Histological sections stained by Masson's Trichrome provided information about the morphology of urothelium in the presence of the metal disks and their corrosion products. The tissue-engineered ureter used in this study consisted of layers of urothelium and stroma (connective tissue) beneath it, which closely mimicked native ureteral tissue. The ability of the urothelial cells to migrate, proliferate, and differentiate to form urothelium is important for maintenance of urothelial integrity and retaining the permeability barrier function [52]. Figure 1d shows the division of the tissue into four areas, from the most proximal area to the metal disk to the most distal one. In all groups, including the control group, which was a disk of the surgical grade stainless steel, in the section stained from area 1 (in direct contact with or under metal disks) no cells or dead cells could be seen (Figure 4a,e,i). In area 2, the layers of urothelium seemed damaged in the control and Zn–1Mg groups, (respectively Figure 4b,j), or consisted of a thin epithelium in the case of the ZM21 group (Figure 4f). The transition from the absence of urothelium to a thin damaged epithelium was gradual, as well depicted in Figure 4j. In area 3, the urothelium seemed unaffected, and the superficial urothelial cells were distinct in the control (Figure 4c) and the Zn–1Mg group (Figure 4k), whereas they formed a wave-like structure in the ZM21 group (Figure 4g). This area is known to be the site of bubbles and thick white structures; it seems obvious that bubbles pushed the urothelium into this shape. The bubbles were made of hydrogen as a result of magnesium corrosion, which can be written as an overall reaction of: Mg + 2H2O → Mg(OH)2 + H2, where the Mg(OH)2 further reacts with Cl− ions to form soluble MgCl2 [38]. Finally, in area 4, the urothelium remained unaffected by the metal disks, whatever the condition. Based on the Masson's Trichrome staining results, the metal alloy groups showed similar morphology compared to the controls.

Differentiated urothelial cells, also known as umbrella cells, assume their role in maintaining the barrier function by producing specific proteins such as zonula occludens (ZO)-1. Other proteins, such as Ki67, allow evaluation of the proliferative state of the tissue, which could be overexpressed in the case of a tissue physiological or pathological response (e.g., injury or cancer). Laminin-5 is also an important molecule because its presence in the basal lamina is an absolute requirement for adequate urothelial cell differentiation, as evidenced by Rousseau et al. [53]. Figure 5 presents the ZO-1, Ki67, and laminin-5 immunofluorescence staining of the histological sections of the UW model. ZO-1 staining was done to evaluate the presence of tight junctions between urothelial cells. Figure 5a–c shows tight junctions separating the superficial urothelial cells in each group. These tight junctions, together with uroplakin, play an important role in urothelium, acting as a barrier from urine leakage, ions, or any other compound in the urine. Cell proliferation processes were not observed for the UW model, as shown by the negative result of the Ki67 staining (Figure 5d–f). Laminin-5 staining confirmed that delamination or detachment of basal layer did not occur in either the control or the metal alloy groups (Figure 5g–i).

‐ ‐ **Figure 4.** Histological analysis of UW tissue with different metal disks, stained with Masson's Trichrome. Nucleus stained as purple and stroma/collagen stained as blue. Tissue sections are divided into four areas: (**1**) under/in contact with metal disks, (**2**) peri-metal site, (**3**) nearest area to peri-metal site, and (**4**) farthest area from peri-metal site. Bar = 100 µm.

‐

‐ ‐ Tight junctions (TJ) between umbrella cells represent the main permeability barrier in combination with another protein, uroplakins [54]. TJs restrict paracellular diffusion and movement, and they contribute to maintenance of the surface polarity of cells by restricting the movement of proteins and lipids between membrane compartments. TJs are composed of cytoplasmic proteins, such as ZO-1, linking TJs to the cytoskeleton, and integral transmembrane proteins, such as occludin, junctional adhesion molecule, and claudins [55–57].

‐ ‐ ‐ ‐ Laminins are a group of proteins that are the most important components of the basal membrane. Laminins play a role in the interaction between urothelial cells and the extracellular matrix. In the ureteral mucosa, laminins are expressed by the basal cells and support the assembly of the basal membrane, contributing to maintenance of cell and tissue integrity. Laminins are also involved in cell adhesion, hemidesmosome formation, migration, proliferation, differentiation, and prevention of programmed cell death [58–61]. The presence of two important proteins, ZO-1 and laminin-5, was noted in our UW model with the direct toxicity method. This means that the urothelium function of producing barrier proteins was not interrupted by the metals and their corrosion products (except in

the area directly under the disk, where the epithelium was destroyed). The urothelium may also have preserved its permeable barrier function in close proximity to the cells.

‐ ‐ ‐ ‐ ‐ **Figure 5.** Expression of urothelial differentiation-associated protein and basal lamina in response to contact with metal disk after 7 days of incubation: (**a**–**c**) immunofluorescence anti-ZO-1 (red), (**d**–**f**) Ki67 (green), and (**g**–**i**) anti-laminin-5 (green). Nuclei were stained with Hoechst (blue).

#### *3.4. Scanning Electron Microscopy Analysis*

‐

‐ ‐ ‐ Evidence of corrosion of metals was observed in alloy groups after 7 days of incubation, while the control group showed the same surface as before the incubation (Figure 6a,b). Corrosion of ZM21 was seen as layers of metal flakes all over the surface (Figure 6d,e), while the Zn–1Mg sample had micro-size holes with attached precipitation on some parts of its surface (Figure 6g,h). Urothelial cells and stroma were observed on the surface of each metal disk (Figure 6c,f,i). Attachment of cells is marked by pseudopodia structures in cells with flattened shape. These results confirmed that the lost or removed parts of urothelium from area 1 in the UW model were actually attached to the metal and were removed together with the metal disks. The cells that were attached 3–13 µm in size, indicating that they were basal and intermediate cells [62]. This indicates that the compression of the urothelium probably inhibited the urothelial cells from differentiating well, which occurred similarly in the control and alloy groups. Basal and intermediate cells alone, i.e., non-differentiated urothelial cells, do not have the ability to produce such proteins able to maintain the impermeable barrier. These very encouraging results pave the way for further studies using a more representative metal structure, e.g., a metal mesh, to closely simulate the ureteral stent.

‐

**Figure 6.** Scanning electron microscope micrographs of the surface of metal disks incubated for 7 days on the UW model: (**a**–**c**) corrosion layer was observed on the alloys' surface, while the control showed no change, (**d**–**f**), urothelial cells and stroma collagen layer attachments were observed on the surface of each disk.

#### **4. Conclusions**

‐ To our knowledge, this is the first study reporting a biological assessment of absorbable metals in view of ureteral stent application using normal human urothelial cells (NHUCs) and a metal direct toxicity test on a flat 3D ureteral wall tissue model that mimics the in vivo conditions. The NHUCs exhibited a survival mode in response to the toxicity of the magnesium and zinc extracts in monolayer culture, shown by reduced viability and unhealthy appearance, but recovered after some days. This result suggests that the cells could rapidly colonize damaged areas to restore mature urothelium, restoring the barrier function to block urine extravasation in the body. The direct toxicity test in the 3D ureteral wall model showed that the corrosion of metals had an effect on the morphology of cell layers, but the urothelium was still able to maintain its function in producing the barrier protein and may have had a permeable barrier function in close proximity to the cells. This indicates that the tissue was healthy despite the presence of the metal disks and their corrosion products after 7 days of exposure. The attached cell layers on the surface of metal disks may indicate that the metal specimens were non-toxic. The zinc alloy group seemed to be more cytocompatible toward human urothelial cells compared to magnesium alloys, and could be further exploited for biodegradable ureteral stent applications. In the future, in vitro test using a flat 3D ureteral wall tissue model could be used to partially replace animal testing in the initial steps of optimization of absorbable metal stents.

**Author Contributions:** Conceptualization, D.P., H.H., S.B., and S.C.; Data Curation, D.P., and S.C.; Formal Analysis, D.P., H.H., S.B., and S.C.; Methodology, D.P., H.H., S.B., and S.C.; Supervision, H.H. and S.B.; Writing—Original Draft Preparation, D.P.; Writing—Review & Editing, D.P., H.H., S.B., and S.C.

**Funding:** This work is supported by the Natural Sciences and Engineering Research Council of Canada (NSERC) and the Fonds de recherche du Québec-Santé (FRQS).

**Acknowledgments:** The authors would like to thank Geneviève Bernard for the help in preparing and conducting the 3D assay, Agung Purnama for the help in the SEM experiment, Sébastien Champagne for the help in corrosion test, Ehsan Mostaed and Maurizio Vedani for providing us with the metals.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Microstructure and Properties of Nano-Hydroxyapatite Reinforced WE43 Alloy Fabricated by Friction Stir Processing**

**Genghua Cao 1 , Lu Zhang 2 , Datong Zhang 2 , Yixiong Liu 1 , Jixiang Gao 1 , Weihua Li 1 and Zhenxing Zheng 1,\***

Received: 10 July 2019; Accepted: 6 September 2019; Published: 16 September 2019

**Abstract:** This research mainly focuses on the successful fabrication of nano-hydroxyapatite (nHA) reinforced WE43 alloy by two-pass friction stir processing (FSP). Microstructure evolution, mechanical properties, and in vitro corrosion behavior of FSPed WE43/nHA composite and FSPed WE43 alloy were studied. The results show that nHA particles are effectively dispersed in the processing zone, and the well-dispersed nHA particles can enhance the grain refine effect of FSP. The average grain sizes of FSPed WE43 alloy and WE43/nHA composite are 5.7 and 3.3 µm, respectively. However, a slight deterioration in tensile strength and yield strength is observed on the WE43/nHA composite, compared to the FSPed WE43 alloy, which is attributed to the locally agglomerated nHA particles and the poor quality of interfacial bonding between nHA particles and matrix. The electrochemical test and in vitro immersion test results reveal that the corrosion resistance of the WE43 alloy is greatly improved after FSP. With the addition of nHA particles, the corrosion resistance of the WE43/nHA composite shows an even greater improvement.

**Keywords:** WE43/HA composite; friction stir processing; microstructure; mechanical properties; corrosion behavior

#### **1. Introduction**

Magnesium and its alloys have several advantages when compared with traditional metal biomedical materials. The Young's modulus and density of magnesium and its alloys are similar to that of natural bone, which can effectively avoid the stress shielding effect [1,2]. Magnesium is biodegradable in vivo, and the corrosion products have proven to be nontoxic. In addition, magnesium-based biomedical materials have been widely reported to positively stimulate the formation of new bone, which is favorable for bone fracture healing [3]. Therefore, magnesium alloys have great potential in applications as biodegradable metal materials [4,5]. However, biomedical magnesium alloys face the urgent issue of controlling corrosion behavior by avoiding local corrosion and controlling corrosion rates, in order to meet the safety and mechanical property requirements for biodegradable metal materials [6].

Hydroxyapatite (HA), the main inorganic component of human bone tissue and teeth, has emerged as a promising bioceramic material for its outstanding biocompatibility and bioactivity [7,8]. However, due to its brittleness and poor strength, HA in biomedical applications is currently limited to non-load bearing parts or low load-bearing parts. In this case, introducing HA particles into magnesium alloys

is considered an effective method to improve the corrosion rate and biocompatibility of magnesium alloys. At present, several processes such as hot extrusion, stirred casting, powder metallurgy, and other methods are performed to prepare HA magnesium matrix composites with uniform corrosion behavior, good mechanical properties, and biocompatibility [9–11].

Friction stir processing (FSP) is an emerging solid-state processing technology for preparing fine-grained metal materials [12]. In recent years, due to its slight interface reaction between matrix and reinforced particles during processing, FSP has been used for metal matrix composite preparation [13–15]. For instance, nano-hydroxyapatite (nHA) particles have successfully been added to pure Mg substrate by multi-pass FSP, and the Mg/nHA composite shows preferable corrosion resistance in simulated body fluid (SBF) or Dulbecco's phosphate buffered saline compared to the substrate. However, the investigations of mechanical properties are not mentioned in these papers [14,15]. As an implant material, the material should have reasonable mechanical properties in order to meet the clinical requirements of implantation, so investigation of these mechanical properties is also important. WE43 magnesium alloy with high strength and low cytotoxicity is suitable for biomedical applications. Therefore, in this research, the casted WE43 alloy was used as matrix, nano-sized HA particles as the reinforcing phase, and FSP was conducted to prepare WE43/nHA composites. The microstructure evolution during FSP, the effects of rotation speed on the distribution of nHA particles, and the effects of dispersed nHA particles on mechanical properties and corrosion behavior of the composites were studied.

#### **2. Experimental Procedure**

#### *2.1. Raw Materials*

Commercially available WE43 magnesium alloy (as-cast) sheets of size 150 mm × 30 mm × 6 mm (length × width × height) were used as the base metal (BM) in this study, and the chemical composition of the BM is shown in Table 1. Nano-sized HA powders (nHA) with a purity of >99% used in this study were purchased from Shanxi Baiwei Biotechnology Co., Ltd. (Xi'an, China). The TEM morphology of nHA particles is shown in Figure 1—the particles were of acicular morphology 20–30 nm in width and 60–120 nm in length.

**Table 1.** Composites of as-cast base metal (wt. %).

**Figure 1.** TEM morphology of nano-hydroxyapatite (nHA) particles.

#### *2.2. Processing*

Figure 2 presents the schematic illustration of the processing steps. To prepare the WE43/nHA composite, FSP was performed on BM sheets with a groove size of 1 mm × 4 mm (width × depth). After nHA particles were filled into the groove, a cylindrical and pin-less FSP tool was pressed down slowly until the shoulder contacted with the material, and it was then processed at a rotation speed of 600 rpm and a traverse speed of 60 mm/min along the groove direction. After processing, a metal sealing layer was formed above the groove, which could avoid the nHA particles escaping from the groove during FSP (Figure 2c). In the next processing step, cylindrical FSP tools with a shoulder diameter of 15 mm, consisting of a tapered cylindrical pin with a diameter varying from 2 to 5 mm over the length of 5 mm, were used to perform FSP on the sealed specimens.

