1 Introduction

Recently, Magnesium has become a famous biomaterial due to its many positive characteristics. Magnesium is a biocompatible element in the body and its presence is a vital necessity for many reactions and metabolisms such as nucleic acid synthesis [1]; also it can be excreted through natural metabolic processes [2]. Magnesium is an appropriate material for medical applications, such as bone implants and fixations, where permanent existence is not desired. These devices support the bone during the healing period and disappear gradually [3, 4]. In addition to these advantages of Mg, it is light but strong and has similar properties to bone such as density and elastic modulus which causes no stress shielding effect when it is in touch with bone [5, 6]. On the contrary, Mg has low corrosion resistance in the body. Fortunately, scientists have figured out several methods to boost the corrosion properties of magnesium alloys. Purification, alloying, coating, and grain refining are the most popular ones [7].

Iron, nickel, and copper, are the most common impurities whose standard reduction potentials are greater than Mg. Galvanic corrosion would occur if they form almost pure metallic inclusions in the Mg matrix. The Mg matrix acts as an anode and corrodes [4,5,6,7,8]. Thus eliminating these impurities lower the corrosion rate undoubtedly which is called purification.

Alloying magnesium with beneficial elements enhances the corrosion properties, omits the effect of impurities, and also refines the grain. Many alloying elements can be added to magnesium because of its hexagonal-closed-packed structure [9], but the elements should be non-toxic for medical applications. For example, aluminum is the most useful alloying element which increases the mechanical properties of Mg alloys but it causes Alzheimer’s and hepatotoxicity [10, 11]. Mg–Ca, Mg–Zn, and Mg–Zn–Ca are typical instances in which nontoxic elements are used for alloying. It is declared that ternary alloys such as Mg–Zn–Ca have introduced better properties in comparison with binary Mg alloys and pure Mg [12].

Ca strengthens the Mg alloy and refines the grain structure [9]. However, Magnesium alloys with more than 1 wt% Ca are not suitable for physiological conditions. They degrade relatively fast and the amount of released Ca exceeds its limits [9,10,11,12,13]. Zn also refines the grains and increases the strength of Magnesium alloys [14, 15]. Gandel et al. explained that Mn addition to an Mg alloy encapsulates Fe impurities which has the most adverse effect on the corrosion rate. Thus, the formation of Fe–Mg galvanic cell is prevented by an Mn barrier leading to less galvanic corrosion [16]. Moreover, Bakhsheshi-Rad et al. found that adding 0.5% Mn increases the corrosion resistance and results in smaller grain size while studying the effect of Mn addition, [17]. Ce and La increase the corrosion and creep resistance respectively in Mg alloys. Ce has a similar standard electrode potential to Mg [18].

In addition to alloying, there are other treatments such as homogenizing, which can refine the microstructure of the Mg alloys. Homogenization treatment abates the composition segregations, non-equilibrium phases, and releases the residual stress of ingots [19].In the present study, the aim is to gather more understandings about the effects of adding Mn on corrosion properties and biocompatibility of a novel biodegradable Mg–1Ca–2Zn–1RE-xMn alloy (x = 0, 0.3, 0.6, 0.9 wt%). This new alloy is expected to have acceptable corrosion properties which can make it a proper material for further analysis to use it as a biodegradable implant.

2 Materials and Methods

The Mg–1Ca–2Zn–1RE–xMn alloys used in this study were prepared from pure Mg (> 99.85%), pure Zn (> 99.95%), pure Ca (> 99.99%) and high purity Mn (> 99.99%) and the rare earth ingots containing 38% La and 62% Ce. All the elements were melted in a stainless steel crucible by an electrical resistance furnace at 720 °C and the melt was held still for 30 minutes to get homogenized. Then the melt was poured into four permanent rectangular steel molds (40 mm × 40 mm × 120 mm), preheated to 200 °C to achieve 4 ingots of the base alloy. Different amounts of Mn were added to these ingots afterward when they were melted again separately in the furnace. All the alloys were homogenized afterward at 425 °C for 16 h in argon atmosphere followed by cooling in the air. Inductively coupled plasma atomic emission spectroscopy (ICP-AES) was used to examine the chemical composition of each alloy. The results are illustrated in Table 1.

