Fabrication and characterisation of MgLi thin films for neurological implants

Fabrication and characterisation of MgLi thin films for neurological implants Dissertation zur Erlangung des akademischen Grades eines Doktors der Ingenieurswissenschaften (Dr.-Ing.) der Christian-Albrechts-Universit¨at zu Kiel vorgelegt von Lisa Hanke aus Hamburg Kiel 2024 Erstgutachter: Prof. Dr. Eckhard Quandt Zweitgutachter: Prof. Dr. Regine Willumeit-R¨omer Drittgutachter: Prof. Dr. Markus Valtiner Datum der Disputation: 25.10.2024 Contents Contents Abstract 1 Zusammenfassung 2 1 Motivation 3 2 Fundamentals 5 2.1 Magnesium and its alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5 2.1.1 Applications in medicine . . . . . . . . . . . . . . . . . . . . . . . . 6 2.1.2 Mg alloys for medical applications . . . . . . . . . . . . . . . . . . . 7 2.1.3 MgLi................................... 7 2.2 Sputteredthinfilms............................... 9 2.2.1 Microstructure of thin films . . . . . . . . . . . . . . . . . . . . . . 9 2.2.2 Mgalloythinfilms ........................... 11 2.3 Corrosion .................................... 12 2.3.1 Mgcorrosion .............................. 13 2.3.2 Influencingfactors ........................... 14 2.3.3 Mg thin film corrosion . . . . . . . . . . . . . . . . . . . . . . . . . 16 2.3.4 MgLicorrosion ............................. 17 2.3.5 Li2CO3formation............................ 19 2.4 Corrosion measurement techniques and feasibility for MgLi thin film measurements .................................... 20 2.4.1 Weight-loss measurements . . . . . . . . . . . . . . . . . . . . . . . 20 2.4.2 Hydrogen evolution . . . . . . . . . . . . . . . . . . . . . . . . . . . 21 2.4.3 Potentiodynamic polarisation . . . . . . . . . . . . . . . . . . . . . 22 2.4.4 In-lineICP-MS ............................. 24 2.4.5 Further measurements . . . . . . . . . . . . . . . . . . . . . . . . . 27 3 Results 28 i Contents 3.1 Publication: Structural characterisation and degradation of Mg-Li thin films for biodegradable implants . . . . . . . . . . . . . . . . . . . . . . . . 28 3.2 Influence of sputtering parameters on the corrosion rate of Mg-3Li . . . . . 41 3.3 Analysis of long-term corrosion in different solutions . . . . . . . . . . . . . 44 3.4 Publication: Investigation of in-situ ion release and surface film formation ofhcpMg-Lithinfilms............................. 48 3.5 Publication: Tailoring of Mg and MgLi thin-film corrosion rates with dielectric barrier discharge plasma treatment . . . . . . . . . . . . . . . . . . 59 3.6 Addition of an alloying element - MgAgLi . . . . . . . . . . . . . . . . . . 69 4 Conclusion and Outlook 70 Appendix lxxii A.1 Supplement Structural characterisation and degradation of Mg-Li thin films for biodegradable implants . . . . . . . . . . . . . . . . . . . . . . . . . . . lxxii A.2 Supplement Investigation of in-situ ion release and surface film formation ofhcpMg-Lithinfilms.............................lxxvi A.3 Supplement Tailoring of Mg and MgLi thin-film corrosion rate with dielectric barrier discharge plasma treatment . . . . . . . . . . . . . . . . . . . . lxxx Full List of Publications lxxxiv Conference Presentations lxxxv Symbols and Abbreviations lxxxvi List of Figures lxxxviii List of Tables lxxxix Bibliography xci Acknowledgements cvi ii Abstract Abstract Due to their mechanical properties, biocompatibility and degradability, magnesium alloys are promising materials for various medical implants, such as bone fixations or stents. Compared to other materials, negative long-term effects of the implants or a second operation to remove the implant can be avoided. Another advantage of the degradation of the material is the potential therapeutic effect of the alloying elements itself. In this work, magnesium-lithium alloys are considered as biodegradable reservoirs for lithium which can be used for neurological treatments. For this purpose, small implants with defined structures need to be fabricated. This manuscript describes the requirements and properties of the thin films produced by magnetron sputtering. Particular attention is paid to the corrosion properties, as these determine not only the lifespan of the implants but also the concentration of released ions. Different possibilities to tailor the properties of the films are also discussed. In a first study, alloys with four different lithium fractions were compared in terms of their structural and mechanical properties and corrosion rate. A subdivision of the properties according to the phases present in the film was found. While the single-phase alloys exhibited a corrosion rate similar to that of unalloyed magnesium films, which only slightly increased with lithium content, secondary phases sometimes resulted in strongly increased corrosion rates due to galvanic corrosion in the absence of passivating surface films. However, the corrosion rate could also be controlled by other factors, such as orientation or grain boundary modification by changing the sputtering parameters during the manufacturing process. The release of lithium during degradation determines the effectiveness of the implants as a reservoir for neurological treatments. It has, therefore, been studied over longer periods as well as at the onset of corrosion. The onset of contact with the solution is of particular interest as the surface films and surface processes significantly influence corrosion, especially for thin films. A lithium carbonate film, which forms in air and dissolves during immersion, was identified on the studied thin films. Subsequent selective dissolution of the lithium leads to a change in concentration throughout the measurement. Since the bulk structure of the film and the surface affect the corrosion resistance, both can be altered to adjust the properties. A plasma treatment with a dielectric barrier discharge plasma resulted in a change in surface chemistry and morphology. The modified surface showed improved corrosion resistance after only a few seconds of treatment. The addition of silver to the alloy increased the tensile strength while not significantly affecting corrosion rate. This study, thus, provides a first insight into magnesium-lithium thin films as neurological implants and different processes and options to consider for optimising the films for the application. 1 Fundamentals (comparison: a= 3.2076 ˚ A, c/a = 1.6227 for 0.65 wt% vs. a= 3.193 ˚ A and c/a = 1.6068 for 5.24 wt% Li) [42]. The change of hcp to a more cubic-like structure or the presence of bcc leads to the availability of more slip systems for mechanical deformation. With a c/a ratio closer to 1, the prismatic and pyramidal slip are more pronounced, leading to a higher ductility and maximum elongation [77, 78]. The elongation varies depending on the Li content. Zhou et al. found, e.g., an increase from around 15 % to 41 % from Mg3.5Li to Mg8.5Li [79]. On the other hand, while low amounts of Li can increase the strain hardening and yield strength, further Li addition and additional slip lowers the possibility for strain hardening and reduces the tensile strength [78]. Furthermore, the plastic anisotropy is reduced, and in comparison to pure Mg, the grain growth during recrystallisation is increased [77]. βMgLi shows an especially high formability, ductility and slightly reduced Young’s modulus but also lower strength and low thermal and creep resistance [72, 80]. Figure 2.2: Phase diagram of Mg-Li. Reproduced and adapted with permission from Springer Nature [76]. While Li could be equally distributed in the single phase material or grains from each phase in the mixed material, a segregation at the grain boundaries is found even for low Li fractions [81]. This can either lead to an enhanced embrittlement or improved grain boundary sliding. For mixed-phase materials after heat treatment and mechanical treatment, grain boundary sliding was found to lead to superplasticity [72]. The Li segregation at grain boundaries possibly influences the diffusion of Li, which is in general lower for α-phase than β-phase [82, 83], leading to an easier transport of Li to the surface over grain boundaries also for the hcp phase. The accumulation of Li could also affect the corrosion properties of the material which will be described later. In addition to the phases, the properties of MgLi are also influenced by microstructure. Different post-treatments such as heat treatments or mechanical treatments, e.g. rolling, change grain sizes and texture [40, 72, 84]. This can lead to hardening of the material. Especially for nanostructured Li-rich materials, an improved strength is found [40]. Other elements such as Zn, Al or REE are often added for improved properties [40, 65, 66]. Working with MgLi also has additional challenges due to the high reactivity of Li. This leads to high oxidation [72] and complicates the processing. Especially for β-MgLi, ad8 Fundamentals ditional ageing effects result in the formation of precipitates and recrystallisation to a partial α-phase even at room temperature [85, 86]. This can lead to age softening and defect formation at interfaces between both phases. Therefore, the properties also depend on the time between sample preparation and experiment. 2.2 Sputtered thin films The definition of thin films differs in the literature. While some describe that a thin film is thinner than 1 µm [87, 88], it can also include structures in the µm-range or even 100s of µm if the material is still small in comparison to the substrate [89]. Thin films can be fabricated by different techniques using solid, liquid or gaseous deposition [90]. The most commonly used techniques are chemical (CVD) or physical (PVD) vapour deposition. In PVD, the material of choice is brought into the gas phase by physical means such as heat or ion bombardment. The atoms will then deposit and form a layer on the chosen substrate without any additional chemical reactions taking place. Sputtering as one of the PVD techniques is carried out in a vacuum chamber which is flooded with an inert gas such as argon for non-reactive sputtering [91]. An applied voltage starts a plasma between the target material and the substrate. Due to the bombardment of the target by compounds from the plasma, atoms of the target will be ejected and deposited on the substrate. For a higher sputtering rate and lower necessary voltage and pressure, permanent magnets can be placed behind the targets for the so-called magnetron sputtering [90, 92]. This leads to a longer, spiral path in the plasma and, thus, further ionising by interaction of gas atoms. Therefore, the amount of ions hitting the target is increased, leading to a higher sputtering rate. A side effect is the formation of specific sputtering trenches due to the small area of attack. The composition of materials including multiple elements can differ from target to deposited material since, depending on the atomic number of the elements, the atoms have a different angular distribution. 2.2.1 Microstructure of thin films During sputtering, parameters of the process, such as the applied power, pressure of gas and gas flow, can be adjusted. This can then increase or decrease the energy of the atoms, therefore influencing the microstructure formed during deposition. The development of microstructure dependent on the temperature of the substrate (Ts) and melting temperature (Tm) as Ts/Tmwas described by Movchan and Demchishin in 1969 [93], indicating different zones 1-3 depending on the substrate temperature since it influences the energy available for activation and, thus, diffusion. Thornton then added the dependence on the pressure and developed the structure-zone model shown in figure 2.3 [94, 95]. The model is mainly described by zone 1-3 and a transition zone T between 1 and 2 [94, 95, 97]. •In zone 1 with the lowest Ts/Tm, the low diffusion leads to a high defect density, shadowing and open grain boundaries. The diameter of the grains increases with 9 Fundamentals Figure 2.3: Structure-zone model developed by Thornton showing the microstructure of sputtered films dependent on the temperature (substrate temperature Tsand melting temperature Tm) and the inert gas pressure. Reproduced with permission from Elsevier [96]. temperature. •In the transition zone T, surface diffusion is already possible, resulting in a denser structure with the formation of less defined grains or fibrous structures. •For zone 2, a smooth surface and dense grain boundaries between columnar grains are reached, leading to better mechanical properties. •In zone 3, volume diffusion and recrystallisation can occur, forming round grains and microstructures similar to bulk material. •On the other axis, the increase in pressure shifts the transformation temperatures to higher values due to a loss in kinetic energy. The revised structure-zone model by Messier et al. exchanged the pressure with a substrate floating potential and claims the T zone as a subzone of zone 1 [98]. Barna and Adamik also added the importance of impurities on the final structure [99], and Anders adjusted the model by switching to a generalised temperature and normalised energy flux to account for other deposition techniques [100]. Thus, the different structure-zone models can give an insight into the approximate structure of thin films but need to be defined and adapted for the system of interest [97]. In comparison to bulk materials, fewer precipitates are found in sputtered films. This is due to the fast cooling rate during deposition, leading to non-equilibrium supersaturated solutions of materials which can possess very different melting points [32]. In addition to the microstructure described in the structure-zone model, changing the sputtering parameters also influences the film stress, specifically, the intrinsic film stress defined by the structural disorder [96]. For the fabrication of stable, possibly freestanding thin films, an optimisation to low film stress is necessary to prevent such films from deformation after lift-off from the substrate. 10 Fundamentals 2.2.2 Mg alloy thin films For medical applications, Mg alloy thin films are of interest in different forms such as coatings or freestanding films for, e.g., stents or scaffolds [26, 61]. Fabrication by sputtering leads to the advantage of an easy change in structure (amorphous or crystalline), adjustment of thickness, production of supersaturated single-phase materials and structuring by other thin-film techniques [27, 61]. Sputtered Mg alloy thin films show a strong preferred orientation and, thus, higher texture than bulk material. This is influenced by a preferred growth in [0001] direction due to the lowest surface energy of the (0001) planes [28], leading for MgAg and WE alloys to a columnar growth with grain sizes in µm range and a grain length equal to the film thickness [28, 29]. The influence of different alloying elements such as Ag, Zn, Ca and REE on the film properties have been tested. While for MgAg typically only 2 wt% Ag are soluble in Mg, sputtering can form metastable solid solutions up to 6 wt% before precipitates occur [28]. By adjusting sputtering parameters, even metallic glasses such as Mg50Zn5Ca, which thus have a high solubility and no precipitates, can be fabricated [30]. The parameters need to be chosen as low power and high pressure to reduce diffusion. For many alloying elements such as Ag, Y and Gd, an increase in lattice parameter with addition of the alloying element is found [27, 31]. The surface roughness is also reduced for sputtered AZ alloys in comparison to sputtered Mg [101]. The texture and addition of elements influence the mechanical properties due to solid solution strengthening, high defect densities and small grain diameters [28, 31]. A higher yield strength is determined for MgAg, MgY and MgGd when including more of the alloying element, while the Young’s modulus stays similar [28, 31]. The maximum elongation is often decreased at the same time [28]. The influence of sputtering parameters for the different alloys was analysed. Studies claim an influence of the sputter pressure on the maximum elongation due to lower diffusion and increasing inert gas atoms addition to the films for higher pressure [29]. A change of pressure and power can change between brittle and ductile films, even leading to amorphous materials [29, 30, 102]. While the adapting of parameters for changing the microstructure is dependent on the specific alloys, a general route for the structuring and fabrication of freestanding thin films in a thickness range of 10-250 µm was developed by Haffner et al. [26] and is described as performed in the studies for this work. The outline is shown in figure 2.4. It includes UV-lithography, magnetron sputtering and etching steps. In the first step, a silicon wafer is coated by spin coating with an image reversal resist (a). After transferring a mask structure onto the wafer by illumination with ultraviolet light and removing the non-illuminated photoresist (b), a layer of Al with a thickness of approximately 150 nm is added by magnetron sputtering (c). The resist is then dissolved in acetone so that only the structure remains as an Al hard mask (d). A Bosch process is then carried out by ICP-RIE (inductively coupled plasma - reactive ion etching) to etch up to a depth of at least 20 µm for the films produced in the following studies, often further to up to 60 µm (e). After cleaning in acetone and isopropanol, AlN is added with a thickness of 500 nm by reactive magnetron sputtering using an Al target and nitrogen gas (f) before the final sputtering of Mg alloy (g). During the lift-off in 20 wt% KOH, the sacrificial AlN layer is dissolved (h). The freestanding thin films can then be cleaned in distilled water and isopropanol before drying them with nitrogen. 