The fabrication defects of the LDED fabricated TiC-TMC coatings on a cross-sectional surface are shown in Figure 7. It could be seen that most fabrication defects were in irregular shapes. In the LDED process, the irregular-shaped fabrication defects on the cross-sectional surface were usually caused by lack of fusion, as demonstrated by Zhang et al. [45]. As the molten pool was enlarged by the assistance of ultrasonic vibration, the lack of fusion at the boundary of the molten pool was reduced. In addition, more powder could be caught by the molten pool during the fabrication, promoting the formation of sufficient overlaps.

**Figure 7.** The effects of ultrasonic vibration on the fabrication defects: (**a**) the molten pool of the TiC-TMC coatings; (**b**) the bonding interface of the TiC-TMC coatings.

With the increased TiC content, the amount of irregular fabrication defects increased due to the lack of fusion. In this study, the laser input energy was constant. The TiC needed to absorb more energy to be melted than titanium. The TiC-TMC coatings with a higher TiC content had more unmelted TiC(O) particles in the molten pool during the fabrication, which aggravated the lack of fusion. Moreover, the large number of solid TiC(O) particles reduced the fluidity of the liquid materials in the molten pool resulting in the aggregation of unmelted TiC(O) particles. In these unmelted TiC(O)-rich regions, a lack of fusion was more likely to happen, as reported in the LDED fabricated Ti6Al4V with trace boron and the selected laser melting of a TiB2 coating on Ti6Al4V [46,47].

#### *3.5. Mechanical Properties*

#### 3.5.1. Microhardness

Figure 8 shows the effects of ultrasonic vibration and TiC content on microhardness. With the assistance of ultrasonic vibration, the microhardness value increased. As discussed in Section 3.1, more refined Ti*x*C*y*(P) particles were precipitated in the titanium matrix with the assistance of ultrasonic vibration. These refined reinforcements could evenly bear the load and increase the resistance of plastic deformation during the microhardness tests. A similar result was demonstrated by Shen et al. through numerical methods [48]. Moreover, as discussed in Section 3.4, with the assistance of ultrasonic vibration, the fabrication defects decreased significantly. The higher density increased the ability to support the load, which could also increase the microhardness. Compared with commercial pure

titanium coating layers, TiC-TMC coatings had a larger microhardness. The major reason was that TiC could support the load and reduce the deformation. However, the excessive TiC content led to the reduction of microhardness. As discussed before, the major reason was that the increase of fabrication defects reduced the microhardness.

**Figure 8.** The effects of ultrasonic vibration and TiC content on microhardness.

#### 3.5.2. Wear Resistance

The sliding width of the material removal trail and wear rate are shown in Figure 9. Both the sliding width and wear rate had a negative relationship with the wear resistance. With the assistance of ultrasonic vibration, the wear resistance increased. The significant reduction of the fabrication defects provided a smoother interface and reduced the friction coefficient of the coating layers [49]. Under the same test condition, the friction force was reduced, resulting in a smaller material removal volume and higher wear resistance. In addition, the reinforcement aggregation was reduced by ultrasonic vibration. It contributed to the formation of the uniform anti-wear protective layer during the dry sliding tests to further increase the wear resistance.

**Figure 9.** The effects of ultrasonic vibration and TiC content on microhardness.

Compared with CP-Ti, TiC-TMC coatings showed a higher wear resistance. In addition, the friction force was measured during the dry sliding tests. The average friction coefficients of each combination of input were calculated and are listed in Table 4. The major reason was that the hardness of TiC was extremely high, which increased the hardness of the TiC-TMC coatings. The adhesive wear mode was changed to abrasive wear, which could significantly increase the wear resistance [50,51]. In addition, during the dry sliding tests, the SiC ball would leave a rough worn surface on the samples. The spherical TiC particles might have the function of self-lubrication [52,53]. With the increase of the TiC content to a high level (10 wt.%), the wear resistance was reduced. The major reason was that the greater number of fabrication defects increased the friction coefficient.

**Table 4.** The average friction coefficient of the CP-Ti and TiC-TMC coatings.

#### **4. Conclusions**

In this study, TiC-TMC coatings with different TiC contents were fabricated by an ultrasonic vibration-assisted LDED process. The effects of ultrasonic vibration and TiC content on the phase composition, TiC aggregation, bonding quality, fabrication defects and mechanical properties were investigated.

With the assistance of ultrasonic vibration, the liquid materials were mixed and stirred by the actions of acoustic streaming and transient cavitation in liquid material solidification, which promoted the dissolution of TiC(O) and the precipitation of refined Ti*x*C*y*(P). In the TiC-TMC coatings fabricated by LDED without ultrasonic vibration, TiC(O) particles preferred to aggregate in a titanium matrix. In contrast, the process with ultrasonic vibration could significantly reduce the TiC(O) aggregation by decreasing the amount of TiC(O) and improving the distribution. In addition, extra heat energy was generated by the ultrasonic vibration assistant, which could increase the bonding quality and reduce the fabrication defects. The improved TiC(O) aggregation and fabrication defects were effective in enhancing the microhardness and wear resistance of the ultrasonic vibration-assisted LDED fabricated TiC-TMC coatings.

With the increase of TiC content, the phase transformation from the TiC(O) phase to the Ti*x*C*y*(P) phase was suppressed. The major reason was that the higher TiC content suppressed the further dissolution of TiC(O) in liquid titanium. At a high content of TiC, there were more undissolved TiC(O) particles existing in the TiC-TMC coatings after the fabrication, which aggravated the reinforcement aggregation. In addition, the higher TiC content exacerbated the generation of a lack of fusion defects as the TiC powders needed more energy to be melted than Ti. The generation of both the TiC(O) aggregation and fabrication defects decreased the mechanical properties of microhardness and wear resistance of the TiC-TMC coatings with a higher TiC content.

**Author Contributions:** Data curation, Y.L. and D.Z.; Formal analysis, Y.L.; Investigation, Y.L.; Methodology, Y.L.; Resources, Y.L.; writing—original draft, Y.L.; Writing—review & editing, Y.L., D.Z., H.W. and W.C. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research received no external funding.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

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

#### **References**

## *Article* **3D-Printed Connector for Revision Limb Salvage Surgery in Long Bones Previously Using Customized Implants**

**Jong-Woong Park 1,2, Hyun-Guy Kang 1,2,\*, June-Hyuk Kim 1 and Han-Soo Kim 3**

**Abstract:** In orthopedic oncology, revisional surgery due to mechanical failure or local recurrence is not uncommon following limb salvage surgery using an endoprosthesis. However, due to the lack of clinical experience in limb salvage surgery using 3D-printed custom-made implants, there have been no reports of revision limb salvage surgery using a 3D-printed implant. Herein, we present two cases of representative revision limb salvage surgeries that utilized another 3D-printed custom-made implant while retaining the previous 3D-printed custom-made implant. A 3D-printed connector implant was used to connect the previous 3D-printed implant to the proximal ulna of a 40-year-old man and to the femur of a 69-year-old woman. The connector bodies for the two junctions of the previous implant and the remaining host bone were designed for the most functional position or angle by twisting or tilting. Using the previous 3D-printed implant as a taper, the 3Dprinted connector was used to encase the outside of the previous implant. The gap between the previous implant and the new one was subsequently filled with bone cement. For both the upper and lower extremities, the 3D-printed connector showed stable reconstruction and excellent functional outcomes (Musculoskeletal Tumor Society scores of 87% and 100%, respectively) in the short-term follow-up. To retain the previous 3D-printed implant during revision limb salvage surgery, an additional 3D-printed implant may be a feasible surgical option.

**Keywords:** extremity; revision; limb salvage surgery; 3D printing; customized; implant

#### **1. Introduction**

Custom-made implants for limb salvage surgery pre-date modular implants; however, their extensive manufacturing duration remains a major drawback [1]. Modular endoprostheses are the most commonly-used implants that have been used to overcome the disadvantages of early custom-made implants [1–3]. They can be used by orthopedic oncology surgeons to provide surgical flexibility, and standardization of the implant pieces facilitates manufacturing. However, there are anatomic sites where a modular endoprosthesis cannot be used. Furthermore, the adjacent normal joint is often sacrificed during reconstruction using the modular implant due to limited bone stock for implant fixation. Recently, three-dimensional (3D)-printing has been introduced for medical applications, wherein custom-made implants can be fabricated in a few days or weeks. These 3D-printed custom-made implants have shown promising short- and mid-term surgical outcomes in orthopedic oncology [4–7].

For limb salvage surgery of the long bones, the modular endoprosthesis system is available for most anatomical locations. However, the 3D-printed custom-made implant has advantages in long bone surgery for preservation of normal adjacent joints, and in limb salvage surgeries for specific anatomical locations without an existing conventional tumor endoprosthesis. When limb salvage surgery with a conventional endoprosthesis including

**Citation:** Park, J.-W.; Kang, H.-G.; Kim, J.-H.; Kim, H.-S. 3D-Printed Connector for Revision Limb Salvage Surgery in Long Bones Previously Using Customized Implants. *Metals* **2021**, *11*, 707. https://doi.org/ 10.3390/met11050707

Academic Editor: Atila Ertas

Received: 2 April 2021 Accepted: 22 April 2021 Published: 26 April 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

an artificial joint is required, intercalary replacement with joint-saving of the long bone is often possible with a 3D-printed custom-made implant.

Limb salvage surgery with an endoprosthesis is more vulnerable to mechanical failure than implants for arthroplasty without replacement of a major bone defect [8–11]. One of the concerns associated with custom-made implants is the higher probability of mechanical failure as the implant design is not standardized [12]. While the mechanical strength of the full solid-body implant fabricated by 3D-printing is comparable to those fabricated by forged, wrought, or casting [13], the lattice structure, that is an intended porotic structure to facilitate tissue integration around the implant, requires mechanical and clinical verification [6,7,14]. In this situation, the biggest concern in practice is overcoming failure of bone union or fracture of implant after performing limb salvage surgery with the 3D-printed implant. As most 3D-printed custom-made implants are not modular but rather a single unit, it is difficult to replace only the broken parts or those where the tumor has recurred nearby.

Herein, we present a surgical method that utilizes an additional 3D-printed implant in revision limb salvage surgery to preserve the previous 3D-printed custom-made titanium alloy implants in long bones.

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

Two representative patients underwent revision surgeries using a 3D-printed connector to connect the previous 3D-implant to the host bone. The reasons for revision surgeries after the 3D-printed limb salvage surgeries were tumor recurrence and implant fracture. The previous implants achieved bony union radiographically at the junctions between the implant and the host bone, and neither patient underwent revision surgery due to non-union at the junction. Rather than removing all implants, the revision surgeries were performed in a manner that preserved the previous implant as much as possible. The 3D-printed customized implants were designed using the Materialise Interactive Medical Image Control System (Materialise; Leuven, Belgium) and Magics 22 (Materialise; Leuven, Belgium) and fabricated using an electron beam melting-type (EBM) 3D-printer (ARCAM A1, Arcam AB, Mölndal, Sweden) with Ti6Al4V (Ti6Al4V-ELI Per ASTM 136). The custom-made implant was created by the MEDYSSEY Company (Jecheon, Korea) and certified by the Ministry of Food and Drug Safety. All study participants provided informed consent, and the study design was approved by the appropriate ethics review board (NCC2017-0129).

For the 3D-printed connector, two junctional parts for the previous implant and remaining host bone were designed first. The connector body was subsequently designed for the most functional position or angle by twisting or tilting. The junction to the previous metal implants and the bodies were mainly solid structure, while the junction to the remaining host bone partly had lattice structure for contact surface with the bone. The lattice structure had 750 um pores [7,15]. The main concept of the 3D-printed connector is that it mimics that of a modular prosthesis, with an assembly mechanism to insert one piece into another (Modular Universal Tumor and Revision System, Implantcast, Buxtehude, Germany; Global Modular Replacement System, Stryker, Kalamazoo, MI, USA). The previous 3D-printed implant was tapered on the inside, and the 3D-printed connector encased the outside of the previous implant. The 1.5-mm gap between the previous implant and the new one was filled with bone cement. The metal artifact around the previous 3D-printed implant that was seen in the computed tomography (CT) images was removed and smoothed as much as possible, though there was a risk of error during the modelling process. For both patients, the design of the implant used in the first surgery was saved as an STL file; in the CT images taken for the new implant, the previous implant part was not modelled by dealing with metal artifacts but was replaced by fusion of the implant design data.

The first patient reported here was a 40-year-old man who underwent treatment for a desmoplastic fibroma, a locally aggressive intermediate tumor, in his forearm [14]. He had undergone orthopedic surgery on the same side of both forearm bones and had used an external fixator for fracture fixation five years ago. It is not clear whether the tumor had originated from the fracture site or had entered the screw holes of the previous external fixator, but the tumor involved both forearm bones. There was no distant metastasis. Bilateral limb salvage surgery for the forearm bones was performed using two 3D-printed custom-made implants for the radius and the ulna (Figure 1A–E). One year postoperatively, local tumor recurrence was detected at the proximal junction of the ulnar implant. The radial implant showed good fixation in both proximal and distal junctions. The first revision surgery was performed to remove the recurred tumor and to make a single forearm bone by fusion of the proximal and distal remnant bones. However, fusion at the proximal part failed and a second revision surgery was planned (Figure 1F–H).

