Cobalt-based alloys, such as cobalt-chromium-molybdenum (CoCrMo), are known for their high mechanical strength and find extensive applications in the biomedical field such as manufacturing of tools, dental components, and orthopedic implants. The longevity of the CoCrMo alloy in service is intricately linked to its resistance to corrosion and wear. Specifically, tribocorrosion can contribute to material loosening; therefore, it is essential to explore surface treatments for cobalt-based alloys as a means to enhance their wear resistance, ensuring the prolonged durability of the material. This study provides novel insights into the bio-tribocorrosion resistance of the borided CoCrMo alloy when immersed in calf serum, emulating the synovial fluid. Two distinct microstructures of boride layers were examined in this research: (1) a CoB–Co2B layer formed through powder-pack boriding and (2) the borided surface underwent diffusion annealing to completely dissolve the CoB, resulting in a monophasic layer (Co2B). Following the ASTM G119-09 procedure, the total material loss (T), encompassing both material loss due to wear (WC) and corrosion (CW), was determined using a linear reciprocating ball-on-flat tribometer equipped with an electrochemical cell. Test results indicated that the presence of CoB–Co2B and Co2B layers on the CoCrMo alloy increased bio-tribocorrosion resistance approximately 2.4 times and 1.3 times, respectively, compared to the non-treated CoCrMo alloy. A dominant wear regime was observed for the borided surface exposed to diffusion annealing and the non-treated CoCrMo alloy, whereas the borided CoCrMo alloy exhibited a corrosion-wear regime. Clearly, these findings highlight the capability of the cobalt boride layer to improve the performance and extend the service life of the CoCrMo alloy in biomedical applications.

The cobalt-chromium-molybdenum (CoCrMo) alloys are commonly used in orthopedic implants for artificial joints due to their excellent mechanical, corrosion-resistant, wear-resistant, and biocompatible properties. However, over their service life, these alloys can experience from detrimental corrosion and tribocorrosion damage, leading to the release of metal ions into the body fluids. This, in turn, can result in undesired toxicity and aseptic loosening.1,2

The corrosion, tribocorrosion, and bio-tribocorrosion tests on CoCrMo alloys using different solutions such as Hank, Tyrode, and artificial saliva have been widely investigated. However, fluids more representative of the human body composition, particularly those rich in proteins—dominant molecules that significantly influence tribocorrosion and bio-tribocorrosion performance, have garnered increasing attention. For this reason, several authors have explored the impact of biofilm, primarily composed of proteins, on the tribocorrosion performance of CoCrMo alloys.3–5 

Namus and Rainforth3 observed a significant reduction in the material loss rate during bio-tribocorrosion in CoCrMo alloy after cathodic polarization of approximately 0.9 V. This decrease was attributed to the interference of the relatively thick adsorbed proteinaceous layer on the surface with the galvanic couple formed between the bare metal surface area and the passivated region. Taufiqurrakhman et al.4 established that a higher protein concentration has a positive effect in reducing the volume loss owing to mechanical wear in CoCrMo alloys. However, the 75% fetal bovine serum (FBS) concentration leads to an increase in volume loss due to corrosion compared to the 25% FBS concentration. In the work of Wang et al.,5 the results demonstrated that proteins have dual effects on the CoCrMo alloy subsurface evolution: forming a multilayered structure and inducing severe subsurface deformation. The tribofilm protected the passive film from scrapping, and the passive film reduced (or even suppressed) stacking fault annihilation, blocking access to the metal surface. Furthermore, the tribofilm, a carbon-rich layer formed by proteins, exhibited a lubricating effect, thereby reducing the friction coefficient (CoF). This alteration shifted the location of the maximum frictional shear stress from the top surface to beneath the surface.

In recent years, researchers have studied surface modifications to enhance the tribocorrosion properties of CoCrMo alloys.6–10 Various techniques, such as plasma electrolytic oxidation, magnetron sputtering, ion implantation, vapor deposition, and others, have been utilized to create layers that improve the surface characteristics of CoCrMo alloys.

Notably, the powder-pack boriding process has been used to enhance the tribocorrosion resistance of CoCrMo alloys when immersed in Hank's solution.11,12 The resulting cobalt boride layer consists of CoB–Co2B (orthorhombic and tetragonal crystalline structures, respectively) combined with CrB, Cr2B, and Mo2B compounds with a well-defined diffusion zone under the layer.13,14 The tribocorrosion findings indicated that the CoB–Co2B layer on the CoCrMo alloy reduced the material loss rate by approximately 16%, primarily due to the wear-corrosion synergism when compared to the non-treated CoCrMo alloy.11 Furthermore, Cuao-Moreu et al.12 demonstrated the impact of laser surface texturing (LST) in combination with boriding on the tribocorrosion behavior of CoCrMo alloys. The textured + borided specimens reduced the wear-corrosion synergism effect by 1.5 times compared to the untreated specimens.

Moreover, in a recent study,15 cytotoxicity tests indicated satisfactory results regarding the survival and proliferative activity of human fibroblasts and Vero cells in contact with the boride layer. Consequently, it is crucial to investigate the tribocorrosion resistance of the boride layer in contact with an oxygen-rich medium containing proteins, enzymes, and other substances.