11 Fundamentals Figure 2.4: Fabrication steps for freestanding Mg alloy thin films using UV-lithography, magnetron sputtering and sacrificial layers as performed for this work, based on the process of Haffner et al. [26]. 2.3 Corrosion Corrosion can occur by chemical or electrochemical means or with the additional influence of biological matter [103 a, 46 b]. It is, in general, defined as the degrading or destruction of a material due to reactions with the environment. While corrosion was defined in the past mainly for metals, it can now include the degradation of all materials [46 b]. The type of corrosion described in the studies of this thesis is electrochemical, more specifically, metal corrosion and oxidation in an aqueous solution. The overall reaction includes an anode, a cathode, an electrolyte and a path for the electrons to transfer between the electrodes, forming a corrosion cell. Other factors, such as temperature, pH, and salt in the solution, can further influence the reactions that take place. [104 a] While the anodic reaction includes the corrosion and dissolution of the sample by oxidation, the cathodic reduction reactions can include different reactions depending on the reaction medium and electrode potential. The probability or strength of corrosion is described by the electrode potential, which is always defined as the potential compared to another electrode. Metals with a lower electrode potential are more prone to corrosion and will act as anodes if they are in contact with metals with a higher potential. [103 b] While corrosion in aqueous solution, in general, is defined by the ion release, this does not have to be a continuous release over the surface. The different types of corrosion are split into homogeneous and local corrosion, which includes, but is not limited to, galvanic, dealloying, intergranular, filiform, pitting and crevice corrosion [103 c, 46 c, 104 b]: •For homogeneous or uniform corrosion, the reaction takes place on the overall surface. It often occurs if no stable passivating layer is formed. While measurements and calculations for corrosion rates only assume uniform corrosion, other corrosion types also take place and need to be considered. •In galvanic or microgalvanic corrosion, two metals or metal phases are in contact 12 Fundamentals with each other, leading to faster corrosion of the compound with the lower electrode potential. •Similar to galvanic corrosion, dealloying corrosion, also called selective corrosion, describes the dissolution process of one specific alloying element due to the lower potential. •If a difference in potential occurs at grain boundaries due to, e.g., preferred accumulation of one compound at boundaries, this could lead to intergranular corrosion. •When corrosion products are formed during corrosion, they can form filaments on the surface, which start from surface defects and lead to an active anodic head of the filament and a growing cathodic area of corrosion products. This process is called filiform corrosion. It often occurs if the corroding material is coated. •Pitting corrosion occurs if the material itself is not homogeneous or the passivation layer forms cracks or other defects, leading to a preferred attack and formation of pits at those places. It can also be influenced by highly reactive anions in solution. •Crevice corrosion takes place if another object limits the volume above an area of the sample and less exchange of solution occurs, leading to localised corrosion. 2.3.1 Mg corrosion Mg has a very low electrode potential of -2.37 V vs SHE [46 a]. The redox reaction consists of the anodic Mg ion release and cathodic hydrogen evolution. Mg ions and hydrogen can then further react to Mg hydroxide. [105] Oxidation: Mg −−→ Mg2+ + 2 e– Reduction: 2 H2O + 2 e–−−→ 2 H2+ 2 OH– Redox: Mg + 2 H2O−−→ 2 H2+ Mg(OH)2 Since Mg often reacts to MgO and forms an oxide layer in air, the reaction to Mg(OH)2 can also be rewritten as 2 Mg + O2−−→ 2 MgO MgO + H2O−−→ Mg(OH)2 The change and reaction of the surface layer are influenced by the stability of such layers on the surface and solubility in water. The stability is typically described by the PillingBedworth ratio (PBR) [106, 107]. The ratio defines the stress of a surface layer on the substrate by calculating the ratio of the corrosion product’s molar volume to the original metal’s molar volume. A value of 1, therefore, defines a stress-free film, and any value of PBR=1-2 is defined as relatively stable. While <1 means that tensile stress occurs, a 13 Fundamentals PBR>1 leads to compressive stress. For MgO, the PBR is <1 with 0.8, but for Mg(OH)2, a value of 1.8 and, therefore, a more stable film is reached [108]. The stability of surface layers is, however, also dependent on an applied potential and the pH of the aqueous solution. Those parameters and the responding products are given in Pourbaix diagrams, showing different areas of reactions [109]. The reaction and formation of Mg(OH)2is just stable for a higher, alkaline pH [109] and, thus, the pH influences the corrosion by reducing the corrosion rate for a higher pH [110, 111]. Figure 2.5: Corrosion rate dependent on the pH in Hank’s balanced salt solution (buffered) for three exemplary Mg materials (high purity Mg HP Mg, AZ91, ZE41 Zn+REE). Reproduced and adapted with permission from Elsevier [111]. At a pH of 7.4, which corresponds to the pH in a human body, neither Mg(OH)2nor MgO are stably formed and therefore do not hinder further reaction and dissolution [112]. However, Mg(OH)2can still form at lower pH due to an increase of pH close to the sample by the reactions taking place [32]. Since the corrosion rate reduction due to the hydroxide layer is higher for a higher pH, the open circuit potential (potential of the sample vs. reference in a certain system if no voltage is applied, EOCV) can also shift pH dependently [113, 114]. An additional point to consider when describing the corrosion of Mg is that the cathodic and anodic reactions cannot be easily separated. For a standard metal, the cathodic reaction - in this case, hydrogen evolution - takes place if a negative potential is applied, while a positive potential leads to more anodic reaction, thus, metal dissolution and a decrease in the cathodic reaction rate. However, for Mg alloys, an enhanced voltage also leads, in many cases, to an increase in hydrogen evolution and a fast Mg dissolution, which already starts in the cathodic region. This is called negative difference effect (NDE) [105, 115, 116]. Certain theories, such as the presence of Mg+ions or the influence of the protective film as cathodic regions for hydrogen evolution, have been proposed. Both hypotheses were contradicted by other studies [34, 115, 117–119]; thus, no final explanation has been determined. 2.3.2 Influencing factors The corrosion process and rate of Mg are very complex and influenced by several factors, such as the structure, including microstructure and texture, the solution in which corrosion 14 Fundamentals takes place and additional elements. The importance of these factors will be discussed in the following sections. Influence of structure The structure of the Mg samples can highly influence the corrosion rate. This is often used to lower the rate by thermal or mechanical treatments [120]. Hereby, two main factors are changed: grain size and orientation. The influence of grain size is widely studied. Grain boundaries, in general, count as defects, increase the area available for corrosion start and, therefore, could increase the corrosion rate, especially if a change of composition occurs between grain volume and boundary [121, 122]. However, it is often found that small grains lead to improved corrosion resistance due to the formation of a more uniform passivating film [121–123], smaller precipitates [32] or a change in potential at the grain boundaries [124]. The formation of a better passivating film also explains why short-term measurements might show higher corrosion rates for small grains since the film has not been stably formed. The orientation of grains and, thus, the texture also changes the corrosion rate. Even though all grains consist of the same phase, each plane can show a slightly different corrosion rate, mainly influenced by the strength of bonds between the atoms. Therefore, theoretically, the densest plane should possess the best corrosion resistance, which is the (0001) plane in the case of Mg [125, 126]. However, planes with the lowest activity for reactions also have the lowest activity towards, e.g., oxidation, therefore forming thinner oxide or protective layers [127–129]. A uniform texture might also be critical for low corrosion additional to the specific plane since it prevents galvanic coupling [130]. Therefore, the preference of a structure for good corrosion resistance has to be determined for each particular system of sample alloy and surrounding. Influence of solution As already described for the pH in section 2.3.1, the solution in which corrosion takes place strongly affects the corrosion rate. This is even more true if additional ions are involved [131]. Kwon et al. showed a strong difference in the corrosion behaviour of a Mg alloy if different ions were present in the solution even though the pH was held constant [132]. Silva et al. found that the oxygen content can also influence the corrosion [133]. When working under physiological conditions, either just in a salt solution or in vitro or in vivo, the system is complex. While added organic components such as molecules and cells can also influence [134], the experiments in these studies are carried out in salt solution; thus, this effect is mainly described in the following. In salt solutions such as Hank’s balanced salt solution (HBSS), hydrocarbonates, carbonates, and carbon dioxide are often present as buffering systems or from air, leading to the formation of MgCO3or more complex carbonates [135]. The formation of carbonates, in general, was found by Santucci et al. [112] to be more stable with increasing pH and increasing concentration, while MgCO3specifically is stable in a broad pH range. Additional salts, including elements such as Ca, can lead to the formation of further carbonates or phosphates. Ca phosphate is especially known to form a stable product which reduces the corrosion rate [131, 136] (table 2.1). It also has a PBR of 1.27 and, thus, forms a 15 Fundamentals surface film with low film stress [108]. A complex system including cells or in vivo measurements can influence the corrosion rate [134, 137], but the influence and process of corrosion product formation are still important. Zhao et al. found, e.g., Mg carbonates and phosphates on Mg alloys after implantation into a mouse model, indicating that the effect of salts in solution is also of interest for the final application [138]. Influence of purity When discussing the influence of additional elements on corrosion, two possibilities have to be considered - impurities and alloying elements. Impurities are, in general, unwanted additions of elements. When comparing ultra-pure Mg with high-purity Mg, the corrosion rate of ultra-pure Mg is significantly lower [139]. The impurities are problematic if they have low solubility and, thus, form additional phases with higher electrode potential which can lead to microgalvanic corrosion. Since nearly all elements have a higher electrode potential than Mg [46 a], this results in faster Mg dissolution. One impurity which is often included is iron. If the concentration exceeds 0.005-0.017 % [140, 141], it can increase the corrosion rate. Techniques such as sputtering might prevent the formation of additional phases and, therefore, reduce such effects [27]. Elements can also be added on purpose as alloying elements to influence the properties. This can include changing the microstructure or surface properties to reduce the corrosion rate, at least in comparison to Mg with low purity [120]. REE have a similar electrode potential to Mg, and the potential of Ca is even lower, leading to a lower overall potential of the intermetallic phase [142]. A decrease in size of the second phase or continuous distribution preventing the propagation of pits can, e.g., reduce the corrosion rate [108, 143–146]. Since Mg alloys can form passivating surface layers, sputtered Mg alloys have also been discussed as self-healing layers, which can prevent galvanic corrosion [120]. 2.3.3 Mg thin film corrosion The structure of thin films produced by sputtering results in two main factors which can reduce the corrosion, namely a strong texture with basal (0001) planes on the surface and a supersaturated solid solution preventing the formation of additional phases up to higher fractions of the alloying element [27, 31, 126, 147]. Therefore, less pitting corrosion and other local corrosion is found for sputtered samples since less galvanic coupling can occur with precipitates [32]. Blawert et al. also discussed the incorporation of alloying elements of thin films into the passivating layers, leading to a better corrosion resistance if distributed homogeneously, which can further be influenced by changes in microstructure [101]. For WE alloys, no effect of the alloying elements on the corrosion rate occurs since no precipitates are found [102]. Even an increase of the alloying element does not lead to changes as studied for up to around 23 at% of Y or 3.6 at% of Gd [31]. For MgAg, the addition of the element even decreases the corrosion rate by increasing the standard potential when forming a solid solution [27]. A reduction is also measured for amorphous MgZnCa, for which the corrosion rate is possibly reduced by reducing the hydrogen evolution via Zn 16 Fundamentals PBR [108] Ksp[149] Li2CO31.35 8.15·10−4 Mg(OH)21.80 5.61·10−12 MgO 0.80 - LiOH 1.26 - MgCO32.04 6.82·10−6 Mg3(PO4)22.29 1.04·10−24 Li3(PO4)21.20 2.37·10−11 Ca3(PO4)21.27 2.07·10−33 Table 2.1: Pilling-Bedworth ratios (PBR) and solubility product constants (Ksp at 25 ◦C of possible corrosion products of MgLi in salt solution. incorporation [30]. For this material, an influence of the sputtering parameters on the corrosion by changing from polycrystalline to amorphous material is described. 2.3.4 MgLi corrosion The electrode potential of Li is even lower than for Mg with -3.05 V vs SHE [46 a]. Thus, Li enhances the reactivity and the corrosion rate. It is found to increase the hydrogen evolution by increasing the cathodic kinetics [35, 148]. While Li is studied to reduce the corrosion resistance short term, studies also show enhanced resistance over longer measurement times [33]. This might be influenced by a stronger increase in pH during the beginning, which stabilises Mg(OH)2on the surface afterwards [73]. Since the influence of corrosion products on the corrosion resistance is influenced by the stability and solubility, PBR and solubility product constants of possible components formed by MgLi in salt solutions are listed in table 2.1. The effect of Li on the corrosion is highly influenced by the exact amount of Li. The main change depends on the MgLi phase. In general, the corrosion resistance is often described to be the highest for β-MgLi, followed by α-MgLi and then α+β-MgLi [150]. The corrosion rates and influences for each phase are described in the following sections. α-phase For MgLi with a hcp structure, the corrosion is generally assumed to be similar to the corrosion of pure Mg [151]. Thus, the formation of MgO and Mg(OH)2will occur [152]. Since Li has a lower electrode potential than Mg, the corrosion potential of the alloy decreases with increasing Li content [153]. During the corrosion, anodic and cathodic regions develop on the surface. Hydrogen and corrosion products are formed in the cathodic regions and the corrosion progresses as filiform corrosion [150, 154]. While Li can lower the overall electrode potential, Li et al. found for α-MgLi with different Li fractions the best corrosion resistance for the highest Li fraction due to the influence on microstructure and surface films [153]. In addition to the corrosion products on Mg, Li-containing components are found. For MgLi-based alloys such as, e.g., LA51 (Mg-5Li-1Al in wt%), Li2CO3and LiOH are formed in air on the surface [155]. 