**Figure 1.** Failure of the 3D-printed implants in both forearm bones [14]. (**A**) A preoperative plain radiography and (**B**) gadolinium-enhanced T1-weighted magnetic resonance image showing desmoplastic fibroma arising from both forearm bones. (**C**) A graphic design and (**D**) photograph of 3D-printed implants. (**E**) A postoperative plain radiograph after limb salvage surgery. Photographs showing (**F**) removal of the ulnar implant with a recurrent tumor and (**G**) a specimen. (**H**) Plain radiographs showing immediate postoperative status after the first revision surgery (**left**) and mechanical failure 4 months later (**right**).

The second representative patient was a 69-year-old woman who had chondrosarcoma at the left distal half of the femur without distant metastasis. She had a limb length discrepancy of 10 cm due to a childhood history of osteomyelitis in the contralateral femur. The patient underwent limb salvage surgery using a 3D-printed customized implant for the distal part of her right femur with retention of the natural knee joint, and acute shortening of the right femur for partial correction of limb length discrepancy (Figure 2A–C). The patient could ambulate independently three months postoperatively. However, six months postoperatively, she slipped, and the proximal part of the implant was broken between the implant body and the fixation plate. The unexpected implant fracture was analyzed and the main causes of the fracture were stress concentration at the broken junction and internal defects of the 3D-implant. The distal junction to the knee joint was well-fixed. The first revision surgery was performed with additional dual plates with expectation of bony

union or extracortical bone bridge to the porous structure of the implant body. However, 21 months later, mechanical failure occurred and a second revision surgery was planned (Figure 2D–F).

**Figure 2.** Failure of the 3D-printed implants in the right distal femur. (**A**) A preoperative plain radiography and gadoliniumenhanced T1-weighted magnetic resonance images showing chondrosarcoma in the right distal femur. (**B**) A graphical design of the 3D-printed implant. (**C**) A preoperative (**left**) and postoperative teleradiogram (**right**) after limb salvage surgery. (**D**) Photographs showing implant fracture at 6 months postoperatively. (**E**) Intraoperative photographs and the removed broken proximal part. (**F**) Plain radiographs showing immediate postoperative status after the first revision surgery (**left**) and another mechanical failure 21 months later (**right**).

#### **3. Results**

#### *3.1. Forearm*

For the patient who underwent single forearm bone surgery, the proximal part of the previous radial implant and the proximal ulna needed to be connected while maintaining the distal part, including the wrist joint. The wrist and hand function in the salvaged forearm was normal; however, thumb extension was limited due to soft tissue adhesion of the thumb flexor muscles. The elbow joint function was originally limited from 0 to 90 degrees due to radial head dislocation secondary to a childhood trauma. The proximal part of the previous radial implant was half the cylindrical plate that originally wrapped the proximal radial shaft. The junctional part of the 3D-printed connector up to the previous radial implant was designed to have a metal insert replace the space of the original radial shaft and the remaining half of the cylindrical plate from the opposite direction of the previous implant. The second junctional part was designed for the proximal ulna near the elbow joint, and this comprised two-thirds of a cylindrical plate with multiple holes for the screws and suture. The functional position of the forearm rotation was set at 15 degrees toward pronation from the neutral position. The body of the 3D-printed connector was twisted to the functional rotational position (Figure 3A–D).

**Figure 3.** Revision limb salvage surgery using 3D-printed connector in the forearm. (**A**) A photograph of the previous 3D-printed implant for the radius. (**B**) Graphic designs of the 3D-printed connector and (**C**,**D**) photographs of the 3D-printed connector for single forearm bone surgery preserving the previous 3D-printed radial implant. Postoperative (**E**) plain radiographs and (**F**) computed tomography reconstruction images after second revision surgery using the 3D-printed connector. (**G**) A photograph showing functional position of the forearm during writing.

In the second revision surgery, the broken screws and the remaining part of the proximal radius were removed. The wrist joint was stable after the first revision surgery, and the host bones near the wrist were firmly bonded to the previous 3D-printed implant. The proximal part of the previous 3D-printed implant and the ulnar remnant near the elbow joint were fixed using the 3D-printed connector. The fixation between the old and new 3D-printed implants was achieved by winding three wires through a matched hole with bone cement. The 3D-printed connector and the proximal ulna showed conforming stability after inserting the bone into the implant, which was designed to sufficiently enclose the bone and fixed using screws (Figure 3E–F). The patient recovered the original range of motion of his elbow (0–90 degrees) at 10 weeks postoperatively and was able to write with his forearm in the 15-degree pronation position (Figure 3G). The wrist and elbow function were well-maintained without mechanical failure until the last follow-up at 26 months after the second revision surgery. The Musculoskeletal Tumor Society (MSTS) score at the last visit was 26 (87%).

#### *3.2. Femur*

For the patient who underwent revision surgery using the 3D-printed connector for the femur, the femoral shaft of the previous implant body and proximal femur needed to be connected while maintaining the distal part, including the knee joint. The knee joint had an extension lag of 10 degrees due to acute shortening of the femur, though independent ambulation without pain was possible. Furthermore, partial correction of the limb length discrepancy led to an improvement in gait, and the left heel that did not touch the ground during ambulation before the limb length correction, could touch the ground while walking after the correction (Figure 4D). The junctional part of the 3D-printed connector to the previous broken 3D-printed custom-made implant was a full cylinder that surrounded the remaining previous implant. Another junctional part was designed as a one-third plate with screw holes to fix the proximal femur near the hip joint. The body of the 3D-printed connector was designed to restore the tilting angle of the femur from the mechanical axis of the lower limb (Figure 4A,B).

**Figure 4.** Revision limb salvage surgery using 3D-printed connector in the femur. (**A**) A graphic design of the 3D-printed connector. (**B**) Photographs of the previous 3D-printed implant (duplicated), broken implant (duplicated), and the 3D-printed connector. (**C**) Intraoperative photographs before and after application of the 3D-printed connector. (**D**) Preoperative teleradiograph and photograph showing the chondrosarcoma of the right distal femur and limb length discrepancy of 10 cm (**left**), and final postoperative pictures (**right**). (**E**) Photographs showing normal knee joint function.

In the second surgery, all the metal plates and screws were removed, leaving the knee joint below the shaft of the broken 3D-printed customized implant used in the first surgery. The 3D-printed connector was connected by inserting the shaft of the previous 3D implant into the 3D-printed connector. The gap between the old and new 3D-printed implant was filled with bone cement. The remaining proximal femur on the hip side was fixed by multiple screws. The patient resumed walking on crutches at 2 weeks after the second operation and was able to walk without crutches without pain at 10 weeks postoperatively. The range of motion of the knee joint was normal without extension lag (Figure 4C–E). At 10 months after the second revision surgery, the MSTS score was 30 (100%).

#### **4. Discussion**

When a megaprosthesis is fabricated by 3D printing, it is often difficult to implement a modular-type; therefore, it is fabricated as a single mass. After limb salvage surgery using a megaprosthesis, revision surgeries are often required for a variety of reasons, including local recurrence and mechanical failure. One of the concerns regarding the use of a 3D-printed customized implant in limb salvage surgery is that it makes revision surgery difficult. In this study, an additional 3D-printed implant was introduced during revision without removing all of the previously inserted 3D-printed implant. The surgical results using the additional 3D-printed connector implant were acceptable in the short-term follow-up.

The metal-to-metal fixation mechanism is still a problem in the 3D-printed implant. The locking screw mechanism has not been introduced in 3D-printed implants as the male and female screw threads are hard to implement due to insufficient precision and

roughness of the electron beam melting-type 3D-printed product. Therefore, cemented fixation to the intended gap between the implants was mainly utilized in these patients.

This study has some limitations. It was a preliminary study based on a small number of patients with a short follow-up. Although the short-term surgical results were promising, they are difficult to generalize. These results imply that the 3D-printed connector implant is a surgical option to consider when revision surgery after previous limb salvage surgery using a 3D-printed customized implant is needed, and retaining the previous 3D-printed implant totally/partially would be beneficial. Regarding the mechanical property of the 3D-printed implants, long-term follow-up clinical data and biomechanical data is required. There was no serious mismatch between the 3D connector and previous implant or remaining bone. However, deformation of the 3D-printed customized implant by residual stress and contraction was not accurately measured in this paper.

In conclusion, to retain the previous 3D-printed implant during revision limb salvage surgery, an additional 3D-printed implant may be a feasible surgical option.

**Author Contributions:** J.-W.P. conceptualized, interpreted the patient data and implant designs and was a major contributor in writing the manuscript. H.-G.K. was a primary implant designer and visualized the illustrations. J.-H.K. and H.-S.K. contributed to validation of the implant designs. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was supported by the Industrial Strategic Technology Development Program ('P0008805 ), funded by the Ministry of Trade, Industry & Energy (MOTIE, Korea).

**Institutional Review Board Statement:** The study was conducted according to the guidelines of the Declaration of Helsinki, and approved by the Institutional Review Board of National Cancer Center, Korea (NCC2017-0129, 1 June 2017).

**Informed Consent Statement:** Informed consent was obtained from all subjects involved in the study.

**Data Availability Statement:** The datasets used and/or analyzed during the current study are available from the corresponding author on reasonable request.

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

#### **References**

## *Article* **In Vitro Physical-Chemical Behaviour Assessment of 3D-Printed CoCrMo Alloy for Orthopaedic Implants**

**Radu Mirea 1,\*, Iuliana Manuela Biris 2, Laurentiu Constantin Ceatra 1, Razvan Ene 3,4, Alexandru Paraschiv 1, Andrei Tiberiu Cucuruz 5, Gabriela Sbarcea 6,\*, Elisa Popescu 3 and Teodor Badea 1**

**Abstract:** In this study, a CoCrMo-based metallic alloy was manufactured using a 3D-printing method with metallic powder and a laser-based 3D printer. The obtained material was immersed in a simulated body fluid (SBF) similar to blood plasma and kept 2 months at 37 ◦C and in relative motion against the SBF in order to mimic the real motion of body fluids against an implant. At determined time intervals (24, 72, 168, 336, and 1344 h), both the metallic sample and SBF were characterized from a physical-chemical point of view in order to assess the alloy's behaviour in the SBF. Firstly, the CoCrMo based metallic sample was characterized by scanning electron microscopy (SEM) for assessing surface corrosion and X-ray diffraction (XRD) for determining if and/or what kind of spontaneous protective layer was formed on the surface; secondly, the SBF was characterized by pH, electrical conductivity (EC), and inductively coupled plasma mass spectroscopy (ICP-MS) for assessing the metal ion release. We determined that a 3D-printed CoCrMo alloy does not represent a potential biological hazard in terms of the concentration of metal ion releases, since it forms, in a relatively short period of time, a protective CoCr layer on its exposed surface.

**Keywords:** 3D printing; powder metallurgy; simulated body fluid; biomaterial

#### **1. Introduction**

"Traditional" orthopaedic implants are titanium (Ti)-based alloys since they provide a series of characteristics suitable for such applications. They can be used in hip joint replacement, tibia rod, clavicle plate, etc. Nevertheless, there are other metallic alloys used in orthopaedic implants that are not Ti based. Cobalt-chromium (CoCr)-based alloys represent a viable solution and molybdenum (Mo) is often added to enhance the inner structure in terms of increasing the ductility and strength of the alloy [1]. Although casting and forging are, and continue to be, the main methods for obtaining both CoCr- and Ti-based alloys for orthopaedic implants, 3D printing is gradually gaining its position within this field. In recent years, 3D printing (both polymer and metal printing) has opened several exciting possibilities to create customized orthopaedic implants [2]. There are three rapid prototyping techniques used to produce metallic implants, i.e., selective laser melting (SLM), electron beam melting (EBM), and laser engineered net shaping (LENS), and each of them having their advantages and disadvantages. The SLM technique uses high power

**Citation:** Mirea, R.; Biris, I.M.; Ceatra, L.C.; Ene, R.; Paraschiv, A.; Cucuruz, A.T.; Sbarcea, G.; Popescu, E.; Badea, T. In Vitro Physical-Chemical Behaviour Assessment of 3D-Printed CoCrMo Alloy for Orthopaedic Implants. *Metals* **2021**, *11*, 857. https://doi.org/ 10.3390/met11060857

Academic Editors: Atila Ertas and Aleksander Lisiecki

Received: 29 April 2021 Accepted: 22 May 2021 Published: 24 May 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

lasers to melt metallic powders in order to achieve the required shape of the implant, while the EBM technique uses a focused high-energy electron beam for direct solidification of metal powders. The LENS technique creates a molten metal pool on a substrate, and then the metal powder is injected [3]. A 3D printing process usually uses digital design software and/or a 3D digital scanner in order to achieve the virtual version of the metallic implant to be produced. Then, the obtained model is transformed into an .STL file format enabling the exact spatial coordinates (xyz) of a model's surfaces. Next, the model is "sliced" by using another software, each "slice" being 25–100 μm thick. By stacking these "slices", the final shape of the implant is obtained. A 3D printer can use different types of materials starting from various plastics to different metallic powder mixtures. About 50 metallic alloys can be used in 3D printing and they are mostly Ti, Ni, Al, stainless steel, CoCr, etc., alloys and more than 80% of them are used as implants [4].