A two-pass FSP with optimized processing parameters, including a rotation speed of 1000 rpm and a traverse speed of 60 mm/min, was employed to obtain WE43/nHA composites with a fine and uniform microstructure. The plunge depth of 0.8 mm and tilt angle of 2.5◦ relative to the normal direction of FSP plane were kept constant. The composite specimens obtained by FSP were coded as WE43/nHA. The same processing parameters and conditions were applied for conducting FSP on BM without the addition of nHA particles, with the obtained specimens being coded as FSP-WE43.

**Figure 2.** Schematic illustration of the processing steps. (**a**) Step 1; (**b**) Step 2; (**c**) Step 3; (**d**) Step 4.

#### *2.3. Characterization of Microstructure and Phase Composition*

The specimens used for microstructure observation were mechanically grinded with emery papers (up to #5000 grade) and polished on a polishing machine. Further, the polished specimens were etched by picric acid solution (picric acid 5 g, alcohol 80 mL, acetic acid 10 mL, and deionized water 10 mL). Optical microscope (DM15000M, Leica, Wizlar, Germany) was used to observe microstructure at lower magnification. The distribution of second phases and reinforcements, as well as the fracture morphology, were observed by a scanning electron microscope (Nova Nano 430, FEI, Hillsboro, OR, USA). The morphology of the nHA particles was characterized by a transmission electron microscopy (JEM-2100F, JEOL, Tokyo, Japan).

#### *2.4. Mechanical Properties Testing*

The microhardness test was conducted on a HVS-1000 digital Vickers microhardness tester (YouHong, Corp., Shanghai, China) with the application of a load of 0.98 N and a loading cycle of 10 s. The indention interval was selected to be 0.5 mm in stirred zone (SZ), and every indentation was measured three times and the average value was calculated. Tensile specimens were machined parallel to the processing direction with the gauge being completely within the stirred zone, the shape and dimension of a tensile specimen is shown in Figure 3. The tensile test was carried out on a SANS CMT5105 universal tensile testing machine (MTS, Eden Prairie, MN, USA) with a strain rate of 2 × 10−3 s −1 . At least five specimens were tested to evaluate the average property values. − −

**Figure 3.** Shape and dimension of tensile specimen.

#### *2.5. Corrosion Behavior*

#### 2.5.1. Electrochemical Test

− Potentiodynamic polarization curve tests were performed on an electrochemical workstation (Vertex.5A. EIX, IVIUM, Eindhoven, the Netherlands) in SBF solution (8.035 g/L NaCl, 0.355 g/L NaHCO3, 0.225 g/L KCl, 0.231 g/L K2HPO4·3H2O, 0.311 g/L MgCl2·6H2O, 0.292 g/L CaCl2, 0.072 g/L Na2SO4, and 6.118 g/L Tris (HOCH2)3CNH2). The reference electrode was a saturated Ag/AgCl electrode and a platinum electrode was used as the counter electrode. One square cm area of the specimens was used as the working electrode. Specimens were exposed to the SBF solution for 30 min prior to the beginning of the experiments to establish open circuit potential. The potentiodynamic polarization was done between the potentials −2.5 and 0.5 V with a scanning rate of 5 mV/s.

#### 2.5.2. Immersion Test

μ

The immersion test was performed per ASTM-G31-72 in SBF at 37 ◦C for 24, 48, and 72 h, respectively. Weight loss specimens of size 6 mm × 4 mm × 2 mm (length × width × height) were cut from the SZ of FSP-WE43 and WE43/nHA composite samples and BM as well, for the purpose of measuring the corrosion rate of specimens in SBF. After immersion, the corrosion products were removed by chromic acid (200 g/L CrO3, 10 g/L AgNO3, and 20 g/L Ba(NO3)2) and then ultrasonically cleaned in distilled water and ethanol, respectively. The weight of specimens was measured before and after immersion. The corrosion rate was calculated by the following equation:

$$\text{CR} = \text{W/At} \text{p} \tag{1}$$

CR = W/Atρ ρ where CR is the corrosion rate (mm/year); W represents the weight loss (g); A refers to the surface area (cm2 ); t is the immersion time; and ρ is the standard density of WE43. A density of 1.83 g/cm3 was used for all specimens in this study.

The corrosion morphology of specimens was observed by the SEM mentioned above.

μ

#### **3. Results and Discussion**

#### *3.1. Microstructure Evolution*

Figure 4 shows the optical microstructure of BM and the stir zone of FSP-WE43 and WE43/nHA specimens. The average grain size of BM is measured ~50.9 µm. Grain refinement is achieved up to ~5.7 and ~3.3 µm in the FSP-WE43 specimen and WE43/nHA specimen, respectively. During FSP, materials in the stir zone will undergo dynamic recrystallization and coarse second phases will break into small particles, resulting from the severe plastic deformation caused by the FSP tool and thermal effect caused by friction [13,16]. This is the main reason for the apparent refinement of grains after FSP. During FSP of magnesium, the peak temperature of SZ is reported lower than 550 ◦C, at which temperature the nHA particles remain stable [17–19]. The incorporated insoluble nHA particles act to stimulate nucleation and impede the migration of grain boundaries [20,21]. As a result, further grain refinement is achieved on WE43/nHA composites.

**Figure 4.** Optical image of (**a**) base metal (BM); (**b**) friction stir processing (FSP)-WE43; and (**c**) WE43/nHA specimens.

#### *3.2. Distribution of HA Particles*

μ Figure 5 shows the distribution of the second phase particles and its corresponding EDS analysis in the WE43/nHA specimen. nHA particles are found well dispersed on the matrix after FSP, joined by only a few clusters with a diameter of less than 10 µm (Figure 5a). The high-angle annular dark field (HAADF) image and the corresponding EDS analyses of the stir zone (Figure 5b,c) confirm that nHA particles are successfully added into the WE43 alloy matrix and most nHA particles remain at nanoscale.

**Figure 5.** (**a**) SEM image; (**b**) TEM image; and (**c**) EDS analysis of the WE43/nHA specimen.

#### *3.3. Mechanical Properties*

#### 3.3.1. Microhardness

The microhardness distribution curves of FSP-WE43 and WE43/nHA specimens are plotted in Figure 6. After FSP, the Vickers microhardness value of the SZ is significantly increased in FSP-WE43 and WE43/nHA specimens. The microhardness value at the SZ of the FSP-WE43 and WE43/nHA specimens is relatively stable, while the microhardness value of the base metal region fluctuates. The region with a higher microhardness value is about 5 mm in width, which is approximately equal to the diameter of the pin on FSP tool. The mean microhardness value of WE43 substrate is ~62.6 HV, which increases to ~79.6 HV by FSP without the addition of nHA powder. In the case of introducing nHA powder during FSP, the mean microhardness value is improved up to ~85.2 HV.

The grain size of as-cast WE43 alloy is coarse and the microstructure is ununiform, so the microhardness value of the BM zone is lower and fluctuates. After FSP, the grains in the stir zone are remarkably refined and grain boundary strengthening is considered the main reason for the significant increase of microhardness value in the stir zone. For the WE43/nHA specimen, the distribution of microhardness values is related to the dispersion of nHA particles, and a relatively stable microhardness value fluctuation for the WE43/nHA specimen indicates that nHA particles are dispersed uniformly on WE43 substrate, which is consistent with the microstructure observation.

**Figure 6.** Microhardness distribution curves of WE43/nHA and FSP-WE43 specimens.

#### 3.3.2. Tensile Properties

Figure 7 shows a comparison of the tensile properties of the base metal and processed WE43 (with or without adding nHA particles) specimens. As plotted in Figure 7, the ultimate tensile strength (UTS), yield strength (YS), and elongation of specimens after processing are improved in different degrees compared with the BM. For BM, the YS and UTS are measured as only ~153.3 and ~193.2 MPa, respectively. After FSP, the YS and UTS of the FSP-WE43 specimen are improved up to ~198.7 MPa and ~255.4 MPa, respectively. Compared with the FSP-WE43 specimen, a slight decline in strength is observed on the WE43/nHA specimen, with the YS and UTS being measured as ~185.1 and ~232.3 MPa, respectively. Furthermore, the value of elongation after fracture is greatly increased by FSP as well. The elongation of BM is only ~5.2%, while in the FSP-WE43 and WE43/nHA specimen it is 20.2% and 10.1%, respectively.

**Figure 7.** Tensile properties of BM, FSP-WE43, and WE43/nHA specimens. UTS: ultimate tensile strength; YS: yield strength.

#### 3.3.3. Fractographic Studies

Figure 8 shows the fractographic images of tensile specimens. The fractographic of BM (Figure 8a) consists of a large amount of cleavage facets and voids defects, which indicates that BM fails in a brittle way. Figure 8b shows the fractographic image of an FSP-WE43 specimen, in which a large amount of fine equal-sized equiaxed dimples can be observed. The fracture morphology exhibits plastic fracture characteristics. This confirms the high ductility of the FSP-WE43 specimen as shown in Figure 7. The fractographic images of the WE43/nHA specimen are represented in Figure 8c,d. Fine dimples

can also be observed in the WE43/nHA specimen and partially agglomerated nHA particles can be seen (as shown by white circles). By observing agglomerated nHA particles at higher magnification, several cracks across the nHA cluster is found (Figure 8d). The loose nHA clusters can be the initiation source of cracks during failure, which is the main reason for the decline of elongation in the WE43/nHA specimen compared with that in the FSP-WE43 specimen.

In general, coarse grains and brittle Mg12Nd networks in as-cast WE43 magnesium alloy cause its poor tensile properties [22]. As shown in Figure 4, remarkable grain refinement is achieved and coarse Mg12Nd phases are broken into fine particles after FSP. Under the combined effects of grain boundary strengthening and dispersion strengthening, the tensile strength of processed specimens (with or without addition of nHA particles) are significantly improved. In addition, FSP eliminates the voids defects in as-cast WE43, which is beneficial to the improvement of strength and ductility. However, the localized nHA agglomerates reduce the tensile properties of the composites to a certain extent compared with those in the FSP-WE43 alloy, while the tensile properties of the WE43/nHA specimen are still improved compared with those in the BM.

**Figure 8.** Fractographic images of (**a**) BM, (**b**) FSP-WE43, and (**c**) and (**d**) WE43/nHA specimens.

#### *3.4. Corrosion Behavior*

#### 3.4.1. Electrochemical Test

− μ − μ − μ The potentiodynamic polarization curves of as-cast WE43 and processed WE43 specimens are demonstrated in Figure 9. The corrosion potentials (Ecorr) and corrosion current density (icorr) of BM are measured as −1.691 mV (vs. Ag/AgCl) and 109.6 µA/cm2 , respectively. For the processed specimens, the Ecorr of the WE43/nHA specimen (−1.661 mV) shifts toward the positive side and the icorr (46.7 µA/cm2 ) is lower than that of BM. The Ecorr and icorr values of the FSP-WE43 sample are measured as −1.678 mV and 53.7 µA/cm2 , respectively. The highest Ecorr value and lowest icorr value indicate that the WE43/nHA composite has the best corrosion resistance in this study.

−

− μ

**Figure 9.** Potentiodynamic polarization curves of BM, FSP-WE43, and WE43/nHA specimens in simulated body fluid (SBF).

As discussed in Section 3.1, grain size is significantly refined by FSP. For magnesium alloy, fine grains are proved beneficial to the formation of a passive layer, as the result of increasing the number of grain boundaries per unit volume and reducing the galvanic couple between grain boundary and grain interior [23,24]. This is the main reason for the increase of corrosion resistance on processed WE43 specimens. Moreover, tiny dispersed nHA particles can also contribute to uniform corrosion behavior [25]. Therefore, the combined effects of grain refinement and dispersion of nHA particles lead to the improvement of corrosion resistance on WE43/nHA composites.

#### 3.4.2. Degradation in Immersion Test

Figure 10 shows the corrosion weight loss curve of BM, FSP-WE43, and WE43/nHA specimens. During the immersion period of 120 h, the corrosion weight loss of BM increases rapidly. After immersion for 120 h, BM specimens are almost completely degraded in the SBF solution. The corrosion weight loss rates of the FSP-WE43 specimen and WE43/nHA specimen are relatively stable. In the first 72 h, the weight loss rates of the FSP-WE43 and WE43/nHA specimen are about the same. After immersion for 72 h, the weight loss rate of the FSP-WE43 specimen increases rapidly, while that of the WE43/nHA specimen maintains a relatively stable value. The corrosion rates are calculated according to Equation (1). After immersion for 120 h, the corrosion rate of the BM specimen is 26.8 mm/year, and the corrosion rates of the FSP-WE43 and the WE43/nHA specimen are 8.1 mm/year and 3.9 mm/year, respectively. Obviously, the FSP process has greatly improved the corrosion resistance of the casted WE43 alloy, and the addition of nHA particles has further improved the corrosion resistance of the material.

Corrosion morphologies (with corrosion products) of specimens after immersion in SBF for 72 h are shown in Figure 11. The SEM images of corrosion morphology show that BM experiences severe localized corrosion after immersion for 72 h and the accumulation of thick corrosion products occur locally, which are reported to be Mg(OH)2 and calcium phosphate bio-minerals [26,27]. For the FSP-WE43 specimen, uniform protective films are observed in most areas, while a small number of protective films fall off locally, which may decrease the protective effect on corrosion attack. For the WE43/nHA specimen, a dense and uniform protective layer generates on the composite surface. These results indicate that the corrosion morphology is changed from local corrosion in as-cast WE43 alloy to uniform corrosion in FSP-WE43 alloy, which is attributed to the fine-grained and homogeneous microstructure by FSP. With the addition of dispersed nHA particles, the uniform corrosion morphology on WE43/nHA composite is more obvious.

**Figure 10.** The weight loss and corrosion rate curves of immersion specimens.

**Figure 11.** Corrosion morphologies of (**a**–**c**) BM, (**d**–**f**) FSP-WE43, and (**g**–**i**) WE43/nHA specimens after immersion in SBF for 72 h at 37 ◦C (with corrosion products).