Table 1 Chemical composition of the Mg–1Ca–2Zn–1RE–xMn alloys in wt% (x = 0, 0.3, 0.6 and 0.9)
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Microstructural investigations were performed using a field-emission scanning electron microscope (TeScan-Mira3) coupled with energy dispersive spectroscopy (EDS). X-ray diffractometry (Inel- EQUINOX3000) was used to analyze significant intermetallic or secondary phases in the alloys. These analyzes were performed using monochromatic Cu-Kα radiation.

Electrochemical tests were performed at room temperature in HBSS (Table 2) using an AUTOLAB electrochemical Interface and specimens with an exposed area of 1 cm2 as working electrodes. A platinum sheet (1 cm × 1 cm) was used as the counter electrode and an Ag/AgCl electrode (in saturated KCl) was used as the reference electrode. To gain a steady-state for testing potentiodynamic polarization, the open circuit potential was considered 2500 s. Then, the electrode potential was scanned from − 1.3  to − 11.7 V versus the reference electrode at a scan rate of 1 mVs− 1. To fit the polarization curves gained from the patterns, Nova software was used with the least square method. EIS tests were performed in the frequency range from 100 kHz to 10 mHz and the perturbation amplitude was 10 mV. The obtained experimental data were fitted with Z view software.

Table 2 Composition of HBSS (Hank’s balanced salt solution)
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The biocompatibility of samples was evaluated by immersing them in HBSS. Each sample was immersed in 50 ml of HBSS and was kept under static conditions inside an incubator at 37 °C. After being immersed for 5 days, the samples were removed from the solution, washed with distilled water and then dried at 40 °C. SEM (TeScan-Mira3) coupled with EDS, was used to observe the morphology and microstructure of the surface after immersion and also Fourier transform infrared spectroscopy (Nicolet-Nexus670) was used to analyze the phase and structure at the surface of the samples.

3 Results and Discussion

The SEM images of homogenized Mg–1Ca–2Zn–1RE–xMn alloys are shown in Fig. 1. Each sample consists of two distinct areas; the matrix (α) and the secondary phases which seem to be placed on the grain boundaries.

Fig. 1
figure 1

SEM images of Mg–1Ca–2Zn–1RE microstructure with; a no Mn, b 0.3 wt% Mn, c 0.6 wt% Mn, d 0.9 wt% Mn

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The secondary phase consists of alloying elements precipitated in the grain boundaries can be seen in the Fig. 2. The alloy with 0.9 wt% Mn has a secondary phase richer of Mg than that of the alloy with no Mn and lower amounts of alloying elements in the precipitations. Thus, by the addition of Mn, the concentration of the solute elements in the boundaries has decreased and it seems that these solute elements have moved toward the matrix. Thus, it helps decreasing the corrosion rate because the structure becomes more uniform with less segregations since the secondary phase get more similar to the matrix.

Fig. 2
figure 2

EDS analyses of secondary phase in Mg–1Ca–2Zn–1RE microstructure with; a no Mn, b 0.9 wt% Mn

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According to the XRD results (shown in Fig. 3), in all samples, the α-Mg phase (Ref. Code: 96-901-3055) corresponding to the matrix and the secondary phases including intermetallic Mg2Ca (Ref. Code: 96-431-3240), MgCe (Ref. Code: 96-900-8819), LaZn (Ref. Code: 96-900-8828), and LaZn5 (Ref. Code: 96-223-3313) are distinguishable. There is not any peak related to Mn containing phases and this implies that the added Mn remains in the matrix.

Fig. 3
figure 3

XRD peaks of Mg–1Ca–2Zn–1RE–xMn alloy with; a no Mn, b 0.3 wt% Mn, c 0.6 wt% Mn, d 0.9 wt% Mn

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The Ca2Mg6Zn3 which is a prevalent phase in the Mg–Ca–Zn alloys can be harmful to the corrosion properties of Mg alloys [20, 21]. However, there is no significant peak related to the Ca2Mg6Zn3 phase in the XRD results. It could be due to Zn/Ca ratio that is more than 1.25 [17]. It seems that Mg–RE phases are formed instead of this ternary phase which leads to better corrosion properties.