According to the ASTM standard F2792-12a [5], the SLM technique is classified as powder bed fusion technology [6]. Therefore, in this paper, we describe the use of this technique and characterize the obtained CoCr-based material used for implants. Thus, an in vitro behaviour assessment of the obtained metallic material used for orthopaedic implants has been performed during 60 days of immersion in a simulated body fluid (SBF) similar to blood plasma.

For a metallic material to be suitable as biomaterial, it should meet a series of desired properties ranging from mechanical (matching the material's elasticity modulus with that of the bone, high strength, high wear, corrosion resistance, etc.) to biocompatibility (not toxic to environmental cells and tissues, quickly integrated in the body, etc.) [7,8].

CoCr-based alloys have higher wear resistance than Ti-based alloys; therefore, they are especially used for implants that, over time are subjected to extensive wear, such as hip joints, etc. Clinically, CoCrMo alloys are used the most, since they are the perfect combination between high strength and high ductility, but the main drawback is their high elasticity modulus, which may lead to greater stress in the bone. This drawback may be overcome by 3D printing of CoCr-based alloys. Although EBM is the most used technique for 3D printing of CoCr alloys [8], the SLM technique can also be used. When prepared by using the SLM technique, Mo must be added to the powder since it enables high temperature gradients during melting and subsequent rapid cooling and, especially, it enables the formation of the fine granular structure of the implant [9].

Thus, it is possible to produce implants that have tailored mechanical properties [10] but the issue of the material's biocompatibility still remains. One of the main problems is metal ion dissociation from the metallic material into the body fluids, which are in close contact with the metal. It should be mentioned that corrosion is the most frequent phenomena that occurs in the case of an implant and one of the most important ones to be avoided. It is well known that, from a biocompatibility point of view, a suitable material performs with an appropriate host response (e.g., minimum disruption of normal body function) and does not cause any allergic and/or inflammatory response when placed in vivo [11].

The first requirement of any material that is placed in the human body is that it should be biocompatible and not cause any adverse reaction. Corrosion and surface oxide film dissolution are the two mechanisms which introduce additional ions into the body. Extensive release of ions from implants can result in adverse biological reactions. Corrosion is the first consideration for a material of any type that is to be used in the body because metal ion release takes place mainly due to corrosion of surgical implants [12]. The two main corrosion types that often appear are spot and pitting corrosion, as described in [13,14].

In the particular case of orthopaedic implants, micro-motions are known to occur at points of fixation, while corrosion is caused by the body fluids, which contain various inorganic and organic molecules [15].

In this paper, we present the results obtained by immersing a CoCr-based 3D-printed material in vitro (being the "next best thing") into an SBF for 60 days at 37 ◦C with the fluid being in relative motion against the material. We assessed the appearance of surface corrosion, metal ion dissociation, and the appearance of a spontaneous protective layer on the exposed surface.

Considering that even a small concentration of metal ion release into the body could be very aggressive due to their migration and accumulation in different organs, sometimes far away from the point of release, it is important to quantify the concentration of ion release in various environments; thus, ICP-MS investigations were conducted [16].

Material corrosion may lead to implant failure as [17] has assessed and may cause an allergic and/or toxic response of the body.

A complete understanding of 3D-printed metallic implants is needed, nevertheless, what has been assessed so far is genuinely promising. Some of the advantages of 3Dprinted implants as compared with the "traditional" manufactured ones, are highlighted in [18], among them, we highlight the following: the final shape of the implant can be obtained in a single phase of production; as many as needed can be produced, even if only one piece; the mechanical properties are better than cast or forged implants; the design and geometry can be as complicated as needed; porosity can be designed as needed for the elasticity module to match that of a bone; and the pores can be designed to be opened or inter-correlated in order to ensure maximum osseointegration.

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

The experimental part of this study was to manufacture a 3D-printed CoCrMo metallic alloy that could be used for an orthopaedic implant. The CoCr alloy was chosen as a good alternative to Ti-based alloys given its wear resistance [7]. Molybdenum was added to the mixture in order to increase the ductility and to increase the strength of the alloy.

#### *2.1. Method for Producing the CoCrMo Alloy*

#### 2.1.1. 3D Manufacturing

The metallic implant was manufactured using a selective laser melting (SLM) TruPrint 1000 (TRUMPF Inc., Ontario, ON, Canada) printer and CoCrMo metal powder (Starbond CoS powder 30 from S&S Schefter GmbH, Mainz, Germany) within the 10–30 μm grain size and with the chemical composition presented in Table 1.

**Table 1.** Chemical composition of the metallic powder.

For 3D manufacturing of the metallic implant, the following process parameters were used: 90 W laser power, 600 mm/s laser speed, 20 μm layer thickness, and 55 μm laser spot. An Ar (99.995%) flow system was used to reduce the oxygen level to less than 300 ppm. The material was mechanically removed from the building platform and was tested in a simulated body fluid solution without any post-processing operations. Moreover, its hardness was measured to be 463.5 (HV 0.5). Then, the sample was polished by using a rotating polishing machine, until its roughness dropped below 0.05 μm (N2). The used polishing paper had a roughness varying from 240 to 1200. Then, the sample was surfaced using diamond powder suspensions of 3 and 1 μm, respectively.

#### 2.1.2. SEM Assessment

The micro-structural analysis was performed by using an FEI F50 Inspect (FEI Company, Brno, Czech Republic) with a magnification of 4000×, HV = 20.00 kV, WD 11.6 mm, spot 3.0, and dwell of 30 μs. The method has been used to emphasize deep corrosion and surface transformations of the metallic samples [17,19].

#### 2.1.3. Mass Variation Assessment

The aim of the mass loss/gain assessment was to determine, the corrosion rate and correlated with XRD the formation of protective layer on sample's surface.

#### 2.1.4. XRD Assessment

The XRD analysis was conducted to establish if and/or what kind of oxides had formed on the sample's surface. A Bruker D8 Discover diffractometer (Bruker, Billerica, MA, USA) was used. The diffractometer settings were as follows: primary optics used a Cu tube (λ = 1.540598 Å) and a Göebel mirror, while secondary optics used a 1D LynxEye detector (Bruker, Billerica, MA, USA). The plots were recorded at 0.04◦ angle and 1s/step scanning speed. They were indexed using the ICDD Release 2015 database.

#### *2.2. Method for Preparing the SBF*

The SBF solution was prepared, as shown in [20]; it was an updated version of the simulated body fluid submitted, in 2003, with detailed instructions for its preparation to the Technical Committee ISO/TC 150 of the International Organization for Standardization as a solution for in vitro measurements of implant materials. Table 2 shows the concentrations of reagents to be used for 1 L of SBF. The SBF solution was prepared in the lab by mixing extra pure substances procured from the market from WWR Chemicals. The sample was immersed in the prepared SBF in an Erlenmeyer polypropylene (PP) recipient. The recipient was immersed in a thermostatic bath equipped with a movable device that allowed a frequency of 60 movements/minute with a distance of ±1 cm right/left in order to mimic the relative motion of body fluid against an implant. A PP wire was used to ensure that the sample was kept immersed in the SBF, however, did not come in contact with the recipient.

**Table 2.** Reagents for preparing the updated simulated body fluid [20]. (reprinted from ref. [20] with permission from Elsevier).

#### 2.2.1. PH and Electrical Conductivity Assessment

The pH and electrical conductivity characterizations were performed in order to determine the redox phenomena that occurred and the variation of ions in the SBF. A Mettler-Toledo pH/conduct meter (Mettler-Toledo, Greifensee, Switzerland) was used.

#### 2.2.2. ICP-MS Assessment

Inductively coupled plasma mass spectrometry (ICP-MS) [21] was used for determining the concentration of metal ions dissociated in the SBF. A Dirac Elan II ICP-MS spectrometer (Perkin-Elmer, Toronto, ON, Canada) was used.

#### **3. Results and Discussions**

The conducted experiments focused on both the material and the SBF. The surface corrosion of the material was assessed by scanning electron microscopy (SEM), XRD to determine if and/or what kinds of oxides had formed on its surface while immersed, and mass variation. The SBF was assessed by its pH, electrical conductivity variation, and by ICP-MS analysis to determine the concentration of metal ion dissociated from the metallic material.

In order to better assess the ion dissociation and surface corrosion, the protective oxide layer of the sample was removed on one side and the exposed material was polished until its roughness dropped below 0.05 μm (N2). All the above-mentioned assessment that refers to the material was performed on the exposed surface.

#### *3.1. Material Characterization*

#### 3.1.1. SEM Assessment

The SEM technique uses electrons for images; thus, it is considered to be the most suitable technique for high-resolution imaging of surfaces. The basic principle is similar to OM, but electrons are used instead of light, allowing a much higher magnification (higher than 100,000×).

The aim of this type of assessment was to investigate even deeper the corrosion that can occur on the surface. A magnification of 4000× was considered to be adequate in this case, and therefore the SEM assessment was performed by means of this magnification and the covered area was 80 μm × 80 μm. The analysis was performed on the exposed surface, aimed at highlighting the formation of the protective layer as a replacement for the layer removed.

The SEM images (Figure 1) show surface corrosion of the sample as pitting in the first two weeks of exposure and, during the next two weeks, the formation of porous structures. According to [22], this porous structure may not be pitting (a type of corrosion that often appears) but porous single crystals of Cr2O3 and Co3O4. This hypothesis is strongly supported by ICP-MS that does not show any increase in metal ions into the SBF and pH and EC plots that show an increase in their values.

(**a**) Initial (**b**) 15 days

**Figure 1.** *Cont*.

(**c**) 30 days (**d**) 60 days

**Figure 1.** SEM images showing surface corrosion of 3D-printed CoCrMo metallic implant during SBF exposure. (**a**) Initial (0 h); (**b**) after 15 days; (**c**) after 30 days; (**d**) after 60 days.

> Thus, the proposed formation mechanisms consist of Cr and Co ions, under the influence of the acidic environment of the SBF (mainly due to HCl) forming porous single crystals of Cr2O3 and Co3O4, as has been studied by [22].

> The surface analysis was performed, as in [23], and showed that, during the first two weeks of exposure, the material became porous due to the metal ion release into the SBF and, as a direct consequence, the porous surface increased almost 150% (Figure 1b). In the next two weeks, the formation of the oxide layer can be observed since most of the pores shrink almost 60% (Figure 1c). In the second month of exposure, the formation of porous single crystals of Cr2O3 and Co3O4 can be observed on the sample's surface (Figure 1d).

#### 3.1.2. XRD Assessment

An XRD assessment is usually used in materials science for determining the crystallographic structure of a material, and it is one of the most used techniques in the field. It uses X-rays to irradiate the material. Firstly, incident X-rays are sent towards the material and the reflected X-rays are measured in terms of intensities and angles.

In this study, X-ray diffraction (XRD) was used to establish if and/or what kind of oxides formed on the sample's surface. As stated before, a D8 Discover diffractometer (Bruker, Billerica, MA, USA) was used for sample analysis.

Figure 2 shows the existence of a protective layer on the polished surface during sample exposure to the SBF. On the basis of each sample's composition, different oxide layers form on the surface.

As shown in Figure 2, the cobalt-chromium (CoCr) protective layer, which is actually an alloy, is formed on the sample's surface. The layer does not have a crystalline structure (peak dimensions are rather modest) but it is well represented leading to the conclusion that it takes a while, but after its formation, the layer offers a protective interface between the SBF and the metallic implant.

**Figure 2.** XRD plot for the 3D-printed CoCrMo metallic implant.

#### 3.1.3. Mass Variation Assessment

The mass variation assessment as shown in Figure 3 was conducted by a gravimetric comparison of the sample's mass. Thus, the sample was weighed at its initial state, and then after 15, 30, and 60 days of exposure in the SBF. The aim of the assessment is to highlight the mass loss and/or gain while the sample is in the SBF. On the one hand, mass loss can be correlated with metal ion release, but this assessment is made by using ICP-MS, as described in Section 3.2.2. On the other hand, mass gain may occur due to the formation of metal oxides on the sample's surface, as emphasized in Section 3.1.2.

**Figure 3.** 3D-printed CoCrMo metallic implants mass vs. immersion time variation.

#### *3.2. SBF Characterization*

3.2.1. SBF pH and Electrical Conductivity Assessment

The pH value refers to the hydrogen ion concentration in a solution. The hydrogen concentration is closely related to corrosion phenomena, since a low pH value indicates an acidic environment suitable for surface corrosion. In addition, electrical conductivity shows the ion concentration in a solution; the higher the value, the higher the number of ions that are dissociated.

Therefore, pH and electrical conductivity monitoring represent an important assessment in order to evaluate the redox phenomena that occur while the sample is immersed in the SBF.