Corrosion morphologies (without corrosion products) of specimens after immersion in SBF for 72 h are shown in Figure 12. It can be seen from Figure 12a that the BM specimen experiences extremely severe corrosion attack and a large amount of material is dissolved in SBF. Deep and large etch pits can be observed (Figure 12b), proving that the material has been eroded by SBF and the corrosion products cannot prevent further corrosion. However, the FSP-WE43 specimen still keeps a relatively complete surface morphology, although parts of materials are dissolved in SBF (Figure 12c,d). For the composite specimen, the original shape is almost maintained after immersion in SBF for three days and only shallow corrosion pits can be observed locally (Figure 12e,f). The corrosion morphology observation indicates that the corrosion resistance of the WE43/nHA composite is superior to that of the FSP-WE43 alloy and much superior to that of the as-cast WE43 alloy, which is in accord with the electrochemical test results.

**Figure 12.** Corrosion morphologies of (**a**,**b**) BM, (**c**,**d**) FSP-WE43, and (**e**,**f**) WE43/nHA specimens after immersion in SBF for 72 h at 37 ◦C (without corrosion products).

#### **4. Conclusions**

Fine-grained WE43/nHA composite was successfully prepared through friction stir processing. Microstructure evolution and mechanical properties as well as in vitro corrosion behavior of WE43/nHA composite were studied. The main findings are summarized as follows:

**Author Contributions:** Methodology G.C. and D.Z.; Supervision Z.Z. and J.G.; resources Z.Z. and D.Z.; Analysis G.C.; W.L. and Y.L.; Writing—original draft G.C. and L.Z.; Writing—reviewing and editing G.C. and J.G.

**Funding:** This research is funded by the Natural Science Foundation of Guangdong Province (2017A030310630), Guangdong Science and Technology Planning Project (2017A070715012), and Guangdong education department project (2017GCZX003, 2017KTSCX115).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

*Article*

## **Biomechanical Loading Comparison between Titanium and Unsintered Hydroxyapatite**/**Poly-L-Lactide Plate System for Fixation of Mandibular Subcondylar Fractures**

#### **Shintaro Sukegawa 1,4,\* , Takahiro Kanno 2 , Norio Yamamoto 3 , Keisuke Nakano 4 , Kiyofumi Takabatake 4 , Hotaka Kawai 4 , Hitoshi Nagatsuka 4 and Yoshihiko Furuki 1**

Received: 5 April 2019; Accepted: 8 May 2019; Published: 13 May 2019

**Abstract:** Osteosynthesis absorbable materials made of uncalcined and unsintered hydroxyapatite (u-HA) particles, poly-l-lactide (PLLA), and u-HA/PLLA are bioresorbable, and these plate systems have feasible bioactive osteoconductive capacities. However, their strength and stability for fixation in mandibular subcondylar fractures remain unclear. This in vitro study aimed to assess the biomechanical strength of u-HA/PLLA bioresorbable plate systems after internal fixation of mandibular subcondylar fractures. Tensile and shear strength were measured for each u-HA/PLLA and titanium plate system. To evaluate biomechanical behavior, 20 hemimandible replicas were divided into 10 groups, each comprising a titanium plate and a bioresorbable plate. A linear load was applied anteroposteriorly and lateromedially to each group to simulate the muscular forces in mandibular condylar fractures. All samples were analyzed for each displacement load and the displacement obtained by the maximum load. Tensile and shear strength of the u-HA/PLLA plate were each approximately 45% of those of the titanium plates. Mechanical resistance was worst in the u-HA/PLLA plate initially loaded anteroposteriorly. Titanium plates showed the best mechanical resistance during lateromedial loading. Notably, both plates showed similar resistance when a lateromedially load was applied. In the biomechanical evaluation of mandibular condylar fracture treatment, the u-HA/PLLA plates had sufficiently high resistance in the two-plate fixation method.

**Keywords:** mandibular condylar fracture; unsintered hydroxyapatite/poly-l-lactide composite plate; bioactive resorbable plate; biomechanical loading evaluation; fracture fixation

#### **1. Introduction**

Mandibular condylar fractures constitute 25%–35% of all mandibular fractures [1] and are the most common form of mandible fracture.

The selection between closed (conservative treatment) and open (surgical treatment) techniques are linked to the type of fracture, the patient's age, and the degree of functional impairment caused by the fracture. Growing patients with certain undisplaced mandibular condylar fractures can undergo closed treatment, but severely displaced fractures are indications for open surgical reduction and fixation with internal rigid devices, generally with plates and screws [2]. In adult patients, the condylar head can be adjusted for functional adaptation of the condyle without restoration of the anatomy caused by displaced base fractures with a loss of ramus height [1]. Therefore, high condylar fractures with little bone available for fixation are mostly treated non-surgically. On the other hand, lower condylar fractures that affect the subcondyle, condylar base, and condylar neck are generally treated by open reduction and internal fixation [3,4]. The surgical treatment of mandibular condylar fractures involves fixation of the fractured stumps with the use of plates and screws. There are several types of fixation for fractures of the mandibular condyle [5]; the most common fixation technique is to use two four-hole straight plates [6]. In general, a titanium metal plate system has been used as a standard osteosynthesis material.

However, titanium metal plates have been reported to produce long-term problems related to implantation in the human body, such as intracranial migration, growth retardation of the craniofacial skeleton, and hypersensitivity to cold. In recent years, as an alternative to metal plates, resorbable plates, which do not require plate removal, have been used as osteosynthesis material in maxillofacial surgery [7]. Bioresorbable osteosynthesis devices have been evolving, with improvements in bioresorbability and marked bioactivity with new material compositions, for different and better in situ behavior. Currently, many osteosynthesis absorbable materials are commercially available; these include OSTEOTRANS-MX (known in Japan as Super FIXSORB MX®; Teijin Medical Technologies Co., Ltd. Osaka, Japan). These bioactive and bioresorbable devices are made from composites of uncalcined and unsintered hydroxyapatite (u-HA) particles and poly-l-lactide (PLLA). This u-HA/PLLA osteosynthesis material is an excellent bioactive device that has an osteoconductive ability in the short-term and absorbs safely and reliably in the long-term [8,9]. However, both the strength and stability of this plate system for fixation to a mandibular subcondylar fracture are still unclear. Thus, the aim of this in vitro study was to assess the biomechanical strength of u-HA/PLLA bioactive resorbable plate systems after internal fixation of mandibular subcondylar fractures.

#### **2. Materials and Methods**

#### *2.1. Materials*

We used the bioresorbable Super FIXSORB MX® (Teijin Medical Technologies Co., Ltd., Osaka, Japan) osteosynthesis system. Forged composites of u-HA/PLLA were processed by machining or milling treatments to form various miniscrews and miniplates, which contained 30- and 40-weight fractions of u-HA particles (raw HA, neither calcined nor sintered material), respectively, in composites (hereinafter referred to as u-HA 30 miniscrews and u-HA 40 miniplates).

A 2.0 mm miniplate system (MatrixMANDIBLE™ Adaption Plate; DePuy Synthes, Raynham, Mass.) (Johnson & Johnson, New Brunswick, USA) was used as the titanium metal plate.

#### *2.2. Evaluation of Tensile and Shear Strength of Titanium and Bioresorbable Plate Systems*

#### 2.2.1. Sample Preparation

We prepared mechanical strength models by affixing the titanium and u-HA/PLLA plates to polyetherketoneketone (PEKK) plates (thickness 3 mm) with screws. The PEKK plates were held in place by different osteosynthesis systems. Plates and screws were used to form the following groups:

In this study, the bioresorbable osteosynthesis material consisted of 1.4 mm-thick plates and screws (diameter 2 mm; length 6 mm). In comparison, the titanium osteosynthesis material consisted of 1.0 mm-thick plates and screws (diameter 2 mm; length 6 mm). The experimental fixed models were mounted on an Autograph AG-20kNXD test frame (Shimadzu Co., Kyoto, Japan) across the chuck. The maximum stress until the plate or screws were destroyed and the stress at the time of 1 mm movement were measured, and the load was applied at a test speed of 10 mm/min. Two types of strength tests (the tensile and shear strength tests) were performed five times each.

#### 2.2.2. Strength Measurement

Using the method illustrated in Figure 1A according to the Japanese Industrial Standard K7113, we measured tensile strength. The peak value of the load profile attained by the Autograph AG-20kNXD (Shimadzu Co., Kyoto, Japan) was considered the tensile strength. Using the method illustrated in Figure 1B according to the Japanese Industrial Standard K7113, we measured shear strength. The peak value of the load profile attained by the Autograph AG-20kNXD was considered the shear strength.

**Figure 1.** Mechanical strength models were prepared by fixing the plate with screws to the polyetherketoneketone plate. (**A**) Tensile strength; (**B**) shear strength.

#### *2.3. Biomechanical Loading Evaluation*

#### 2.3.1. Sample Preparation

We used 20 polyurethane replicas of human hemimandibles (Mandible, Code #8900, SYNBONE AG, Laudquart, Tardisstrasse, Switzerland). Although polyurethane mandible replicates the property of spongy bone [10], this model is useful to obtain preliminary results concerning the stability of the investigated osteosynthesis systems. An ordinary subcondylar fracture model was investigated by using the osteotomy line connecting the mandibular notch to the midpoint of the mandibular ramus' posterior border. We created a fracture in the right condyle of each model, mimicking the experimental method of Meyer et al. [11]. Initially, these fracture models were created with the use of a cutting guide

by a computer-controlled program. A partial cut was made in each hemimandibular model with a diamond disk (KG Sorensen, Cotia, São Paulo, Brazil). A silicone mold of the subcondylar area of the osteosynthesis was made on the mandible and drilled to form guide holes, and then, using the mold as a guide, we completed the cut. This method was used to create identical cuts and perforations in all the hemimandibular models, thus guiding the positions for plate fixation according to each study group.

'

#### 2.3.2. Fixation Method for Mandibular Condylar Fracture

We used two straight miniplates along the ideal line of the mandibular condyle as proposed by Meyer et al. [12]. The technique is to plate below the mandibular notch and below the mandibular posterior border, respectively. This technique with monocortical screws had proved to be the most reliable and functionally stable osteosynthesis procedure for subcondylar fractures. The bone segments were held in place by different commercially available osteosynthesis methods (a titanium miniplate and bioresorbable plate) and monocortical screws. With both types of plates, bone fixation could be performed without plate bending. Thus, we created two conditions (Figure 2A):

**Figure 2.** (**A**) The mandibular subcondylar fracture replicas were held in place by double titanium (Ti) straight plates (thickness: 1.0 mm) with four monocortical screws (2.0 mm in diameter and 6 mm long screws; group A) or by double unsintered hydroxyapatite (u-HA)/ poly-l-lactide (PLLA) straight plates (thickness: 1.4 mm), each with four monocortical screws (2.0 mm in diameter and 6 mm long; group B). A linear load was applied at a displacement speed of 1 mm/min. (**B**) Anteroposterior (vertical) linear loading. (**C**) Lateromedial (lateral) linear loading.

MatrixMANDIBLE™ (1) In each bone segment in group A, double titanium straight plates (MatrixMANDIBLE™ Adaption Plate, 1.0 mm thick) with four monocortical screws (2.0 mm in diameter and 6 mm long) were installed; and (2) in each bone segment in group B, double-u-HA/PLLA straight plates (Super FIXSORB MX ®, 1.4 mm thick), each with four monocortical screws (2.0 mm in diameter and 6 mm long), were installed.

We used a total of 20 hemimandibular replicas, and 10 were allocated to each of the two groups.

#### 2.3.3. Biomechanical Loading Test

After bone segment fixation, the replicas were mounted on a testing machine (AG-20KNX; Shimazu) (Shimadzu Co., Kyoto, Japan), which was based on a biomechanical cantilever-bending model that simulates masticatory forces. The mandibular body and angle areas of each replica were then stabilized. Adaptation of the polyurethane hemimandible to the machine was guaranteed through a metallic support. In reference to past biomechanical evaluation methods, the replicas were then subjected to linear loading in two directions: from anterior to posterior (vertically; Figure 2B) and from lateral to medial (horizontally; Figure 2C). These forces simulated the muscular forces applied to an actual fractured condyle. The material testing unit created linear displacement at a rate of 1 mm/min, and loading continued until the maximum load was reached. The peak load and displacement for each replica were recorded. All replicas were analyzed for 0.5, 1.0, 1.5, 2.0, 3.0, and 5.0 mm displacements by loading and for the amount of displacement by the maximum load. Means and standard deviations were derived and assessed for statistical significance.

#### *2.4. Statistical Analysis*

Data were recorded and entered into an electronic database during the course of the evaluation by means of Microsoft Excel (Microsoft, Inc., Redmond, USA). Means and standard deviations were used for normal data distributions. The database was transferred to JMP version 11.2 for Macintosh computers (SAS Institute, Inc., Cary, USA) for statistical analysis. Then, t-tests for independent samples were performed to investigate whether there were significant differences between the mean values of the groups. The level of significance was set at *p* < 0.05.

#### **3. Results**

#### *3.1. Tensile and Shear Strength Evaluation*

The results are shown in Figure 3. In the tensile test, the mean maximum test forces were 806.1 N (standard deviation 7.9 N) in group A (titanium plates fixed with titanium screws) and 208.8 N (standard deviation 10.3 N) in group B (u-HA/PLLA bioresorbable plates with bioresorbable screws). The mean test forces at 1 mm displacement were 382.1 N (standard deviation 19.5 N) in group A and 181.8 N (standard deviation 4.2 N) in group B. The mean maximum tensile test force in group B was 25.9% that of group A, and the test force at 1 mm displacement in group B was 47.6% that of group A. These differences were significant (*p* < 0.05) (Figure 3A).

**Figure 3.** (**A**) Comparison of titanium plates and unsintered hydroxyapatite (u-HA)/poly-l-lactide (PLLA) bioresorbable plates with regard to maximum tensile test force and test force at 1 mm displacement in the tensile test. (**B**) Comparison of titanium plates and u-HA/PLLA bioresorbable plates with regard to maximum tensile test force and test force at 1 mm displacement in the shear test.

In the shear test, the mean maximum test force and the test force at 1 mm displacement were 390.5 N (standard deviation 11.0 N) and 130.6 N (standard deviation 12.2 N), respectively, in group A. In group B, these mean forces were 93.2 N (standard deviation 4.6 N) and 58.1.8 N (standard deviation 12.2 N), respectively. The maximum shear test force in group B was 23.9% that of group A, and the test force at 1 mm displacement in group B was 44.5% that of group A.