In this study, novel insights into the bio-tribocorrosion behavior of both borided CoCrMo alloys and borided CoCrMo alloys subjected to a diffusion annealing process were obtained. Bio-tribocorrosion tests were conducted using a linear reciprocating tribometer coupled with the standard three-electrode corrosion cell employing a calf serum solution to emulate synovial fluid. The experimental procedures followed the ASTM G119-09 standard. The results showed that the microstructural properties of the borided surface, characterized by the presence of CoB–Co2B or Co2B layers, along with the experimental conditions of corrosion in the absence of wear (C0), corrosion in the presence of wear (CW), wear in the absence of corrosion (W0), and total material loss due to wear and corrosion synergy (T) had varying effects on the bio-tribocorrosion performance. Finally, the bio-tribocorrosion results of borided specimens were compared with those obtained from the non-borided CoCrMo alloy.

Commercial CoCrMo alloy specimens with 20 mm of diameter and 4 mm of thickness were prepared for the powder-pack boriding (C1) and the powder-pack boriding subjected to a diffusion annealing (C2). The nominal chemical composition of the specimens complying with the Micro-Melt BioDur Carpenter CCM Alloy standard is depicted in Table I. Prior to the boriding process, the samples were ground to 2000-grit finish using SiC papers, polished with alumina powder of 0.3–0.05 μm, and cleaned with ethyl alcohol in an ultrasound bath.

TABLE I.

Nominal chemical composition (wt. %) of CoCrMo alloy.

0.14 max 
Cr 26–30 
Mo 5–7 
Ni 1.0 max 
Mn 1.0 max 
Fe 0.75 max 
Si 1.0 max 
0.25 max 
Co balance 
0.14 max 
Cr 26–30 
Mo 5–7 
Ni 1.0 max 
Mn 1.0 max 
Fe 0.75 max 
Si 1.0 max 
0.25 max 
Co balance 

The C1 and C2 experimental conditions followed the established powder-pack boriding procedures for the CoCrMo alloy.11,13,14,16 In the case of C1, the specimens were placed inside a closed cylindrical AISI 304 L steel container filled with a powder mixture consisting of 20% B4C, 10% KBF4, and 70% SiC. Boriding was carried out at 1000 °C for 5 h to obtain a 33 μm thick CoB–Co2B layer. Subsequently, the container was removed from the furnace and allowed to cool to room temperature. After boriding, the specimens underwent a diffusion annealing process inside an electrical muffle with an argon atmosphere at 1000 °C for 2 h. Once the exposure time was completed, the borided specimens were cooled to room temperature inside the muffle. This procedure was carried out with the aim of achieving a single Co2B layer with a thickness of approximately 30 μm.

The C1 and C2 specimens were then prepared metallographically to reveal the microstructure of the cobalt boride layer. First, the specimens were ground with SiC emery sandpapers and polished with alumina powder of 0.3–0.05 μm until a mirror finishing was achieved. Finally, the polished specimens were etched with 3% HCl in H2O2 during 5 min. The cross-section measurements of the CoB–Co2B (C1) and Co2B (C2) layers were acquired by optical microscopy (GX51 Olympus, USA) with the aid of image pro plus software (MEDIA CYBERNETICS, USA). The phases on the borided surfaces for C1 and C2 were inspected by x-ray diffraction (D8 ADVANCE Bruker, USA) using a glancing angle of 0.5° with Cu-Kα radiation (λ = 0.179 nm) and 2θ angle ranging from 20° to 80°.

The bio-tribocorrosion experiments were conducted using calf serum in three conditions: C1, C2, and non-treated CoCrMo alloy (RM). In accordance with the ISO 14242-1:2014 procedure (Implants for surgery—Wear of total hip-joint prostheses—Part 1: Loading and displacement parameters for wear-testing machines and corresponding environmental conditions for test), the calf serum (In Vitro, MEX) was diluted by mixing it with deionized water at a molar ratio of 1:1. This procedure aimed to achieve a protein mass concentration in the solution ranging from 23 to 28 g L1. To prevent bacterial growth, sodium azide (NaN3 ∼ 0.2 wt. %) was added, and disodium ethylenediaminetetraacetic dihydrate, EDTA disodium salt (∼7.45 g L1) was introduced to mitigate calcium precipitation in the calf serum solution.17 The nominal composition of calf serum is summarized in Table II.

TABLE II.

Nominal elemental composition of calf serum.

Bicarbonate 25–30 mM 
Calcium 2.12–2.72 mM 
Chloride 100–108 mM 
Phosphorus 2.87–4.81 mM 
Potassium 3.5–4.7 mM 
Sodium 134–143 mM 
Amino acids 20–51 mg ml–1 
Glucose 650–966 mg ml–1 
Uric acid 30.5–70.7 mg ml–1 
Water 930–955 mg ml–1 
Albumin 37.6–54.9 mg ml–1 
lgG 6.4–13.5 mg ml–1 
Fibrogen 2–4 mg ml–1 
Bicarbonate 25–30 mM 
Calcium 2.12–2.72 mM 
Chloride 100–108 mM 
Phosphorus 2.87–4.81 mM 
Potassium 3.5–4.7 mM 
Sodium 134–143 mM 
Amino acids 20–51 mg ml–1 
Glucose 650–966 mg ml–1 
Uric acid 30.5–70.7 mg ml–1 
Water 930–955 mg ml–1 
Albumin 37.6–54.9 mg ml–1 
lgG 6.4–13.5 mg ml–1 
Fibrogen 2–4 mg ml–1 

For the bio-tribocorrosion assays on C1, C2, and RM specimens, the ASTM G119-09 procedure (Standard Guide for Determining Synergism Between Wear and Corrosion) was used. The arithmetic surface roughness (Ra) of C1, C2, and RM specimens was measured to be approximately 0.05 μm using non-contact profilometry (Contour GT-K 3D Bruker, USA).