It can be observed in Figures 4 and 5 that the variation of pH values and electrical conductivity are closely interdependent, as well as with mass variation shown in Figure 3. Thus, observing the curve's allure which indicates pH and EC variation during the immersion period, in the first week of exposure, the SBF's pH varies significantly while the electrical conductivity increases. During this period, the sample gains mass due to the formation of an oxide layer on its surface. Redox phenomena during the first week of exposure are intense, as seen in Figure 4, and the concentration of metal ions is increased, as can be observed in Table 3.

**Figure 4.** SBF pH vs. immersion time variation.

**Figure 5.** SBF electrical conductivity vs. immersion time variation.

**Table 3.** Measured metal ion concentrations in the SBF.

Furthermore, this trend continues during the second week of exposure when the sample's mass decreases as metal ions are dissociated in the SBF, correlated with a continuous increase in electrical conductivity and pH.

After the second week of exposure, the amplitude of redox phenomena diminishes, as the protective layer forms on the sample's surface.

The most probable hypothesis regarding the allure of the plots shown in Figures 3–5 can be described as follows: The newly formed oxide layers of Cr2O3 and Co3O4 are still reacting to each other in order to form the protective layer of Co-Cr highlighted by the XRD plot. Thus, O2 is released back into the SBD in its ion form and quickly reacts with ionic H2 leading to a decrease in pH, observed in the last section of the plot shown in Figure 4. Nevertheless, metal ions do not dissociate since ICP-MS measurements show this aspect. In fact, Cr continues to be consumed since its concentration decreases (as shown in Table 3). The sample's mass decreases in the final interval of the plot shown in Figure 3, since O2 is released from the initial oxide layers of Cr2O3 and Co3O4 and also due to the continuous (even at slower rate) release of Co and Mo ions into the SBF.

#### 3.2.2. Metal Ion Release Assessment

ICP-MS is a powerful method for determining trace metal in given environments, due to its capability to measure very low concentrations of metal ions dissolved in fluids. This approach is well known, widely used by various research teams, and reported in frequently quoted scientific journals.

Basically, the ICP-MS method consist of determining the concentration of metal ions in a given solution. Thus, as required by the experimental design, 50 mL of the SBF used for sample immersion was collected and the solution was used for determining the nature and concentration of metal ions that might have dissociated from the sample into the SBF.

Given the data provided by Table 1, the following metal ions were monitored: Co, Cr, Mo, and Mn, since it is well known that large concentrations can lead to allergic and/or toxic reactions.

Table 3 shows the variation of metal ions dissociated from the sample in the SBF during the exposure period. Given the fact that the sample's exposed surface is relatively large, being almost impossible for a metallic implant to lose such a large area of protective layer, the immediate effect is the high concentration of metal ions released into the SBF. The aim of this assessment is to highlight the formation of a new protective oxide layer soon enough after the implant's surface is damaged (for various reasons, e.g., poor manipulation during implant surgery, mechanical shock, etc.) and to assess the concentration of metal ions dissociated into the SBF.

It can be observed that the formation of the oxide protective layer occurs after just 1 month of exposure. The correlation with pH, electrical conductivity, and mass variation plots highlights the fact that 3D-printed CoCr-based alloy may be used as orthopaedic implant. More mechanical and biocompatibility testing must be performed in the near future, but steps forward have been made.

Even though the metallic sample used in this study had a protective film on its surface, sample preparation removed it before immersing the sample in the SBF. It should be mentioned that the exposed surface was 909.65 mm2, and based on this, the corrosion rate and ion release rate could be calculated.

Table 4 shows the metal ion release rate from each sample and their corrosion rate. The values can be calculated using the following formula:

$$x\_i = \frac{m\_{i,\text{sol}}}{m\_{i,\text{sample}} \times S \times 16} \tag{1}$$

where *ci* is the ion concentration, *mi,sol* is the ion mass in solution, *mi,sample* is the ion mass in sample, *S* is the exposed surface and, 16 is the number of weeks of exposure.

**Table 4.** Calculated corrosion and ion release rates.

The corrosion rate can be calculated by using the following formula:

$$C\_{I} = \frac{m\_{m}}{m\_{i} \times S \times 16} \tag{2}$$

where *mm* is the measured mass of the sample, *mi* is the initial mass of the sample, *S* is the exposed surface and, 16 is the number of weeks of exposure

The corrosion rate is almost insignificant (e−5) since its measurement is grams over mm2 and week. Most implant materials are kept in the body for a maximum of 2 years (almost 100 weeks in bad cases), and the overall corrosion rate can be declared as 4.3‰. The more important aspect is the release rate of Mn, but as Table 3 shows, this release is only in the first 72 h of exposure, then, the trend is constant, and therefore it can be considered that Mn turns passive.

As a comparison, Table 5 lists the maximal allowed concentrations in the human body and the biological effects.

**Table 5.** Biological effects and maximal concentrations of metal ion accumulation in the human body.

#### **4. Conclusions**

After being immersed for 60 days in a blood plasma SBF and having a large area depleted from its protective oxide layer, the sample shows pitting corrosion in the first two weeks of exposure which is visible by SEM.

Although the material has a large area in direct contact with SBF, a new protective layer of CoCr alloy is formed relatively quickly, as highlighted by the XRD assessment. Nevertheless, during this period, metal ions are released in the SBF and the more concerning ones are Co and Cr.

The pH, electrical conductivity, and sample mass variation give a general idea about redox phenomena that occur.

A 3D printing technique was used for producing a CoCrMo-based alloy to be used as an orthopaedic implant by using the SLM technique and the incipient conclusion is that the produced CoCrMo alloy is suitable for such applications.

The concentration of metal ions dissociated from the biomaterial into the SBF solution was assessed by using the ICP-MS technique, which highlighted that the concentration of some metal ions varied during SBF solution exposure. The main conclusion is that those ions form a protective layer on the sample's surface mainly due to their reactivity and lead to the passivation of the exposed surface, thus, minimizing the concentration of ions that can dissociate into real body fluid.

**Author Contributions:** Conceptualization, R.M. and I.M.B.; methodology, R.M., L.C.C., R.E., and G.S.; software, A.P. and G.S.; validation, R.M., I.M.B., and R.E.; analysis, L.C.C., A.T.C., G.S., T.B., and E.P.; investigation, R.M., L.C.C., A.T.C., A.P., G.S., E.P., and T.B.; resources, R.M.; writing—original draft preparation R.M., R.E., and G.S.; writing—review and editing, R.M., R.E., and G.S., funding acquisition, R.M. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was carried out within POC-A1-A1.2.3-G-2015, ID/SMIS code P\_40\_422/105884, and the "TRANSCUMAT" Project, grant no. 114/09.09.2016 (subsidiary contract no. 1/D.1.5/114/24.10.2017), a project supported by the Romanian Minister of Research and Innovation.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Data sharing not applicable.

**Conflicts of Interest:** The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

#### **References**

## *Article* **Fatigue Assessment of Selective Laser Melted Ti-6Al-4V: Influence of Speed Manufacturing and Porosity**

**Unai Segurajauregi 1, Adrián Álvarez-Vázquez 2, Miguel Muñiz-Calvente 2,\*, Íker Urresti 1 and Haydee Naveiras 1**

**Abstract:** Additive Manufacturing represents a promising technology as an alternative to the conventional manufacturing process, with rapid and economic product development, as well as a significant weight reduction and a freeform design.  This paper presents a study of the influence of speed manufacturing and porosity in the fatigue behaviour of a Ti-6Al-4V alloy. To this aim, a numerical simulation of the expected porosity at different laser velocities is performed, together with a simulation of the residual stresses. These numerical results are compared with experimental measurements of residual stresses and a qualitative analysis of the porosities. Then, fatigue strength is experimentally obtained for two different laser speeds and fitted by a probabilistic model. As a result, the probabilistic S–N fields for different laser velocities are found to be similar, with scatter bands nearly coincident, drawing the conclusion that this effect is negligible in comparison with other concurrent ones, such as roughness or surface defects from manufacturing conditions, promoting crack initiation and premature fatigue failure.

**Keywords:** additive manufacturing; fatigue; titanium; selective laser melting

#### **1. Introduction and Motivation**

Additive Manufacturing (AM), formerly known as Rapid Prototyping (RP) (see ASTM F2792-12a [1]), has undoubtedly been increasing over the last two decades as a technology that is disrupting current manufacturing processes, and attracts interest from both industrial and academic perspectives [2–5]. Known also as 3-D printing, AM consists of a progressive consolidation of raw materials, such as powder or wire, in a layer-by-layer fashion, in an opposite approach to traditional manufacturing processes, which are typically based on machining block parts, that is, the subtraction or removal of material [6]. Moreover, this novel technology has several important advantages compared with traditional methods: an agile development product from Computer Aided Design (CAD) to fabrication; a significant reduction of weight in the final design (with potential reductions of up to nearly 50% [6,7]); and a geometric freedom that allows the production of parts otherwise not possible with conventional methods. Additionally, AM may lead to the reduction of carbon emissions compared with traditional manufacturing processes, due to the use of lighter weight parts [6,8].

Polymeric materials were originally preferred for producing additively manufactured parts [5], but nowadays both non-metallic (composites, ceramics) and metallic materials are usually employed. The titanium (Ti) alloys, in particular, are of great interest because

**Citation:** Segurajauregi, U.; Álvarez-Vázquez, A.; Muñiz-Calvente, M.; Urresti, Í.; Naveiras, H. Fatigue Assessment of Selective Laser Melted Ti-6Al-4V: Influence of Speed Manufacturing and Porosity. *Metals* **2021**, *11*, 1022. https://doi.org/10.3390/met11071022

Academic Editors: Atila Ertas, Adam Stroud and Matteo Benedetti

Received: 3 May 2021 Accepted: 22 June 2021 Published: 25 June 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

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of their increasing use in the aerospace industry due to their weight saving, operating temperature, corrosion resistance and compatibility with biological and composite materials. Unfortunately, their higher cost compared with other alternatives hitherto represents the most important limitation [6,9–12].

Three different technologies were developed to produce metal additively manufactured parts: Laser Beam Melting (LBM), Laser Metal Deposition (LMD) and Selective Laser Melting (SLM), the latter being the most commonly used [13,14]. In this process, the energy of the laser source is applied to melt powder, as a raw material, within a powder-bed layer. Then, the 3D geometrical design is built up by recoating a new powder-layer and subsequent melting [5]. Nowadays, these manufacturing technologies for developing structural alloys are particularly useful, since the intrinsic heat can be directly used to trigger the chemical reactions, such as those implied in precipitation hardening alloys [15,16].

There are different works in the literature that have focused on researching the influence of different additive manufacturing parameters on the fatigue performance of Ti-6Al-4V, such as the microstructure, the build direction, the residual stresses and the porosity. In the first case, Nalla et al. [17] have investigated the influence of the microstructure on both bimodal and coarser lamellar types, concluding that the latter improved the fatigue behaviour in the HCF zone, whereas Thijs et al. [18] studied the influence of the scanning parameters and scanning strategy on the microstructure during the SLM process. In the second case, Edwards et al. [6,19] presented a study on the effect of the build direction, revealing that the cracks oriented perpendicular to the build layers provide enhanced fatigue crack growth behaviour. Regarding the residual stresses, several researchers [20,21] have found that a high temperature pre-heating during the additive manufacturing process may reduce thermal gradients. Lastly, [22] evidenced that the failure initiation in SLM or EBM manufactured titanium alloys is governed by porosity and lack of fusion. Nevertheless, previous works have not investigated the influence of speed manufacturing, which would be conducted with different porosities and could imply different fatigue behaviours.

The aim of this paper is to study the influence of different laser velocities on the porosity of additively manufactured specimens of the Ti-6Al-4V alloy and to evaluate the fatigue performance associated with those porosities. Firstly, numerical simulations were developed to study the expected porosity considering different laser velocities, together with residual stresses inherent to the manufacturing process. Secondly, these numerical results were compared with experimental measurements of residual stresses and a qualitative analysis of the porosities. Thirdly, a tensile test was conducted on specimens produced at two different laser velocities in order to evaluate its influence on the mechanical properties. After that, a fatigue experimental campaign was carried out on specimens at two different laser speeds and the results were evaluated according to a probabilistic S–N model developed by [23], contrary to the fatigue assessment methodologies performed in the literature, based on deterministic S–N models [24–26], through the inherent and non-negligible scatter titanium fatigue tests [6,27–31].

The paper is structured as follows: in Section 2, the material and methods employed in this study are detailed, including the material and geometry selected (Section 2.1), the manufacturing conditions (Section 2.2) and the experimental procedures followed in the testing (Section 2.3). Section 3 details the numerical study of the porosity and the residual stresses together with the experimental results obtained. Section 4 is focused on both the tensile (Section 4.1) and fatigue characterization of two different laser velocities (Section 4.2). Section 5 presents an interpretation of the experimental results and, finally, Section 6 summarises the main conclusions drawn from this work.