17 Fundamentals corrosion. The measurement generally does not show corrosion in a non-disturbed system and varies significantly from the corrosion rate of weight loss or hydrogen evolution [180]. One problem is also the assumptions for the Tafel extrapolation, which only describes one reaction occurring for each branch. In praxis, this does not hold true, especially considering the NDE, which, together with the formation of oxides and hydroxides, often prevents a crossing of both linear extrapolations at EOCV [118, 180–182]. The hydrogen evolution and formation of bubbles on the sample surface can further influence the results [183], especially considering small measurement areas as often used for the thin film samples. However, an extrapolation from the cathodic branch can be used to estimate a corrosion rate and compare different samples placed in the same set-up [180, 184]. 2.4.4 In-line ICP-MS During the corrosion in solution, the sample itself gets dissolved. Therefore, additional information regarding the corrosion rate and, even more importantly, about the specific released element can be determined by analysing the dissolution product in the electrolyte. Samples of solution with dissolved species can be analysed by techniques such as inductively coupled plasma - mass spectrometry (ICP-MS) or inductively coupled plasmaatomic emission spectrometry (ICP-AES, also called ICP-OES, OES = optical emission spectrometry) to determine, e.g., catalyst loading, stability of partially metallic catalysts or the concentration in solution after a defined duration of corrosion [66, 79, 185–187]. However, this only gives the integral information over time between measurements. Based on previous measurement techniques [188], an in-line technique was presented by Ogle and Weber in 2000 [189]. The AESEC (atomic emission spectroelectrochemistry) combines ICP-AES with an electrochemical flow cell upstream in which the sample is added as the working electrode. Thus, the dissolved elements can be directly analysed downstream and during electrochemical measurements. Similar set-ups with ICP-MS have been developed [185, 190, 191], and both types are used for the study of time-resolved corrosion or catalysis [185, 188, 192]. Both will be discussed; however, the main focus is on the technique used for studies in this thesis, thus, a combination of the flow cell with ICP-MS. In both cases, the system consists of the flow cell - a cell with defined sample surface in contact with solution and defined volume over the sample through which the electrolyte is pumped with a defined speed -, the spectrometer and the data acquisition system. Multiple important parameters have to be taken into account during the measurement. While, e.g., the flow rate can be determined to compare the ion release of different samples measured with slightly differing rates, this rate will also directly influence the corrosion process by, e.g., hindering precipitation, changing passivation layers or pH changes [188, 193–195]. The flow cell includes a three-electrode set-up with a counter electrode and reference electrode. In figure 2.9, the set-up of the flow cell [196, 197] described in this work is shown. In the ICP-MS system, the electrolyte is converted into an aerosol in the nebuliser by Ar flow and then transported into the spray chamber, from which only a few droplets will be transferred into an Ar plasma to bring the elements into the gas phase and ionise them. Afterwards, the ions are passing into a vacuum chamber for mass determination. Here, at first, any remaining neutral particles are separated from the ions, and the beam is focused before it reaches a collision/reaction cell. In this cell, any ions 24 Fundamentals Figure 2.9: Schematic of flow cell as described in [196, 197] and inline ICP-MS measurement. a) Flow cell as used for the measurement of Mg alloy samples, including a Pt counter electrode (CE), reference electrode (RE) and the sample as a working electrode (WE), connected by copper tape. b) Cross section of flow cell. c) Overall set-up with solution flow through the flow cell to the ICP-MS. including multiple atoms are split, in the case of the machine used in this work (Agilent 7900 ICP-MS) by collision with He atoms. The quantification dependent on the mass takes place using a quadrupole mass spectrometer in high vacuum to avoid recombination. While ICP-MS allows the quantification of isotopes and detection down to a limit of ppt (parts per trillion), the electrolyte needs to be much more diluted than for ICP-AES. [185, 188] The dissolution rate jmis then defined by the concentration c(t) after calibration, the flow rate fof the electrolyte and the area of sample Ain contact with said electrolyte [185, 188]: jm(t) = f·c(t)/A (2.6) To calculate the dissolution current density i, the Faraday constant F, electrons released per atom nand the molar mass Mare taken into account i(t) = jm(t)·n·F/M (2.7) While this current density (if it is added up for all elements) can, in theory, be compared to the results of, e.g., potentiodynamic polarisation if a voltage sweep is performed in the flow cell (e.g. [182]), the influence of side reactions, signal broadening and transfer time from flow cell to ICP-MS need to be taken into account [182, 188, 189]. By applying a potential or current, a change in composition during and after passivation regions can also be determined. The combination with the flow cell and identification of each element individually, thus, makes the identification of the stoichiometry of dissolved elements possible. Therefore, selective dissolution due to dealloying, preferential corrosion, film formation during corrosion, or catalytic reactions can be measured time resolved. The combination with electrochemistry also helps to receive information not directly available from ICP-MS, such as the oxidation state of the elements after corrosion [185, 188]. Several studies by both in-situ ICP-MS and ICP-AES have been carried out for further understanding the corrosion process and protective layers during corrosion. The system also allows to analyse the ion dissolution during other electrochemical measurements [198] or reactions under additional influences such as, e.g. light or scratching [185, 191, 196, 25 Fundamentals 199, 200]. The techniques were used for several studies on Mg and MgLi. For Mg corrosion, studies by Thomas et al. investigated the influence of impurities of Fe, proving an increase in activity [201]. The influence of a pH change and applied anodic potential was tested by Rossrucker et al. [110, 182], showing a lower NDE for low polarisation and high pH. Compared to ex-situ measurements, the dependence on time and change after applying anodic polarisation can be directly observed. For Mg corrosion, earlier studies suggest the possibility of the formation of Mg+additional to Mg2+ as a reason for the NDE occurring for Mg, however, analysis by ´ Swiatowska et al. [117] and Lebouil et al. [115] showed only Mg2+ and rather suggested the growth of cathodes by corrosion product formation if an anodic potential is applied. This also leads to an improved stability of the dissolution after anodic polarisation when the sample is held at EOCV. The formation of hydrogen gas can be identified as noise in the measurements [117]. For MgLi, specifically with high Li content (β-phase), the composition of a passivation layer formed during corrosion was analysed by determining the change in stoichiometry when applying an increasing voltage and scratching the surface [162, 166, 200], showing, e.g., a preferred Li dissolution after the scratch in the study by Yan et al. [200]. Hou et al. found a greater Mg:Li ratio in the dissolved ions than in the material itself and determined the presence of Li2CO3by a change to a ratio similar to bulk when the voltage is increased over a certain threshold [166]. However, Yan et al. [162] did not find the same difference in ratio but rather found a Mg-rich layer determined by Mg-rich particles released after scratching. For the studies of thin films, the possibility to gain insight into the start of the corrosion in-situ is especially important since the surface affects the overall sample properties significantly due to the high surface-to-volume ratio. Additionally to thin films on substrate, free-standing thin films can also easily be placed into the flow cell, however, special care has to be taken to prevent corrosion mainly as crevice corrosion when the thin film is bent around the O-ring [197]. Since the dissolved ions are directly detected, no effects such as the NDE can influence the assumed corrosion rate compared to the actual ion release. Since the detection limit is very low, there is also no effect of concentrations that are too low due to the thin films, especially since, in any case, the technique analyses in short time frames, and thus, no long-term accumulation of corrosion products is necessary. It also allows for the identification of light elements such as Li, therefore allowing the studying of nearly any Mg alloys. While the technique can be used for longer in-situ experiments, the measurement times for thin films are limited to prevent crevice or pitting corrosion through the whole film to avoid contact between the electrolyte and the substrate material or sample holder additional to the sample. Thus, no long-term corrosion properties are analysed. Since the machine is very sensitive, the electrolyte needs to be adjusted. While the corrosion rates were mainly determined in HBSS (155 mmol) at a pH of 7.4±0.2 and a temperature of around 37±1◦C (see 2.4.1, 2.4.3) to simulate physiological conditions, for the in-situ measurements, the concentration was adjusted to 15 mmol HBSS during the experiments for this work. The pH was close to 7.4 but no temperature control was applied. This is, however, possible by, e.g., a water bath heating [188]. In-line ICP-MS might not be necessary to determine the corrosion rate, which can be done with less effort with other techniques, but it provides a further understanding of corrosion processes and element-dependent dissolution of MgLi thin films. 26 Fundamentals 2.4.5 Further measurements While several techniques have been discussed to determine the corrosion rate and analyse the corrosion process, those often need to be combined with additional measurements [105]. Those can be used for studying the surface preand post-corrosion to investigate a change in surface composition by, e.g., formation of corrosion products (e.g. SEM/EDX (scanning electron microscopy/energy dispersive X-ray spectroscopy), XRD (X-ray diffraction), XPS (X-ray photoelectron spectroscopy)), or change in surface morphology (e.g. AFM (atomic force microscopy), SEM) for defining the overall change during corrosion. Additional measurements such as EIS (electrochemical impedance spectroscopy) can give further insight into the electrochemical processes taking place at different time scales and help to identify the formation of corrosion layers. Thus, since the corrosion of Mg and Mg alloys in aqueous solution and especially in more complex electrolytes is a complex system, no single technique can give all results necessary to understand the ongoing process. 27 Results 3 Results To analyse the viability of MgLi thin films as a local drug reservoir for the treatment of neurological disorders, in the following studies and chapters with additional experimental results, the fabrication, structural properties and corrosion rate of MgLi thin films with varying Li content will be discussed. In addition to the general test of the properties of such alloys, the studies opt to further understand the corrosion, the influence of participating parts such as phases, Li content, microstructure, orientations or surface films on the corrosion rate and a more detailed corrosion process and ion release will be described. This not only provides a more thorough understanding of the underlying mechanisms occurring but also allows the understanding of adjustments to the corrosion rate by changing the sputtering parameters or applying additional post-treatments. 3.1 Publication: Structural characterisation and degradation of Mg-Li thin films for biodegradable implants To develop MgLi thin films for medical treatments, freestanding films of the alloys have to be prepared and studied. In this work, thin films in αand α+β-phase were fabricated by magnetron sputtering with low film stress. The microstructure and phases were analysed, showing not only β-phase for higher Li fractions but also additional Li2CO3and a change in orientation. In mechanical tests, a lower tensile strength was found for all MgLi compositions compared to pure Mg, and the maximum elongation was reduced except for samples with the lowest Li fraction. While the corrosion rate increased for an increase in Li fraction in the α-phase due to higher reactivity and the influence of orientation changes, a lower corrosion rate was found at high Li fraction in the mixed phase, indicating a possible effect of protective surface films. Own contributions to the following article (concept - 50 %, planning - 75 %, experiments - 75 %, analysis - 90 %, writing - 90 %): •Sample preparation and preparation optimisation •XRD, SEM and EDX investigation •Tensile testing, weightloss and potentiodynamic polarisation measurements •Interpretation and discussion of the results •Writing of the manuscript 28 Results The paper Hanke, L., Jessen, L.K., Weisheit, F., Bhat, K., Westernstr¨oer, U., GarbeSch¨onberg, D., Willumeit-R¨omer, R., Quandt, E. Structural characterisation and degradation of Mg–Li thin films for biodegradable implants. Scientific Reports 13, 12572(2023) [202] published by Springer Nature is open access and the use is permitted by the Creative Commons Attribution 4.0 International License (http://creativecommons.org/ licenses/by/4.0/). 29 1 Vol.:(0123456789) Scientific Reports | (2023) 13:12572 | https://doi.org/10.1038/s41598-023-39493-9 www.nature.com/scientificreports Structural characterisation and degradation of Mg–Li thin films for biodegradable implants Lisa Hanke 1, Lea K. Jessen 1, Felix Weisheit 1, Krathika Bhat 2, Ulrike Westernströer 3, Dieter Garbe‑Schönberg 3, Regine Willumeit‑Römer 2 & Eckhard Quandt 1* Freestanding thin films of Mg–Li (magnesium–lithium) alloys with a Li mass fraction between 1.6% (m/m) and 9.5% (m/m) were prepared and studied with respect to their structure and degradation properties. With increasing Li content, the microstructure deviates from hexagonal Mg–Li with strict columnar growth and preferred orientation, and additional cubic Mg–Li and Li2CO3 occur. The corrosion rate was measured in Hanks’ balanced salt solution by potentiodynamic polarisation and weight loss measurements to investigate biodegradation. Influences of the orientation, phase and protective layer formation lead to an increase in corrosion from 1.6 to 5.5% (m/m) from 0.13 ± 0.03 to 0.67 ± 0.29 mm/year when measured by potentiodynamic polarisation but a similar corrosion rate for 9.5% (m/m) and 3% (m/m) of Li of 0.27 ± 0.07 mm/year and 0.26 ± 0.05 mm/year. Magnesium and its alloys are widely studied as materials for applications in the medical field due to their biodegradability. Different elements such as, e.g., Ca, Zn or rare earth elements (REE) are included to improve mechanical properties or tailor the degradation rate to adjust them for applications as, e.g., stents or bone implants1–4. Additional to the advantage of having an implant which degrades after it is no longer required, the possible therapeutic effects of the implants are explored by, e.g., loading stents with drug-eluting layers5 or using the corrosion process and changes in the environment such as pH and hydrogen evolution directly for its antibacterial properties6. In line with the idea to use the implant itself as treatment, an alloy including the therapeutically active element lithium will be analysed in this study. Lithium is used in treatments for mood disorders, in particular bipolar disorder, and is also studied to have effects on Alzheimer’s and Parkinson’s disease7–10. Magnesium itself shows also neurological effects11. Therefore, the degradation and, thus, continuous release of both the magnesium and additional elements would allow a local treatment in the brain. If a controlled and local release is achieved by understanding the degradation of the material, side effects which can occur during the treatment with Li12 could be reduced. For Mg–Li, the structure in bulk materials differs from pure Mg by a reduction of the distance in c-direction and a phase change to a body centred cubic (bcc) phase (β phase) for higher Li fractions (Mg–Li phase diagram, Fig.1 13). This change leads to additional non-basal slip on the prismatic planes, twinning, and more ductile properties even in hcp Mg–Li alloys. For higher Li fractions, the addition of the second phase can significantly change the mechanical properties. Li etal. showed that cracks are preferably formed at the phase boundaries which is facilitated by the difference in the number of gliding systems present in both phases and, thus, a difference in stress accumulation14. Additionally, the ageing of the second phase and change from bcc to hcp phase even at room temperature influences the properties over time14,15. The corrosion rate of Mg–Li alloys is influenced by several factors such as the high activity of Li, change of microstructure and surface films. The low electrochemical potential of Li leads to an increase in the cathodic kinetics and a more significant shift of the pH. Filiform corrosion is found to be one of the main corrosion processes occurring for Mg–Li alloys in the α or α + β materials16–18. For films with mixed phases, micro-galvanic coupling is found as a main factor for an increase in corrosion rate with preferred corrosion and pitting at the phase boundaries19,20. However, the diversity of microstructure and protective layers formed during corrosion makes a clear indication of the influence of the different factors on the corrosion rate difficult. Li etal. showed that the corrosion rate decreases from α + β > α > β18. The lower corrosion rate of the bcc phase is assumed to be OPEN 1Inorganic Functional Materials, Institute for Materials Science, Faculty of Engineering, Kiel University, Kiel, Germany. 