–

The strength of the titanium-based fixation system was obviously higher than that of the u-HA/PLLA bioresorbable system in both the tensile test and the shear test. In the shear test, the titanium plates were significantly stronger (*p* < 0.05) than the u-HA/PLLA bioresorbable plates at maximum shear test force and at 1 mm displacement test force (Figure 3B).

#### *3.2. Biomechanical Loading Evaluation*

The results of this experiment show that the mechanical resistances among the osteosynthesis devices remained proportional to the amount of displacement. In the vertical loading test, the titanium plates had a significantly higher load value than did the u-HA/PLLA bioresorbable plates at displacements of 0.5 and 1 mm. In the displacements of 1.5–5 mm, there was no significant difference between the two types of plates (Figure 4). As the load in the anteroposterior direction increased, although the bone fragments fixed by the anterior plate were separated, the position of the posterior bone fragments did not change (Figure 5).

**Figure 4.** Load values of the titanium plates and unsintered hydroxyapatite (u-HA)/poly-l-lactide (PLLA) bioresorbable plates according to the amount of displacement in the anteroposterior loading test.

In contrast, when the forces were applied to the condyle in the lateromedial direction, there was no significant difference between the u-HA/PLLA bioresorbable plates and the titanium plates at all displacements (Figure 6). As the load in the lateromedial direction increased, there was no excessive load on either titanium or u-HA/PLLA bioresorbable anterior and posterior plates, and both types of plates in both locations were flexed (Figure 7).

**Figure 5.** Change in subcondylar fracture segments under anteroposterior loading. Ti, titanium (plates); u-HA/PLLA, unsintered hydroxyapatite/poly-l-lactide (bioresorbable plates).

**Figure 6.** Load values of the titanium and unsintered hydroxyapatite (u-HA)/poly-l-lactide (PLLA) bioresorbable plates according to the amount of displacement in the lateromedial loading test.

**Figure 7.** Change in subcondylar fracture segment under lateromedial loading. Ti, titanium (plates); u-HA/PLLA, unsintered hydroxyapatite/poly-l-lactide (bioresorbable plates).

#### **4. Discussion**

– Bioresorbable materials have been used clinically for osteosynthesis in various fields of maxillofacial surgery such as orthodontic surgery [13], repair of craniomaxillofacial fractures [7], bone augmentation for dental implantation [14], and reconstruction in situations of maxillofacial cysts and tumors. In traumatology, bioresorbable plates are used mainly in the treatment of orbital and zygomatic fractures [15–17] and in the treatment of mandibular fractures [8,18]. However, there are very few clinical studies of the treatment of mandibular subcondylar fractures with resorbable plates, and the number of cases reported is often low [19]. Bioresorbable material, unlike titanium plates, generally cannot resist high mechanical loads. Therefore, for mandibular condylar fracture treatment, it is necessary to know how much load the u-HA/PLLA bioresorbable plates can withstand in comparison with titanium plates. We aimed to clarify the strength of the bioresorbable plate system and the titanium plate system and to examine the possibility of the treatment choice of the osteosynthesis material for mandibular subcondylar fracture.

First, the mechanical strengths of the titanium and u-HA/PLLA bioresorbable plate systems were examined. The strength of the u-HA/PLLA bioresorbable system, in either tension or shear, was only approximately 45% that of the titanium system. The u-HA/PLLA bioresorbable plate strength was improved slightly when the form of the plates was modified. This is comparable to previous research on the strength of titanium plates when the same material of the bioresorbable plate was used. However, our results were similar to those of previous research [20]. These were very important findings. For pure plate strength, these results suggest that the u-HA/PLLA bioresorbable plate system was much weaker than the titanium plate system.

However, the results of our biomechanical loading tests were very interesting. In the loading tests, we reproduced mandibular subcondylar fracture treatment with two straight miniplates along the ideal line of the mandibular condyle. In the lateromedial loading test, there was no significant difference between the u-HA/PLLA bioresorbable plates and the titanium plates. One of the most important clinical indications for surgical treatment of mandibular condylar fracture is the medial displacement

of the condyle [21], which is caused by traction of the lateral pterygoid muscle. The primary function of the lateral pterygoid muscle is to pull the head of the condyle out of the mandibular fossa along the articular eminence to make the mandible protrude [22]. The lateral pterygoid muscle allows lateral movement of the mandible. This test simulates the lateral movements of mastication, and the strength of the bioresorbable plate system was sufficient. In contrast, in the anteroposterior loading test, the load value at the initial displacements of 0.5 and 1.0 mm was significantly larger in the titanium plates than in the u-HA/PLLA bioresorbable plates; the load in the latter was approximately 80% that of the titanium plates. The u-HA/PLLA bioresorbable plates were much weaker than the titanium plates. However, proper plate placement for mandibular subcondylar fracture treatment greatly improved the strength of u-HA/PLLA bioresorbable plates.

Champy and Lodde [23] advocated "functionally stable osteosynthesis" or "dynamic osteosynthesis," according to which the plates must be placed along physiological tension lines that appear during the fracture. This concept was based on the theory that stabilization depends mainly on the compression of the fracture site until function is restored. Meyer et al. [11] demonstrated the presence of the compression line on the mandibular posterior border and the tension line on the mandibular notch in the mandibular condylar process. One posterior plate is located on the compression strain lines when placed along the condylar neck axis. This plate enables functionally stable osteosynthesis. The other plate, which enables dynamic osteosynthesis, should be on the traction line. To conform to dynamic osteosynthesis principles, the plate should be placed higher and more obliquely, parallel to the mandibular notch [11]. In short, the plate located at the traction line resists the dynamic force of the occlusion. The key to anteroposterior direction load testing is the presence of plates on the traction line.

In our experiment, as the anteroposterior load increased, the bone opening on the traction line increased. This opening was influenced by the strength of the plate on the traction line and was thought to cause the difference in load values of the titanium plates. Load values at displacements of 1.5 mm or more in the bioresorbable plates did not differ significantly from those in the titanium plate group. Therefore, we believe that the bioresorbable plate can serve clinically in osteosynthesis of mandibular subcondylar fractures. However, the mechanical model used for this study was a polyurethane model, which is softer than the actual mandible [9]. Hence, considering a slight decrease in the biomechanical strength of the bioresorbable model compared with the titanium equivalent, in clinical use of the bioresorbable plate system, early overloading soon after surgery should be avoided. A young patient with strong occlusal force also needs dietary instruction, such as eating soft foods after surgery.

At present, resorbable plate systems are used widely, and various resorbable plate systems have been developed on the basis of PLLA. Like the u-HA/PLLA plate system used in this study, these plate systems composed of various materials, including pure PLLA [15] or polyglycolic acid/PLLA [24], are commercially available. In particular, the u-HA/PLLA bioresorbable plate system that we studied has a very large biological advantage. Resorbable plates composed of pure PLLA are very strong [25], however, the strength of plates is reduced when PLLA is mixed with other materials. Unfortunately, PLLA osteosynthetic devices have several disadvantages, including lower dynamic strength, an inability to bond directly to bones, unstable resorption and decomposition processes, and long replacement times [26–28]. In contrast, u-HA/PLLA osteosynthesis devices significantly improved biological activity. It has been reported that u-HA/PLLA material induced no apparent inflammatory or foreign body reactions after implantation, and bonded directly to the human bone quickly [14]. Clinical studies have also clearly shown the effective osteoconductivity of the u-HA/PLLA plate system in maxillofacial treatment [8,29]. In addition, the reduction in strength due to the use of composite material was overcome by the unique compression molding and machining treatment in this osteosynthesis device [30]. The early bioactivity, such as osteoconductivity and direct bone bonding capacity, can produce early functional improvements in maxillofacial osteosynthesis of mandibular condylar fractures. However, to clarify the relationship between loading and mandibular condylar

fracture healing with this plate system, it is necessary to evaluate the tissue over time. Further research reports are expected in the future.

The approach to mandibular condylar fracture is generally an extraoral approach [4]. Therefore, when it is necessary to remove the plate, there is a risk of recurrent facial nerve damage [31]. This is greatly stressful for both the surgeon and the patient. The bioresorbable system applied clinically may be an important option for the treatment of mandibular condylar fractures in the future.

The search for improved methods of subcondylar fracture fixation devices has been the subject of several studies. Mechanical loading tests are used to evaluate the behavior of fixation devices and methods, allowing the study of different osteosynthesis constructs. However, the present study did not involve a real mandible, and it will be necessary to conduct further research on mechanical experiments, as well as studies from various clinical viewpoints.

#### **5. Conclusions**

The titanium-based fixation system was much stronger than that based on bioresorbable u-HA/PLLA. In the biomechanical evaluation of mandibular condylar fracture treatment, however, our results showed that the u-HA/PLLA bioresorbable plate had sufficiently high resistance in the two-plate fixation method. The titanium plates were more resistant than the bioresorbable plates in subcondylar fractures when the initial force was applied in the anteroposterior direction; conversely, in the medial direction, there were no statistically significant differences. These results suggest that titanium and bioresorbable plate fixation systems have similar mechanical resistance within the limitation of our mandible model study.

**Author Contributions:** S.S. participated in the design of the study and experiments, acquisition of data, and drafted the manuscript. T.K. performed and coordinated the study and contributed to the drafting of the manuscript. N.Y. analyzed the data and contributed to the drafting of the manuscript. H.N. and K.N. performed and coordinated the study and contributed to the drafting of the manuscript. K.T. and H.K. made many contributions to the experiment and the correction of manuscript of this paper. Y.F. conceived the study and participated in its design. All authors approved the final version of the manuscript prior to submission.

**Funding:** This work was supported by JSPS KAKENHI Grant Number 20837661.

**Acknowledgments:** We declare no financial support from Teijin Co., Ltd., Osaka, Japan. Some analyses, such as those of tensile and shear strength and biomechanical loading evaluations, were performed at the Teijin Medical Institute, Okayama, Japan.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **In Vitro Degradation of Absorbable Zinc Alloys in Artificial Urine**

**Sébastien Champagne 1,2 , Ehsan Mostaed 3 , Fariba Safizadeh 1 , Edward Ghali 1 , Maurizio Vedani 3 and Hendra Hermawan 1,2,\***

Received: 24 December 2018; Accepted: 16 January 2019; Published: 18 January 2019

**Abstract:** Absorbable metals have potential for making in-demand rigid temporary stents for the treatment of urinary tract obstruction, where polymers have reached their limits. In this work, in vitro degradation behavior of absorbable zinc alloys in artificial urine was studied using electrochemical methods and advanced surface characterization techniques with a comparison to a magnesium alloy. The results showed that pure zinc and its alloys (Zn–0.5Mg, Zn–1Mg, Zn–0.5Al) exhibited slower corrosion than pure magnesium and an Mg–2Zn–1Mn alloy. The corrosion layer was composed mostly of hydroxide, carbonate, and phosphate, without calcium content for the zinc group. Among all tested metals, the Zn–0.5Al alloy exhibited a uniform corrosion layer with low affinity with the ions in artificial urine.

**Keywords:** absorbable; corrosion; degradation; magnesium; ureteral stent; zinc

#### **1. Introduction**

In 2008, up to five in 10,000 Canadian adults suffered from acute urinary tract obstruction, a kidney-related disease contributing to the economy burden of Canada by \$3.8 billion [1,2]. It is a blockage that limits the flow of urine out of the body, which can be caused by a variety of factors such as a kidney stone, a swollen prostate, or a tumor [3]. To quickly relieve the kidney of the built-up of urine, urologists will mechanically open the urinary tract using a stent. An ideal ureteral stent will maintain an excellent urine flow to optimize upper tract drainage, be resistant to infection and encrustation (calcification), and be absorbable [4]. The majority of urologists consider the placement of stents after ureteral dilation to be routine. Some studies showed that about 60–85% of urologists maintained the stent for less than seven days [5,6]. This temporary need has brought the concept of biodegradable ureteral stent. Introduced in the late 1980s, their purpose is especially for temporary treatment such as for patients waiting for the effect of a medical therapy [7]. Poly(glycolic acid), poly(lactic acid), and poly(lactic-*co*-glycolic acid) are among the biodegradable polymeric materials used for ureteral stents [4,8]. However, polymeric stents possess a limited ability to resist external compression forces such as those created in a malignant extrinsic ureteral obstruction. This is due to their intrinsic low strength, which often results in their reinforcement by a metal skeleton made of non-degradable alloys [9,10]. At this point, metals that can corrode would constitute the ideal materials for biodegradable ureteral stents.

Nowadays, with the recent development in new biomaterials technology, namely absorbable metals, this ideal ureteral stent should be feasible. Absorbable metals, also known as biodegradable metals, are metals that corrode gradually in vivo with an appropriate host response, then dissolve completely while assisting tissue healing [11]. Iron (Fe), magnesium (Mg), zinc (Zn), and their alloys are among the studied absorbable metals, mostly for cardiovascular and orthopaedic applications. Coronary stents made of Fe and its alloys have shown their safety and efficacy when tested in animals [12], while those made of Mg alloys were clinically tested in humans, showing a continuous favorable safety profile up to 12 months [13]. Zn and its alloys have been more recently proposed as alternatives to Fe and Mg. In recent in vivo studies, pure Zn wire showed a steady corrosion rate for up to 20 months post-implantation in murine artery without causing local toxicity [14], while pure zinc stents corroded in rabbit abdominal aorta without obvious accumulation of corrosion products even after 12 months of implantation [15].