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

This section describes the material and geometry of the specimens to be used in the experimental campaign in this work. Then, the manufacturing conditions are also detailed, distinguishing different batches of samples fabricated for porosity, tensile and fatigue characterization. Finally, the experimental testing procedures are exposed.

#### *2.1. Material and Geometry*

The specimens were produced using the SLM technique employing a titanium-based alloy, Ti-6Al-4V. The dimensions of the specimen are indicated in Figure 1. Note that the build direction *Z*-axis is indicated, starting at the support.

**Figure 1.** Dimensions (in mm.) of the Ti-6Al-4V specimens with build direction in *Z* and the support.

#### *2.2. Manufacturing Conditions*

All the specimens were provided by the manufacturer Optimus3D (Vitoria–Gasteiz, Spain) in the same orientation. The SLM parameters selected are detailed in Table 1. Three different batches of additively manufactured Ti-6Al-4V samples were produced with the following purposes:

*2.3. Testing and Characterization Procedures*

2.3.1. Porosity Measurement Procedure

The samples from *Batch 1* were sliced using a diamond disc cutter by removing a layer thickness of no less than 1 mm, in order to avoid the effect of the cutting process on the microstructure. Silicon carbide papers with different grades of 80, 240, 600, 1200 and 2500 (in this order) were used for the grinding with a continuous water stream for flushing the loose and abrasive particles. Then, a manual polishing process was applied to the surfaces using a diamond suspension of 9, 3 and 1 microns of particle size on a Remet LS1.

It is worth mentioning that neither the heat treatment nor the machining process were conducted on specimens before the experimental campaign, since the aim of this work was to study only the laser velocity, that is, without any additional varying concomitant effect that could mislead the interpretation of the experimental results. After that, the porosity was qualitatively analysed using both optical and scanning electron microscope (SEM). The entire surface of the specimens was analyzed with a resolution of 500 μm in order to identify the zones with larger pores. After that, different scanning zooms (10 and 50 μm) were applied to focus on the zones where pores had been observed. In cases where the resolution of 500 μm was not enough to identify pores in any part of the specimen, 50 μm was used to check the entire surface. It is important to remark that smaller pores could not be identified by the applied resolution. Still, the authors assumed that the influence of those pores on the fatigue life could be disregarded, compared to the pores identified in this study.

#### 2.3.2. Residual Stresses Measurement Procedure

The measurement of the residual stresses was performed by way of the hole drilling strain gage method according to ASTM E837–13a [32] using an MTS-300 RS measurement machine supplied by SINT Technology, as can be seen in Figure 2. The parameters selected included a drilling speed of 0.2 mm/min, a drill delay of 2–3 s and an acquisition delay set to 5–10 s. The RESTAN software (SINT Technology) was used to obtain the residual stresses.

**Figure 2.** Strain gage method in hole drilling for measurement of residual stresses.

2.3.3. Tensile and Fatigue Characterization Procedure

The tensile tests were developed according to EN 2002 [33]. A strain rate of 0.05 mm/min was used for yield stress *σys* at 0.2 % and 2 %/min for tensile strength *σr*. The displacement was measured using DIC equipment.

The fatigue tests were conducted for sinusoidal load at *R* = 0.1 and at ambient temperature, according to ASTM E466-07 [34], at a frequency of 6 Hz with stress ranging from 100 to 600 MPa, in a servohydraulic MTS Bionix. As previously mentioned, a total of 7 samples were tested for laser velocity *v* = 1200 mm/s and 8 samples for laser velocity *v* = 1900 mm/s.

#### **3. Study of Porosity and Residual Stresses**

This section presents the results obtained from numerical and experimental studies related to the estimation and measurement of specimens porosity and residual stresses.

#### *3.1. Numerical Study: Expected Porosity and Residual Stresses*

Two different numerical studies were conducted: a porosity simulation for each of the different laser velocities considered in *Batch 1*, and a finite element simulation for the estimation of the residual stresses.

In the first case, the tool known as "Additive Science", from the Additive Suite developed by ANSYS [35], was used to simulate the expected porosity with different laser velocities. The following constraints in the dimensions of the melt pool were introduced as an input (see Figure 3):

**Figure 3.** Dimensions of the melt pool.

Based on the constraints on the melt pool, the geometrical mean results provided by Additive Science for each velocity are summarised in Table 2.

**Table 2.** Numerical results of the geometrical dimensions of the melt pool for different velocities.

Once the dimensions of the melt pool had been estimated, the software provided the expected porosity for each of the different laser velocities considered, as shown in Table 3, together with the corresponding energy density in the SLM process. The porosity was directly estimated by ANSYS as a function of the energy density applied, which depends on the power, velocity and the dimensions of the melt pool. It is also worth mentioning that this software only considered the porosity that was due to a lack of fusion, which was predominant at high velocities, and the porosity that was due to spherical vaporization was discarded. Then, in order to identify the kind of porosity to be experimentally found, the causal relationship between the energy density and the kind of porosity proposed by Dilip et al. [36] was used: spherical pores correspond with energy density >60 J/mm3, a lack of

pores (or completely net) corresponds with energy density in the interval 55–60 J/mm3 and sharper pores correspond with energy density <55 J/mm3. Then, according to these limits, only the velocity 800 mm/s was expected to produce spherical pores, but with a lower percentage of porosity. On the contrary, those velocities higher than 1300 mm/s were expected to give sharper pores, with the porosity percentage increasing as much as the velocity increases. According to the simulation performed and the results obtained, middle velocities ranging from 1100 to 1200 were not expected to produce pores, but completely net structures. Finally, it is worth mentioning that the values reported in Table 2 regarding the percentage of porosity are not related to experimental measurements, but to results obtained by the Additive Suite tool developed by ANSYS, in combination with the work reported by Dilip et al. [36], which studied the influence of processing parameters on the evolution of melt pool and porosity in Ti-6Al-4V alloy parts fabricated by selective laser melting.

In the second case, the Additive Science tool allowed the residual stress to be numerically simulated, as can be seen in Figure 4, where the equivalent von Mises stress is depicted along the sample.

**Figure 4.** Residual stresses simulated upon the removal of the support.

#### *3.2. Experimental Results*

#### 3.2.1. Porosity

The porosity was qualitatively analysed using both optical and scanning electron microscopes (SEM) for each of the laser speeds considered in Batch 1, as can be seen in Figure 5. As is well known, the geometrical forms of the pores are expected to be heterogeneous depending on the laser velocity, which occurred in this case; ranging from spherical (Figure 5a) at low velocities and high laser power, to irregular and sharper (Figure 5d–f) for large velocities and low laser power, which is in accordance with the simulated results in Section 3.1. The former are known to be caused due to improper settings or processing parameters [12,37,38], while the latter are usually related in the literature to the argon gas entrapped during the manufacturing process [39–41]. Furthermore, the middle velocities were expected to produce a negligible porosity in comparison with the other velocities (see Table 2), which is corroborated in the micrographs in Figure 5b,c).

**Figure 5.** SEM micrographs for different laser velocities: (**a**) *v* = 800 mm/s, (**b**) *v* = 1100 mm/s, (**c**) *v* = 1200 mm/s, (**d**) *v* = 1500 mm/s, (**e**) *v* = 1700 mm/s, (**f**) *v* = 1900 mm/s.

#### 3.2.2. Residual Stress Measurement

Figure 6 illustrates the experimental results of the maximum and minimum residual stresses along the distance for both velocities considered in Batch 1. The maximum values of the residual stresses for the higher velocity evolve steadily along the distance, while for the lower velocity, a peak is present at the 0.1 mm distance. The same behaviour is observed in the case of minimum residual stress but at a lower order of magnitude. In general terms, there is an inverse trend between the development of residual stresses and the laser speed, that is, the values for both maximum and minimum residual stresses are higher for the larger velocity until a certain distance of almost 0.6 mm, where both trends tend to be equal.

**Figure 6.** Experimental values of maximum and minimum residual stresses for two different laser velocities: *v* = 1200 mm/s and *v* = 1900 mm/s.

#### **4. Study of Tensile and Fatigue Behaviour**

Once the study of the expected porosity and residual stresses and its comparison with the experimental results was performed, the tensile and fatigue characterization was conducted, and is now described in this section.

#### *4.1. Tensile Behaviour*

Engineering stress–strain curves for both velocities considered in Batch 2 are illustrated in Figure 7. As can be seen, the linear-elastic regime is approximately the same in both cases, but with a larger yield strength the lower the laser speed. In the plastic zone, however, the samples manufactured at a lower laser velocity produce a better tensile performance for the same strain value. Table 4 summarises the mechanical constants of the additively manufactured specimens for both velocities.

**Figure 7.** Engineering stress vs. strain curves for Ti-6Al-4V for two velocities: 1200 and 1900 mm/s.

**Table 4.** Tensile properties for additively manufactured Ti-6Al-4V obtained from the stress–strain curves in Figure 7.

#### *4.2. Fatigue Behaviour*

Finally, the fatigue assessment of manufactured Ti-6Al-4V samples from Batch 3 was conducted according to the probabilistic model developed by [23]. In this model, the *p*-percentile curves in the S–N field are given as the following Weibull distribution:

$$p = 1 - \exp\left[-\left(\frac{(\log \Delta \sigma - \mathbb{C})(\log N - B) - \lambda}{\delta}\right)^{\beta}\right],\tag{1}$$

with *B* as the horizontal asymptote for the stress, that is, the fatigue strength, *C* as the vertical asymptote for the lifetime, that is, no fatigue failure will occur below this limit, and *λ*, *δ* and *β* as the location, scale and shape Weibull parameters, such that (log Δ*σ* − *C*)(log *N* − *B*) > *λ*.

Figure 8 depicts the estimated S–N fields for both velocities considered in this batch, which were estimated easily with ProFatigue software [42]. As can be seen, the inherent scatter of fatigue results is non-negligible, thus a probabilistic model is more suitable than a deterministic one. The fatigue performance on both velocities exhibits the same behaviour with no differences in the lifetime cycle for different given stress ranges, and the scatter bands are approximately similar. For this reason, speed manufacturing seems to not have a significant effect on the fatigue performance of additively manufactured Ti-6Al-4V.

**Figure 8.** Resulting S–N curves from the Castillo–Canteli model for Ti-6Al-4V with velocities 1200 mm/s and 1900 mm/s.

Finally, the experimental campaign retrieved from [6], corresponding with a lower laser velocity of *v* = 200 mm/s, was also estimated according to the Castillo–Canteli model and is superposed in Figure 9 for comparison purposes. As a result, a wide range of laser speeds were considered from 200 to 1900 mm/s, providing robustness to the final conclusions drawn from this work. Indeed, though having a lower laser velocity, the resulting

S–N field is approximately the same with both previous velocities at 1200 and 1900 mm/s, enhancing the conclusion that there was a negligible effect on the fatigue performance.

**Figure 9.** Comparison of S–N fields for different laser velocities: *v* = 200 mm/s (data from Edwards and Ramulu [6]), *v* = 1200 mm/s and *v* = 1900 mm/s.

#### **5. Discussion**

In this work, different laser speeds were considered in additively manufactured Ti-6Al-4V alloy to evaluate their influence on the fatigue performance, where any other additional concurrent effect was avoided (heat treatment, machining surfaces, etc.). The experimental fatigue results were estimated according to a probabilistic model developed by [23], and the resulting S–N curves were approximately the same with the scatter bands being nearly coincident. In other words, the laser speed effect is negligible and no influence is found on the fatigue lifetime of Ti-6Al-4V specimens. Then, in order to enhance this conclusion, an external experimental campaign retrieved from [6]—at a lower velocity from those previously considered, that is, *v* = 200 mm/s—was also estimated and compared with previous results. Though having a wide range of laser velocities, from 200 to 1900 mm/s, the resulting S–N fields are approximately the same and the scatter bands are nearly coincident. However, this evidence does not allow for the conclusion that the effect of laser speed is negligible in terms of the fatigue lifetime of the specimens, since other concurrent effects could be covering the effects associated with the laser velocity and misleading the conclusions. The authors postulate that the high surface roughness obtained in the additive manufacturing process (*Ra* = 3.27, *Rz* = 15.83), together with the few differences in porosities depending on both laser speeds, lead to all cracks originating from the surfaces of the specimens, misleading conclusions about the potential effects of the porosities on the fatigue lifetime. Future work could consider the prior polishing of the specimens in order to improve the surface roughness, thus inducing the failure to occur from the pores.

The authors want to remark that the tensile and fatigue properties presented in this paper are related to only one SLM building direction. Taking into account that the material properties of metals manufactured by SLM cannot be considered isotropic that is, dependent on the testing direction—the conclusions of this paper must be taken into consideration only for the manufacturing direction in which the specimens were made, since different results could be achieved for other directions. Furthermore, the relationship between speed manufacturing and the resulting mechanical properties is not straightforward in additively manufactured samples, as other variables are implied, such as the melt pool depth and temperature required for complete melting or evaporation; the

effects of these variables were not considered in this study but are proposed as the basis for further work.