Mg2Ca reduces the corrosion properties of Mg alloys [17], but MgCe has a positive effect on decreasing the corrosion rate of the alloys since the corrosion potential of Ce is similar to that of Mg and it lowers the galvanic corrosion in the alloys [18].

Figure 4 illustrates the electrochemical polarization curves obtained in HBSS, which indicates the decrease in the current density of Mn-containing Mg–1Ca–2Zn–1RE alloys. As it is illustrated, the addition of Mn to the Mg–1Ca–2Zn–1RE alloy has clearly moved the corrosion potential toward more active values but has decreased the current density by the Tafel extrapolation method. There is an exception for the alloy with 0.6% Mn which has a bigger current density in comparison with the alloy with 0.3% Mn. This can be due to the formation of films on the alloys which can be analyzed better using further examinations such as EIS measurements.

Fig. 4
figure 4

Potentiodynamic polarization curves of Mg–1Ca–2Zn–1RE–xMn alloys in HBSS

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The exact data about this test is demonstrated in Table 3 where Ecorr and Icorr represent corrosion potential and corrosion density respectively. βa and βc indicate the Tafel anodic and cathodic slopes and corrosion rate is calculated using Eq. (1).

Table 3 The fitting results of polarization curves of Mg–1Ca–2Zn–1RE–xMn alloys
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In Eq. (1), corrosion density, icorr (µA cm− 2), density, d (g cm− 3) and equivalent weight, E.W. (g) are involved:

$${\text{Corrosion}}\;{\text{Rate (mpy) = }}\frac{{{{0}}{{.13}}\;{\text{i}}_{{{{\rm corr}}}} \;({\text{E}}.{\text{W}}.)}}{{\text{d}}}$$
(1)

EIS measurements were carried out (Fig. 5) to examine the protection gifted by the film formed from corrosion on Mg–1Ca–2Zn–1RE–xMn alloy in HBSS since the results from potentiodynamic polarization cannot explain the whole effect of Mn addition on the corrosion resistance. The electrochemical response recorded from these alloys to EIS tests was simulated regarding the equivalent circuits pictured in Fig. 6 .In these circuits, the solution resistance between the reference and working electrodes is represented with Rs. Rfilm and Rct indicate the resistance of the corrosion product film and the charge transfer resistance at the interface between the solution and the specimen respectively. CPEfilm is the constant phase element of the film produced by corrosion, and CPEdl is the capacitance of the double layer and charge transfer. In this model, the inductive behavior refers to L and RL. L represents the induction and the resistance is indicated with RL [22, 23].

Fig. 5
figure 5

Nyquist plot from EIS results of Mg–1Ca–2Zn–1RE–xMn alloy immersed in HBSS for 2 h

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In Fig. 6 capacitive loops obtained in the high and low-frequency regions show that all the Mn-containing Mg–1Ca–2Zn–1RE alloys have bigger capacitive loops in all frequency ranges compared to the Mg–1Ca–2Zn–1RE alloys with no Mn. (bigger capacitive loop means that the alloy is more resistant). It is seen that the Mn addition has led to larger capacitive loops, increasing the corrosion resistance in the alloys. As long as there is the L element in equivalent circuits of alloys with 0, 0.3% and 0.9% Mn, it can be derived that although alloys with 0.3% and 0.6% Mn have better corrosion resistance in comparison with alloy having no Mn, the corroded Mg surface film is vulnerable to the attack by corrosive chloride ions and other species which diffuse through the structure. However in the alloy with 0.9% Mn, larger Rct and Rfilm demonstrate that the surface oxide film of the alloys defends the Mg substrate against attacks by chloride ions in HBSS. There is no diffusion and no sign of any induction.