Application of absorbable metals in urology is a novel avenue. In 2014, Lock et al. [16] conducted a study on the degradation and antibacterial properties in artificial urine of Mg, Mg–Y, and Mg–3Al–1Zn (AZ31) that was the first to show the potential of absorbable metals for urological applications. More recent work by Zhang et al. [17,18] reported that implantation of Mg, Mg–6Zn, and Mg–5.4Zn–0.5Zr (ZK60) showed no adverse effect in the surrounding rat's bladder tissue and no sign of toxicity towards the liver and the kidneys. Their degradation tests demonstrated that the pure Mg consistently corroded faster than the alloys in simulated body fluids and artificial urine and in the bladder of rats. However, those previous corrosion studies did not assess the formation of the corrosion layer. In reaction with the ionic content of urine, the surface of the ureteral stent can be easily calcified. Known as encrustation, this is one of the main problems currently faced with respect to ureteral stents [6]. Therefore, this work aims at revealing a detailed in vitro degradation behavior of Zn alloys under simulated urinary tract conditions by using an electrochemical method and advanced surface characterization techniques.

#### **2. Material and Methods**

#### *2.1. Material and Specimen Preparation*

A group of binary Zn alloys (Zn–0.5Mg, Zn–1Mg, and Zn–0.5Al) was selected based on their excellent mechanical properties as determined in our previous study [19]. Hot rolled pure Zn (99.995%) and pure Mg (99.94%) and commercially available extruded Mg–2Zn–1Mn alloy were used as comparative materials. The Zn alloys were melted at 500 ◦C in a cylindrical steel mold inside a resistance furnace (Nabertherm, Lilienthal, Germany) and cast as cylindrical billets of 80 mm length and 15 mm diameter. The billets were annealed at 350 ◦C for 48 h, water-quenched, and extruded to a final diameter of 8 mm as described in detail elsewhere [19]. Specimens for corrosion tests were cut from the extruded metals at a thickness of 4 mm using a slow speed diamond rotating blade (Isomet 1000 Precision Saw, Buehler, Uzwil, Switzerland). The specimens were then polished using SiC paper from #600, to #800, and to #1200 (Carbimet 2 Abrasive Paper, Buehler, Uzwil, Switzerland).

#### *2.2. In Vitro Degradation Testing*

Four methods of electrochemical corrosion tests were conducted: open circuit potential (OCP), electrochemical impedance spectroscopy (EIS), potentiodynamic polarization (PDP), and electrochemical noise (EN). The metal specimens were mounted as working electrodes in a waterproof chamber with an exposed area of 0.096 cm2 (0.35 cm of diameter), and all tests were done in triplicate for statistical analysis. The first three methods were performed one after another on all metal specimens using a three-electrode cell configuration. The metallic samples served as the working electrode, a graphite rod as the counter electrode, and a saturated calomel electrode (SCE) (E = +0.241 V saturated) as the reference electrode. The electrodes were connected to a VersaSTAT3 potentiostat and monitored using the VersaStudio software (Ametek Princeton Applied Research, Oak Ridge, TN, USA). The OCP experiments were conducted by monitoring the potential versus the SCE without applying any outside source for 3600 s, until equilibrium was reached at the corrosion potential Ecorr. Once the equilibrium was achieved, EIS was started at the amplitude of 10 mV RMS and a frequency scan from 100 kHz

to 1 Hz for all specimens. The interpretation of EIS results was made with the help of ZSimpWin software (Ametek Princeton Applied Research, Oak Ridge, TN, USA). The PDP experiments were done following the EIS, using a scan rate of 0.6 V/h, from −0.25 to +0.6 V vs. OCP.

The EN experiments were conducted on pure Mg, pure Zn, and Zn–0.5Al specimens only. Two identical working electrodes (mounted in resin) and a saturated silver chloride electrode (E = +0.197 V saturated) reference electrode were used and connected to a PC4/750 galvanostat/potentiostat (Gamry Instruments, Warminster, PA, USA). Data processing was done using an electrochemical signal analyzer (ESA400, Gamry Instruments, Warminster, PA, USA). The EN analysis was performed independently from the other electrochemical tests, as different equipment was used. The EN experiments were conducted for 48 h and at a frequency of acquisition of 10 Hz. The two identical coupled WEs were connected to a zero resistance ammeter (ZRA) to maintain the bias potential at 0 V. Each set of EN analysis, containing 1024 data points, was recorded with a data sampling rate of 0.1 s.

The artificial urine solution was made from analytical grade chemicals (Sigma-Aldrich, St. Louis, MO, USA) following the composition presented in Table 1. The pH of the artificial urine solution was adjusted to 6.0 using a solution of 1 N NH4OH and monitored with a pH meter (Accumet pH meter 25, Fisher Scientific, Portsmouth NH, USA). To mimic the ureteral condition, the solution was kept at a normal human body temperature of 37 ± 1 ◦C using a water heater and a jacketed beaker. It was stirred with a magnetic stirrer at 80 rpm. The artificial urine was chosen over natural urine to allow more reproducible results. The natural variability of the chemical composition of urine is greatly inconsistent between individuals because of the different diet and physical activity level, and it also depends on the time of the day, with more concentrated urine in the morning [20].

**Table 1.** Artificial urine composition [21].

Brought to 1 L by adding distilled water.

#### *2.3. Surface and Corrosion Layer Characterization*

The surface morphology observation and chemical composition analysis of the specimens after PDP were done using a scanning electron microscope (SEM, Quanta 250 FEI, Hillsboro, OR, USA) coupled with an energy dispersive X-ray spectroscope (EDS, SwiftED 3000, Oxford Instruments, Concord, MA, USA). The specimens were rinsed with distilled water and dried in desiccator for one day before being subjected to SEM observation. An X-ray photoelectron spectroscope (XPS, PHI 5600-ci spectrometer, Physical Electronics, Chanhassen, MN, USA) was then used to further determine the nature of the corrosion products at the surface of the corroded samples. A neutralizer was used to negate the charging effect at the surface of the metals. An XPS spectra survey was recorded from 1400 eV and ending at 0 eV using a standard aluminum anode (1486.6 eV) at 300 W. High-resolution spectra were then collected using a standard magnesium anode (1253.6 eV) at 300 W aimed at the following elements: carbon (300–280 eV), oxygen (545–525 eV), zinc (1037–1017 eV), phosphorus (147–127 eV), magnesium (64–44 eV), and, finally, aluminum (88–68 eV). The C1s peak was chosen as a reference at 285.0 eV to negate the effects of charging. Both survey and high-resolution scans were done in triplicate for each metal.

#### **3. Results and Discussion**

Once immersed in the artificial urine solution, all metals tended to stabilize after 1 h (Figure 1a). Pure Mg and Mg–2Zn–1Mn showed a more negative open circuit potential (−1.76 and −1.52 V, respectively) compared to all Zn groups (between −1.15 and −1.05 V with pure Zn having the highest value). Under potentiodynamic polarization, the Zn group shows a more noble behavior compared to the Mg group (Figure 1b). Corrosion rates vary from 2.16 mm/year for pure Mg to 0.87 mm/year

for Zn (Table 2). The Mg–2Zn–1Mn corrodes slightly slower than pure Mg at 1.90 mm/year due to alloying effects. The corrosion of magnesium is highly dependent on the impurities that can act as active cathodic sites to accelerate corrosion. Zn is nobler than Mg and may act as an impurity, but the addition of Mn increases the tolerance limit of impurity's content by reducing the grain size of the alloys, thus lowering the corrosion rate of Mg–2Zn–1Mn compared to that of pure Mg [22–24]. In the Zn group, corrosion rate increases from pure Zn to Zn–0.5Al and Zn–Mg (Table 2). The addition of Mg into Zn produces an Mg2Zn11 phase that creates microgalvanic action with the Zn matrix, which in turn increases the corrosion rates of Zn–Mg alloys [25]. Meanwhile, the high solubility of Al in Zn produces a single phase microstructure in the Zn–0.5Al, so its corrosion rate only slightly increases [19].

**Figure 1.** Typical open circuit potential (OCP) and potentiodynamic polarization (PDP) results: (**a**) potential vs. time, (**b**) potential vs. current.

−

**Table 2.** Corrosion parameters derived from PDP results.

‐ ‐ After the polarization test, the surface of Mg group specimens is covered by a fairly uniform but not seemingly compact corrosion layer as indicated by the presence of cracks (Figure 2a,b). A pure Zn surface presents a uniform and dense corrosion layer (Figure 2c), as it tends to passivate in a near neutral aqueous solution, such as artificial urine [26]. Both the investigated Zn–Mg alloys show a non-uniform corrosion layer with the appearance of a flower-like crystal on the surface. The flower size is larger in Zn–0.5Mg (Figure 2d) with a seemingly compact underlying layer compared to that of Zn–1Mg (Figure 2e), which could explain the slightly faster corrosion rate of the latter. The flower is not observed on the Zn–0.5Al surface, but a rather uniform, smooth, and dense corrosion layer (Figure 2f).

**Figure 2.** SEM images of the specimens' surfaces after polarization: (**a**) pure Mg, (**b**) Mg–2Zn–1Mn, (**c**) pure Zn, (**d**) Zn–0.5Mg, (**e**) Zn–1Mg, and (**f**) Zn–0.5Al.

‐ The distinct difference on the corrosion layer's composition between the Mg group and the Zn group is the calcium content, which was only detected on that of the former group with a concentration range of 13–15 wt % (Table 3). The Mg group's composition has similarity to that of Mg alloys reported in many non-urinary in vitro and in vivo studies, and it generally forms hydroxide, phosphate, and carbonate [23,27–31]. Apart from the lack of calcium, the EDS spectra of the Zn group corrosion layer

‐

showed a lower content of oxygen and phosphorous, but a higher carbon content as compared to the Mg group. A focused EDS analysis on the flower-like formation on the surface of pure Zn and Zn–Mg alloys (Figure 2c–e) reveals a high content of zinc and oxygen, which supposedly forms ZnO, as later confirmed by XPS analysis. On the very top layer, as measured by the XPS, both groups have a similar content of carbon and oxygen, but a higher phosphorous content for the Zn group. Apart from the aluminum content, the surface of the Zn–0.5Al alloy has the lowest content of those elements, as measured by both techniques. Previous non-urinary in vitro and in vivo corrosion studies of pure Zn generally confirmed the formation of zinc oxide, zinc carbonate, zinc phosphate, and CaP on the very top layer [14,15,19,32,33]. The presence of calcium on the corrosion layer may initiate calcification and thus encrustation, a main current problem of ureteral stents [6], leading to the formation of urinary stones. These stones are composed mainly of calcium oxalate (CaOx), struvite (NH4MgPO4·6H2O), calcium phosphate (CaP), and uric acid (C5H4N4O3) [34,35].

**Table 3.** Elemental composition of the corrosion layer detected by EDS and XPS.

Further XPS analysis allowed for the observation of two Mg2p peaks at 51.15 and 49.66 eV on pure Mg and at 51.19 and 50.07 eV for Mg–2Zn–1Mn (Figure 3a), which correlate to MgCO3 and Mg(OH)2, respectively [36,37]. Three C1s peaks (Figure 3b) are also detected on these specimens at 285 eV for C–C, 286.6 eV for C–O(H) [38], and 289.7 eV for carbonate [37]. Three O1s peaks (Figure 3c) are present at 530.15 eV related to MgO or Mg(OH)2, 531.7 eV for the double binding of P=O [39], and at 533.3 eV related to O–C=O of CO3 species such as MgCO3 or CaCO3 [40]. Additionally, two P2p peaks were observed at 133.6 and 135 eV (Figure 3d), which correspond to MgHPO4 and Mg3(PO4)2, respectively [41].

Regarding the Zn–0.5Al alloy, the XPS analysis demonstrates the presence of three C1s peaks related to C–C at 285 eV, carbonate at 289.15 eV, and C–O(H) at 286.89 eV (Figure 4a). The first O1s peak was observed at 532 eV and corresponds to Zn(OH)2 [42], the second at 529.6 eV is related to ZnO [42], and the third at 533 eV fits to carbonate (Figure 4b). Two P2p peaks were observed at 133.47 and 134.31 eV (Figure 4c), characterized as PO3− 4 and HPO2− 4 [41], which confirms the presence of a phosphate compound in the corrosion layer. One peak was observed for Zn2p3 at 1022.8 eV (Figure 4d), which corresponds to Zn(OH)2 [42]. An additional peak of Al2p (not shown) was found at 74.76 eV for Al2O3 [43]. Overall, the XPS analysis confirms that no calcium-related compound was formed on the surface of Zn–0.5Al, which makes this alloy less susceptible to calcification compared with pure Mg and the Mg–2Zn–1Mn alloy.

**Figure 3.** High resolution XPS spectra for pure Mg specimen: (**a**) Mg2p peaks, (**b**) C1s peaks, (**c**) O1s peaks, and (**d**) P2p peaks.

**Figure 4.** High resolution XPS spectra for Zn–0.5Al specimen: (**a**) C1s peaks, (**b**) O1s peaks, (**c**) P2p peaks, and (**d**) Zn2p3 peak.

‐

‐

The corrosion layers formed on the metals' surface exhibit a different impedance behavior, as indicated by their Nyquist plots (Figure 5a). The Mg group's plots form a semi-circle (one time constant), whilst those of the Zn group show two semi-circles (two time constant). Looking at the Bode plots, in the Mg group the phase angles increase from about 10◦ to 55◦ at a high frequency and then decrease to about 25◦ at a lower frequency (Figure 5b), indicating a capacitive loop [44]. Meanwhile, the plots for the Zn group show two peaks, one at a high frequency and another at a low frequency, except for the Zn–0.5Mg where the second peak is not that apparent. In the other Bode plots, all specimens show similar evolution from low impedance at high frequency to high impedance at low frequency, with the Mg group having the highest impedance (Figure 5c), indicating the corrosion layer exerts an electric effect between a resistor and a capacitor.

**Figure 5.** Typical electrochemical impedance spectroscopy (EIS) results: (**a**) Nyquist plots, (**b**) Bode phase angle plots, (**c**) Bode impedance plots, and (**d**) proposed equivalent electrical circuit models. Note that at very low frequency the Mg specimens show some negligible scatter.

