynthesis.
Acknowledgements
The authors are grateful to Christoph Mareischen, B.Sc., for hiswork in the laboratory. We would also like to thank the electronmicroscopy centre at ETH Zurich (EMEZ) for providing the facili-ties for TEM work and for general support during TEM studies. Wehighly appreciate the financial support of the Austrian Institute ofTechnology (AIT), the state of lower Austria, the European Fund forRegional Development (EFRE) (within the framework of the “Bio-compatible Materials and Applications” (BCMA) project) and theStaub/Kaiser Foundation (Switzerland).
References
[1] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37–45.[2] B. Heublein, R. Rohde, V. Kaese, M. Niemeyer, W. Hartung, A. Haverich, Heart
89 (2003) 651–656.[3] F. Witte, V. Kaese, H. Haferkamp, E. Switzer, A. Meyer-Lindenberg, C.J. Wirth,
H. Windhagen, Biomaterials 26 (2005) 3557–3563.[4] M.P. Staiger, A.M. Pietak, J. Huadmai, G. Dias, Biomaterials 27 (2006)
1728–1734.[5] A. Lambotte, Bull. Mem. Soc. Nat. Chir. 28 (1932) 1325–1334.[6] T. Honma, N. Kunito, S. Kamado, Scripta Mater. 61 (2009) 644–647.[7] Y. Kawamura, M. Yamasaki, Mater. Trans. 48 (2007) 2986–2992.[8] D.H. StJohn, M. Qian, M.A. Easton, P. Cao, Z. Hildebrand, Metall. Mater. Trans. A
36 (2005) 1669–1679.[9] E.O. Hall, Proc. Phys. Soc. B 64 (1951) 946–953.
[10] N.J. Petch, J. Iron Steel Inst. 174 (1953) 25–28.[11] J. Koike, Mater. Sci. Forum. 449–452 (2004) 665–668.[12] J. Koike, T. Kobayashi, T. Mukai, H. Watanabe, M. Suzuki, K. Maruyama, K.
Higashi, Acta Mater. 51 (2003) 2055–2065.[13] A.C. Hänzi, F.H. Dalla Torre, A.S. Sologubenko, P. Gunde, R. Schmid-Fetzer, M.
Kuehlein, J.F. Löffler, P.J. Uggowitzer, Phil. Mag. Lett. 89 (2009) 377–390.[14] A.C. Hänzi, A.S. Sologubenko, P.J. Uggowitzer, Int. J. Mater. Res. 100 (2009)
1127–1136.[15] P.A. Manohar, M. Ferry, T. Chandra, ISIJ Int. 38 (1998) 913–924.[16] C.S. Smith, Trans. Am. Inst. Min. Metall. Pet. Eng. 175 (1948) 15–51.[17] C.L. Mendis, K. Oh-Ishi, Y. Kawamura, T. Honma, S. Kamado, K. Hono, Acta Mater.
57 (2009) 749–760.[18] F.H. Dalla Torre, A.C. Hänzi, P.J. Uggowitzer, Scripta Mater. 59 (2008) 207–210.[19] Q. Li, Q.D. Wang, H.T. Zhou, X.Q. Zeng, Y. Zhang, W.J. Ding, Mater. Lett. 59 (2005)
2549–2554.[20] S.-Y. Chang, H. Tezuka, A. Kamio, Mater. Trans. 38 (1997) 526–535.[21] I.J. Polmear (Ed.), Light Alloys – Metallurgy of the Light Metals, Edward Arnold,
Metallogr. 103 (2007) 64–71.[23] P. Gunde, A. Schiffl, P.J. Uggowitzer, Mater. Sci. Eng. A 527 (2010) 7074–7079.[24] Pandat Software, PanMg8 database, software available at
http://www.computherm.com/pandat.html.[25] M. Yamasaki, N. Hayashi, S. Izumi, Y. Kawamura, Corr. Sci. 49 (2007) 255–262.[26] A. Drynda, N. Deinet, N. Braun, M. Peuster, J. Biomed. Mater. Res. A 91A (2009)
360–369.[27] K. Oh-ishi, C.L. Mendis, T. Homma, S. Kamado, T. Ohkubo, K. Hono, Acta Mater.
57 (2009) 5593–5604.[28] A.C. Hänzi, T. Ebeling, R. Bormann, P.J. Uggowitzer, Magnesium Technol. 2009
(2009) 521–526.[29] W. Yu, Z. Liu, H. He, N. Cheng, X. Li, Mater. Sci. Eng. A 478 (2008) 101–107.[30] X. Zeng, Y. Zhang, C. Lu, W. Ding, Y. Wang, Y. Zhu, J. Alloys Compd. 395 (2005)