The bio-tribocorrosion experiments were performed using a ball-on-flat tribometer coupled with a three electrode-electrochemical cell (UMT-2 Bruker, USA) controlled with a potentiostat (Bruker, USA) and leaving a specimen exposed area of 1.5 cm2. A schematic representation of the tribocorrosion system11,12,15,18 is presented in Fig. 1, while Table III summarizes the experimental parameters employed during the bio-tribocorrosion tests.

FIG. 1.

Schematic representation of the tribocorrosion system.

FIG. 1.

Schematic representation of the tribocorrosion system.

Close modal
TABLE III.

Experimental parameters for the bio-tribocorrosion tests.

Normal load 20 N 
Relative wear distance 100 m 
Track distance 2.5 mm 
Sliding rate 10 mm/s 
Ball diameter (Al2O35 mm 
Potential rate EOCP ± 300 mV 
Scan rate 1 mV/s 
Electrolyte Calf serum 
Temperature 15–30 °C 
Reference electrode Ag/AgCl 
Counter electrode Pt wire 
Work electrode C1, C2, and MR 
Initial contact pressure max. Al2O3 – RM ∼2.25 GPa 
Initial contact pressure max. Al2O3 – CoB ∼2.57 GPa 
Initial contact pressure max. Al2O3 – Co2∼2.42 GPa 
Normal load 20 N 
Relative wear distance 100 m 
Track distance 2.5 mm 
Sliding rate 10 mm/s 
Ball diameter (Al2O35 mm 
Potential rate EOCP ± 300 mV 
Scan rate 1 mV/s 
Electrolyte Calf serum 
Temperature 15–30 °C 
Reference electrode Ag/AgCl 
Counter electrode Pt wire 
Work electrode C1, C2, and MR 
Initial contact pressure max. Al2O3 – RM ∼2.25 GPa 
Initial contact pressure max. Al2O3 – CoB ∼2.57 GPa 
Initial contact pressure max. Al2O3 – Co2∼2.42 GPa 

From the experimental results, the material loss rates C0 (corrosion in the absence of wear), CW (corrosion in the presence of wear), W0 (wear in the absence of corrosion), and T (total material loss due to wear-corrosion synergy) were estimated. A brief summary of each procedure is presented in Fig. 2.

  1. Prior the C0 and CW tests, the C1, C2, and RM specimens were immersed in the calf serum at room temperature for approximately 3 h to stabilize the open circuit potential (EOCP), whereas the polarization resistance experiments were performed within the EOCP ± 300 mV versus (Ag/AgCl) with a scan rate of 1.0 mV s−1 (the scan rate was chosen due to the approached steady-state and small disturbance of charged current).

FIG. 2.

Overview of the key steps in the bio-tribocorrosion tests.

FIG. 2.

Overview of the key steps in the bio-tribocorrosion tests.

Close modal

Similarly, during CW tests, sliding occurred continuously from the beginning to the end of polarization with the normal load and sliding rate specified in Table III.

For C0 and CW tests, the electrochemical parameters were obtained connected the tribocorrosion system to a computer running a CETR software (Bruker, USA). The material loss results were obtained by replicating three experiments for each specimen and using Faraday's Law,
C = K j corr EW ρ ( m m 3 m m 2 y r 1 ) .
(1)

Here, C = C0 or CW, K is a constant ( 3272 mm g / A cm yr ), jcorr is the corrosion current density (μA cm−2), the equivalent weight (EW) corresponding to C1 (∼12.89), C2 (∼12.43), and RM (∼24.73), and ρ is referred to the density ( C 1 7.37 g c m 3, C 2 8.10 g c m 3, RM 8.29 g c m 3).

Notably, after C0 tests, the corrosion products on the surfaces of C1, C2, and RM specimens were characterized by x-ray diffraction (D8 ADVANCE Bruker, USA) technique, operating in a Bragg-Brentano configuration, using a glancing angle of 0.5° with Cu-Kα radiation and 2θ angle ranging from 0° to 90°. The diffraction peaks on the surfaces were analyzed by x'pert highscore plus V.2.2d software (PANalytical, UK).

  • A potentiostatic polarization test (W0) at −1 V of the EOCP (recorded by a potentiostat; UMT-2 Bruker, USA) was performed to avoid the influence of corrosion on wear. The cathodic protection was initiated at the beginning of sliding and continued until the end of sliding (refer to wear variables in Table III), while the friction coefficients were obtained connected the tribocorrosion system to a computer running a CETR software (Bruker, USA). The tests were carried out at least thrice for each specimen.

The material loss rate for C1, C2, and RM was estimated based on the following:
W 0 , T = [ Wrem At ] × 8760 ( m m 3 m m 2 y r 1 ) ,
(2)
where Wrem is the volume removal over the worn track, A is the specimen exposed area during the test (1.5 cm2), and t is the test duration (∼4 h). The material loss rate was evaluated by measuring the dimensions (width and depth) at center of the worn track employed a non-contact profilometer (Contour GT-K 3D Bruker, USA).