#### **6. Conclusions**

**Author Contributions:** Conceptualization, U.S., M.M.-C., Í.U. and H.N.; methodology, U.S., M.M.-C., Í.U. and H.N.; software, U.S., M.M.-C. and H.N.; validation, M.M.-C. and A.Á.-V.; formal analysis, U.S., M.M.-C. and A.Á.-V.; investigation, U.S., M.M.-C. and Í.U.; resources, U.S., Í.U. and H.N.; data curation, U.S. and H.N.; writing—original draft preparation, U.S. and A.Á.-V.; writing—review and editing, A.Á.-V., U.S. and M.M.-C.; visualization, A.Á.-V. and M.M.-C.; supervision, All authors; project administration, I.U. and M.M.-C.; funding acquisition, U.S. and Í.U. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research received no external funding.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

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

#### **References**

## *Article* **Deposition of Nickel-Based Superalloy Claddings on Low Alloy Structural Steel by Direct Laser Deposition**

**André Alves Ferreira 1,2,\*, Rui Loureiro Amaral 2, Pedro Correia Romio 1, João Manuel Cruz 3, Ana Rosanete Reis 1,2 and Manuel Fernando Vieira 1,2,\***

**Abstract:** In this study, direct laser deposition (DLD) of nickel-based superalloy powders (Inconel 625) on structural steel (42CrMo4) was analysed. Cladding layers were produced by varying the main processing conditions: laser power, scanning speed, feed rate, and preheating. The processing window was established based on conditions that assured deposited layers without significant structural defects and a dilution between 15 and 30%. Scanning electron microscopy, energy dispersive spectroscopy, and electron backscatter diffraction were performed for microstructural characterisation. The Vickers hardness test was used to analyse the mechanical response of the optimised cladding layers. The results highlight the influence of preheating on the microstructure and mechanical responses, particularly in the heat-affected zone. Substrate preheating to 300 ◦C has a strong effect on the cladding/substrate interface region, affecting the microstructure and the hardness distribution. Preheating also reduced the formation of the deleterious Laves phase in the cladding and altered the martensite microstructure in the heat-affected zone, with a substantial decrease in hardness.

**Keywords:** direct laser deposition; Inconel 625; parametrisation; microstructure; microhardness; preheating

#### **1. Introduction**

The laser-based additive manufacturing (LBAM) technologies applied in the production and repair of industrial components emerged in the late 1990s. Their use continues to extend to many industrial sector applications for components that operate in extreme conditions. LBAM technologies are unique and versatile in the manufacturing of parts with complex geometry, functionally graded or customised, producing an improvement in properties that can be used for a variety of industrial applications, such as within the aerospace, metallurgy, energy, and automotive industries [1,2]. Direct laser deposition (DLD) is an LBAM technology used for the additive manufacturing of metal parts, reconstructions, and repairs. DLD consists of the supply, through a nozzle, of metallic powder (or wire) processed by a focused laser, creating a melt pool on the surface of a metallic substrate. Several processing variables directly or indirectly affect the quality and structural integrity of components, dictated by solidification and metallurgical bonding [3].

DLD involves interactions between the laser beam, powder, and substrate in an environment with local protection from inert gas. Laser power, scanning speed, beam size, and powder feed rate are parameters that play a dominant role in cladding geometry (height, width, and length), dilution, and metallurgical properties. Clad overlapping, gas flow rate, powder flow profile, powder quality (size, shape, and density), and preheating are important secondary parameters [4,5].

**Citation:** Ferreira, A.A.; Amaral, R.L.; Romio, P.C.; Cruz, J.M.; Reis, A.R.; Vieira, M.F. Deposition of Nickel-Based Superalloy Claddings on Low Alloy Structural Steel by Direct Laser Deposition. *Metals* **2021**, *11*, 1326. https://doi.org/10.3390/ met11081326

Academic Editors: Atila Ertas and Adam Stroud

Received: 16 July 2021 Accepted: 20 August 2021 Published: 22 August 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

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The success of DLD depends on the selection of processing conditions that guarantee an effective bonding of the deposited material. This proper bonding produces adequate thermal delivery control, dense layers, a small heat-affected zone (HAZ), low dilution, minimal distortion, and good surface quality, with an attractive set of mechanical properties as well as resistance to wear and corrosion [6–8]. A controlled DLD process can replace conventional processes (i.e., electric arc welding and thermal spraying) in order to repair industrial components. Traditional approaches present drawbacks in component repairs, such as the time required, the limited thickness of the deposition layers, the low metallurgical bond, the formation of porosities and cracks, and the distortion of substrates (caused by overheating of the components). Therefore, it is of industrial interest to develop high-efficiency and -precision repair technologies to increase component life.

This additive manufacturing technology is considered the best strategy for reconstructing and repairing damaged components in terms of environmental benefit and economic feasibility. However, it is regarded as a complex process due to uncertainties in the quality and reliability of recovered industrial components [9], requiring further investigation to consolidate the results reported in this area. Although the equipment cost is high, DLD has successfully repaired dies, moulds, turbines, and gears. Adaptability for automation, ease of assembly of the laser on a CNC machine or robotic arm, and lower post-processing requirements are additional advantages of the DLD process [10,11].

DLD still has a way to go for broader industrial applications. Theoretical and experimental studies have developed relevant information about DLD; however, there are still many challenges, such as process optimisation, 3D reconstruction of highly complex structures, and substrate preheating effects, which need to be clarified. The production of wear-resistant claddings on low and medium carbon steel substrates is an application that can have many industrial applications, both in component repair and protective coating with a thick resistant layer.

Nickel-based superalloys are an excellent option for producing this wear-resistant layer. These alloys have been adopted in multiple applications due to their properties, such as mechanical behaviour at high temperatures, hardness, mechanical resistance, and good fatigue resistance, creep, and corrosion [12,13]. These properties are conferred by the structure and chemical composition of the alloy, mainly by elements such as molybdenum (Mo) and niobium (Nb), which form a solid solution in a nickel–chromium matrix [14]. While conventional manufacturing with these high-performance alloys has been difficult due to excessive tool wear and low material deposition rates, LBAM technologies can overcome these constraints, improving delivery times, and reducing manufacturing costs [15].

The use of nickel-based superalloys in DLD must consider the high cooling rates of the process, promoted by the localised thermal delivery induced by the laser beam, which can lead to the formation of metastable phases and the segregation of elements. These microstructural effects reduce the toughness and hardness of the coated components [16–19]. Preheating (PHT) is essential to control the cooling rate, minimising this effect. Increasing PHT temperature promotes the growth of the melt pool (depth and width), melting more substrate, thus increasing dilution [20–24]. PHT also prevents cladding delamination or cracking and reduces distortion and residual stresses due to the lower thermal gradient between the cladding and substrate [7,25,26].

In this study, Inconel 625, a nickel-based superalloy, powder was deposited on 42CrMo4 steel, while the process parameters were varied. 42CrMo4 steel is often used to produce components, such as gears and main shafts, and Inconel 625 is employed in the repair/remanufacturing of these components by SERMEC Group. Single layers were produced to evaluate the metallurgical bonding with a substrate; the influence of several processing parameters, such as laser power, scanning speed, and powder feed rate on the cladding quality, was evaluated considering the absence of cracks and structural imperfections. Preheating was performed on an optimised cladding condition in order to moderate

the microstructure and mechanical responses. Microhardness profiles of claddings were obtained and correlated with the microstructures.

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

*2.1. DLD System Setup*

A laser system, LDF 3000–100, was used to produce DLD claddings. The system has a high-power fibre-coupled laser diode (wavelength 900–1030 nm, depending on the power), with a nominal beam power of 6000 W. The machine concept is a KUKA KR90 R3100 industrial robot, based on a 6-axis industrial robot. All axes are connected to the robot and laser control units, which command the temperature of the melt pool as well as the laser power. The powder was fed during the deposition process by a coaxial feeding system, as illustrated in Figure 1. The substrate was preheated (PHT) to 300 ◦C with a manual gas system. The temperature control of the preheated substrate was done by a digital pyrometer, for verification of the uniformity of the substrate surface temperature distribution. Preheating is intended to decrease the cooling rate in the melt pool and HAZ regions as well as eliminate moisture. Tests were performed on substrates after the production of clads, with and without PHT, to evaluate susceptibility to cracking and eventual formation of metastable phases.

**Figure 1.** (**A**) Schematic representation of the direct laser deposition system with two feeders; (**B**) system in operation.

#### *2.2. Feedstock Powder and Substrate*

A nickel-based superalloy (MetcoClad 625 from Oerlikon), similar to Inconel 625, produced by the gas-atomised process, was used in this study. This powder was developed specifically for laser processing and presents a spherical morphology as well as particle sizes ranging between 45 and 90 μm. Figure 2 shows scanning electron microscopy (SEM) images of the morphology of the MetcoClad 625 powder.

**Figure 2.** SEM images of MetcoClad 625 powder.

42CrMo4 steel, in the quenched and tempered condition, was used as a substrate for the DLD deposition. Specimens with dimensions of 100 mm × 120 mm × 15 mm were prepared for depositions. This steel is classified as a low alloy structural steel with high mechanical strength and toughness as well as good fatigue resistance and machinability, being widely used in the manufacturing of critical industrial components, such as gears, automotive components, wing generators, and drilling joints for deep wells [27]. Its mechanical and chemical properties are described in standard EN 10,269 [28]. The chemical composition of the MetcoClad 625 (M625) powder and 42CrMo4 steel are shown in Table 1.

**Table 1.** Chemical composition of M625 powder and 42CrMo4 steel (wt.%).

#### *2.3. Process Parameters*

The main process parameters (laser power, scanning speed, and feed rate) were considered to evaluate the effect of processing conditions on clads. The evaluation of the clad quality depends on clad characteristics, namely, absence of cracks or pores, good metallurgical bond (interfaces without microcracks, pores, and fragile phases, exhibiting good wettability), and dilution between 15 and 30%. The values of the tested parameters are shown in Table 2. All results are representative of single-layer samples. The terminology ILP\_SS\_FR was used to identify the samples, being I—M625 powder (Inconel powder); LP—laser power (kW); SS—scanning speed (mm/s); FR—feed rate (g/min). Eighteen combinations of different processing conditions were tested, with and without preheating (see Table 3). Before deposition, the substrates were cleaned with pure acetone.

**Table 2.** Process parameters tested for M625 deposition on 42CrMo4 substrate.

In all tests, a spot size of 2.5 mm and an offset in the *Z*-axis of 0.2 mm were used. High-purity argon (99.99%), with a 5.5 L/min flow rate, was used as the shielding gas for minimising contamination of the melt pool during the DLD process. Samples with and without PHT were cooled in air.

#### *2.4. Mechanical and Microstructural Characterisation*

Samples from each deposition were cut for microstructural and mechanical characterisation using a metallographic cut-off machine with refrigeration in order to avoid substrate and cladding overheating. Samples were mounted in resin and polished down to a 1 μm diamond suspension. Kalling's N◦. 2 chemical etching (CuCl2—5 g, hydrochloric acid—100 mL, and ethanol—100 mL) was used to reveal the microstructures.

The measurements of the height, depth, and width of the claddings produced by the DLD technique were performed using a Leica DVM6 A 2019 digital microscope (DM) (Wetzlar, Germany). The Leica DM 4000 M optical microscope (OM) (Wetzlar, Germany) allowed for a microstructural analysis at low magnifications to evaluate, for example, the size of the heat-affected zone. A scanning electron microscope, FEI Quanta 400 FEG (ESEM, Hillsboro, OR, USA), equipped with energy-dispersive X-ray spectroscopy (EDX) (EDAX Genesis X4M, Oxford Instrument, Oxfordshire, UK) and electron backscatter diffraction (EBSD) (EDAX-TSL OIM EBSD, Mahwah, NJ, USA) was used for higher magnification observation and phase identification. For EBSD evaluation, the samples went through an additional polishing step, using a 0.06 μm silica colloidal suspension, for a superior surface

finish and to remove polishing-induced plastic deformations, allowing Kikuchi patterns to be obtained [29]. EBSD allows for the obtaining of information on microstructural characteristics with a small interaction volume and a high resolution, for which TSL OIM Analysis 5.2 software was used. For all raw data obtained by EBSD, a dilatation clean-up routine was performed, with a grain tolerance angle of 15◦ and a minimum grain size of 10 points.

Quantitative image analysis was employed on optical images using the ImageJ software, version 1.51p (National Institutes of Health, Bethesda, MD, USA).

Vickers microhardness tests gave the mechanical characterisation. The tests were performed using a test force of 300 g for 15 s in a Struers Duramin 5 (Struers Inc., Cleveland, OH, USA) Vickers hardness tester. Each hardness value corresponds to the average of three indentations.

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

#### *3.1. Processing Effects*

The quality of cladding was first evaluated by inspection with a digital microscope (DM). The microstructural analysis of all the claddings produced did not detect cracks, pores, or inclusions of significant dimensions, an essential requirement for obtaining highperformance deposits.

SEM characterisation confirmed the observations made by the DM. Figure 3 shows an SEM image of the cross-section of an M625 clad deposited on 42CrMo4 steel, representing the geometric aspects of cladding: height (h), width (w), depth (d), clad area (AC), melting area (AM), and wetting angle (θ). These geometric aspects were measured on all claddings using the ImageJ software. The results are shown in Table 3.