‐ ‐ ‐ By applying a single-frequency voltage to the interface, EIS allows a measurement of the electric properties of the corrosion layer, including a phase shift, a real and imaginary part of the impedance and the amplitude of the resulting current at that particular frequency [45,46]. All these properties can be simplified as an electrical circuit with each electrical component representing one of the elements of the electrochemical cell. For the Mg group, their semi-circle plots indicate the formation of a typical uniform oxide layer (Figure 2a,b). Its characteristic can be fitted with a simple electrical circuit (Model 1 (Figure 5d)) composed of the Rs (solution's resistance), Rct (charge transfer resistance), and CPEdl (constant phase element) of the double layer at the electrolyte–surface interface. The CPE represents a non-ideal electrical element in the circuit; in this case, it behaves as an imperfect capacitor and is very useful for assessing its dispersive effect at the surface of the electrode [47]. For a corroding metal, this effect is attributed to the microscopic increase in the surface roughness of the electrode [48]. When Rct

‐

‐

is the only resistance in the circuit, it can often be substituted with Rp (polarization resistance) and it is inversely proportional with the corrosion rate. For the Zn group, its oxide layer characteristic is fitted with the Model 2 circuit that includes a CPEol for the capacitive properties of the oxide layer, considering that Zn is less reactive than Mg and formed a rather non-uniform corrosion layer (Figure 2c–f). A more complex electrical circuit can be fitted to all specimens, representing a bi-layered surface with outer CPEol and Rol (Model 3). The calculated value of each electrical component from the equivalent circuits is presented in Table 4.

**Table 4.** EIS-derived parameters (R in Ω cm2 and CPE in µF).

The solution's resistance, Rs*,* is in the constant range of 5–10 Ω throughout all the experiments. As shown in the Model 1 and Model 3 circuits, the Mg group exhibits a higher Rct compared to the Zn group, conversely to their CPEdl. The addition of Mn may play a role in improving the coating resistance of the Mg–2Zn–1Mn by increasing the passivity of the corrosion layer [28]. Both Zn–0.5Mg and Zn–1Mg exhibit similar electrical values, so in EIS analysis they are treated as one metal. The presence of a large value of CPEol (Model 2) indicates that a preferential corrosion occurred on the surface of the Zn–Mg alloys, leading to the creation of micropores on the corrosion layer (Figure 2e). The CPEdl value of the Zn group is higher than those of the Mg group due to the higher capacitive properties of Zn, which produces fewer active ions than magnesium, and this could be the reason for the slower degradation rate of the former [49]. The value decreases from pure Zn to Zn–0.5Al and to Zn–Mg, meaning more charge should be accumulated at the electrode-electrolyte interface of pure Zn since no element acts as a captor of charged ions, i.e., Mg in Zn–Mg and Al in Zn–Al in the process of localize corrosion. In the Model 3 circuit, Rct and CPEdl are affected by the interaction of different layers, which in turn changes the behavior of both components. The Rol value is much higher for the Mg group, which could be correlated to a thicker oxide layer (Figure 2a,b). The capacitive properties of the oxide layer are bound with CPEol. In this bi-layer model, the oxide layer of pure Zn and Zn–0.5Al has dielectric properties, which could be influenced by a thick and dense corrosion layer governed by the diffusion phenomenon (Figure 2c). However, this may not be the case for Zn–Mg alloys as the layer becomes less uniform due to the localized corrosion (Figure 2d,e).

The electrochemical noise data were analyzed in the time domain with the calculation of the noise resistance Rn. It is well established that Rn is inversely proportional to the corrosion rate [50]. This helps to monitor the corrosion rates after immersion of the specimens in the urine solution. Pure Mg, pure Zn, and Zn–0.5Al were used as samples for the EN investigation. Up to 12 h of immersion, pure Mg exhibited the highest corrosion rate followed by Zn–0.5Al and pure Zn (Figure 6), a similar trend was obtained by the accelerated PDP test. Pure Mg experienced a stabilization after 12 h, so, for longer periods, pure Mg possesses the lowest corrosion rate than the other two metals.

**Figure 6.** Electrochemical noise (EN)results presented as 1/Rn vs. time for pure Mg, pure Zn, and the Zn–0.5Al alloy over a period of 48 h.

In artificial urine, pure Mg dissolves quickly. Due to the high affinity of Mg ions with the ions in the solution, an insoluble corrosion layer (i.e., Mg3(PO4)3) forms and protects the metals underneath from further corrosion [18]. This thick layer (Figure 7a) acts as a more effective barrier than Mg(OH)2 or Zn(OH)2 on pure Zn [17]. A rather smooth surface was observed on the latter (Figure 7b). The Zn–0.5Al is prone to microgalvanic corrosion due to the potential difference of the Zn and Al phases (Figure 7c).

**Figure 7.** Surface morphology after 48 h of the EN test of (**a**) pure Mg, (**b**) pure Zn, and (**c**) Zn–0.5Al.

#### **4. Conclusions**

‐ ‐ ‐ ‐ In artificial urine, the studied pure Mg and its alloy show a corrosion behavior distinct from that of pure Zn and its alloys. Based on the accelerated potentiodynamic polarization test, the Mg group exhibits a higher corrosion rate than does the Zn group. The corrosion layer of the Zn group has a more capacitive behavior than that of the Mg group. A flower-like structure of Zn oxide forms on the polarized surface of the Zn group, except for Zn–0.5Al, while precipitation of Ca and Mg was observed in the Mg group only. In the long-term electrochemical noise test, pure Mg exhibits stabilization after 12 h and thus has an overall lower corrosion rate than does pure Zn and Zn–0.5Al. The continuous metal dissolution of Zn–0.5Al and the low affinity to ions in artificial urine prevents the formation of a thick corrosion layer. In view of ureteral stent applications, a continuous dissolution of metal in urine, such as the case of the Zn–0.5Al alloy, is ideal to prevent calcification (stent encrustation) and bacterial adhesion on the surface.

**Author Contributions:** Conceptualization, H.H. and S.C.; Data Curation, S.C.; Formal Analysis, F.S., H.H. and S.C.; Methodology, F.S., H.H. and S.C.; Resources, E.M.; Supervision, H.H.; Writing—Original Draft Preparation, S.C.; Writing—Review & Editing, E.G., E.M., F.S., H.H., M.V. and S.C.

**Funding:** This research was funded by Natural Sciences and Engineering Research Council of Canada (NSERC) via Discovery grant.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

*Review*

## **The Effect of Surface Treatments on the Degradation of Biomedical Mg Alloys—A Review Paper**

**Marcjanna Maria Gawlik 1,\*, Björn Wiese 1 , Valérie Desharnais 1,2, Thomas Ebel 1 and Regine Willumeit-Römer 1**

Received: 19 November 2018; Accepted: 14 December 2018; Published: 16 December 2018 -

**Abstract:** This report reviews the effects of chemical, physical, and mechanical surface treatments on the degradation behavior of Mg alloys via their influence on the roughness and surface morphology. Many studies have been focused on technically-used AZ alloys and a few investigations regarding the surface treatment of biodegradable and Al-free Mg alloys, especially under physiological conditions. These treatments tailor the surface roughness, homogenize the morphology, and decrease the degradation rate of the alloys. Conversely, there have also been reports which showed that rough surfaces lead to less pitting and good cell adherence. Besides roughness, there are many other parameters which are much more important than roughness when regarding the degradation behavior of an alloy. These studies, which indicate the relationship between surface treatments, roughness and degradation, require further elaboration, particularly for biomedical Mg alloy applications.

**Keywords:** surface treatments; roughness; Mg-alloys; degradation behavior

#### **1. Introduction**

The study of Mg as degradable biomaterial for implants is an advanced research area. A second operation to remove the implant after bone healing can be avoided [1–6]. Mg is naturally available as trace element in the body, and is thus non-toxic and biocompatible [7–9]. Implant processing is feasible due to the ductility and workability of Mg [10]. Strength and toughness are higher than of polymer implants, which is beneficial for load-bearing implants [11,12]. Mg alloys are reported to show improved osseointegration and bone implant strength compared to permanent Ti alloys [13,14]. In particular, Mg alloys are suitable as biodegradable implant materials [1,15–19]. Mg is able to degrade in aqueous solutions with the formation of magnesium hydroxide and hydrogen [20–24]. In particular, aqueous salt solutions containing ions including chlorides or sulphates, with the exception of alkali metals or alkaline metal containing solutions, are able to dissolve the protective magnesium hydroxide layer, leading to enhanced degradation [24–28]. In order to improve the mechanical properties of Mg, elements are added to tailor, for example, its tensile strength and ductility. Thus, it is possible to produce implants that have tailored mechanical properties to use it as temporary bone fixation. However, when alloying and processing the material, impurities like Fe, Ni, and Cu or phases with a high electrochemical potential difference are found at or near to the surface of the material, which increases the degradation rate through galvanic corrosion [23,29]. For the application of biodegradable Mg implants to become feasible in the future, two different objectives must be met in order to achieve usable degradation behavior.

One objective is limiting the degradation rate of the initial state of the alloy, which, as explained later, is related to the amount of hydrogen evolution. The deeper and rougher the surface morphology, the more hydrogen gas will be produced [9]. An overly fast degradation with gas evolution in the initial state leads to degradation of the mechanical integrity. Excessive gas evolution can also modify the bone remodeling process and impair the consolidation of bones [30]. However, relatively strong hydrogen evolution is crucial for cell adherence and implant-bone integration [21,31,32]. Aqueous salt solutions including chloride ions, like those found in the human body fluids [33], increasing decomposing of Mg(OH)2, release OH− and raise the pH [34]. Besides hydrogen production, a local alkalization might provoke necrosis [9].

The second objective is to control the degradation rate of implants during the healing time. The required degradation rate depends on the application with lifetime and stability of the implant and the potential of the surrounding tissue to tolerate pH changes and high ion concentrations. It is reported that the properties of the material, e.g., crystallographic orientation [35–37], microstructure [21,38–48], grain size [41,49–53], secondary phases [51,54,55], contamination [38,40,56], and deformation [38,41,57–60], affect the degradation behavior, as well the environment, e.g., the immersion medium [61–64]. It is possible to control the degradation behavior of Mg alloys using chemical, physical, and mechanical surface treatments [27,32,65–69]. Additionally, surface uniformity has been shown to decelerate degradation [70]. Surface morphology can differ despite identical roughness parameters, and also affects the degradation process [39]. Studies have shown that surface roughness can affect the initial degradation [71], the degradation rate [38,71–77], degradation resistance [73,78–82], pitting behavior [38,71,72,83], bone integration [84–86], cell adherence [21,74,87,88], cell proliferation [88–91], and cell differentiation [92]. Besides roughness, surface unevenness can also influence the adhesion of cells [73]. In some cases, a smoother surface will reduce the degradation rate [72,74,93]. However, this behavior has been contradicted in other studies [73,79,81,82,94,95].

The aim of this review is to show the correlation between surface treatment, roughness, and the degradation behavior of Mg alloys in order to define meaningful roughness values and suitable surface treatments for biodegradable Mg implants. An overview of studies mentioning surface treatments, roughness, and degradation is given in Tables 1–5.

#### **2. Mechanical Surface Treatments**

#### *2.1. Grinding and Polishing*

The degradation behavior of sand-cast, ground, and polished AZ91 alloys were investigated by Walter and Kannan [96]. The use of a grinding paper with increased grit size decreased the surface roughness (Table 1, Ref. [96]). Three methods were used to evaluate the relationship between pitting and roughness: a 24 h immersion test in a 0.5 wt.% NaCl solution, 1 h of potentiodynamic polarization (PDP) and 1 h of Electrochemical Impedance Spectroscopy (EIS). For the polished samples, no inductive loop after EIS was observed. A low inductive loop is related to a low or negligible amount of surface pitting. This was found for all ground samples. Thus, it is suspected that no pitting will occur on polished surfaces due to a higher passivation. As a consequence, passivation is reduced for higher surface roughness values. The polarization curves in Ref. [96] show that a higher anodic current, as indicated by the current density icorr, is produced with greater surface roughness (Table 1, Ref. [96] and Figure 1).

In particular, the polarization curve of material ground with 320 grit paper in Ref. [96] exhibits a strong increase in anodic current, which suggests a high number of pits being formed. The surface appearance was analyzed by scanning electron microscope (SEM) after immersion for 24 h and after galvanostatic testing. Numerous pits were observed after immersion when the 320 grit size paper was used, which confirms the results from the electrochemical testing. Less pitting was seen to occur when using a finer grinding paper. In the case of paper with a 1200 grit size, no localized pitting was found after testing. Surprisingly, more pits were observed on polished samples after galvanostatic testing. This is likely to have been caused by a high anodic current which reduces passivation for all surface treatments. Walter and Kannan [96] concluded from these experiments that roughness does affect the passivation layer, but does not directly affect the likelihood of pitting. After removing the passivation layer, pitting occurred for all surface roughness values as seen by SEM [96].

**Figure 1.** Graph shows a linear relationship between Sa and icorr after grinding and polishing. R2 is the coefficient of determination, which assess the linear mutual dependence of x and y. R2 = 1 defines the highest linearity [97]. Sa and icorr values were obtained from [96].

Walter et al. [71] also investigated the correlation between the degradation and surface roughness Sa (the arithmetic mean height within a sample area, three dimensionally determined roughness) [98] for samples which were ground using 120 SiC grit size paper (Sa = 973 nm) and samples that had been ground using 2500 grit size paper, followed by polishing with a 3 µm diamond paste (Sa = 22 nm). The samples were cleaned with acetone and ethanol. The degradation behavior was characterized under simulated body fluid (SBF) using EIS. The results for both surface finishes exhibited similar tendencies. The ground and the polished samples showed a mid-frequency capacitive loop at the beginning of testing [71]. A mid-frequency capacitive loop corresponds to a passivation layer [96]. For the ground sample, a mid-frequency capacitive loop was observed for the first 2 h and was then followed by an inductive loop at low frequencies. The polished samples had inductive loops at low frequencies after 4 h, which confirms that passivation layers on smoother surfaces last longer [71]. In general, an inductive loop implies surface pitting [99]. The polarization resistance was present for a maximum of 3 h for the polished sample in contrast to a maximum of 2 h for rougher ground surface. This observation agrees with the assumption that polished samples have a higher passivation. Thus polishing samples reduces the degradation behavior, especially at the beginning of immersion. The SEM results support these findings. After 2 h immersion, a general degradation for both finishes was observed. After 6 h, the rougher surface clearly showed more pitting, while the few pits seen on the polished surface indicated the start of pit formation. Pitting was studied for both surfaces after 12 h immersion. The initial pitting of the ground surface had progressed further compared to the polished finish. Walter et al. [71] explained this observation as a local pH drop caused by deep valleys in the rough surface. Additionally, the passivation layer of the rough surface broke up earlier than the smooth surface [71].