2Institute of Metallic Biomaterials, Helmholtz Centre Hereon, Geesthacht, Germany. 3Marine Climate Research, Institute of Geosciences, Faculty of Mathematics and Natural Sciences, Kiel University, Kiel, Germany. *email: [email protected] Results 30 2 Vol:.(1234567890) Scientific Reports | (2023) 13:12572 | https://doi.org/10.1038/s41598-023-39493-9 www.nature.com/scientificreports/ due to the high density and stability of a formed protective layer. While for the Mg rich phase, mainly a porous Mg(OH)2 layer is expected to form during corrosion21,22, the layer structures of surface films formed on Mg–Li including the bcc phase are complex. Xu etal., e.g., analysed the structure of the natural film formed in air as a Li2CO3 film on the surface, Mg oxide and Li oxide film underneath and Mg-rich film before the bulk material23. Other studies claim for films formed in air or during corrosion several compounds including carbonates, oxides and hydroxides of both Li and Mg, often separated in a layer structure16,24,25. Previous studies assumed or suspected that the formed Li2CO3 has the main influence on the higher corrosion resistance of the bcc phase23,26,27. The Pilling–Bedworth ratio (PBR), which is a measure for film stress and, thus, identifies a stable film for 1 < PBR < 2, is > 1 for all Mg:Li ratios for Li2CO3. Thus, it could already be formed for lower Li mass fractions in the hcp phase18. Yan etal. suggest another possible influence as Li doping and therefore strengthening the MgO and hindering the formation of the more porous and less protective magnesium hydroxide. Since the critical Li fraction for forming a stable layer of MgO is calculated to be around 15–18at.% (4.8–5.9%(m/m)), this is in agreement with the formation of the layer only on Mg–Li with β or α + β28. Therefore, with the assumption of the formation of a stable MgO layer by Li doping, the higher Li fractions leads to a decrease in corrosion rate by changing the film stress. However, it has to be noted that the corrosion of Mg–Li with different phases is a complex system, leading to other studies showing the lowest corrosion rate for the hcp phase19. The main influences on the actual corrosion rate are therefore not just Li content and phases but also the microstructure of the material. For neurological implants, small freestanding films and structures in the sizes of a few µm–mm are needed. Previous to this paper, no extensive studies of Mg–Li thin films regarding the growth and properties are available to our knowledge. As found for thin films of other Mg alloys, significant influences on the properties in comparison to bulk materials are expected. For different Mg alloys such as Mg–Ag or Mg–REE, the structure of thin films was studied29–33. For sputter-deposited films, the hexagonal close packed (hcp) magnesium phase (α phase) is strongly textured with a preferred growth direction of [001] and a columnar grain structure is apparent31,33–35. The strong texture influences the deformation by, e.g., reducing work hardening and influences therefore the mechanical properties33. The corrosion of thin films in comparison to bulk is shown to be more homogeneous with less effect of pitting35,36. Additionally, the corrosion resistance and oxidation of different planes are different due to the packing and binding energies, thus, the corrosion rate is influenced by the texturing of the films37. The (001) plane is found to show the lowest corrosion rate due to the densest packing, however, since the faster oxidation of other planes could lead to a protected surface, the actual influence on the corrosion rate cannot be directly predicted38,39. Since the properties are highly dependent on structure and microstructure, they can be influenced for the same alloy by changing the sputtering parameters, leading to higher densities or differences in the film growth depending on the energy available for diffusion of the atoms29,32,40–42. In this study, Mg–Li thin films with a pure hcp or an α + β structure (Mg–Li phase diagram, Fig.1) are prepared via magnetron sputtering. To gain insight into the corrosion process dependent on the specific structure present in the thin films, studies with respect to their growth and microstructure are carried out to allow a correlation with influences on the corrosion rate. Additionally, the possibilities to influence and tune those to meet specific requirements given for applications are discussed. Figure1. Mg–Li phase diagram, adapted from13. The region with a Li fraction from 0 to 50% (n/n) is depicted and the concentration of sample films (Li mass fraction of 1.6% (m/m), 3% (m/m), 5.5% (m/m) and 9.5% (m/m)) are marked. Low Li concentrations lead to a Mg rich α-phase with a hcp structure and high Li concentrations to a β-phase with a bcc structure. In the mixed phase regions, α-phase and β-phase are formed. Results 31 3 Vol.:(0123456789) Scientific Reports | (2023) 13:12572 | https://doi.org/10.1038/s41598-023-39493-9 www.nature.com/scientificreports/ Results Composition and microstructure. The compositions of the prepared Mg–Li thin films are marked with red dashed lines at mass fractions of 1.6% (m/m), 3% (m/m), 5.5% (m/m) and 9.5% (m/m) in the phase diagram in Fig.1 (fractions given in Supplementary Table1). As shown, two sample types are prepared with Li mass fractions leading to pure hexagonal phase, while the other two theoretically consist of α and β phase with approximately 89% (Mg–5.5Li) or 22% (Mg–9.5Li) α phase. Very low iron contaminations are measured for all film compositions. Representative XRD diffractograms for all compositions are displayed in Fig.2a in comparison to pure Mg prepared by the same procedure. While pure Mg exhibits a strong texture with a main peak of (002) at 34.3°, including 1.6% (m/m) Li leads to less preferred growth and for Mg–3Li, a random orientation is identified (Supplementary Fig.1). For both materials in the α + β phase, the hexagonal phase shows a preferred orientation of (110). Additional bcc can be identified as small peaks in the diffractogram, and additional studies of the reciprocal space allow the identification of strong (110) peaks for Mg–9.5Li at an angle of χ = 32°–40°, indicating a strongly textured β-phase. While the lattice parameter a is only slightly decreased from Mg–1.6Li to higher Li mass fractions, c is reduced, especially for the increase of Li from 1.6% (m/m) to 3% (m/m) (peak shift in Fig.2a, calculated lattice parameters in Supplementary Table2). When the second phase is present, the parameters do not decrease further since the added Li is included in the additional phase. Of interest are as well the peaks at, e.g., 21.3°, 23.3°, 29.4° and 34.1° for Mg–5.5Li and Mg–9.5Li which indicate the existence of Li2CO3 (Supplementary Table3). Because of the formation of this additional phase, the amount of β-phase is reduced. Investigations via EDX show a layer including carbon and oxygen on the surface of Mg–9.5Li, thus, the Li2CO3 identified by XRD can mainly be assigned to a film formed on the surface of the samples when those are stored in air. The formation of Li2CO3 in humid air for β-phase Li is also shown in previous studies23,28. A layer of Li2O is formed on the surface of alloys with high Li content and can react further to Li2CO3 if CO2 is present in the surrounding atmosphere23. There is no significant compositional change over the layer thickness for, e.g., Mg–1.6Li (Fig.2b). Cross-sectional images of the different Mg–Li alloy freestanding thin films with a thickness of 10µm are given in Fig.3a for analysis of the microstructure. Mg–1.6Li exhibits a columnar growth with a constant diameter of approximately 500nm to 1µm over the whole film thickness. This structure is also identified for pure Mg films with the strong (001) texture prepared via magnetron sputtering30. For Mg–3Li, smaller grains are formed close to the substrate while columns start after a few 100nm with increasing diameter up to 1.5µm to 2µm with a few columns exhibiting a diameter of around 4µm. Less columnar growth is visible for Mg–5.5Li and cannot be identified for the highest Li mass fraction. The surface, however, still exhibits a structure which leads to the identification of grain sizes of approximately 1.5–2µm. The difference in the cross-sectional images can not only be assigned to a change of columns to a different microstructure because of additional phases but also to less preferred fracture at the grain boundaries during bending. This is influenced by, e.g., voids formed due to the self-shadowing of the columns. For Mg–9.5Li, oxidation of the samples also plays a major role in the visible structure since the samples are highly affected. While a thin oxide film is formed for all films and is apparent in the surface images in Fig.3a, only for films with higher Li content (Mg–9.5Li) the oxide grows significantly until the film is completely oxidized (Fig.2c). To classify the microstructures further, they can be compared with the structure-zone model41,43. The substrate temperature during the deposition was (49–54)°C for Mg–1.6Li, (54–66)°C for Mg–3Li and (60–66)°C for Figure2. Structure and composition of freestanding thin films (a) XRD diffractograms for Mg, Mg–1.6Li, Mg–3Li, Mg–5.5Li, Mg–9.5Li thin films. The positions of the hcp and bcc Mg–Li phases and Li2CO3 are marked. Additionally, the orientation of the corresponding planes for the hcp phase are indicated, (b) EDX line scans of cross sections of Mg–1.6Li and Mg–9.5Li thin films as used for corrosion measurements, (c) EDX line scans of cross sections of Mg–9.5Li freestanding thin films after 1day and 8–9 months. Results 32 4 Vol:.(1234567890) Scientific Reports | (2023) 13:12572 | https://doi.org/10.1038/s41598-023-39493-9 www.nature.com/scientificreports/ Mg–5.5Li and Mg–9.5Li samples. This leads to T/Tm = 0.36 for Mg–1.6Li and T/Tm = 0.39 for Mg–9.5Li as the highest and lowest possible values, leading in theory to structures in the T-zone, changing into zone 2. For higher zones, the grain boundaries get denser, and the defect density decreases, therefore fewer voids occur which could influence the fracture at grain boundaries. Additionally, a shift from straighter fibres to more complex structures at lower layers for Mg–3Li is in good agreement with the microstructural cross-sections. Film growth. To identify and understand the differences in microstructure, the growth of Mg–1.6Li, Mg– 3Li and Mg–9.5Li films is additionally analysed. The sputter times were chosen according to the sputtering rates for 10µm to result in approximately 10nm, 100nm, 1µm and 20µm. The results are displayed in Fig.3b. The growth for the lowest Li mass fraction starts with island growth, similar to pure Mg44. The film covers the whole surface for a thickness > 10nm, resulting in columnar growth. The energy of the particles from the sputtering process itself and the low temperature of the substrate do not allow sufficient diffusion for a more homogeneous growth45. The columnar growth is additionally in good accordance with the structure formed with the preferred orientation of (001) because the fastest growth for hexagonal faces is in direction of the c-axis45. The film growth process is studied to be highly influenced by including alloying elements44. For Mg–3Li, the layer formation starts with layer growth, including only few defects in a 10 nm thick layer. After approximately (400–500)nm, the growth changes to columnar growth. This change can be attributed to film stress which accumulates over the layer and changes the energetically favourable growth. Pores and voids can already be identified after the column growth for a layer thickness of 1µm. Even though columnar growth is visible, the homogeneity of the signal on the reciprocal space mapping from XRD shows that a random orientation is formed throughout the whole film (Supplementary Fig.1). Since the growth is not as strictly orientated as for Mg–1.6Li, it leads to an increase in column diameter and void formation. For Mg–9.5Li, the film growth at the beginning cannot be directly identified as island growth, however, at a film thickness of approximately 100nm, a grain-like surface with a high roughness is visible. Even though these samples were measured directly after preparation and the oxidation of the samples is thus minimal, a colour change of the samples showed a slight oxidation even for the fastest possible measurement and therefore Figure3. SEM images (a) Surface and cross-section of freestanding Mg–Li thin films. (b) Side view or crosssection of Mg–Li films (Li: 1.6% (m/m), 3% (m/m), 9.5% (m/m)) on Si substrate with thicknesses of 10 nm, 100 nm, 1 µm and 20 µm. Results 33 11 Vol.:(0123456789) Scientific Reports | (2023) 13:12572 | https://doi.org/10.1038/s41598-023-39493-9 www.nature.com/scientificreports/ 58. Haffner, D., Zamponi, C., de Miranda, R. L. & Quandt, E. Micropatterned freestanding magnetron sputtered Mg-alloy scaffolds. BioNanoMaterials 16, (2015). Acknowledgements The authors thank Dr. Heike Helmholz from the Institute for Metallic Biomaterials of the Helmholtz Centre hereon for her support with the AAS measurements. This work was supported by the DFG in the framework of the research training group 2154—Materials for Brain (project 270394294). Author contributions L.H. and L.K.J. conception and design of the study, L.H. writing-original draft, analysis of results; L.H., F.W., K.B. and U.W. performing experiments, E.Q, R.W.-R. conception of general project, E.Q., R.W.-R. and D.G.-S. supervision of the work. All authors reviewed the manuscript. Funding Open Access funding enabled and organized by Projekt DEAL. Competing interests The authors declare no competing interests. Additional information Supplementary Information The online version contains supplementary material available at https:// doi. org/ 10. 1038/ s4159802339493-9. Correspondence and requests for materials should be addressed to E.Q. Reprints and permissions information is available at www.nature.com/reprints. Publisher’s note Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. Open Access This article is licensed under a Creative Commons Attribution 4.0 International License, which permits use, sharing, adaptation, distribution and reproduction in any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons licence, and indicate if changes were made. The images or other third party material in this article are included in the article’s Creative Commons licence, unless indicated otherwise in a credit line to the material. If material is not included in the article’s Creative Commons licence and your intended use is not permitted by statutory regulation or exceeds the permitted use, you will need to obtain permission directly from the copyright holder. To view a copy of this licence, visit http:// creat iveco mmons. org/ licen ses/ by/4. 