Additionally, the role of corrosion-wear synergism in the total material loss (T) was evaluated, by considering the variations of the EOCP (recorded by the potentiostat; UMT-2 Bruker, USA) before, during, and after sliding. The test began with approximately 1 h stabilization period in the calf serum solution before the start of sliding period. During sliding, the friction coefficients were acquired by the CETR software (Bruker, USA), and then, the test ended with a stabilization period of approximately 1 h. There were three replications of the test performed for each specimen. The total material loss due to wear and corrosion (T) for the tested materials was determined using Eq. (2); the Wrem was determined by measuring the depth and width from five 2D profiles at the center of the worn area using a non-profilometer (Contour GT-K 3D Bruker, USA).

The failure mechanisms, debris products, and chemical composition over the surface of worn tracks developed after T and W0 tests were analyzed using scanning electron microscopy (SEM) and energy dispersive x-ray spectrometry (EDS) at different accelerated voltages (JEOL JSM IT-100, JPN).

The OM cross-sectional images for C1 and C2 are presented in Figs. 3(a) and 3(c), respectively. The microstructure and composition of cobalt boride layers on C1 and C2 have been previously reported.11–14,16,19 A brief summary follows. Figure 3(a) illustrates the microstructure of the CoB–Co2B layer on the surface of C1. The total boride layer thickness was around 33 μm (CoB = 20 ± 1 μm, Co2B = 13 ± 1 μm) with the presence of a diffusion zone (∼25 μm) beneath the layer. The presence of alloying elements in the substrate, such as Cr and Mo, hinders the cobalt boride layer growth; these elements precipitate into the layer, competing with Co and forming boron compounds (CrB and Mo2B), as evident in the x-ray diffraction pattern [Fig. 3(b)].

FIG. 3.

Cross-sectional OM images of borided CoCrMo alloy: (a) C1 and (b) C2. The XRD patterns of borided surfaces: (c) C1 and (d) C2.

FIG. 3.

Cross-sectional OM images of borided CoCrMo alloy: (a) C1 and (b) C2. The XRD patterns of borided surfaces: (c) C1 and (d) C2.

Close modal

The diffusion annealing changed the boride layer's microstructure for C2 [Fig. 3(c)]. In the absence of boriding media, the B diffusion occurs in CoB–Co2B and Co2B–substrate interfaces that progress toward the equilibrium state. The diffusion annealing temperature enhances the B diffusion providing the energy needed to break the bonds in the CoB crystalline structure, which disappears almost completely due to the chemical reaction between Co and B to form Co2B.20,21 Similarly, in the Co2B-diffusion zone, the B atoms moving into the substrate interact with existing boron compounds, extending the diffusion zone. The resulting boride layer microstructure consisted mainly of Co2B, Cr2B, and Mo2B according to the x-ray pattern of Fig. 3(d); the Co2B layer thickness was around 29 μm, while the depth of the diffusion zone was approximately 34 μm.

Before the C0 and CW tests, the EOCP as a function of time was measured for the C1, C2, and RM: the EOCP value of RM shifted to more positive value (0.071 ± 0.010 V) in comparison to the value estimated for C2 (−0.126 ± 0.004 V) and C1 (−0.360 ± 0.003 V). The EOCP of RM and C2 revealed more electropositive values, which denoted a minor corrosion susceptibility than C1.

For C0 and CW tests, the over potential (η) against the current density (j) was plotted to obtain the linear polarization resistance (Rp) for the entire set of experimental conditions as shown in Fig. 4(a). In addition, the electrochemical parameters such as Ecorr (corrosion potential) and jcorr (corrosion current density) were obtained from the Tafel extrapolation method [Fig. 4(b)]. Notably, jcorr is obtained as follows:
j corr = B Rp ,
(3)
B = b a b c 2 .303 ( b a + b c ) ,
(4)
where B is the Stern-Geary constant and ba and bc are the anodic and cathode Tafel slopes.
FIG. 4.

(a) Linear polarization plots and (b) Tafel curves after C0 and CW tests.

FIG. 4.

(a) Linear polarization plots and (b) Tafel curves after C0 and CW tests.

Close modal

During C0 tests, the chemical interactions between the calf serum and the C1 and C2 surfaces led to the formation of a protein film. Initially, electrons released by B and Cr attracted water molecules from the calf serum to the calf serum-borided surface interface [Fig. 5(a)]. Subsequently, B and Cr ions interacted with hydroxides ( O H ) to initiate an oxidation-reduction process [Fig. 5(b)]. Thus, the following simultaneous reactions22 were carried out on the surfaces of C1 and C2 [Eqs. (5)–(13)]:

FIG. 5.

Schematic representation of oxides and hydroxides formation [(a) and (b)] on the borided surface and [(c) and (d)] proteins' adhesion on the borided surface.

FIG. 5.

Schematic representation of oxides and hydroxides formation [(a) and (b)] on the borided surface and [(c) and (d)] proteins' adhesion on the borided surface.