Alvarez et al. [94] found that polished AE44 samples encouraged more pitting compared to semi-polished samples in an immersion test. At the beginning of degradation, polished samples exhibited a higher pit volume compared to semi-polished samples. However, semi-polished samples had higher pit radii. The smoother surface of the polished sample and its related chlorine absorption capacity is given as a possible explanation. This behavior is distinct from reports that report that rougher surfaces on steel [75,76] and aluminum [77] lead to faster degradation and more pitting. Walter and Kannan [96] suggested a change of the passivation layer provoked by shifting the local pH, initiated by aeration of the solution, as a reason for this behavior [96]. Pitting for both conditions leads to intergranular degradation after several hours. However, the start of intergranular degradation of the polished samples started earlier than for semi-polished samples [94].

Lorenz et al. [74] showed that surface roughness does not only influence the degradation resistance of pure magnesium; it also affects the cell (HeLa cells/GSP-C12 mouse fibroblasts) adhesion on the surface of Mg. For this study, discs were prepared with 600 paper grit size, a combination of 6 µm diamond paste, and an ethanol/glycerol solution. Sample cleaning was performed using an ultrasonic bath filled with ethanol for 3 min. In order to analyze the effects of surface morphology on cell adherence, one series of samples was immersed in 1 mol NaOH for 24 h and another series in modified simulated body fluid (M-SBF) at 37 ◦C for 5 d. Afterwards, the samples were flushed with ethanol and dried in air. The roughness increased after immersion in both solutions, but especially for the modified simulated body fluid (M-SBF) solution. pH measurements were also carried out on Mg samples degrading in a Minimum Essential Media (MEM) that included fetal bovine serum (FBS). A pH of 8.96 was observed for the M-SBF treated Mg samples and was higher compared to the other treatments after 2 h (Tables 1 and 4, Ref. [74]). The thicker Ca/Mg phosphate layer after M-SBF immersion does not protect the Mg sample due to its porosity, but the corrosion resistance increased by a factor of five compared to the untreated samples. In contrast, the cell density is higher compared to the polished and NaOH treated samples. The increase of the roughness by immersing in M-SBF improved the cell adhesion. The medium alkalization of the M-SBF samples is only suitable for short term applications. The smooth surface of the polished samples exhibits nearly no cell adherence and degrades very quickly. The passivation of Mg with NaOH reduces degradation, but cell adhesion is lower compared to M-SBF immersion [74].

Liu [61] compared the cell adherence of rolled pure Mg foils with an oxide layer and on ground foils without an oxide layer. He also studied the effects of roughness and degradation in Dulbecco's Modified Eagle's Medium (DMEM) and deionized (DI) water. The smoother ground samples varied by only 1.2% (oxide layer 13.6% and on ground foils 14.8%) in cell density from the rough oxidized samples (Tables 1 and 2, Ref. [61]). As such, it can be assumed, that surface roughness did not affect the cell density. No correlation between roughness and degradation rate was found, though ground samples in DMEM showed a slower degradation rate compared to oxidized samples. The opposite behavior was observed in DI water [61].

In contrast to the studies carried out by Liu [61] and Lorenz et al. [74], it was reported by Johnson et al. [21] that ground Mg-4Y samples demonstrated a better cell attachment than samples with an electrical discharged machined (EDM) surface. The roughness of the surfaces could be a possible explanation, as it was found that rough surfaces degrade faster than smooth surfaces [21,72,93]. Mg-4Y exhibits a contrary degradation behavior as pure Mg [61]. The ground surface leads to a lower mass loss in DI water which is opposite to the higher mass loss in DMEM (Tables 1 and 2, Ref. [21]). This effect is not thought to be due to roughness, but rather, from a different evolution of the pH under the different testing conditions.

Song and Xu [38] investigated the effect of tempering (HT), sandblasting, grinding, and etching on the degradation resistance of the alloy AZ31. Tempering and sandblasting reduced degradation resistance, while grinding or acid etching as a cleaning procedure decreased weight loss and hydrogen evolution. Heat treatment led to the precipitations of large Al-Mn-Fe particles which deteriorated the degradation resistance [38]. In addition to impurities, it is also known that roughness influences the degradation rate [100]. The roughness Ra (two-dimensionally determined roughness, arithmetic mean deviation of the roughness profile) and the hydrogen evolution of ground samples were very low compared to sandblasted samples (Tables 1 and 2, Ref. [38]). Sandblasting led to a very rough surface, accompanied by micro stresses in the surface layer. The Fe impurities rather than this surface roughening increased the degradation rate. Grinding the surface removes contaminations and leads

to a slower degradation rate. Though ground surfaces are smoother than etched, the effect on the degradation is not as significant as removing a significant amount of Fe [38].

Zhao and Zhu [6] investigated, in addition to collagen monomer concentration, pH, and assembly time, the influence of ground surfaces on collagen fibril formation. They tested different surface finishes and the collagen formation with subsequent cell attachment. They ground Mg and AZ31 discs with 180 (Ra = 1.89 µm), 800 (Ra = 0.29 µm) and 1200 (Ra = 0.15 µm) SiC paper and apply 50 µL of 200 µg/mL D-phosphate-buffered solution (DPBS) diluted collagen solution for 2 h on the samples. By SEM they observed the morphology of collagen fibers for every surface finish and alloy. A clear difference of collagen formation was visible from roughest to smoothest surface for both alloys. While the collagen formation on both alloys for both smoothest surfaces was comparable, the roughest surface of AZ31 showed less dense structure in contrast to Mg. The roughest surface of both alloys adsorbed the highest amount of collagen after 2 h, while the smoothest surface showed the lowest adsorbed amount. This trend was more distinct for Mg compared to AZ31. Cell attachment observations after 2 h lead to the assumption that cells were more attached on collagen treated and smoother samples in contrast to the roughest surface finish. The roughest surface finish with a more fiber woven structure and highest collagen adsorption also showed in another Fluorescent live/dead cell analysis that, after one day, the collagen structure on a rough surface is more detrimental to cell density, independent of the alloy [6]. Nudelman et al. [101], reported a correlation between collagen and cell attachment [101]. In contrast to Nudelman et al. [101], Zhao and Zhu [6] evidenced a decrease in cell density with higher collagen adsorption. For this reason, it is assumed, that more collagen does not always result in a stronger cell attachment. In reference to roughness, this study shows an effect on the collagen formation which influences the cell density indirectly [6].

#### *2.2. Burnishing*

A comparison between ground and burnished sample degradation was performed using hydrogen evolution, PDP, and white light interferometry on the AZ31B alloy [39]. Ground and burnished (dry and cryogenic) samples had a very similar roughness before degradation. Burnishing was performed using a severe plasticity burnishing (SPB) process. Cryogenic burnishing is distinguished from dry burnishing by the use of liquid nitrogen. After degradation, the morphology of the ground samples differed from the burnished samples. Thus, roughness will not be the only factor to influence the degradation behavior. In addition, grain size and basal texture had an influence on the uniformity and amount of degradation. Dry and cryogenic burnishing decreased hydrogen evolution over a 7 h period with respect to ground surfaces. Both burnishing processes lead to a smoother finish with reduced pit depth and pit volume compared to the 4000 grit size paper treatment immersion test. The PDP analysis shows the same trend for both burnished surfaces with a higher degradation resistance, indicated by a broader capacitive loop [39]. The crystallographic orientation and grain refinement has to be considered, as well as the surface roughness [38]. In this report, the surface roughness did not affect the degradation, in disagreement with the prediction of Song and Xu [38]. Moreover, a small grain size and a strong basal texture led to a higher degradation resistance [39].

**Table 1.** *Cont*.

**Table 1.** *Cont*.

1 MEM: Minimum Essential Media; 2 FBS: fetal bovine serum; 3 CD: Cell density; 4 DMEM: Dulbecco's Modified Eagle's Medium; 5 DR: degradation rate; 6 DI: deionized; 7 PDP: potentiodynamic polarization; 8 EIS: Electrochemical Impedance Spectroscopy; 9 PBS: phosphate buffered saline; 10 CR: corrosion rate; 11 α-MEM: MEM alpha modification Media; 12 SBF: simulated body fluid.

#### *2.3. Machining*

Turned, threaded and sandblasted Mg-0.8Ca samples were examined and tested in vivo [93]. Smooth (Ra = 3.65 µm) turned and threaded samples exhibited the best interlocking connection between the bone and implant. Rough (Ra = 32.7 µm) sandblasted rods degraded most rapidly with the highest number of visible gas bubbles. Turned surfaces led to the lowest gas evolution and decomposition in these studies [87,93]. Despite a similar integration of threaded and turned implants into the bone tissue, threaded implants showed a non-uniform bone resorption at the thread edges [93]. This is in agreement with the findings of Walter et al. [71], which may be explained by local variations in pH.

Mhaede et al. [102] reported a relationship between roughness and corrosion resistance. For the degradation test in 0.9 wt.% NaCl solution, eight different conditions of AZ31 alloy were prepared. Samples were either ground or shot-peened with 3 different Almen intensities (saturation value of residual arc height of an Almen strip, established by John Almen) [102,106,107], ground and coated, or shot-peened with 3 different Almen intensities and coated without prior grinding. Shot peening was performed with ceramic shot which had a diameter of 850 µm (Z850). The dicalcium phosphate dihydrate (DCPD) coating used was produced by electro-deposition of samples in a 0.1 mol Ca(NO3)2 4H2O + 0.06 mol NH4HPO4 solution. In Table 2, Ref. [102] it is shown that the current density icorr for the shot-peened (SP) samples was increased compared to the other conditions in Table 4, Ref. [102]. It has been shown that having a rough surface after shot-peening affects icorr (Table 2, Ref. [102] and Figure 2), as the resulting greater surface area increases the surface reactivity [102]. However, it is not possible to relay Ra linear to icorr (Tables 1, 2 and 4, Ref. [102] and Figure 2) for all surface finishes due to the protective properties of the DCPD coating compared to only shot peened samples. The linear relationship between Ra and icorr for the shot-peened and shot-peened/coated samples (Figure 2) agrees with the study of Walter and Kannan [96], whereas linear correlation was observed for only the ground samples. However, it should be noted that higher deformation and internal stress, arising from higher Almen intensities, could also affect the degradation behavior.

**Figure 2.** Diagram current density icorr against roughness Ra for ground, shot peened, and shot peened + DCPD coated samples. A non-linear R2 relationship between Ra and icorr is shown by comparing all conditions (ground/SP/SP+DCPD) together. Linear R2 is plotted for only shot peened or only shot peened and coated samples. A trend of linearity can be seen only for roughness values arising from same surface treatments. Ra and icorr values were obtained from [102].

Denkena and Lucas [108] studied the surface and subsurface properties after turning and deep rolling a Mg-3Ca alloy. Three different conditions for each machining process were investigated. With

regard to turning, the roughness Rz (distance from deepest valley to highest peak within sample length from a linear measurement) [109] decreased from around Rz ~ 4.48 µm to Rz ~ 3.75 µm after increasing the cutting speed from 10 m/min to 100 m/min at constant cutting depth and feed rate. By reducing the feed rate from 0.1 mm to 0.05 mm at a constant cutting speed (100 m/min) and cutting depth (0.5 mm), the roughness (Rz ~ 2.17 µm) was reduced (Table 2, Ref. [108]). After deep rolling with different rolling forces (F = 50 N, 200 N, 500 N) and constant rolling speed (25 m/min) and feed rate (0.1 mm), no significant change in roughness (Rz ~ 0.91–1.26 µm) occurred (Table 2, Ref. [108]). Degradation tests were performed in 0.9 wt.% NaCl solution and the hydrogen gas evolution was measured. The mass loss was calculated from the amount of hydrogen produced and a correction factor. The degradation rates for turning with higher roughness were greater compared to the degradation rates after rolling. However, for turning, the condition with the highest roughness showed the lowest mass loss compared to the smoother samples. For deep rolling, the condition with the lowest rolling force led to the highest, while not signifying mass loss (calculated from hydrogen generation) after around 240 h exposure time compared to conditions with higher rolling forces and comparable Rz values. No significant correlation between roughness and mass loss was found. High residual compressive stress was reported to reduce the degradation rate by about 100 times [108], and the degradation results were comparable with the results from high speed dry milled Mg-0.8Ca with the lowest roughness [110].

The influence of machining and deep rolling on Mg-3Ca and Mg-0.8Ca was analyzed by Denkena et al. [111], and the results were compared to those of Denkena and Lucas [108]. Only 3 conditions per alloy were tested. The turning was carried out with a cutting speed of 100 m/min, cutting depth of 200 µm, and a feed rate of 0.1 mm. Two deep rolling conditions were studied with rolling forces of 50 N and 200 N and the same cutting speed and feed as described by Denkena and Lucas [108]. The roughness Rz after turning and deep rolling was comparable for each alloy for every machining process. The roughness Rz after turning was about 4 µm, while Rz for deep rolling resulted in a lower Rz of between 0.44–0.76 µm (Table 2, Ref. [111],[108]) compared to Denkena and Lucas [108]. The corrosion behavior was tested by hydrogen evolution and performed in a 0.9 wt.% NaCl solution and µ-CT. It was shown that turned Mg-3Ca with the highest Rz resulted in the highest hydrogen evolution (~ 20.2 mL/cm2 after 29 h) and greatest degradation in µ-CT compared to deep rolled samples (~0.76–1.27 mL/cm2 after 29 h). For the Mg-0.8Ca alloy the hydrogen evolution (~5.42–6.22 mL/cm2 after 29 h) showed no significant dependence on the method of machining. From these investigations, it is possible to say that roughness had no influence on the degradation behavior. Rather than roughness, a high compressive stress and the Mg2Ca phase in the Mg-3Ca alloy was reported to affect the degradation behavior [111].