0/. © The Author(s) 2023 Results 40 Results 3.2 Influence of sputtering parameters on the corrosion rate of Mg-3Li As discussed in the publication in chapter 3.1, the corrosion rate can not only be influenced by the Li content itself but also by the structure, e.g., the orientation. To investigate this influence further, the impact of a structure change by adjustment of sputtering parameters and, thus, the direct tuning of the corrosion rate during the sputtering process, is discussed in the following study. Power [W] Pressure [10−3mbar] Sputtering rate [nm/s] 50 2.3 1.78 100 1.5 3.19 2.3 4 6 150 2.3 4.70 Table 3.1: Set of sputtering parameters for magnetron sputtering of Mg-3Li to analyse the influence of power and pressure change on the structure and corrosion rate of the thin films. The properties of thin films prepared by magnetron sputtering can be influenced by changing the parameters such as pressure, power or gas flow. In this study, a variation of power (50-150 W) and pressure (1.5-6·10−3mbar) for Mg-3Li is exemplarily investigated to identify possible influences on the structure of the film and the resulting corrosion rate (sputtering parameters table 3.1). Figure 3.1 shows the corrosion rates for all tested parameter combinations. 50 100 150 0.15 0.20 0.25 Corrosion rate [mm/yr] Power [W] Constant pressure 2.3*10 -3 mbar (a) 2 4 6 0.10 0.15 0.20 0.25 Corrosion rate [mm/yr] Pressure [10 -3 mbar ] Constant power 100 W (b) Figure 3.1: Comparison of corrosion rates for Mg-3Li thin films on substrate sputtered with different sputtering parameters (for power change: a) constant pressure of 2.3·10−3mbar, for pressure change: b) constant power of 100 W). Since the film stress varies, all corrosion studies were performed on substrate. Different corrosion potentials between -1.70 V and -1.79 V are measured for all parameter sets, but 41 Results no direct dependence on power or pressure can be assigned. The lowest corrosion rate is determined for samples prepared with a power of 100 W and pressure of 4·10−3mbar. Especially for the highest pressure of 6·10−3mbar, the potentiodynamic polarisation measurement is more unstable with high fluctuations, including steps from possible pitting and depassivation, therefore leading to higher errors. Since not all parameter combinations are tested, this study does not aim to determine the lowest or highest possible corrosion rate but instead to identify possible correlations between parameters, microstructure and corrosion rate, therefore showing tuning possibilities. Film stress analysis by producing freestanding thin films and comparing the rolling ratios shows that changing the parameters leads to significant film stress (figure 3.2). Figure 3.2: Comparison of freestanding Mg-3Li sputtered at different pressure and power (for power change: constant pressure of 2.3·10−3mbar, for pressure change: constant power of 100 W). The film stress can be identified by the rolling of the films. For a pressure of 2.3·10−3mbar, the lowest film stress is exhibited for 50 W samples; higher power leads to compressive stress. At 150 W, both compressive and tensile stress samples can be found, depending on the position on the wafer during sputtering. The surface of all samples, but especially for 100 W, are shiny and, therefore, smooth. At 100 W, a change of pressure leads to flat and stable films for a pressure of 4·10−3mbar with compressive stress for the other films, thus not allowing a direct trend. For all pressures sputtered at 100 W, the surface is smooth. However, a colour change for the lowest pressure indicates a change in the sample surface (figure 3.2). Factors such as grain size, roughness, and grain boundaries could also influence the corrosion rate. As discussed by Qu et al., smaller grains can decrease the biodegradation rate of pure Mg [124]. The improved corrosion resistance could be influenced by an improved oxide layer or grain boundaries acting as corrosion barriers [124, 203]. However, if the surroundings do not allow for sufficient passivation due to high corrosion rates [122], the corrosion rate can also be increased. Grain boundaries offer possible attack points for corrosion if the structure is not dense due to increased surface area. Here, the grain size decreases with increasing sputtering power, and the structure is denser, with smaller grain structures apparent on the surface. This reduces the area of easy attack. Due to the additional energy at a higher power, more diffusion of atoms is possible, hindering a strong shadowing. At much higher energies, this can additionally lead to a shift from the T zone to zone 2 with stronger columnar growth [204]. However, in the samples discussed here, microstructural analysis of cross-sections of all sample types does not show a significant change in the microstructural type (figure 3.3). Increasing the pressure can lead to lower energy of the incident sputtered atoms and, thus, more lattice defects [29]. However, other studies describe that a pressure change does not 42 Results Figure 3.3: Microstructure of Mg-3Li thin films sputtered at different pressure and power (for power change: constant pressure of 2.3·10−3mbar, for pressure change: constant power of 100 W). influence the corrosion rate [102] or can lower the corrosion rate by leading to finer granular structures [205]. For samples prepared at the lowest pressure of 1.5·10−3mbar, the SEM images confirm the change in the surface properties, which was already apparent due to the colour change. The sputtering at such pressures is less stable, and smaller particles or droplets can be found on the surface and in the film. However, since no significant change of the corrosion rate to the pressure of 2.3·10−3mbar is measured, this cannot be the main influence for a corrosion rate change. For samples sputtered at the pressure of 4·10−3mbar, thin columnar structures are formed. Even though the surface is not flat, no significant gaps and no rough sub-structure on the grains are formed. Since the Mg-3Li samples for a pressure of 4·10−3mbar and a power of 100 W showed the lowest corrosion rate and low film stress, freestanding thin films were prepared for corrosion measurements, leading to corrosion rates of 0.14±0.02 mm/yr in comparison to 0.27±0.07 mm/yr for the standard sputtering parameters at 50 W described in the chapter 3.1. For both samples, a similar Li fraction (3.5±0.5 wt% vs 3.9±0.1 wt%, measured by AAS) is found. Therefore, this cannot be a main influence on the change in corrosion rate. The main difference shown in cross-sectional images of the microstructure is analysed to be a denser structure and smaller columns, indicated by fewer voids at grain boundaries, and a flatter surface structure with less pronounced grain boundaries (figure 3.4 a). The comparison of exact XRD peak heights for freestanding thin films is difficult since they cannot be fixed completely flat. However, no significant change in orientation can be observed and stronger signals and, thus, more low-indexed planes parallel to the surface can be found for the 100 W sample (figure 3.4 b). Thus, by changing the sputtering parameters, the same corrosion rate as for Mg-1.6Li in freestanding thin films can be achieved for Mg-3Li. As a result, no direct trend, such as a decrease in corrosion rate with increasing power or decreasing pressure, can be identified. Still, the possibility of influencing the corrosion rate is confirmed. Multiple factors, including a preferred orientation, grain size, void formation, surface roughness and low film stress, affect the corrosion rate. Thus, the parameters have to be chosen accordingly for each alloy type. To allow a controlled tuning of the properties, the influence of the parameters on the plasma and, therefore, film growth needs to be further analysed. 43 Results Figure 3.4: Comparison of SEM and XRD analysis of Mg-3Li sputtered with two sets of sputtering parameters (50 W, 2.3·10−3mbar and 100 W, 4·10−3mbar) producing stress-free films. 3.3 Analysis of long-term corrosion in different solutions The corrosion experiments of the studies are carried out mainly in HBSS over shorter periods of time due to the measurement set-up (potentiodynamic polarisation or weight loss in sample holder). To identify further influencing factors and see the corrosion long term, freestanding samples of Mg, Mg-1.6Li and Mg-3Li are placed in HBSS without heating (18-23 ◦C) or pH control (pH range of 7-8.5). The pictures of Mg-1.6Li are shown in figure 3.5. After a colour change due to the formation of corrosion product on the surface, holes occur due to inhomogeneous corrosion. Figure 3.5: Exemplary photos of corrosion of Mg-1.6Li over 182 days in HBSS. As can be seen in figure 3.6, for all three sample types, the corrosion does not take place homogeneously over the whole surface but holes are formed due to defects and pitting. Due to the rolling of mainly Mg-3Li and Mg, the corrosion rate cannot be directly identified but after around 55-60 days, Mg-1.6Li has the largest remaining area of thin film (dark 44 Results film area), indicating a lower corrosion rate as already described in the paper in 3.1. For Figure 3.6: Corroded samples of a) Mg (55 days), b) Mg-1.6Li (62 days), c) Mg-3Li (57 days) in HBSS. all samples, the formation of a white corrosion product partially retaining the shape of the thin film is found. To further identify change of the film during the corrosion, XRD and SEM analysis on the corroded samples were performed. In figure 3.7, the formation of a corrosion product layer can be easily identified by a change in microstructure. This surface layer, however, is not dense and is cracked throughout the whole film. In the area of the white corrosion product, the microstructure is completely changed, only leaving a porous material. Figure 3.7: SEM of Mg-1.6Li samples after 63 days in HBSS. a) Cross section of white, corroded area, b) cross section of dark area, c) surface. To identify the change of the film, XRD analysis was carried out on both the dark (possibly remaining metal alloy) and white (pure corrosion product) area of Mg, Mg-1.6Li and Mg3Li samples. In figure 3.8, both are directly compared to the spectra of uncorroded thin films. In the grey regions, the Mg or MgLi can still be identified for all samples but is much lower or not present for the white regions due to the corrosion. Instead, MgO and carbonate containing compounds (MgCO3or MgxCa(CO3)x+1) can be found on both MgLi alloys with additional influences of salts such as NaCl since the samples could not be completely cleaned due to instability after corrosion. It can be assumed that the white area is not only a salt accumulation of NaCl singe it mimics the shape of the metal film before (see figure 3.6), thus, corrosion products of the thin films are included. The signal in general is lower due to the reduced amount of material. In addition, XRD only detects crystalline material, thus, any amorphous components cannot be determined. A possible inclusion of Ca into the corrosion product hints to the importance of the solution on the corrosion process as discussed in section 2.3.2. HBSS is chosen since it includes the many elements which are present in the human body such as Na, Ca, Mg, K, Cl, S or P and glucose (H1387, Sigma-Aldrich with added sodium bicarbonate). Other studies for Mg corrosion are carried out in simple NaCl salt solutions (e.g. [156, 165]) or other simulating body fluids (e.g. [163, 187]). To see the effect of the salts on the corrosion of MgLi, samples 45 Results (a) (b) (c) Figure 3.8: XRD of a) Mg after 56 days, b) Mg-1.6Li samples after 66 days, c) Mg-3Li after 74 days corrosion in HBSS on white and grey sample area. Additionally, XRD diffractograms of the samples before corrosion are shown. of Mg-1.6Li were placed in HBSS, DMEM (Dulbecco’s modified Eagle’s medium, DMEM, 31966047, Thermo Fisher, with 10 % fetal bovine serum, S0615, Merck), aCSF (artificial cerebrospinal fluid) and distilled water as a comparison without temperature and pH control. In figure 3.9, images of the samples after short immersion and after longer immersion are shown. For samples in both distilled water and DMEM, the sample is nearly completely corroded before 10 days in solution while samples in both aCSF and HBSS are more stable with large areas intact or more homogeneously corroded after over 50 days. An estimation therefore leads to a corrosion rate <0.6 mm/yr in DMEM and water, <0.1 mm/yr in aCSF and <0.05 mm/yr in HBSS. Since the film is, however, not corroding homogeneously, this only gives an estimation of the influence of corrosion and no direct corrosion rates. The composition of aCSF is similar to HBSS, containing mainly salts and glucose including carbonates, phosphates and Na, Ca, Mg, Cl. The DMEM used in this studies overall has a higher concentration of salts than in HBSS and additionally includes amino acids and vitamins. The concentration of Ca and carbonates and phosphate containing components 46 Results is higher, thus, an easier formation of such components on the formation of a passivating layer cannot be the deciding factor. The organic components can in theory also decrease the corrosion rate [135]. One possible explanation is a change in pH. DMEM is found to buffer a change in pH better than HBSS [135], thus, due to the lack of pH control, an increase in pH can be more pronounced in HBSS, leading to lower corrosion rates. Figure 3.9: Mg-1.6Li samples in different solutions after short (1-2 days) and longer corrosion (for fast corroding samples 6 days, for slow corroding samples around 60 days). 47 Results 3.4 Publication: Investigation of in-situ ion release and surface film formation of hcp Mg-Li thin films In addition to the influences of the microstructure on the corrosion rate of the hcp α-MgLi alloys (Mg-1.6Li and Mg-3Li), which are discussed in 3.1, the ion release of the samples can give further insight into changes of the corrosion with a change in Li content. Therefore, both long-term (3 days) and short-term studies (in-situ study during the first 30 min-1 h of immersion) were carried out to analyse the release of both Mg and Li over time. For comparison, thin films with 5.5 wt% of Li were added with a composition very close to the boundary between hcp and mixed phase. The difference in corrosion process previously suspected by the change in corrosion rate to an increased corrosion rate with Li fraction and microgalvanic coupling for Mg-5.5Li were confirmed, and a process of preferred ion release over time was proposed. Further, a Li-rich layer on the surface of the α-phase samples stored in air was found, which is released directly after contact in solution and, thus, increases the Li concentration during the start of the treatment if not pre-treated. Therefore, the formation of Li2CO3on the surface of α-MgLi confirms that the corrosion rate reduction for β-MgLi might not only be influenced by Li2CO3as discussed in section 2.3.5. Own contributions to the following article (concept - 75 %, planning - 80 %, experiments - 75 %, analysis - 90 %, writing - 90 %): •Sample preparation •Set-up and corrosion for preparation of long-term study •Online ICP-MS measurements •Interpretation and discussion of the results •Writing of the manuscript The paper Hanke, L., Kalchgruber, L., Westernstr¨oer, U., Garbe-Sch¨onberg, D., Quandt, E., Valtiner, M. Investigation of in-situ ion release and surface film formation of hcp Mg-Li thin films. Corrosion Science 238, 112361(2024) [206] published by Elsevier is open access and the use is permitted by the Creative Commons CC-BY License (https: //creativecommons.org/licenses/by/4.0/). 48 Investigation of in-situ ion release and surface film formation of hcp Mg-Li thin films Lisa Hanke a , Lukas Kalchgruber b , Ulrike Westernstr¨ oer c , Dieter Garbe-Sch¨ onberg c , Eckhard Quandt a , Markus Valtiner b , * a Inorganic Functional Materials, Institute for Materials Science, Kiel University, Kiel, Germany b Applied Interface Physics, Institute of Applied Physics, Vienna University of Technology, Vienna, Austria c Marine Climate Research, Institute of Geosciences, Kiel University, Kiel, Germany ARTICLE INFO Keywords: A Magnesium A Sputtered films B ICPMS B XPS C Surface films ABSTRACT In this work, the dissolution process of magnetron sputtered Mg-Li thin films was investigated by in-situ flow cell/ICP-MS measurements and ex-situ ICP-MS measurements after longer immersion and additional XPS measurements. High Li concentrations are released due to a Li rich carbonate layer formed in air. The depletion of Li leads to preferred Mg release before preferred Li release occurs due to the higher activity of Li and incorporation of Mg in corrosion products. This data provides a baseline for developing release profiles for medical application, more generally, it unravels details of the corrosion mechanism of lightweight MgLi alloys. 