Close modal
Chemical reactions developed between the calf serum and B,
B B 3 + + 3 e ,
(5)
B 3 + + 3 O H B ( O H ) 3 ,
(6)
B ( O H ) 3 + 1 2 B 2 O 3 B 2 O 3 ,
(7)
B 2 O 3 + H 2 O B ( O H ) 3 .
(8)
Chemical reactions formed between the calf serum and Cr,
Cr C r 2 + + 2 e ,
(9)
C r 2 + + 1 4 O 2 + 1 2 H 2 O C r 3 + + O H ,
(10)
C r 3 + + 3 O H Cr ( O H ) 3 ,
(11)
Cr ( O H ) 3 + 1 2 C r 2 O 3 C r 2 O 3 ,
(12)
C r 2 O 3 + H 2 O Cr ( O H ) 3 .
(13)

It is worth noting that additional reactions can take place between the calf serum and the boride layer, such as those involving Co and Mo. However, B and Cr exhibited a higher affinity for reacting with O H than Mo and Co, primarily due to their atomic properties.

The interaction between proteins and surface can involve various charge bond combinations, including hydrophilic, hydrophobic, positively, and negatively charged interactions, due to the heterogeneous nature of the surface. Despite the original structural arrangement of the protein, its rearrangement can lead to the development of more active sites for protein-surface contacts. The rate of protein adsorption (thickness growth) depends on the charges of the proteins and the surface. While interactions between negatively charged proteins and negatively charged surfaces do occur, the adsorption kinetics is generally slower when compared to other charge combinations.23 

Then, a rich layer of oxides and hydroxides was formed on the surface of C1 and C2 [Fig. 5(c)]. The outer hydroxide layer on these surfaces exhibited an affinity for the electrical charges of proteins, such as albumin, based on factors like molecular orientation and surface characteristics, including chemical composition, hydrophobicity, and topography. As time progresses, the interactions between the surface and proteins can lead to the development of organometallic bonds of various types, including covalent, ionic, and coordination links, forming an adsorbed biofilm [Fig. 5(d)].

Moreover, the chemical reactions between the RM and calf serum are depicted in Eqs. (9)–(13); a schematic representation for the development of biofilm on the surface of RM is presented in Fig. 6.

FIG. 6.

Schematic representation of oxides and hydroxides formation [(a) and (b)] on the RM and [(c) and (d)] proteins' adhesion on the RM.

FIG. 6.

Schematic representation of oxides and hydroxides formation [(a) and (b)] on the RM and [(c) and (d)] proteins' adhesion on the RM.

Close modal

In this sense, after C0 tests, the presence of glucose (C6H12O6) was detected on the surfaces of C1, C2, and RM as part of the biofilm [refer Figs. 7(b), 7(d), and 7(f), respectively].

FIG. 7.

SEM images and XRD patterns after C0 tests on the surfaces of [(a) and (b)] C1, [(c) and (d)] C2, and [(e) and (f)] RM.

FIG. 7.

SEM images and XRD patterns after C0 tests on the surfaces of [(a) and (b)] C1, [(c) and (d)] C2, and [(e) and (f)] RM.

Close modal

Table IV presents the electrochemical parameters obtained after C0 and CW tests. In the C0 tests, both C1 and C2 exhibited increased jcorr values compared to RM. Based on the XRD patterns after C0 tests on the surfaces of C1 and C2 [Figs. 7(b) and 7(d), respectively], the reduction of the corrosion resistance was attributed to the presence of B2S3 and CrPO4 species in a crystal cluster form, resulting from the interaction of the calf serum and the cobalt boride layers,11,15,24 while the recorded measurements of RM corrosion resistance were related to the presence of a well-developed passive layer,25 consisting of CrO, CrO3, and Cr2O3 as illustrated in Fig. 7(f).

TABLE IV.

Electrochemical parameters obtained by means of the potentiodynamic polarization technique during C0 and CW tests.

C0CW
SpecimenEcorr (mV)jcorr (μA cm−2)Rp (kΩ cm2)Ecorr mVjcorr (μA cm−2)Rp kΩ cm2
C1 −504 ± 14 1.63 ± 0.13 8 ± 0.6 −472 ± 10 3.63 ± 0.2 4 ± 0.3 
C2 −281 ± 20 0.04 ± 0.002 322 ± 13 −319 ± 29 0.13 ± 0.01 100 ± 8 
RM −229 ± 19 0.01 ± 0.001 1085 ± 62 −321 ± 13 0.14 ± 0.01 90 ± 6 
C0CW
SpecimenEcorr (mV)jcorr (μA cm−2)Rp (kΩ cm2)Ecorr mVjcorr (μA cm−2)Rp kΩ cm2
C1 −504 ± 14 1.63 ± 0.13 8 ± 0.6 −472 ± 10 3.63 ± 0.2 4 ± 0.3 
C2 −281 ± 20 0.04 ± 0.002 322 ± 13 −319 ± 29 0.13 ± 0.01 100 ± 8 
RM −229 ± 19 0.01 ± 0.001 1085 ± 62 −321 ± 13 0.14 ± 0.01 90 ± 6 

On the other hand, the C2 specimen revealed a significant decrease of the jcorr (∼0.04 μA cm−2) compared to C1 specimen (∼1.63 μA cm−2). Previous findings indicate that diffusion annealing applied to borided CoCrMo alloy redistributes and enriches the Cr and Mo content in the Co2B surface layer. This redistribution contributed to heightened protection against corrosive ions and promoted the enhancement of the passive film. The passive film, in turn, offered protection and less formation of the B2S3 specie.26 