**Table 3.** Dimensional analysis of claddings produced by DLD.

This analysis of the geometry of the single-track deposits is critical as it provides information on process yield and cladding performance. For example, the contact angle is an essential parameter in assessing the quality of the cladding [30,31]. Higher beads can promote low wettability during the deposition of multi-tracks, hindering overlapping and generating discontinuities in the hatch spacing of the overlapping deposits, thus facilitating crack propagation. Typically, a contact angle greater than 90◦ is associated with lowerquality claddings [32]. Table 3 shows that depositions with higher heights have a high contact angle, as samples I1\_2\_15, I2\_2\_15, and I3\_2\_15 demonstrate.

**Figure 3.** Cross-section of a single layer of M625 deposited on 42CrMo4: d—depth; h—height; w—width; θ—contact angle; AC—clad area; and AM—melting area.

The results presented in Table 3 show that the analysed processing parameters (laser power, scanning speed, powder feed rate, and preheating) strongly influence the production of the cladding and its bonding to the substrate. The selection of a processing window that guarantees a metallurgically bonded clad with good material yield is a fundamental task. This selection is difficult since the mutual interaction of the various parameters is complex. Thus, it is common to apply combined parameters in the DLD process to obtain a more accurate relationship between processing parameters and the clad characteristics [33].

One of the most used complex parameters is powder deposition density (PDD), which expresses the combined effect of feed rate, scanning speed, and laser spot size (ϕ) (Equation (1)) [34–36]. Figure 4 exposes the linear growth of the cladding area with the increase in the PDD parameter.

$$\text{PDD} \left( \text{g}/\text{mm}^2 \right) = \text{FR} / \left( \text{SS} \times \text{\textquotedblleft} \right) \tag{1}$$

This parameter shows that by increasing the feed rate or decreasing the scanning speed, we can obtain claddings with a larger area, which was expected since both situations result in more powder supply in the same period of time. PDD is a valuable parameter, but this relationship is only valid for the cladding area and not for the total area of the deposit, including the area of the substrate that has been melted. This last area is vital because it affects the quality of the cladding.

**Figure 4.** Dependence of cladding area on the value of the PDD complex parameter.

Additionally, with the measured areas (Table 3), it was possible to calculate the dilution, which quantifies the relative amount of melted substrate during laser processing, according to Equation (2).

$$\text{Dilution} \left( \% \right) = \text{AM} / \left( \text{AC} + \text{AM} \right) \times 100 \tag{2}$$

A dilution ratio between 15 and 30% is sufficient to allow a good metallurgical bond between the substrate and the cladding. Higher dilution, which means a greater melting of the substrate, is undesirable since it reduces the deposition yield and induces considerable changes in the chemical composition of the deposited material, modifying the expected properties of the cladding.

The results obtained, and presented in Table 3, indicate that dilution increases with increased laser power, keeping the other processing variables constant. A laser power of 1.5 kW is enough to guarantee a dilution higher than 15% for almost all conditions (the only exception is the I1.5\_10\_15 sample).

Despite this apparent direct relation between laser power and dilution, the effect of other critical processing variables, namely the scanning speed and the feed rate, makes the establishment of relationships between processing condition and dilution complex. To overcome this difficulty, it is common to apply complex parameters, empirically adjusted, to the clad/substrate set under analysis to define the processing window [33,37,38].

Figure 5 shows a process window map, which associates laser power with the scanning speed and feeding rate ratio, as well as the dilution that is correlated with the laser power through complex parameter LP (SS/FR)0.5; it is represented by two curves, one for 15% and the other for 30%. Additionally, as was also considered in the map, a vertical line that corresponds to the acceptable limit for the wetting angle is present.

As seen in Figure 5, the shaded area, delimited by the previous conditions (dilution range, acceptable wetting angle, and process parameters), reveals the desirable practical manufacturing processing window of Inconel 625. In addition, the dilution increases proportionally with the increase in laser power and decreases with the ratio between scanning speed and feeding rate [39]. The flow of liquid metal in the melt pool is dominated by Marangoni's convection effect, caused by the surface tension gradient. As the temperature of the substrate increases, this effect becomes more evident. However, in practice, the surface tension gradient (γ) dγ/dT depends on both temperature (T) and composition. In this case, the most significant influence factor is the thermal gradient promoted between laser beam and substrate, as Le et al. [40] demonstrated, where the increase in the substrate temperature becomes more evident.

**Figure 5.** DLD process window map of M625 for additive manufacturing of single tracks. Red point represents the experimental I2\_6\_20 condition.

Laser power control allows the required good metallurgical integrity and dilution to be achieved [41], producing a cladding with a good metallurgical bond and uniform properties. Increasing the power of the laser promotes increased energy density as well as greater dilution and mixing between the substrate and the deposited powder, which is characteristic of the laser alloy process [42].

Considering the analyses performed, the I2\_2\_15, I2\_4\_15, I2\_6\_15, I2\_6\_20, I2\_10\_10, I3\_2\_15, and I3\_4\_15 conditions showed dilutions within the established range, between 16.6% and 29.9%, but a lower dilution value is preferable. Nevertheless, the I2\_2\_15, I2\_4\_15, and I3\_2\_15 samples presented poor wettability angles (107◦, 88◦, and 99◦, respectively) that may lead, in future, to claddings with overlapping defects between strands. On the other hand, the four remaining conditions presented good wetting angles, but taking into account not only the quality of the process but also its efficiency, condition I2\_6\_20 (Figure 5 red point) is the only one that allows for the possibility of manufacturing a larger cladding area and consequently a higher cladding high, which has an inverse linear relationship with the SS/FR ratio. Considering this evaluation, the I2\_6\_20 condition will be used to perform the analyses throughout the following sections.

#### *3.2. Microstructures and Mechanical Characterisation of the DLD Samples*

Figure 6 shows the cladding microstructure formed adjacent to the substrate by the deposition of M625 on 42CrMo4 steel. This microstructure consists of columnar grains, mainly dendrites, and cellular morphologies in a few regions.

As shown in Figure 6B, on solidification a continuous thin layer, < 10 μm, consisting of planar grains formed in the vicinity of the substrate. This morphology evolves into columnar grains with continued solidification of the cladding. This microstructure is consistent with the results of a similar study in which martensitic stainless steel is deposited [43]. The very high thermal gradient in the contact zone of the melt pool with the cold substrate contributed to the formation of planar grains. The microstructure evolves into a columnar/dendritic structure due to a rapid decrease in the thermal gradient when more material solidifies. Moreover, columnar grains grow perpendicular to the substrate/solidified material, i.e., in the opposite direction of the primary heat dissipation source, as is usual in DLD solidification [44]. As seen in Figure 7, this columnar/dendritic structure is the characteristic microstructure of the cladding.

**Figure 6.** Microstructure of cladding produced by the I2\_6\_20 cladding condition on the preheated substrate, showing (**A**) substrate interface cladding zone. (**B**) A higher magnification illustration of the interface.

**Figure 7.** SEM image of the I2\_6\_20 cladding microstructure.

Figure 8 shows a high-magnification SEM image and elemental mapping of the microstructure of the I2\_6\_20 cladding. This microstructure consists of the γ matrix, bright zones (surrounded by a segregation zone), and rounded dark particles. The elemental mapping indicates that the bright zones are rich in Mo and Nb, indicating the formation of the Laves phase located at the interdendritic region. Rounded dark particles are complex oxides dispersed in the γ matrix. Table 4 shows the chemical composition obtained by EDX analysis of the zones indicated in Figure 8.

**Figure 8.** SEM image and the elemental maps obtained by EDX microanalysis showing segregation zones associated with the Laves phase and complex oxides.

**Table 4.** Chemical composition (%) of the zones indicated in Figure 8.

The γ dendrites are formed during the solidification of nickel-based superalloys processed by DLD, segregating Nb and Mo into the liquid, thus creating the local conditions for forming the Laves phase. The final stage of non-equilibrium solidification thus gives rise to this microstructure consisting mainly of the γ matrix and Laves phase. A similar microstructure has been found in other nickel-based superalloy solidification studies [45–51].

Laves phase formation promotes the initiation and propagation of cracks, with a detrimental effect in mechanical response, reducing ductility, ultimate tensile strength, fracture resistance, and fatigue life [52,53]. Thus, this phase reduces the performance of Inconel 625, requiring control of morphology and distribution in the cladding.

A similar analysis was performed at the cladding/substrate interface, Figure 9. In the continuous thin layer adjacent to the substrate, characterized by planar grains (see Figure 6), the thermal gradient and growth rate are significantly different from those of the dendritic region, and the Laves phase was not detected.

Figure 9 shows that DLD deposition of M625 on the 42CrMo4 substrate led to cladding with a flat interface, with a thin continuous layer of planar γ grains, followed by γ dendrites and a dispersed Laves phase. As already mentioned, the appearance of the Laves phase in this region has a detrimental effect and should be minimized. Therefore, another cladding was performed with the substrate preheated to 300 ◦C. This procedure leads to a decrease in the thermal gradient of the deposited layer, influencing the microstructure. Preheating (PHT) increases the interdiffusion of constituents of M625 and 42CrMo4 at the bonding interface and, by decreasing the substrate cooling rate, can also affect the microstructure and properties in the heat-affected zone.

**Figure 9.** SEM image and elemental maps obtained by EDX microanalysis of M625 deposited on 42CrMo4 substrate. An EBSD image shows the phase distribution in this region: FCC phases in green and BCC phases in red.

The influence of PHT application on the cladding/substrate interface is shown in Figure 10. The thin layer of planar grains was not formed. Substrate heating significantly reduced the thermal gradient in the initial solidification phase of the melt pool, leading to the formation of dendrites throughout the cladding. Furthermore, it appears that PHT slightly increases the interdendritic spacing from 5–7 μm to 6–9 μm, measured by ImageJ software in Figure 6 and Figure 10, respectively. This observation confirms the decrease in the cooling rate allowing the growth of the interdendritic spacing. This feature is consistent with a study that indicates that lower cooling rates promote dendrite growth and decrease cellular grain formation in the initial stages of solidification [54].

**Figure 10.** Microstructure of cladding produced by the I2\_6\_20 cladding condition on the preheated substrate, showing (**A**) substrate interface cladding zone. (**B**) A higher magnification illustration of the interface.

Besides, the application of PHT has led to the reduction in Laves phases, this effect highlighted at the interface zone, as seen in Figure 11. A lower cooling rate effectively maintained the Nb and Mo in the γ matrix, avoiding segregation. Moreover, the elemental mapping and EBSD image in Figure 11 confirm a more intense interdiffusion at the interface, which is not as plane as it was without PHT, Figure 9. Diffusion of nickel, which is a gammaphase stabilizer, to steel is associated with a greater amount of residual austenite in this region. Thus, the application of preheating seems helpful for reducing the deleterious Laves phase and for enhancing metallurgical bonding.

The HAZ of 42CrMo4 steel is also a critical region, as substrate heating by the laser followed by rapid cooling leads to martensite formation, which can create cracks and allow for rapid crack propagation. Microstructural differences in HAZ caused by PHT were analysed by EBSD, as illustrated in Figure 12. This figure shows that PHT affects the HAZ microstructure, with larger (longer and wider) martensite laths caused by a slower cooling rate. Martensite with wider laths is associated with lower mechanical strength which, together with the more significant amount of residual austenite determined when using PHT, can reduce the brittleness of this region.

Figure 12 also shows that no preferential crystalline orientation was formed in the HAZ of claddings without or with PHT.

**Figure 12.** EBSD images of the I2\_6\_20/substrate interface region. (**A**) Without PHT; (**B**) with PHT.

#### *3.3. Microhardness Measurements*

The microhardness profile shown in Figure 13 shows that the hardness distribution in the cladding zone is uniform and independent of PHT, with an average hardness of 258 ± 2 HV and 253 ± 2 HV for the samples with and without preheating, respectively. A sharp transient zone with a pronounced hardness increase is measured in the HAZ of the sample without PHT (maximum hardness of 491 ± 23 HV). The application of PHT to the substrate before deposition produced a more uniform distribution of hardness in the HAZ (368 ± 25 HV), with a less sharp transient near the interface. Furthermore, the hardness peak has been eliminated, and the hardness values show less dispersion. These differences in hardness are attributed to changes in the microstructure induced by PHT, as seen in Figure 12, and its influence on chemical composition distribution, seen in Figures 9 and 11, and indicate that the HAZ region is less prone to cracking.

**Figure 13.** Microhardness profile of cladded samples applying the I2\_6\_20 condition with and without PHT.

#### **4. Conclusions**

The deposition of Inconel 625 claddings onto a 42CrMo4 steel substrate was performed using direct laser deposition (DLD) and varying processing conditions: laser power, scanning speed, feed rate, and preheating. Macro- and microstructural analysis, in addition to the hardness measurements, led to the following main conclusions:

**Author Contributions:** A.A.F., M.F.V., and A.R.R. proposed the methodology and results data analyses followed in this research; A.A.F., J.M.C., P.C.R., and R.L.A. carried out the experimental tests. All authors participated in the discussion of results and writing of the manuscript. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by FEDER through the Operational Programme for Competitiveness and Internationalization (COMPETE 2020), Projetos em Copromoção (project POCI-01-0247- FEDER-039848).