1. Introduction Biodegradable materials such as magnesium (Mg) are of interest for several medical applications to reduce permanent implants’side effects, improve bone healing or release bioactive components [1–4]. A key point of interest is the understanding and control of the degradation rate of such implant materials to verify the lifetime of the implants, avoid adverse effects from fast corroding materials and allow for the adjustment of the material properties to facilitate a better treatment [5–7]. Understanding the degradation process becomes even more crucial if the degradation products, e.g., corrosion products and pH change [4] or ion release [8], are used as therapeutically active species. A material of interest for ion release studies is magnesium-lithium. While the release of lithium (Li) from different materials is already studied [9–12], it was recently proposed to use Mg-Li thin films as a reservoir implant releasing Li for the local treatment for neurological applications [8,13] such as the treatment of bipolar disorder, Alzheimer’s or Parkinson’s disease [14–17]. The therapy with Li is associated with several side effects and has a small therapeutic window [15,18]. Thus, the control of the concentration is of high importance. Mg can be used as the base material for implants not only because it is already widely studied as a biodegradable and biocompatible material and used in clinical trials and medical applications [2,19] but also because it shows possible positive therapeutic effects [19,20]. The treatment with Mg-based materials is already established for particles loaded with additional drugs or using the degradation products, such as hydrogen. Here, both the benefit of a biodegradable carrier and the effect of Mg itself are taken advantage of [21–24]. Mg-Li alloys can be differentiated in α -Mg-Li, α +β-Mg-Li and β-Mg-Li with α as the hexagonal closed packed (hcp) Mg-rich phase and βas the body centred cubic (bcc) Li-rich phase. Generally, the highest corrosion rate is found in the mixed phase α +βdue to the microgalvanic coupling of those phases, while single phases show a lower corrosion rate [25,26]. For hcp α -phase Mg-Li with low Li content, the corrosion process is discussed to be similar to pure Mg with the formation of oxide and hydroxide layers [27,28] and filiform corrosion occurring [26,29]. Li carbonate containing protective layers were also found for Mg-5Li-1Al samples in the hcp phase [30]. For neurological applications, implants need to be produced in small sizes. The properties of such thin films are highly affected by surface effects and can differ from bulk material [31, 32]. As previously shown [8], similar to bulk material, Mg-Li thin films with low Li fraction show a corrosion rate similar to Mg thin films in a medium chosen to simulate physiological conditions, while the corrosion rate increases for higher Li fractions. The ion release of Mg-Li-Zn alloys [33] and Mg-Li-(Al)-(RE) [34] was tested to identify the corrosion rate and influence on cells after longer immersion time for applications as stent material. Zhou et al. found an increase in both Li and Mg release for Mg-8.5 wt% Li in comparison to * Corresponding author. E-mail address: [email protected] (M. Valtiner). Contents lists available at ScienceDirect Corrosion Science journal homepage: www.elsevier.com/locate/corsci https://doi.org/10.1016/j.corsci.2024.112361 Received 13 May 2024; Received in revised form 29 July 2024; Accepted 8 August 2024 Corrosion Science 238 (2024) 112361 Available online 13 August 2024 0010-938X/© 2024 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY license ( http://creativecommons.org/licenses/by/4.0/ ). Results 49 decreases and even though the passivation is not strong, additional compounds such as Li 2 CO 3 or LiOH have been found on the surface of Mg-Li based samples [29,30]. In general, the Li release is expected to increase with the Li content available in the film, which was found for both short-term and long-term measurements. However, the release is influenced by not only the Li content but also the degradation process, in the simplest case, the corrosion rate. In a previous study, the corrosion rate of Mg-1.6Li was determined to be the lowest, while the corrosion rate increased due to the change in orientation, higher activity of Li and galvanic coupling because of the formation of a second phase for Mg-5.5Li [8]. A higher Li content leads to a higher corrosion rate and, thus, to a higher release in both Mg and Li ions. This can be even more pronounced if the increase in Li content also changes the orientation of the material and the microstructure [8]. This is in agreement with the higher overall ion release and, thus, corrosion rate measured herein. Both α -Mg-Li alloys show a preferred Mg release after the first peak of Li release in the first hour, and there is a nearly direct start of Mg release after immersion. Thus, after the dissolution of Li-rich compounds, Li is no longer released preferentially, and a larger amount of Mg is available, possibly due to the depletion after Li reacting on the surface. In the 3-day measurements, both α -phase materials show a similar trend of release with a higher release of Mg at the beginning, as already mentioned for in-situ measurements. However, the results show a preferred dissolution of Li over longer immersion times, possibly due to the lower electrochemical potential [50]. The preferred dissolution of Li is also observed in the ion release studies at constant voltage. While the Mg release is nearly completely suppressed for both materials at cathodic polarisation, Li release still occurs, especially for Mg-3Li. Thus, the release is not a continuous release of both materials present in the phase but preferred of one material even though no two different phases can be found. Since the Li was found to be accumulating at the grain boundaries [53], it is possible that these Li-rich areas are corroding preferentially; however, further studies would be needed to prove the effect on the ion release directly. The corrosion of α -Mg-Li, thus, changes from surface to bulk of the thin films and can be described by three phases (Fig. 8): 1. Li release from Li containing surface film formed in air, 2. Preferred release of Mg from Mg-rich area under the surface, 3. Continuous release from the bulk of the film with preferred Li dissolution. The release of Mg is reduced over time while the Li release is less reduced; thus, in addition to the higher activity of Li, the incorporation of Mg in corrosion products might additionally reduce the Mg release and lead to a higher concentration of Li in comparison to Mg. The anodic polarisation by sweeping to values above passivation and the start of pitting shows an increase of Mg release, thus indicating the presence of a Mg-rich corrosion product. The formation of the corrosion products can then change decrease corrosion over time [54] together with additional components such as phosphates (Ca, P) formed due to the additional salts in the solution. Li-containing components have not been found in the outer layer after corrosion without storage in air, and after 1 h of immersion, only a thin layer of oxides and hydroxides is formed. Thus, the Mg metal underneath is still detectable. Further measurements would be necessary to identify the composition of the corrosion products in detail which is not Fig. 8. Schematics of possible processes of the surface layer formation and degradation of hcp Mg-Li thin films a)-c) in air and d)-f) in solution. a) Formation of a Li 2 CO 3 and MgO containing surface layer in air on as sputtered samples, b) formation of Mg-rich surface layers in air after short storage time and low Li content after corrosion and c) development of Li 2 CO 3 on corroded samples after longer storage time in air for films with higher Li content. d) The start of the contact with solution leads to a dissolution of the Li-rich carbonate layer, followed by e) preferred, strong Mg release in short-term corrosion, before f) the release switches to preferred Li release, possibly due to preferred Li release and formation of corrosion products. L. Hanke et al. Corrosion Science 238 (2024) 112361 8 Results 56 part of this study. For the application of such films, the Li release needs to be determined over time. While the therapeutic range is discussed to be around 0.4–1.2 mM [55], this cannot be directly compared to the concentrations reached in the 3-day studies due to the influence of cells or placement in vivo on the corrosion rate [56], and the actual volume of solution and the flow on the final concentration. However, this study and the preferred release of Li identify that the Li concentration available cannot be directly derived from the degradation rate of the films and will change, especially during the beginning of placement and after long immersion times when Li is depleted. Especially the first Li-rich layer might lead to higher Li concentrations, and removal of such layer by, e.g., previous immersion in solution might be necessary. 5. Conclusions Mg-Li thin films in the hcp phase were studied with respect to the surface chemistry and degradation by a combination of in-situ and exsitu techniques for determining the ion release and additional analysis of the surface composition by XPS. A process of the formation of surface layers in air and during corrosion and preferred ion release is described: •It was found that both sputtered hcp alloys (Mg-1.6Li and Mg-3Li) formed Li carbonate containing compounds on the surface after long-term exposure to air. This layer is also formed if corroded samples are exposed to air. However, the development is mainly visible for Mg-3Li films after 7 days due to the higher Li content and, thus, availability of the element. •During degradation in Hanks’balanced salt solution, this surface layer is dissolved, monitored by a high release in Li and reduction in both carbonate and Li signal in XPS. The layer does, therefore, not contribute to the passivity and influence on the corrosion resistance of the hcp thin films. •While a higher Mg concentration can be found during the first hour of immersion, a higher Li concentration than present in the film is determined over a longer time due to the preferred release of Li and the formation of Mg-rich corrosion products (containing oxides and hydroxides), which form a passivating layer and decrease the corrosion rate. The corrosion process and ion release are time-dependent and change from direct contact in solution to long-term corrosion. The combination of in-situ short-term and studies with longer immersion times, as described in this work, is therefore beneficial to gain a more comprehensive picture of the complex degradation behaviour of the thin films. Understanding the degradation process and reactions in air is also of interest for other applications of MgLi in fields such as lightweight construction. For the medical field specifically, the ion release results also highlight the importance of carefully analysing the ion release when discussing such thin films for application as biodegradable reservoirs for treatments with dissolving elements. For an optimisation of the treatment, it is not sufficient to determine the Li release by degradation rate since a change over time occurs. The change in the ratio of both ions must also be considered if both can influence the therapeutic effect of the other element involved. Furthermore, a treatment or solution dipping before tests in vitro or in vivo might be beneficial to avoid the impact of the high Li release at the start. CRediT authorship contribution statement Ulrike Westernstr¨ oer: Investigation. Dieter Garbe-Sch¨ onberg: Supervision, Resources. Lisa Hanke: Writing –review &editing, Writing –original draft, Investigation, Formal analysis, Conceptualization. Lukas Kalchgruber: Writing –review &editing, Methodology, Investigation. Eckhard Quandt: Writing –review &editing, Supervision, Funding acquisition. Markus Valtiner: Writing –review & editing, Supervision, Resources, Methodology. Declaration of Competing Interest 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. Data Availability Data will be made available on request. Acknowledgements This work was supported by the DFG in the framework of the research training group 2154 –Materials for Brain (project 270394294). We also acknowledge the CzechNanoLab project LM2018110 funded by MEYS CR for the financial support of the XPS measurements at CEITEC Nano Research Infrastructure. Appendix A. Supporting information Supplementary data associated with this article can be found in the online version at doi:10.1016/j.corsci.2024.112361. References [1] J.-W. Lee, H.-S. Han, K.-J. Han, J. Park, H. Jeon, M.-R. Ok, H.-K. Seok, J.-P. Ahn, K. E. 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While chapter 3.2 already discusses the possibility of changing the corrosion rate during the sputtering itself, this only decreases the corrosion rate if still a microstructure with a lower corrosion rate is possible and can be formed. In general, tuning the corrosion rate for Mg-based alloys means decreasing it since the structures need to be present in the body for a longer time so that the implant can serve its job. To further reduce the corrosion rate of Mg and MgLi thin films, an additional treatment by dielectric barrier discharge plasma was tested. The set-up for the treatment in ambient air is simple, cheap and could be easily upscaled. The surface change by reducing roughness and formation of a carbonate-containing layer reduced the corrosion rate by around 50 % without significantly changing the film’s thickness. Own contributions to the following article (concept - 50 %, planning - 70 %, experiments - 40 %, analysis - 70 %, writing - 90 %): •Sample preparation •XRD, SEM/EDX investigation •Partially potentiodynamic polarisation measurements •Interpretation and discussion of the results •Writing of the manuscript The paper Hanke, L., Hartig, T., Weisheit, F., Tjardts, T., Pogoda, T., Faupel, F., Quandt, E. Tailoring of Mg and MgLi thin-film corrosion rates with dielectric barrier discharge plasma treatment. Journal of Vacuum Science and Technology A 41, 053109(2023) [207] is reproduced with permission from the American Vacuum Society (AVS) and licensed under a Creative Commons Attribution (CC BY) license (https://creativecommons. org/licenses/by/4.0/). 59 Tailoring of Mg and MgLi thin-film corrosion rates with dielectric barrier discharge plasma treatment Cite as: J. Vac. Sci. Technol. A 41, 053109 (2023); doi: 10.1116/6.0002783 View Online Export Citation CrossMar k Submitted: 22 April 2023 · Accepted: 28 July 2023 · Published Online: 18 August 2023 Lisa Hanke, 1 Torge Hartig, 2 Felix Weisheit, 1 Tim Tjardts, 2 Tim Pogoda, 2 Franz Faupel, 2 and Eckhard Quandt 1,a) AFFILIATIONS 1 Chair for Inorganic Functional Materials, Institute for Materials Science, Faculty of Engineering, Kiel University, Kiel, Germany 2 Chair for Multicomponent Materials, Institute for Materials Science, Faculty of Engineering, Kiel University, Kiel, Germany Note: This paper is part of the Special Topic Collection including papers from the Pacific Rim Symposium on Surfaces, Coatings and Interfaces (PacSurf 2022). a) Author to whom correspondence should be addressed: [email protected] ABSTRACT Magnesium and magnesium alloys such as magnesium-lithium are of great interest for the application as biodegradable implants. To control the degradation, a tailoring of the corrosion rate is needed. In this study, the effect of a short (5–20 s) dielectric barrier discharge plasma treatment in ambient air on the corrosion rate of magnetron sputtered Mg and MgLi thin films is presented. The treatment with atmospheric plasma of as sputtered samples leads to a decrease of the corrosion rate of 45%−50% in Hanks’balanced salt solution. The higher corrosion resistance is influenced by a change in surface structure and a formation of an MgCO 3 containing film. © 2023 Author(s). All article content, except where otherwise noted, is licensed under a Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/). https://doi.org/10.1116/6.0002783 I. INTRODUCTION Magnesium and magnesium alloys are widely studied biodegradable materials as candidates for medical applications. Applications of interest reach from biodegradable orthopedic implants and stents up to the possible field of therapeutical treatment. 