**Table 2.** Overview of different studies investigating SFF and machining and their influence on the degradation behavior. \* Values were determined from the diagram with the corresponding reference.

**Table 2.** *Cont*.

**Table 2.** *Cont*.

1 PDP: potentiodynamic polarization; 2 HBBS: Hank's Balanced Salt Solution; 3 HEPES: Biological buffer for cell culture media; 4 SFF: indirect solid free-form fabrication; 5 DMEM: Dulbecco's Modified Eagle's Medium; 6 FBS: fetal bovine serum; 7 P/S: Penicillin/Streptomycin; 8 DR: degradation rate; 9 DI: deionized; 10 HT: heat treated at 450 ◦C for 10 min (tempering); 11 SB60: sandblasting with glass bead type MS-6 at 60 psi; 12 SP: shot peening; 13 CR: corrosion rate; 14 SB40: sand blasting by #40 SiO2; 15 EDM: wire electrical discharge machining; 16: Deep Rolling: with vr = 25 m/min and Fr = 0.1 mm.

#### **3. Chemical Surface Treatments and Coatings**

#### *3.1. Acid Etching*

The treatments reported by Supplit et al. [95] indicated that it was possible to improve the degradation resistance of rolled AZ31 alloy by pickling with different acids like acetic acid, phosphoric acid, nitric acid, and hydrofluoric acid [95]. Acid pickling, especially with acetic acid decreased the degradation rate from 7.17 mg cm−2 d −1 (rolled condition) to 0.70 mg cm−2 d −1 . The second best etching method was found to be phosphoric acid. The degradation rates were determined by measuring hydrogen gas evolution in 5 % NaCl. A rougher surface after pickling with acetic acid was observed when compared to the other etching solutions. The samples with the lowest degradation rates had rougher surfaces, an observation that contradicts the findings of Nguyen et al. [72].

Organic acids like acetic, citric, or oxalic, and inorganic acids like phosphoric, nitric, and sulfuric acid were used to treat AZ31 alloy by Nwaogu et al. [40]. The aim was to remove contamination and impurities from resulting from rolling. After etching, 1–20 µm was removed from the surface. It was observed that more material was removed as the etching time increased. A roughness analysis showed that the roughness value Ra after etching is higher than Ra of rolled samples. Removing 5 µm of material generally reduced the number of Ni impurities. However, Fe impurities still remained at the surface even after material had been removed. To determine the degradation behavior, a 48 h salt spray test was used as a screening test. The lowest degradation rates were obtained from samples with the lowest impurity levels which had the greatest amount of material removed. Acetic acid-etched samples had the slowest degradation rates. EIS measurements supported the finding that acetic acid etching leads to the resistance, due to having the highest polarization resistance (Rp) [40]. This finding is in agreement with the results obtained by Supplit et al. [95] and Nwaogu et al. [40], who showed that a low impurity level and a 5 µm etching depth improved the degradation behavior. When more than 5 µm of material was removed, the surface became rougher [40].

The change in roughness after inorganic acid etching confirmed this finding [56]. For sulfuric acid etching, Ra (>2 µm) was much higher compared to other inorganic etching solutions when 7 µm of material had been removed. In addition, sulfuric acid etching leads to a lower degradation resistance in spite of the resulting low level of impurities. Degradation of sulfuric acid-etched material mostly results from galvanic degradation initiated by second phases. Though the effect of roughness were not the main focus in these investigations, nitric acid etching showed a high degradation resistance for a surface with an initially uniform roughness distribution and low roughness value [56]. Thus, the roughness of a sample after etching could also be a parameter which has to be considered in order to determine the full degradation behavior.

Gawlik et al. [57] measured the roughness after acetic acid etching with various combinations of acid concentration and immersion time. The surface roughness Sa increased after etching compared to the milled surface and varies with different conditions. After 30 days immersion, the same degradation rate was determined for all etched conditions (Table 3, Ref. [57]), in spite of different Sa and Sq (root mean square value of surface deviations) [98] values after etching. This leads to the conclusion that the initial roughness of the sample has no long-term effect on degradation. The varying surface morphology and near-surface deformation arising from milling also affected the degradation rate [57].

Similarly to Nwaogo et al. [40,56], Song and Xu [38] described that Fe impurities accelerate the degradation of the AZ31 alloy. As such, as-received samples and heat-treated samples have lower degradation resistance compared to ground and sulfuric acid etched samples, due to Fe particles remaining on the surface. Sulfuric acid etching roughens the surface much more than grinding, but both conditions lead to a similar degradation rate. Thus, the roughness of the etched samples itself does not contribute to the degradation rate. Acid cleaning removes contamination and the deformation zone arising from processing, and thus directly impacts the degradation behavior [38].

1 SST: salt spray test; 2 H2SO4: sulfuric acid; 3 CR: corrosion rate; 4 HNO3: nitric acid; 5 H3PO4: phosphoric acid; 6 CH3COOH: acetic acid; 7 C2H2O4: oxalic acid; 8 C6H8O7: citric acid; 9HF:hydrofluoricacid;10HT:heattreatedat450◦Cfor10min(tempering);11SBF:simulatedbodyfluid;12DMEM:Dulbecco'sModifiedEagle'sMedium;13FBS:fetalbovineserum;

 14 P/S: Penicillin/Streptomycin.

Gray-Munro et al. [70] also tested acetic treatments on AZ31 alloy. Gray and Luan's study [68] described that a strong passive oxide layer was formed during the etching process when compared to the as-received state. Gray-Munro et al. [70] found that as-received samples have a greater non-uniform morphology in comparison with phosphoric acid etched samples. Phosphoric acid treated samples showed a lower degradation rate of 8.27 mg/d compared to non-etched samples (~31 mg/d). Additionally, the modified surface after etching improves adhesion and minimizes the porosity of coatings [70].

#### *3.2. Coatings*

Gray-Munro et al. [70] discovered that biomimetic calcium phosphate coatings (Ca/P) and polymer coatings after phosphoric etching led to a uniform morphology, which in turn led to a uniform degradation over the surface of the AZ31 alloy. The degradation rates of 6.17 mg/d (Table 4, Ref. [70]) for a polymer poly(L-lactic acid) (PLA) coating and of 3.83 mg/d for a poly (desaminotyrosyl tyrosine hexyl) (DTH) carbonate coating are compared to etched samples with rates of 8.27 mg/d (Table 3, Ref. [70]) (sample size: 1 mm thick foil, 10 mm × 20 mm). The degradation rate (7.27 mg/d) resulting from the Ca/P coated samples did not strongly differ from the only etched alloys [70]. However, the Ca/P-coated Mg alloy exhibited non-toxic and biocompatible properties. Ca/P enhanced the osseointegration and bioresorption of the alloy in a physical environment [112–116], which is why a Ca/P coating is more favorable compared to phosphoric etching.

Bakhsheshi-Rad et al. [104] performed potentiodynamic polarization (PDP) tests and immersion tests in SBF (Kokubo solution) on polished Mg-0.5Ca-6Zn samples with and without a coating. The coatings tested were a fluoride conversion coating, a dicalcium phosphate dihydrate/magnesium fluoride (DCPD/MgF2) coating, and a nano-hydroxyapatite/magnesium fluoride (nano-HA/MgF2) coating. A higher root mean square roughness (Rq) was measured for polished Mg-0.5Ca-6Zn samples with either a DCPD/MgF2 coating (Rq = 395 nm) or a nano-HA/MgF2 coating (Rq = 468 nm). The Rq for polished Mg-0.5Ca-6Zn (Rq = 210 nm) samples without coating and polished Mg-0.5Ca-6Zn samples with a fluoride coating (Rq = 280 nm) were somewhat lower. As seen in Table 4, Ref. [104], hydrogen evolution, icorr and the degradation rate declined as the surface roughness increased, in contrast to studies of Walter and Kannan [96] and Mhaede et al. [102]. As shown in Table 4, Ref. [104], the degradation rate after coating compared to non-coated Mg alloy in Table 1, Ref. [104] is about a factor of 60 smaller, even though Rq only differs by 70 nm [104]. Thus, the protective coatings have a greater influence on corrosion than the Rq values.

Pompa et al. [103] investigated the morphology, surface roughness, cell viability, and degradation rate of ground and anodized AZ31B, AZ91E, and ZK60A alloys. Grinding was performed with a 1200 grit size grinding paper. Anodization was carried out using a mixture of alcohol and organic acid. The surface roughness increased dramatically from Sa = 29.76 nm to Sa = 204.81 nm after anodization for the AZ91E alloy. The anodization of AZ31B (Sa = 48.58 µm) and ZK60A (Sa = 78.30 µm) did not change the roughness significantly. It was shown that anodizing decreased the degradation rate compared to a ground surface. No correlation between surface roughness and degradation rate was found. This may be due to corrosion resistances being similar for all anodized surfaces despite their variation in roughness [103].

In the study of Chiu et al. [80], AZ31 plates were arc sprayed and hot pressed. Anodizing with oxalic acid was then performed [80]. The current density icorr decreased after a combination of spraying and hot pressing or additional anodizing (Table 4, Ref. [80]) as compared to uncoated sandblasted samples (Table 2, Ref. [80]). Hot pressing decreased the surface roughness, which seems to improve the acid treatment afterwards. No correlation between roughness and degradation resistance was reported [80].

**Table 4.** Overview of different Coatings studies involving roughness and degradation behavior. \* Values were determined from the diagram with the corresponding reference.

**Table 4.** *Cont*.

 1 MEM: Minimum Essential Media; 2 M-SBF: modified simulated body fluid; 3 FBS: fetal bovine serum; 4 CD: Cell density; 5 H3PO4: phosphoric acid; 6 CR: corrosion rate; 7 PLA: polymer poly(L-lactic acid); 8 DTH: desaminotyrosyl tyrosine hexyl; 9 PDP: potentiodynamic polarization; 10 DCPD: dicalcium phosphate dihydrate; 11 SP: Shot peening; 12 PBS: phosphate buffered saline; 13 SB: Sand blasting; 14 ASC:arc-spray coating; 15 PHP: post hot pressing; 16 PEO: plasma electrolytic oxidation; 17 KF: potassium fluoride; 18 HF: hydrofluoric acid; 19 HA: hydroxyapatite.

Yoo et al. [81] studied the effect of roughness on the degradation resistance of a plasma electrolytic oxidation (PEO) layer on a AZ91 alloy. Surfaces with various Ra were prepared by grinding and polishing. Afterwards, all samples were coated using the same PEO process. Due to the differing roughness values of the ground and polished surfaces, the coating differentiated in pore size as well. This affects the degradation process. The current density increased with higher initial Ra (Table 4, Ref. [81]). A salt spray test also showed that the amount of pitting increased with higher Ra after 120 h, which indicates that the surface roughness before PEO indirectly influences the degradation resistance [81].

Cho et al. [73] compared the degradation resistance of PEO coatings on AZ91 alloy for different amounts of potassium pyrophosphate in the electrolyte. The size of the pores increased as the amount of potassium pyrophosphate was increased. It was also reported that the surface roughness increased with increasing pore size. There was a trend between pore size, surface roughness, and icorr for additions of potassium pyrophosphate (Table 4, Ref. [73]). The PEO coating for the potassium pyrophosphate-free electrolyte exhibited the lowest degradation resistance compared to the rougher coatings [73].

In contrast to Cho et al. [73], Hwang et al. [79] compared PEO coatings on AZ91 alloy with and without potassium fluoride in the electrolyte. In addition to varying the surface roughness, the evolution of the degradation resistance with coating time was also examined. The surface roughness increased for longer coating times. The roughness of the coated samples, after dipping in the potassium fluoride containing electrolyte, was higher compared to coatings dipped into potassium fluoride free electrolyte. The roughness increases due to pore size enlargement as reported by Cho et al. [73]. The degradation resistance of the coatings when exposed to potassium fluoride-containing electrolyte was higher than for potassium fluoride-free electrolyte. Hwang et al. [79] explained that the oxide thickness is the reason for the improved degradation resistance, and did not assess the influence of roughness on the degradation resistance, as investigated in Hwang et al. [82]. The effect of the PEO coating roughness on the degradation behavior was examined in Hwang et al. [82] with PDP and three different coating surface roughness Ra values. The surface roughness increased with increased pore size, as was also seen in Hwang et al. [79] and by Cho et al. [73].

#### *3.3. Ion Implantation*

Jamesh et al. [105] implanted Si ions from a plasma on polished WE43 plates. Atomic force microscope (AFM) measurements after polishing and Si implantation showed that the surface became smoother after the ion implantation process. The smoother Si-implanted surfaces had improved degradation resistance. However, the roughness did not vary enough between the polished and Si-implanted surface types to obtain a correlation between roughness and degradation [105].

Zhao et al. [88] reported a slower degradation rate for Mg-Ca and Mg-Sr alloys after ion implantation (Zr and O ions) onto their surfaces. After the surfaces were implanted, measurements determined that the surfaces were uniformly rough. The roughness increased after implantation for both alloys (compare Tables 1 and 5, Ref. [88]). The current density icorr decreased for surfaces with higher roughness, and the cell adherence and proliferation improved [88].

**Table 5.** Summary of different studies concerning the influence of ion implantation on the degradation behavior. \* Values were determined from the diagram with the corresponding reference.

 1PDP: potentiodynamic polarization; 2 SBF: simulated body fluid; 3 DMEM: Dulbecco's Modified Eagle's Medium; 4 EIS: Electrochemical Impedance Spectroscopy

#### **4. Summary of the Influence of Roughness on Degradation**

#### *4.1. Mechanical Surface Treatments*

Using grinding papers with higher grit size and/or polishing reduced surface roughness and reduced pitting during the degradation tests [71,96]. These results differ from those in studies by Alvarez et al. [94], where polished samples encouraged pitting compared to semi polished surfaces [94]. Some papers reported that roughness affects cell adherence to the surface [21,74]. However, one study showed that cell adherence was not influenced by surface roughness [61]. Some studies showed that roughness did not affect degradation [39,102,108,111]. In contrast, it was demonstrated that a linear relationship existed between roughness and degradation if the surface treatments were comparable (Figures 1 and 2) [72,96,102].

#### *4.2. Chemical Surface Treatments and Coatings*