Fig. 6
figure 6

Equivalent circuits of EIS results of Mg–1Ca–2Zn–1RE–xMn alloy immersed in HBSS for 2 h: a without and with 0.6% Mn, b 0.3% Mn and c 0.9% Mn alloys

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The FTIR spectra of the immersed Mg specimens in HBSS are shown in Fig. 7. It can be seen that the absorption bands at 1636 cm− 1 and 3416 cm− 1 are due to O–H stretching vibration which proves that water molecules are present in the structure [24, 25].

Fig. 7
figure 7

FT-IR spectra of Mg–1Ca–2Zn–1RE–xMn alloys after immersion in HBSS

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The absorption bands at 433 cm− 1 and 3698 cm− 1 are related to Mg–O and Mg–OH respectively showing the oxidation of Mg and stretching and bending vibration of the hydroxyl (OH−) groups due to the formation of hydroxide film in the structure. The characteristic peaks that appeared at 1080 cm− 1 and 592 cm− 1 are assigned to the phosphate groups and are considered to be the IR fingerprint of a hydroxyapatite (HAP) structure [26, 27].

In particular, the peaks at 1459 cm− 1 and 874 cm− 1 can be related to CO32− which replaces the PO43− in HAP, thus forming the B-type carbonated HAP (CHAP) [28]. The CO32− originated from CO2 in air, which is dissolved in the electrolyte solution, and the B-type CHAP can normally be obtained from precipitation reaction [29].

It seems that the HAP is formed on the specimen and its structure consists of CO32−, PO43− and water molecules. This coating has elevated the corrosion resistance of the magnesium alloys and its formation indicates the bioactivity of the Mg alloys as they were immersed in HBSS.

Figure 8 shows the corroded surface of Mg–1Ca–2Zn–1RE–xMn specimens after immersion in HBSS for 72 h. Figure 8a shows that the surface of the alloy with no Mn is covered by a film and cracks with a heterogeneous distribution. The crack formation can be due to water loss of the corrosion products and surface shrinkage. The existence of cracks let the solution to contact the matrix easier, accelerating the corrosion of the matrix. Thus the low corrosion resistance of this alloy is explained.

Fig. 8
figure 8

SEM images of microstructure after immersion of Mg–1Ca–2Zn–1RE alloy with; a no Mn, b 0.3 wt% Mn, c 0.6 wt% Mn, d 0.9 wt% Mn

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Figure 8b reveals that heterogeneous coatings are formed on the specimen with different corrosion resistances and the cracks decreased but there are still cracks making the contact between the solution and the matrix. In Fig. 8c the specimen is covered with a uniform film which is less compact than the film formed on Mg alloy with 0.3% Mn.

Figure 8d indicates that the alloy with 0.9% Mn is covered with a compact and more homogeneous film which is made of phosphate and carbonate groups and the morphology is very likely to be HAP. By analyzing the EIS measurements and the SEM picture of this specimen, it can be derived that there has been no contact between the solution and the matrix.

Both FT-IR analyses and SEM pictures indicate that alloy specimen with 0.9% Mn has the best biocompatibility among the analyzed Mg alloys.

4 Conclusion

After investigating the effect of Mn addition on corrosion properties of Mg–1Ca–2Zn–1RE-based alloys, it could be concluded that Mn does not produce any intermetallic phase or compound in any added amount and increasing Mn content moves the solute elements to the final solidification zone and makes the grain boundaries more vacant. The presence of phases such as MgCe has led to better corrosion properties as long as corrosion potential of Ce is similar to that of Mg and it lowers the galvanic corrosion.

Mn addition has resulted in decreasing the current density and lowering the corrosion rate and the alloy with 0.9 wt% shows a desired corrosion resistance. Also, it was observed that the Mg–1Ca–2Zn–1RE-base alloy is biocompatible and hydroxyapatite was formed on the specimen. Mn containing alloys have also good biocompatibility. The alloy with 0.9 wt% shows the best biocompatibility among the other alloys. The morphology of the formed hydroxyapatite on this alloy is more homogeneous and compact which has not let any ion to diffuse and acts like an aappropriate barrier. Also this alloy has a more similar secondary phase to its matrix among other alloys which has led to better corrosion properties.