1–5 To ensure the duration of the integrity of the implant required by the application or the therapeutic activity, the degradation rate needs to be adjustable. Since the corrosion rate of Mg and Mg alloys is often too high for the applications, 6–9 several strategies to reduce the rate have been developed: These include alloying of the bulk material, 10–12 influencing the microstructure, 12–14 surface treatments, and coatings. 15–17 The surface treatments range from deposition of inorganic 18–20 and organic 21–23 deposition coatings over chemical conversion by immersion in solution 24,25 to ion implantation 26,27 or plasma electrolytic oxidation. 28–30 In this study, the focus of the application is on miniaturized implants fabricated by MEMS (micro-electro-mechanical systems) technology which can serve as, e.g., biodegradable materials for brain implants, either as structural implants, substrates for additional materials or reservoirs for therapeutically active ions. As an example of degradable films that release ions which can act as treatments, MgLi thin films are studied since Li is used in treatments for mood disorders such as bipolar disorder. 31,32 The possibility to include additional ions in a coating to reduce the corrosion rate specifically for the application as reservoirs is limited since they could influence the therapeutic effect if they are released during the degradation. While small Mg structures with a thickness of 10– 100 μm can be prepared by thin-film deposition techniques such as sputtering, 33–35 the addition of a thick coating would additionally counteract the effort to reduce the implant size. Thus, a treatment that leads to a thin layer only consisting of the alloy elements itself, possibly also in the form of, e.g., oxides and carbonates, is beneficial. Possible techniques to form those layers include, e.g., chemical treatment in solution to form MgO or Mg(OH) 2,36,37 ion implantation of nitrogen, 38 or plasma treatments. 39–42 Kocijan et al. showed that plasma treatment in O 2 and H 2 leads to pin–hole free oxide layers and, thus, lowering the corrosion rate. 39 The formation of an oxide film is also studied for plasma treatment with Ar/O 2 by Tiyyagura 40 while Nakazawa et al. examined the implementation of nitrogen and oxide after the treatment of a Mg surface with an atmospheric plasma jet. 41 Additionally, carbon contaminations on ARTICLE pubs.aip.org/avs/jva J. Vac. Sci. Technol. A 41(5) Sep/Oct 2023; doi: 10.1116/6.0002783 41, 053109-1 ©Author(s)2023 15 April 2024 10:22:05 Results 60 the samples are removed by plasma processes. 39 Li et al. show that a dielectric barrier discharge (DBD) plasma treatment can reduce or even prevent the corrosion of MgLi alloys with a Li content of 14.2% (m/m) significantly by the formation of Li 2 CO 3 and oxidecontaining layers. 42 Atmospheric pressure plasmas are highly interesting to influence surface properties in an economic way. In comparison to other techniques, DBDs can easily be scaled up to surface treatments in industrial scales. 43,44 By DBD treatment oxygen groups or compounds can form on the surface by the interaction with elevated oxygen species from the plasma. 45 In the literature, DBDs are also discussed to clean the surface of Mg from remaining contaminations 46 or to lead to the reaction of carboncontaining components to carbonates additional to the formation of an oxide layer. 47 Since for thin films no prior grinding and polishing can be easily performed, the process can, thus, provide advantages additionally to the reduction of the corrosion rate due to the formation of protective layers including oxides and carbonates. This study aims to evaluate the effect of a DBD plasma treatment to tailor the corrosion resistance of thin films of Mg and MgLi. Since the MgLi samples in this study have a Li content of only 1.6% (m/m), it does not lead to additional Li rich phases which could lead to high Li concentrated protective surface films as seen in previous studies, 42 but might influence the activity and microstructure. Additionally, the nature of thin films with μm thickness does not allow the implementation of thick protective layers. Thus, the effect of the treatment forming layers below 1 μm thickness on the corrosion rate is studied in Hanks’balanced salt solution to simulate the environment in medical applications. Additionally, the layer formed during the treatment is analyzed regarding the structure and chemical composition. By including Mg and MgLi [1.6%(m/m)], the influence of alloying and varying microstructure can provide a detailed understating of the overall process. II. EXPERIMENT A. Sample preparation Magnesium and magnesium-lithium films were prepared by magnetron sputtering (Von Ardenne CS730S), using targets of pure Mg and MgLi [2.5%(m/m) Li] from FHR. As a substrate, 4 in. silicon wafers were cut into 15 × 15 mm 2 samples and coated with aluminium (Al) and aluminium nitride (AlN) for comparability to freestanding thin films. The sputtering was carried out with a base pressure of < 5 × 10 −7 mbar and an Ar pressure of 2.3 −2.6 × 10 −3 mbar with a gas flow of 25 SCCM Ar. A final sample thickness of 10–20 μm of Mg and MgLi thin films with a Li mass fraction of 1.6% (m/m) was reached. For one measurement set, all samples were prepared in the same sputtering process, thus, with the same thickness for untreated and treated samples. Additionally, freestanding thin films were prepared with the same sputtering parameters, following the process described by Haffner et al. 34 After structuring on the wafer by UV-lithography and etching, a sacrificial AlN layer was added before the final layer of Mg or MgLi was deposited. For the etching of the sacrificial layer, the samples were afterwards immersed in 20 wt. % KOH solution. B. DBD treatment The thin-film samples were treated in the gap of a self-built symmetric volume dielectric barrier discharge setup. All treatments were performed in stagnant ambient air. The freestanding thin films were placed on a Si chip (15 × 15 mm 2 ) and fixed on the edges to ensure a flatter surface during the treatment. A laboratory power supply (SM7020-D, Delta Elektronika) and a function generator (DDS function generator 4025, Peak Tech) were connected to the high frequency high voltage power supply (Minipuls 4, GBS Elektronik). The resulting sinusoidal voltage signal of the high frequency high voltage power supply with an output of 1:2000 was monitored via an oscilloscope (UTD2025CL, UNIT). The treatment time of the plasma was controlled by an inhibiting signal by a microcontroller board (Arduino nano every) connected to the high frequency high voltage power supply. The DBD setup connected to the power supply can be seen in Fig. 1. The parameters of the plasma power supply were chosen to allow the formation of filaments all over the sample surface during the complete treatment time, giving a state of saturation (Table I). An overlay of the plasma filaments at all times of the treatment can be seen in Fig. 1(b). The parameters were adjusted if the setup had to be adapted in between measurements of different sample sets due to functioning reasons to ensure a homogeneous surface treatment. For each sample and measurement type, treatments with different treatment times of 5–25 s were carried out with the same parameters to ensure comparability. C. Sample analysis The surface of the thin films was imaged and analysed using a Zeiss Ultra 55 Plus scanning electron microscope (SEM) and an Oxford Instruments ULTIM MAX 65 energy-dispersive x-ray spectroscope (EDX). An accelerating voltage of 3 kV for imaging and 10 kV for EDX measurements was used. Additionally, the structure of the thin films was analysed by x-ray diffraction (Smart Lab 9 kW, Rigaku) with a parallel beam and monochromatic Cu Kα radiation on a θ/2 θ-scan with a range of 20°−90° with a speed of 5−10°/min and a step size of 0.03°. The chemical composition of the sample surface was characterized by x-ray photoelectron spectroscopy (XPS). For this purpose, an XPS UHV system from Omicron Electron Spectroscopy Ltd. with a 240 W Al anode was used. Survey scans to screen for the elements present at the surface were conducted at a pass energy of 100 eV, a step-size of 0.5 eV, and averaged over three sweeps. High-resolution scans of the characteristic core hole-level spectra used for the chemical analysis were conducted at a pass energy of 30 eV, 15 sweeps, and a step-size of 0.05 eV. For data analysis, the software CASA XPS (Version 2.3.23PR1.0) was utilized and charge correction was done by shifting the C 1 s main peak to 284.8 eV and adjusting all the corresponding spectra accordingly. The corrosion rate was determined by potentiodynamic polarisation measurements in a 155 mmol Hanks’balanced salt solution (H1387, Sigma-Aldrich with added sodium bicarbonate) at a pHof 7.4 ± 0.2 (CO 2 regulation) and a temperature of 37 ± 1 °C. A VersaSTAT 3–300 potentiostat (AMETEKSI) and a three-electrode setup with an Ag/AgCl reference electrode, a Pt mesh counter electrode, and the sample included into a sample holder with an exposed area of 0.916 cm 2 ARTICLE pubs.aip.org/avs/jva J. Vac. Sci. Technol. A 41(5) Sep/Oct 2023; doi: 10.1116/6.0002783 41, 053109-2 ©Author(s)2023 15 April 2024 10:22:05 Results 61 were used. After 5 min of measuring the open circuit potential (E OCV ), a linear voltage sweep from −0.3 V vs E OCV to +0.3 V vs E OCV was performed. Further information can be found in Ref. 48. Additionally, the same measurement was carried out after 1h of immersion time. III. RESULTS AND DISCUSSION A. Thin-film characterization Representative cross sections and surface images of the sputtered thin films of Mg and MgLi are shown in Figs. 2(a)–2(d). FIG. 1. (a) Scheme of the self-built dielectric barrier discharge system. The electrodes are connected to the high frequency high voltage power supply. The top electrode is a transparent FTO-coating on glass (Sigma Aldrich, 100 × 100 × 2.2 mm 3 ,13Ω/sq) acting as the top dielectric, making the observation of the plasma treatment from the top possible. The bottom electrode is made from aluminum with the bottom dielectric being made from 6 mm thick Al 2 O 3 . (b) Photo of the DBD treatment of an Mg thin film seen through the top transparent electrode. For the inset, photos (video frames from a video with 30 fps) over a treatment time of 15 s were combined to show all filaments during a treatment, ensuring area saturation. ARTICLE pubs.aip.org/avs/jva J. Vac. Sci. Technol. A 41(5) Sep/Oct 2023; doi: 10.1116/6.0002783 41, 053109-3 ©Author(s)2023 15 April 2024 10:22:05 Results 62 While for both materials a columnar growth is visible, the surface exhibits different structures for both sample types with a more structured surface for Mg and larger grainlike areas for MgLi. EDX analysis of the surface of the thin film identify an oxide and carbonate signal for both sample types with no significant difference between Mg and MgLi [2%–3%(m/m) C and 0.5%–1%(m/m) O]. No further quantification of oxide or carbonate components is carried out due to the possibility of the influence of contaminations on the exact intensity. In Fig. 2(e), XRD diffractograms show signals of the thin film itself and the substrate. Mg and MgLi have a hexagonal closed packed (hcp) structure with a strong preferred orientation for pure Mg. Samples of Mg and MgLi on the substrate and freestanding films of MgLi were treated with an atmospheric pressure dielectric barrier discharge plasma in air with a saturation of filamentary discharges, as seen in Fig. 1(b). B. Corrosion measurements Exemplary potentiodynamic polarization curves for an untreated and 15 s treated sample after 5 min of immersion are shown for Mg and MgLi in Fig. 3(a). The corrosion current densities and corrosion rates can be determined via Tafel extrapolation. 48,49 The cathodic branch was used for the estimation of the corrosion rate for the following studies due to the larger linear area. The current density of the anodic branch is influenced by additional hydrogen evolution, film formation, and passivation regions, 50–52 thus, a corrosion rate determined on the anodic branch may differ from the rate determined from the cathodic branch. MgLi samples were treated for 5, 10, 15, and 20 s. To exclude the effect of Li on the influence of the plasma treatment, Mg samples were treated for 5 and 15 s for comparison. The corrosion rates for the measurements are shown in Fig. 3(b). A significant decrease in the corrosion rate for both material type is apparent, lowering the corrosion rate (CR) during a treatment of 15 s from CR untreated = 2.19 mm/yr to CR treated = 0.45 mm/yr for MgLi and from CR untreated = 1.69 mm/yr to CR treated = 0.38 mm/yr for Mg. Since the effect for both sample types is similar and Li or the different surface structure does not seem to influence the effect of the treatment significantly, further studies were only carried out with MgLi samples. MgLi thin films with a lower corrosion rate were measured with the corresponding corrosion rates given in Fig. 3(c). The main decrease in corrosion rate already occurs after a treatment time of 5 s with a reduction of the corrosion rate of about 46% (0.23 ± 0.03 mm/yr to 0.13 ± 0.01 mm/yr). Since the corrosion resistance differs only slightly for longer treatments, the times were set to 5 and 15 s. Since the potentiodynamic polarization measurements show the degradation of samples only for the short term, TABLE I. The parameters of DBD treatment. Listed are the distance dbetween both electrodes, the frequency f, and the discharge voltage U discharge. d (mm) f(kHz @5Vpp) U discharge (kVpp) Mg, MgLi [Fig. 3(b)] 3.9 20 24 MgLi [Fig. 3(c)] 1.9 17 25.3 MgLi [Fig. 3(c),1h immersion] 1.9 19 22.8 FIG. 2. Thin films as sputtered on the substrate (Si wafer with the added Al + AlN layer) (a) and (b) SEM cross section and image of the surface of MgLi, (c) and (d) SEM cross section and image of the surface of Mg, (e) XRD diffractogram, signals of Mg, Al, AlN, and Si are marked. Peaks without symbol are kβ, WLa, and edge effect signals of the Si substrate. ARTICLE pubs.aip.org/avs/jva J. Vac. Sci. Technol. A 41(5) Sep/Oct 2023; doi: 10.1116/6.0002783 41, 053109-4 ©Author(s)2023 15 April 2024 10:22:05 Results 63 additional measurements are necessary to confirm the improvement of corrosion resistance over longer time periods. Weight-loss measurements over longer terms are difficult due to the low sample weight of thin films and the cleaning resulting in the removal of corrosion products and possible products formed during the treatment, thus, not allowing the final identification of the corroded mass. Therefore, additional potentiodynamic polarization measurements were carried out after 1h of immersion to allow the prior formation of corrosion products on the surface of the film which can protect the film from further corrosion. 53,54 As shown in Fig. 3(c), the corrosion rate still decreases after the DBD treatment significantly to approximately half of the corrosion rate. Thus, the treatment not only leads to a passivation decreasing the first corrosion before a protective corrosion layer is formed but also leads to a stronger protective layer throughout the degradation. No significant change in the corrosion potential E corr can be identified for MgLi thin films after the treatment for 15 s in comparison to the untreated samples. While an E corr of −1.85 ± 0.08 V is measured for the untreated samples directly after immersion, E corr for samples with a 15 s treatment measured is −1.85 ± 0.01 V, thus, the potential varies for untreated samples, while it is more stable for the treated samples. After 1h of immersion, the potential is slightly increased to −1.78 ± 0.02 V for untreated samples and −1.79 ± 0.03 V for samples after 15 s of treatment. The increase after longer immersion time can be assigned to the lower activity of the material, possibly due to the depletion of Li on the surface 55 and the formation of other products during corrosion. 