Furthermore, the significant differences between the jcorr values of C1 (CoB–Co2B) and C2 (Co2B) can be explained by the texture fiber orientations on the outer surface of the boride layers. The presence of {020} and {021} planes (loose planes) on the CoB layer (C1) led to decreased corrosion resistance due to the low atomic density and high surface energy in these atomic planes. Conversely, the preferred texture fibers {002} orientation in C2 resulted in low surface energy on its closed-packed plane, contributing to increased corrosion resistance.26,27

From the CW results summarized in Table IV, the physical interaction of surfaces in relative motion, which generates phenomena of adhesion, abrasion, erosion, etc., can lead to particle and passive film detachment, resulting in material loss due to wear accelerated corrosion. In all cases, the jcorr values were higher compared to those obtained in C0. The change in jcorr suggests that under rubbing conditions, the passive film on the surfaces of C1, C2, and RM was locally removed, exposing the materials to the calf serum solution. Consequently, the rubbed areas experienced an increase in oxidation rate due to the loss of passivity. Furthermore, as per Wimmer et al.,28 the larger size of albumin hinders the repassivation processes during corrosion in the presence of wear.

Figure 8(a) displays the CoF curves obtained on the surfaces of C1, C2, and MR when slid against the alumina ball during the W0 and T tests.

FIG. 8.

(a) CoF behavior during W0 and T experiments. (b) Wear depth profiles after W0 and T tests. (c) The EOCP behavior before, during, and after the T experiments.

FIG. 8.

(a) CoF behavior during W0 and T experiments. (b) Wear depth profiles after W0 and T tests. (c) The EOCP behavior before, during, and after the T experiments.

Close modal

For the entire set of experimental conditions, the CoF exhibited a gradual increase with sliding distance, which coincides with static friction preceding the motion of the bodies and with the interlocking of asperities in the very early period of sliding. Subsequently, some particularities were observed between the friction behavior during W0 and T for the different specimens as shown in Table V.

TABLE V.

Wear depth, Wrem, and CoF results after T and W0 tests.

W0T
SpecimenWear depth (μm)CoFWrem (mm3 × 10–4)Wear depth (μm)CoFWrem (mm3 × 10–4)
C1 4.35 ± 0.4 0.35 ± 0.01 16.2 ± 1.1 2.66 ± 0.1 0.32 ± 0.02 7.1 ± 0.7 
C2 2.77 ± 0.2 0.34 ± 0.01 8.37 ± 0.8 4.53 ± 0.4 0.33 ± 0.01 14.0 ± 1.0 
RM 3.90 ± 0.1 0.37 ± 0.01 10.1 ± 1.0 6.40 ± 0.1 0.35 ± 0.02 20.2 ± 1.5 
W0T
SpecimenWear depth (μm)CoFWrem (mm3 × 10–4)Wear depth (μm)CoFWrem (mm3 × 10–4)
C1 4.35 ± 0.4 0.35 ± 0.01 16.2 ± 1.1 2.66 ± 0.1 0.32 ± 0.02 7.1 ± 0.7 
C2 2.77 ± 0.2 0.34 ± 0.01 8.37 ± 0.8 4.53 ± 0.4 0.33 ± 0.01 14.0 ± 1.0 
RM 3.90 ± 0.1 0.37 ± 0.01 10.1 ± 1.0 6.40 ± 0.1 0.35 ± 0.02 20.2 ± 1.5 

During W0, slightly higher CoF values in the steady state were obtained in C1 (∼0.35), C2 (∼0.34), and RM (∼0.37) compared to the CoF results in the T conditions (C1 ∼ 0.32, C2 ∼ 0.33, and RM ∼ 0.35). This phenomenon can be explained by the absence of an oxide film during W0 tests, resulting in decreased biofilm adsorption on the C1, C2, and RM surfaces. Consequently, there is an increase in material transfer to the counterpart (alumina ball) when compared with the balls slid on C1, C2, and RM surfaces during T tests [refer to Figs. 9(a)9(f)].

FIG. 9.

OM images of the alumina balls rubbing against the surfaces of (a) C1, (b) C2, and (c) RM after the T experiments. OM images of the alumina balls under W0 tests sliding on the surfaces of (d) C1, (e) C2, and (f) RM.

FIG. 9.

OM images of the alumina balls rubbing against the surfaces of (a) C1, (b) C2, and (c) RM after the T experiments. OM images of the alumina balls under W0 tests sliding on the surfaces of (d) C1, (e) C2, and (f) RM.

Close modal

Figure 8(b) illustrates the wear depths on the surfaces of C1, C2, and RM after W0 and T tests obtained by non-contact profilometry technique. In addition, the wear depths and Wrem results are summarized in Table V.

According to the results in Table V, the wear depths and Wrem estimated for C2 and RM decreased when subjected to cathodic polarization (W0) compared to those under bio-tribocorrosive (T) conditions. This occurred due to a significant reduction in the oxidation reaction, suppressing corrosion and resulting in less corrosive wear. However, in W0 tests, the wear depth and Wrem results for C1 showed an opposite behavior compared to C2 and RM.

As demonstrated by Akonko et al.,29 the application of cathodic potential leads to the hardening of the surface specimens due to hydrogen embrittlement, and this effect is more pronounced when wear is involved. Although hydrogen embrittlement exists in C2 and RM, the susceptibility to failure through cracking during sliding is heightened in C1 due to the presence of a brittle CoB on its outer surface.21,30 During W0 tests, this vulnerability is further exacerbated by hydrogen embrittlement, leading to increased material removal in C1 compared to C2 and RM.