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

**Acknowledgments:** The authors give their thanks to the CEMUP (Centro de Materiais da Universidade do Porto) for expert assistance with SEM.

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

#### **References**

## *Article* **Comparative Study on Microstructure and Corrosion Resistance of Al-Si Alloy Cast from Sand Mold and Binder Jetting Mold**

**María Ángeles Castro-Sastre 1,\*, Cristina García-Cabezón 2, Ana Isabel Fernández-Abia 1, Fernando Martín-Pedrosa 2 and Joaquín Barreiro 1**

**Abstract:** This investigation is focused on the corrosion evaluation of an as-cast Al-Si alloy, obtained by two different casting methods: traditional sand casting and three-printing casting, using a binder jetted mold. The experimental results are discussed in terms of chemical composition, microstructure, hardness, and corrosion behavior of two different casting parts. The microstructure and composition of the sample before and after the corrosion tests was analyzed using light microscopy (OM), scanning electron microscopy (SEM), and energy-dispersive X-ray spectroscopy (EDX) and X-ray diffraction (DRX). The corrosion of the two processed castings was analyzed using anodic polarization (PA) test and electrochemical impedance spectroscopy (EIS) in an aerated solution of 3.5% by weight NaCl, similar to the seawater environment. After the corrosion process, the samples were analyzed by inductively coupled plasma/optical emission spectrometry (ICP/OES); the composition was used to determine the chloride solution after immersion times. The sample processed by binder jetting mold showed higher corrosion resistance with nobler potentials, lower corrosion densities, higher polarization resistance, and more stable passive layers than the sample processed by sand casting. This improvement of corrosion resistance could be related to the presence of coarse silicon particles, which decrease of cathodic/anodic ratio and the number of micro-galvanic couples, and the lower amount of intermetallic β Al-Fe-Si phase observed in cast alloy solidified in binder jetting mold.

**Keywords:** binder jetting; sand casting; aluminum alloy; microstructure; corrosion

#### **1. Introduction**

Techniques based on layer-by-layer fabrication have expanded enormously over the past two decades due to the availability of a variety of materials (metals, ceramics, polymers, and sands) and post-treatment procedures [1]. These technologies began for the manufacture of prototypes due to the geometric freedom and validation of physical prototypes during the product development cycle, reducing delivery time. [2,3]. However, several sectors (aerospace, automotive, foundry) have included this new technology in their production to make functional parts. In particular, casting industries immediately identified the benefits of additive manufacturing for the production of molds [4]. The ability to make parts layer by layer allows the designer to achieve one-piece molds, integrating sprue, runner, gate, and feeder into the mold design, and these unique geometries and complex designs are difficult to achieve through traditional casting processes. In addition, manufacturing time is reduced since the production of patterns, use of tools, and machining processes is avoided [5]. These aspects generate enormous advantages considering that more than 70% of all metal castings are produced by the sand casting process [6]. Moreover, sand consumption can be significantly reduced with the additive manufacturing process, which can be helpful for energy saving and reducing environmental pollution as it allows the selective deposition of an acid catalyst (e.g., furan).

**Citation:** Castro-Sastre, M.Á.; García-Cabezón, C.; Fernández-Abia, A.I.; Martín-Pedrosa, F.; Barreiro, J. Comparative Study on Microstructure and Corrosion Resistance of Al-Si Alloy Cast from Sand Mold and Binder Jetting Mold. *Metals* **2021**, *11*, 1421. https:// doi.org/10.3390/met11091421

Academic Editor: Atila Ertas

Received: 26 July 2021 Accepted: 31 August 2021 Published: 8 September 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

Binder jetting is one of the most suitable additive manufacturing technologies for the manufacture of casting molds. Several studies have focused on improving this process, applied to different materials. Tang et al. [7] optimized the binder jetting process for inkjet printing Ti6Al4V parts with excellent properties. Mariani et al. [8] demonstrated the feasibility of binder jetting technology to print WC–Co parts. This study revealed the importance of the characteristics of the initial powder and its influence on the mechanical properties of the printed parts. Other studies have focused on improving inkjet technology. Cheng et al. [9] investigated the process of driving waveform design for the multi-nozzle piezoelectric printhead. The developed procedure allows reducing the number of experimental tests to find suitable waveform parameters for different 3D printers and new materials. Other studies [10] have focused on improving the roller-type leveling mechanism in order to achieve desired layer thickness, smooth surface, and good color quality. Li et al. [11] conducted a study to solve the jamming problem of the printhead and to improve the dimensional accuracy of binder jetting technology. For this purpose, several water-based binders and gypsum composite powders were prepared, and the optimum binder-powder assembly was then determined through elementary adhesive testing and roller paving testing.

In the bibliography, there are several works focused on validating the advantages of applying additive manufacturing technologies to the manufacture of casting molds. Sneling et al. [12] took advantage of using 3D printed molds to cast parts with complex geometries, such as cellular structures and sandwich panels. In this way, they would avoid joints in the final part that would act as stress concentrators. Walker et al. [13] used binder injection technology to manufacture molds with built-in sensors that allowed monitoring the casting process by recording different parameters (temperature, pressure, humidity, gas chemistry) in different parts of the mold during the casting process.

Despite the advantages of applying additive manufacturing in the foundry industry, the process is not yet ready to completely replace the traditional casting process. A deeper understanding of the properties of printed molds and their effect on the quality of castings is still needed. The quality of castings depends on various factors related to the casting process, as well as the properties of the mold used. With regard to the casting process, it is necessary to consider the rate of heat removal from the metal to the mold, the pouring temperature, and the chemical composition of the starting alloy. Regarding the properties of the mold, it is necessary to know the mechanical properties, the surface finish, and the permeability fundamentally. Most of the studies have focused on studying the properties of printed molds using gypsum powder and sand in order to obtain castings of light alloys such as aluminum. McKenna et al. [14] analyzed the effects of temperature and curing time on the mechanical properties of the printed mold. The authors concluded that due to the loss of volatile compounds, permeability increased with increasing curing time. Vaezi et al. [15] studied the effect of layer thickness on mold properties. They observed that by decreasing the thickness of the layer and increasing the amount of binder added between particles bond, an improvement in tensile strength was achieved, although worse results were produced in the surface quality of the mold. Edwin et al. [16] studied the feasibility of printing plaster molds to produce aluminum metal castings. To improve the mechanical and thermal properties of the mold, they performed post-processing operations such as infiltration. The infiltrated molds had good resistance to heat, although the mold did not have enough mechanical strength to withstand the metallostatic pressure due to the molten metal.

There have been some studies that analyzed the quality of the castings obtained from 3d printed molds but the number of studies is limited. Snelling et al. [17,18] conducted a comparative study using two different materials to produce the 3DP mold—ZCorp's plaster vs. ExOne's silica—in order to obtain A356 aluminum alloy casting. In this study, they focused on analyzing various properties of the casting, including roughness, density, hardness, porosity, and microstructure. They concluded that the material of the mold significantly influences the quality of the final part. Different factors, such as the amount

of binder used to bind the particles of the mold, the grain size, the permeability, and the composition of the mold material, modify the material properties of resultant castings. These properties are related to the microstructural aspects of the final casting matrix (secondary dendrite arm separation, grain size), secondary intermetallic phases (size, morphology, distribution, and quantity), and porosity [19]. Specifically, the microstructure of the Al-Si alloys, the object of study in this article, present in their microstructure a matrix of α-aluminum dendrites as the main component and several phases. Among the phases present are eutectic Si particles, several secondary intermetallic phases such as Al2Cu, Mg2Si, and several Fe-bearing phases (AlFeSi, AlFeSiMn). The secondary phases (eutectic Si, Al2Cu, Mg2Si, AlFeSi, and AlFeSiMn) provide specific properties to the alloy and can affect the mechanical and corrosion behavior of the material.

Al-Si-Cu alloys are passive materials, however, they show high susceptibility to pitting corrosion in chloride solution; this has been related to the presence of secondary phases because most of these particles are more noble than Al-matrix [20]. These micro- galvanic cells cause the dissolution of the aluminum matrix.

Part of this secondary phase are Si particles. The low solubility of silicon in the aluminum matrix originates the precipitation of Si particles. The size, morphology, and distribution of these particles play an important role in the mechanical and corrosion behavior of Al-Si-Cu alloys [21]. The cathodic behavior of Si within the aluminum-rich matrix acting as an anode results in the formation of micro-galvanic couples that cause localized corrosion [22]. The influence of silicon morphology and its effect on the area ratio in galvanic corrosion has been analyzed by Tahamtan et al. [23]. Thick flake Si particles with sharp edges that increase the cathode/anode ratio greatly increase the density of the corrosion stream [24].

Also, the electrochemical properties of Al-Si-Cu alloys are affected by morphologies, type, and distribution of intermetallic phases that are, in turn, a function of alloy composition and cooling rate [25,26]. Intermetallic particles such as Al2CuMg, and Mg2Si, are anodic with respect to the Al matrix and preferably corrode with respect for the surrounding Al matrix [27]. However, the particles of intermetallic phases such as Al2Cu and especially the iron-rich phases (β-Al-Fe-Si and α-Al-Fe-Mn-Si) are mostly cathodic to the Al and cause the dissolution of surrounding eutectic Al matrix adjacent to these intermetallic phases [28]. The iron content determines the type and morphology of these iron-rich intermetallic phases. The plate shape phase, such as βAl5FeSi, is the most abundant, and the greater the amount of iron, the longer and wider the needles are. This phase is very hard and fragile and induces a high susceptibility to localized corrosion. The presence of manganese favors the "Chinese script" form, such as αAl15(Fe,Mn)3Si2, which improves corrosion resistance compared with those that are manganese-free [29–31].

The studies at the microstructural level, mechanical properties and corrosion behavior of Al-Si alloys produced by the traditional sand casting process are known. However, little information on these aspects has been studied for castings using molds fabricated by additive manufacturing technology. This work aims to focus on this gap to investigate the microstructure and corrosion behavior of the AlSiCu alloy. For this purpose, the alloys were solidified in traditional sand molds and binder jetting molds.

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

#### *2.1. Materials and Molds Manufacturing*

For this study, the aluminum alloy Al-Si-Cu (EN-AC 4600) was used. The chemical composition was analyzed by optical emission spectroscopy (OES), and the results are shown in Table 1.

The metal pouring conditions were the same for both molds. The aluminum smelting was carried out in a Hobersal PR/400 (Hobersal, Barcelona, Spain) furnace at 770 ◦C, and the pouring temperature was 750 ◦C.

The material used to print the binder jetting mold was a commercial powder plaster, CaSO4.1/2H2O (VisiJet PXL Core, from 3DSystems, Rock Hill, SC, USA) with a purity of 80–90%, according to the manufacturer and a commercial binder solution (VisiJet PXL Clear) based on water with 2% Pyrrolidone. The characterization of the powder and the binder were carried out by the authors in a previous work [32]. The mold was manufactured using binder jetting technology with a Project CJP 660Pro machine (3D Systems, Rock Hill, SC, USA). Binder Jetting technology is a powder-based process that uses a liquid binder to selectively bind the powder particles that are spread on a work platform. This process is repeated layer by layer until the part is completed according to the CAD design.

To manufacture the sand casting mold, commercial silica sand (Petrobond) was used. A pattern, including the part and filling system, was manufactured by fused deposition modeling (FDM) technology, using Ultimaker 2+ machine (Ultimaker B.V., Utrecht, Dutch). Figure 1 shows the geometry of the part and the methodology used for the manufacture of both molds, obtaining the casting parts, and metallographic and corrosion study. The part has a rounded geometry with a sufficient area-volume ratio to analyze changes in the internal structure due to heat transfer and cooling rate.

**Figure 1.** The methodology used for the casting process and analysis of the casting samples (metallographic and corrosion study).

#### *2.2. Microstructural Characterization and Porosity*

The metallographic study of both castings was carried out. First, the castings were longitudinally cut and then smaller samples were cut from two different zones of the casting parts, the center and the outer zone. The samples obtained were cold encapsulated and polished (as standard). To reveal the microstructure of the samples, the polished specimens were then etched using Keller's reagent (2.5 mL HNO3; 1.5 mL HCl, 1 mL HF and 95 mL H2O) for 12 s. Once polished and attacked, the study of different parameters was carried out, resorting to the use of different techniques.

For the microstructure and porosity study, optical and electronic microscopy was used. The images were taken from each sample (center sample and outer sample) corresponding to the castings obtained from both molds (sand mold and binder jetting printed mold). For the porosity study, a total of 30 images were taken with light microscopy (LM) randomly distributed on the surface of each sample, and then ImageJ 2016-0205 software (National Institute of Health, Bethesda, MD, USA) was used to determine the area occupied by the pores.