53,54 Thus, the DBD treatment itself does not influence the potential of the material. For possible applications of Mg or Mg-based alloys as implants for small-size applications and not coatings on other materials, the thin films need to have a thickness in the μm range without an additional substrate. The process to produce the thin films used in this study is described by Haffner et al. 34 Since the sacrificial layer of Al and AlN is dissolved in KOH, the films are also exposed to KOH for the duration of the lift-off, resulting in a changed surface. A treatment with KOH is reported to lower the corrosion rate even without additional treatment if the sample is anodized due to the formation of more stable MgO and Mg (OH) 2;36 however, this effect is not observed here, possibly due to the insufficient thickness and density of the layer formed during the simple immersion. The effect of the plasma treatment is lower on the freestanding samples; however, a decrease in the corrosion rate is still visible [Fig. 3(c)]. Additional to the change of surface structure, the samples on substrate also had a flatter surface than freestanding thin films which were only attached to Si chips, thus, ensuring a more homogeneous treatment. Thus, an optimization of the treatment and sample fixation for freestanding thin films could improve the degradation rate decrease further. The corrosion studies prove that only very short DBD plasma treatments of Mg and MgLi are required to significantly influence FIG. 3. Potentiodynamic polarization measurements in Hanks’balanced salt solution (pH= 7.4 ± 0.2 and T = 37 ± 1 °C). (a) Exemplary Tafel plots for Mg and Mg-1.6Li as sputtered on the substrate and after 15 s DBD plasma treatment. (b) Corrosion rate (CR) of as-sputtered Mg and Mg-1.6Li on the substrate. (c) Corrosion rates (CR) for Mg-1.6Li thin films as sputtered, after 1 h immersion and freestanding thin films (without preimmersion). FIG. 4. SEM image of MgLi thin films, (a) cross section (inset shows the surface film with higher magnification) and (b) surface image of the surface film formed after 15 s of DBD treatment. ARTICLE pubs.aip.org/avs/jva J. Vac. Sci. Technol. A 41(5) Sep/Oct 2023; doi: 10.1116/6.0002783 41, 053109-5 ©Author(s)2023 15 April 2024 10:22:05 Results 64 FIG. 5. XPS spectra of untreated and treated MgLi thin films (a) full spectrum, (b) Mg 2p and Li 1 s, (c) C 1 s (positions of carbonates, C–C and carboxylates are marked, see the supplementary material at for spectra with a full description of C compounds, fitted according to Fotea et al. (Ref. 62, Fig. S1) and (d) O 1 s regions. See the supplementary material Tables SII and SIII for the corresponding peak positions and full-width-at-half-maximum. ARTICLE pubs.aip.org/avs/jva J. Vac. Sci. Technol. A 41(5) Sep/Oct 2023; doi: 10.1116/6.0002783 41, 053109-6 ©Author(s)2023 15 April 2024 10:22:05 Results 65 Appendix Appendix A.1 Supplement Structural characterisation and degradation of Mg-Li thin films for biodegradable implants lxxii Structural characterisation and degradation of Mg-Li thin films for biodegradable implants - Supplementary information Lisa Hanke1, Lea K. Jessen1, Felix Weisheit1, Krathika Bhat2, Ulrike Westernströer3, Dieter Garbe-Schönberg3, Regine Willumeit-Römer2, Eckhard Quandt1* 1Inorganic Functional Materials, Institute for Materials Science, Faculty of Engineering, Kiel University, Kiel, Germany 2Institute of Metallic Biomaterials, Helmholtz Centre hereon, Geesthacht, Germany 3Marine Climate Research, Institute of Geosciences, Faculty of Mathematics and Natural Sciences, Kiel University, Kiel, Germany Mg-Li alloy Li (%(m/m)) Fe (%(m/m)) Target Film, ICP-MS Film, AAS Film, ICP-MS Mg-1.6Li 2.5 1.60±0.06 1.55±0.03 0.0029±0.0008 Mg-3Li 5 3.07±0.12 3.13±0.08 0.0027±0.0003 Mg-5.5Li 9 5.15±0.69 6.12±0.32 0.0028±0.0004 Mg-9.5Li 14 9.31±0.84 9.88±0.07 0.0033±0.0006 Table 1 Mass fractions of Li in targets (nominal) and prepared freestanding thin films analysed by ICP-MS and AAS. Additionally, the Fe contamination measured by ICP-MS is given. The measurement uncertainties given are statistical deviations from minimum 3 samples. Figure 1 Reciprocal space maps measured with 2Ddetector and sample tilting to χ=0°,15°,30°,45° for a) Mg-1.6Li b) Mg-3Li c) Mg-5.5Li d) Mg-9.5Li and zoomed in area around the (002) and (101) hcp and (110) bcc peak for e) Mg-5.5Li and f) Mg-9.5Li. Appendix lxxiii Mg-1.6Li Mg-3Li Mg-5.5Li Mg-9.5Li a [Å] 3.208±0.002 3.202±0.002 3.196±0.004 3.200±0.002 c [Å] 5.197±0.005 5.175±0.007 5.156±0.010 5.161±0.008 Table 2 Average lattice constants a und c and standard deviations calculated from XRD results for Mg-Li alloys (Li: 1.6 %(m/m), 3 %(m/m), 5.5 %(m/m), 9.5 %(m/m)). Mg-1.6Li Mg-3Li Mg-5.5Li Mg-9.5Li hcp 100 32.16±0.03 32.25±0.03 32.35±0.08 32.27±0.06 002 34.48±0.03 34.67±0.06 34.78±0.07 34.75±0.07 101 36.63±0.03 36.75±0.03 36.87±0.08 36.78±0.05 102 47.90±0.02 48.09±0.03 48.28±0.10 48.21±0.04 110 57.44±0.01 57.55±0.03 57.69±0.09 57.62±0.05 103 63.26±0.02 63.51±0.04 63.79±0.09 63.75±0.05 200 67.40±0.02 67.56±0.03 67.71±0.07 67.57±0.01 112 68.79±0.02 69.00±0.02 69.16±0.09 69.11±0.06 201 70.11±0.02 70.28±0.04 70.42±0.10 70.33±0.03 004 72.85±0.07 73.21±0.09 73.45 73.48±0.01 202 78.24±0.05 78.33±0.01 104 81.90±0.04 82.23±0.06 82.64±0.12 82.65±0.10 bcc 110 36.10±0.07 36.03±0.06 200 51.66±0.10 51.48 211 64.89±0.06 64.84±0.07 Li2CO3 110 21.39±0.09 21.29±0.06 200 23.38±0.06 23.32±0.07 111 23.49 111 29.48±0.04 29.38±0.03 202 30.67±0.09 30.52±0.06 002 31.77±0.06 31.65±0.05 112 34.11±0.01 34.02±0.06 020 36.13±0.03 36.06±0.03 021 39.65±0.01 39.64±0.10 310 39.93±0.02 39.91 221 42.60 112 42.60 220 43.55±0.02 43.45 130 56.77±0.04 331 65.41 Table 3 Average 2θ angles (°) from XRD diffractograms for Mg-1.6Li, Mg-3Li, Mg-5.5Li and Mg-9.5Li with corresponding miller indices for hcp and bcc Mg-Li and Li2CO3 with. A minimum of three samples was measured and standard deviations are given if the peak was present in multiple diffractograms. Appendix lxxiv Figure 2 XRD diffractograms for Mg-1.6Li thin films with orientations of (002) and (110). The main orientations are indicated. Figure 3 EDX line scans of cross sections of Mg-1.6Li and Mg-9.5Li thin films after corrosion in HBSS over 3-5 days. The inset shows the area of the corrosion layer. MgLi alloy Power (W) Pressure (10-3 mbar) Sputtering rate (nm/s) Mg-1.6Li 50 2.3 1.45 Mg-3Li 50 3.3 1.75 Mg-5.5Li 50 2.5 1.17 Mg-9.5Li 50 2.3 0.88 Table 4 Sputter parameters for Mg-Li sputtering at standard parameters for stress-free freestanding thin films and for Mg-3Li at different pressure and power. Appendix lxxv Appendix A.2 Supplement Investigation of in-situ ion release and surface film formation of hcp Mg-Li thin films lxxvi Supplementary - Investigation of in-situ ion release and surface film formation of hcp Mg-Li thin films Lisa Hanke1, Lukas Kalchgruber2, Ulrike Westernströer3, Dieter Garbe-Schönberg3, Eckhard Quandt1, Markus Valtiner2 1Inorganic Functional Materials, Institute for Materials Science, Kiel University, Kiel, Germany 2Applied Interface Physics, Institute of Applied Physics, Vienna University of Technology, Vienna, Austria 3Marine Climate Research, Institute of Geosciences, Kiel University, Kiel, Germany 1. Calculation of dissolved mass during 3-day corrosion, measured by ICP-MS Time cMg [µg/L] cLi [µg/L] Blank 17910 <1 Mg-1.6Li, sample 1 1 h 19139 12 4 h 19539 35 1 day 24593 142 3 days 36108 433 Mg-1.6Li, sample 2 1 h 22117 16 4 h 19786 43 1 day 27067 226 3 days 43626 661 Mg-3Li, sample 1 1 h 18132 32 4 h 20292 67 1 day 25677 360 3 days 29598 812 Mg-3Li, sample 2 1 h 20210 47 4 h 19141 98 1 day 27323 379 3 days 39602 888 Mg-5.5Li, sample 1 1 h 17134 111 4 h 20852 213 1 day 21810 847 3 days 37944 2120 Mg-5.5Li, sample 2 1 h 19140 98 4 h 20840 224 1 day 26640 1598 3 days 41950 4714 Table 1. Measured Mg and Li concentrations by ICP-MS in a blank HBSS solution (155 mmol, Hanks´ balanced salts H1387, Sigma-Aldrich with added sodium bicarbonate (0.35 g/L)), and after 1 h, 4 h, 1 day and 3 days of corrosion of Mg-1.6Li, Mg-3Li and Mg-5.5Li thin films in the solution at 37±1 °C and pH 7.4±0.2. The mass of the element (Mg or Li) in the solution at a certain measurement step mi can be directly determined from the concentration ci determined (listed in table 1) and the volume adjusted by the Appendix lxxvii evaporation loss Vi. However, to calculate the overall mass released into the solution mg,i, the amount of extracted solution (in the experiments of this study 15 ml) needs to be taken into account. Thus, mg,i can be described by 𝑚𝑔,𝑖 = 𝑚𝑔,𝑖−1 + 𝑚𝑖− 𝑚𝑙,𝑖−1 With ml,i-1 as the remaining mass of the element in the overall solution after sample extraction for the step before. c0 is the concentration of element in the blank solution which is added in the same amount as solution is extracted. ml,i =mi Vi ⋅(Vi−15ml)+ c0⋅15mL 2. Additional results Figure 1. Current I measured by the potentiostat in the flow cell at constant voltages of -0.5 V vs EOCV, 0.2 V vs EOCV and -0.9 V vs Ag/AgCl during 15 min after 10 min at EOCV for two samples per alloy (Mg1.6Li and Mg-3Li). iMg,CV/iLi,CV (EOCV) iMg,CV/iLi,CV (-0.9 V vs Ag/AgCl) iMg,CV/iLi,CV (+0.2 V vs EOCV) Mg-1.6Li 44.3±2.0 64.4±5.2 55.3±1.5 Mg-3Li 20.7±2.6 211.1±45.5 29.1±3.7 Table 2. Ratio of averaged Mg and Li current density during 15 min at a constant potential of EOCV, anodic polarization of -0.9 V vs Ag/AgCl or +0.2 V vs EOCV after 10 min at EOCV for Mg-1.6Li and Mg-3Li thin films. Appendix lxxviii Figure 2. Dissolution rate for two samples per alloy (Mg-1.6Li and Mg-3Li) over LSV from EOCV until - 0.5 V vs Ag/AgCl, measured in situ by ICP-MS. The linear sweep is performed after 10 min OCV, 15 min conditioning and further 5 min OCV. Figure 3. XPS spectra for a) Ca 2p and b) P 2p for Mg-1.6Li and Mg-3Li thin films after 1 h corrosion and storage in ethanol. Appendix lxxix Appendix A.3 Supplement Tailoring of Mg and MgLi thin-film corrosion rate with dielectric barrier discharge plasma treatment lxxx Tailoring of Mg and MgLi thin-film corrosion rate with dielectric barrier discharge plasma treatment – Supplementary Information Lisa Hanke 1, Torge Hartig 2, Felix Weisheit 1, Tim Tjardts 2, Tim Pogoda 2, Franz Faupel 2 and Eckhard Quandt 1 1 Chair for Inorganic Functional Materials, Institute for Materials Science, Faculty of Engineering, Kiel University, Kiel, Germany 2 Chair for Multicomponent Materials, Institute for Materials Science, Faculty of Engineering, Kiel University, Kiel, Germany TABLE SI. Atomic concentrations of the surface of untreated MgLi thin films and thin films after 15 s of DBD determined by XPS. Untreated Sample Treated Sample Element Peak used for Analysis Relative Amount (at%) Element Peak used for Analysis Relative Amount (at%) C C 1s 31.16 C C 1s 13.00 O O 1s 43.94 O O 1s 61.74 Mg Mg 2p 0.45 Mg Mg 2p 19.59 Li Li 1s 24.45 Li Li 1s 4.12 Appendix lxxxi List of Figures List of Figures 2.1 Slip systems in a hexagonal unit cell of Mg. . . . . . . . . . . . . . . . . . 5 2.2 Phase diagram of Mg-Li. Reproduced and adapted with permission from SpringerNature[76]. .............................. 8 2.3 Structure-zone model developed by Thornton showing the microstructure of sputtered films dependent on the temperature (substrate temperature Tsand melting temperature Tm) and the inert gas pressure. Reproduced with permission from Elsevier [96]. . . . . . . . . . . . . . . . . . . . . . . 10 2.4 Fabrication steps for freestanding Mg alloy thin films using UV-lithography, magnetron sputtering and sacrificial layers as performed for this work, based on the process of Haffner et al. [26]. . . . . . . . . . . . . . . . . . . 12 2.5 Corrosion rate dependent on the pH in Hank’s balanced salt solution (buffered) for three exemplary Mg materials (high purity Mg HP Mg, AZ91, ZE41 Zn+REE). Reproduced and adapted with permission from Elsevier [111]. ...................................... 14 2.6 SEM images of Mg-3(wt%)Li thin films before and after treatment with chromicacid. .................................. 21 2.7 Schematic set-up for potentiodynamic polarisation measurements with a three-electrode set-up including the sample in the sample holder as the working electrode (WE), a reference Ag/AgCl electrode (RE) and a Pt mesh counter electrode (CE). . . . . . . . . . . . . . . . . . . . . . . . . . 22 2.8 Schematic of a potentiodynamic polarisation measurement, depicted as a Tafelplot..................................... 23 2.9 Schematic of flow cell as described in [196, 197] and inline ICP-MS measurement. a) Flow cell as used for the measurement of Mg alloy samples, including a Pt counter electrode (CE), reference electrode (RE) and the sample as a working electrode (WE), connected by copper tape. b) Cross section of flow cell. c) Overall set-up with solution flow through the flow celltotheICP-MS................................ 25 3.1 Comparison of corrosion rates for Mg-3Li thin films on substrate sputtered with different sputtering parameters (for power change: a) constant pressure of 2.3·10−3mbar, for pressure change: b) constant power of 100 W). . 41 lxxxviii List of Figures 3.2 Comparison of freestanding Mg-3Li sputtered at different pressure and power (for power change: constant pressure of 2.3·10−3mbar, for pressure change: constant power of 100 W). The film stress can be identified by the rolling of the films. . . . . . . . . . . . . . . . . . . . . . . . . . . . 42 3.3 Microstructure of Mg-3Li thin films sputtered at different pressure and power (for power change: constant pressure of 2.3·10−3mbar, for pressure change: constant power of 100 W). . . . . . . . . . . . . . . . . . . . . . . 43 3.4 Comparison of SEM and XRD analysis of Mg-3Li sputtered with two sets of sputtering parameters (50 W, 2.3·10−3mbar and 100 W, 4·10−3mbar) producing stress-free films. . . . . . . . . . . . . . . . . . . . . . . . . . . . 44 3.5 Exemplary photos of corrosion of Mg-1.6Li over 182 days in HBSS. . . . . 44 3.6 Corroded samples of a) Mg (55 days), b) Mg-1.6Li (62 days), c) Mg-3Li (57days)inHBSS................................ 45 3.7 SEM of Mg-1.6Li samples after 63 days in HBSS. a) Cross section of white, corroded area, b) cross section of dark area, c) surface. . . . . . . . . . . . 45 3.8 XRD of a) Mg after 56 days, b) Mg-1.6Li samples after 66 days, c) Mg-3Li after 74 days corrosion in HBSS on white and grey sample area. Additionally, XRD diffractograms of the samples before corrosion are shown. . . . . 46 3.9 Mg-1.6Li samples in different solutions after short (1-2 days) and longer corrosion (for fast corroding samples 6 days, for slow corroding samples around60days). ................................ 47 3.10 Comparison of Mg-1.6Li and Mg-6Ag-1.6Li. a) Diffractograms of thin films. In red, the positions of a MgAgLi phase (cubic, based on [208]) are marked. b) Exemplary stress-strain curves of Mg-1.6Li, Mg-3Li and Mg-6Ag-1.6Li. . 69 lxxxix List of Tables List of Tables 2.1 Pilling-Bedworth ratios (PBR) and solubility product constants (Ksp at 25 ◦C of possible corrosion products of MgLi in salt solution. . . . . . . . . 17 3.1 Set of sputtering parameters for magnetron sputtering of Mg-3Li to analyse the influence of power and pressure change on the structure and corrosion rateofthethinfilms............................... 41 xc Bibliography Bibliography 1. Hammond, C. R. & Lide, D. R. 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Therefore, I would like to thank: My supervisor Prof. Dr. Eckhard Quandt for being able to work in his group on this interesting, interdisciplinary topic and the support and helpful suggestions during the last four years. Prof. Dr. Regine Willumeit-R¨omer for her support during my PhD regarding magnesium and corrosion specific questions as my second supervisor. My colleagues of the group inorganic functional materials for helping me to learn the necessary machines and for the general support both scientific and personal, especially Lea Jessen for teaching me everything about Mg thin films and measurement techniques available, but also helping me to generally feeling welcomed in the group when I started. Thanks to Hanna Lewitz, Lars Thorm¨alen, Dr. Justin Jetter, Felix Weisheit and Duygu Dengiz for being not only a great help by practical support, answering questions and discussing problems but also by moral support. A special thanks goes to Lars for keeping the machines alive so that experimental work was possible. My colleagues from the research training group ”Materials for Brain”, especially Krathika Bhat, for continuous discussions over the time of the PhD and a lot of help to understand the requirements from a biological point of view but also the willingness to learn the material science part so that actual exchange was possible. Prof. Dr. Markus Valtiner for not only hosting me during my research at the TU Vienna, but also being available for any questions and discussions during that stay and motivating me to develop a deeper understanding of my scientific problems. I would like to extend my gratitude to the group members in Vienna, Lukas Kalchgruber, Dr. Laura Mears and Matteo Olgiati for helping me learn the machines, answering all my questions and helping me with further measurements and everyone else for creating a pleasant working atmosphere. Ulrike Westernstr¨oer and by extension Dr. Dieter-Garbe Sch¨onberg for agreeing to measuring my samples by ICP-MS without which many parts of my studies would not have been possible or relevant. Last but not least I would like to thank my family, especially my parents and my sister, and my friends for continuously supporting and motivating me. cvi