On the other hand, during the T tests, the EOCP behavior was monitored before, during, and after sliding, as depicted in Fig. 8(c). The EOCP serves as an indicator of the potential state of both the unworn surface and the wear track surface.31 

At the onset of sliding, a notable sudden decrease in EOCP to more negative values was observed for C1, C2, and RM, indicating damage to the passive film during wear. This could induce a galvanic coupling between the passive layer (and biofilm) and the bare specimen material, subsequently, causing local dissolution of the base material. For the RM, the EOCP dropped to values ranging −310 and −330 mV; in the case of C2, the EOCP reduced to values around−310 mV, while for C1, the average EOCP was approximately −500 mV. Once the sliding was ended, the EOCP for C1, C2, and RM was restored to the values prior to wear because of the formation of adsorption layer (biofilm) and repassivation of the worn surfaces.

Building upon the T results from Table V, the wear depth of the RM was approximately 2.4 and 1.4 times larger than C1 and C2, respectively. This difference was attributed to the presence of the CoB–Co2B and Co2B layers on the CoCrMo alloy.

On the other hand, the Co2B layer on the C2 specimen increased its wear depth during sliding in calf serum (about 1.7 times larger) with respect to that of the CoB–Co2B layer on the C1 specimen. A reasonable explanation for these findings is related to the microstructure and mechanical properties on the outer surface of C1 and C2 specimens.32 The CoB layer on the outer surface of C1 displays a hardness around 27 GPa (Young's modulus ∼374 GPa), while the Co2B layer in C2 exhibits a hardness of 20 GPa (Young's modulus ∼324 GPa). Furthermore, the difference in hardness (and Young's modulus) between CoB and Co2B can be attributed to the chemical characteristics of their crystal structures.33,34 The CoB exclusively forms covalent bonds between B and Co atoms (Co–B) and two B atoms (B–B), while Co2B maintains two types of bonding: (1) covalent involving B and Co atoms (Co–B), and 2 B atoms (B–B) and (ii) metallic which includes two Co atoms (Co–Co).

Following the W0 and T procedures, SEM and EDS analyses were conducted on the wear tracks, as shown in Figs. 10 and 11, respectively.

FIG. 10.

SEM images and EDS spectra in the worn track and unworn surface after W0 experiments: (a) C1, (b) C2, and (c) RM.

FIG. 10.

SEM images and EDS spectra in the worn track and unworn surface after W0 experiments: (a) C1, (b) C2, and (c) RM.

Close modal
FIG. 11.

Representative SEM-EDS images of the worn track and unworn surface under T experiments: (a) C1, (b) C2, and (c) RM.

FIG. 11.

Representative SEM-EDS images of the worn track and unworn surface under T experiments: (a) C1, (b) C2, and (c) RM.

Close modal

The worn tracks on the surfaces of C1, C2, and RM after W0 tests are depicted in Figs. 10(a)10(c). For all the circumstances, failure mechanisms such as grooving and adhered particles were observed, indicative of three-body abrasive wear. The EDS analyses conducted on the worn tracks of C1 and C2 (zones II and IV, respectively) revealed, basically, the presence of C (attributed to the biofilm remnants), B, and alloying elements of the substrate such as Cr, Co, and Mo. The elemental composition of RM's worn track (zone VI) showed the presence of C from the calf serum solution, a little content of O, along with Si, Cr, Co, and Mo from the substrate. Additionally, EDS spectra on the unworn surfaces of C1, C2, and RM (zones I, III, and V, respectively) indicated that the weight percentage of C and O was relatively higher when compared to their contents on the worn tracks. This observation revealed that all surfaces were covered by an adsorbed biofilm and oxide layer before sliding at the cathodic potential.

After the T experiments, the worn track of C1 [Fig. 11(a)] revealed two-body abrasive wear as the primary failure mechanism. This was evidenced by the presence of grooving and the absence of adhering particles within the worn track. Conversely, in C2 and MR, grooves of abrasive wear and attached particles (associated with corrosion and wear products) were observed in various areas of the sliding tracks [Figs. 11(b) and 11(c), respectively], indicative of three-body abrasive wear, accelerated by corrosion.35 EDS analyses were introduced to evaluate the elemental compositions of C1, C2, and RM worn and unworn surfaces and to confirm the presence of organic material following testing. Particularly, the high C peaks and the presence of O in zones I, III, and V was a good indicator of organic material and passive reformation on the unworn surfaces of C1, C2, and C3 [Figs. 11(a)11(c)]. In addition, the EDS revealed elements of calf serum solution combined with alloying elements of the substrate. However, due to the continuous removal of the biofilm during mechanical contact, a reduction in C concentration within the worn tracks (zones II, IV, and VI) was detected for all experimental conditions. This observation indicates that tribochemical reactions occurred during sliding.

The ASTM G119-09 standard procedure establishes a relationship between the total material loss (T), the synergistic component (S), and the damage resulting from pure mechanical wear without corrosion (W0) and the electrochemical corrosion rate without wear (C0). This relationship is expressed as follows:
T = C 0 + W 0 + S ( m m 3 m m 2 y r 1 ) ,
(14)
T = C 0 + W 0 + Δ C w + Δ W c ( m m 3 m m 2 y r 1 ) .
(15)

Here, Δ C w = C w C 0 is the change in the corrosion rate due to wear and Δ W c = W c W 0 is the change in the wear rate due to corrosion. C w is the total material loss rate due to corrosion and W c is the total material loss rate due to wear. The results of the synergistic interactions are summarized in Table VI.

TABLE VI.

Parameters of the synergistic contribution.

mm3 mm−2 yr−1 ( × 10−3)
SpecimenTW0C0CWSWCCW/WC
C1 12 ± 1 23 ± 2 9.35 ± 0.75 17.15 ± 1.08 −20 −5 3.467 
C2 21 ± 2 12 ± 1 0.20 ± 0.01 0.67 ± 0.05 13 0.052 
RM 29 ± 2 15 ± 1 0.11 ± 0.01 1.34 ± 0.09 14 28 0.048 
mm3 mm−2 yr−1 ( × 10−3)
SpecimenTW0C0CWSWCCW/WC
C1 12 ± 1 23 ± 2 9.35 ± 0.75 17.15 ± 1.08 −20 −5 3.467 
C2 21 ± 2 12 ± 1 0.20 ± 0.01 0.67 ± 0.05 13 0.052 
RM 29 ± 2 15 ± 1 0.11 ± 0.01 1.34 ± 0.09 14 28 0.048 

The material loss rate resulting from the synergistic effects of wear and corrosion (T) was approximately 1.7 times lower in C1 compared to C2 and roughly 2.4 times lower than the value observed in RM. Moreover, the material loss due to corrosion and wear was ∼1.3 times lower in C2 in comparison with RM. During T, the results revealed an accelerated corrosion wear in C2 and RM, related to the presence of corrosion and wear products entrapped in the tribological pair. These products acted as abrasive particles.36 In addition, the estimated value of synergistic component (S) for C1 was negative, indicating the presence of an antagonistic synergy due to the corrosion products. These products developed on the surface of C1 provided a better contact surface, resulting in a decrease in the T value compared to the results in W0 (refer to Table VI).

On the other hand, the transition between bio-tribocorrosion regimes was obtained by employing the C w / W c ratio according to the following criteria:37 
Wear dominated regime : C w W c < 0.1 ,
(16)
Wear - corrosion regime : 0.1 C w W c < 1 ,
(17)
Corrosion - wear regime : 1 C w W c < 10 ,
(18)
Corrosion dominated regime : C w W c 10.
(19)

Based on the C w / W c results presented in Table VI, C2 and RM exhibited a wear-dominated regime due to the significantly greater acceleration of wear-induced corrosion (CW ∼ 3.4 C0 and CW ∼ 12.2 C0, respectively) compared to wear accelerated by the effect of corrosion (Wc ∼ 1.1 W0 and Wc ∼ 1.9 W0, respectively). The results of C1 revealed a corrosion-wear regime, attributed to wear accelerated corrosion (CW ∼ 1.8 C0); the formation of corrosion species on the surface of C1 had a more pronounced impact on the reduction of material due to the wear effect, resulting in a negative value for W c.

Recently, the tribocorrosion behavior of borided CoCrMo alloy in Hank's solution was reported.11 The tribocorrosion experiments were conducted with an alumina ball of 5 mm, applying 20 N of load over a 2.5 mm stroke distance. The authors found a total material loss rate resulting from the synergistic effects of wear and corrosion (T) around 25.4 × 10−3 mm3 mm−2 yr−1. In this study, when comparing the total material loss (T) results of C1 with those obtained for borided CoCrMo alloy immersed in Hank's solution, the presence of biofilm reduced wear and corrosion synergy by approximately two times (see Table VI). As demonstrated by Taufiqurrakhman et al.,4 a higher protein concentration in the calf serum positively influences the reduction of material loss due to the synergy of corrosion and wear. During the mechanical mixing and chemical actions, the adsorbed and deposited proteins are denatured and transformed into a tribofilm that protects the passive film from being mechanically removed. Additionally, the passive film and tribofilm can inhibit the corrosion reactions by acting as the mass and electron transfer barriers.5 In this scenario, the biofilm served as a boundary lubricant, contributing to a decrease in material loss,38 complementing the advantageous superficial mechanical properties of CoB–Co2B layer on the CoCrMo alloy.

Kanka et al.39 investigated the role of calf serum on the tribological behavior of borides grown on cobalt-28chromium-6molybdenum (Co–28Cr–6Mo) alloy. Their wear experiments were conducted in a ball-on-disk tribometer using an alumina ball (6 mm in diameter), using an applied load of 20 N and a sliding distance of 400 m. The results indicated that the presence of albumin proteins in the calf serum solution played a crucial role in establishing boundary lubrication, thereby mitigating wear. The borided cobalt alloy showed a reduction in volume loss of approximately 0.003 mm3 compared to the untreated Co–28Cr–6Mo alloy, which experienced a volume loss of 0.008 mm3.

Finally, the results of this study showed that the bio-tribocorrosion resistance can be enhanced by the presence of the CoB–Co2B or Co2B layers on CoCrMo alloy. However, for use in biomedical applications, it is essential to assess the proliferative activity and survival of various human cells on the surface of the borided CoCrMo alloy through in vitro cytotoxicity assays. Likewise, evaluating the mineralization ability of cells, such as osteoblasts adhered to the borided surface, is important for determining the osteoconductive and osteoinductive properties of cobalt boride layer, essential for classifying it as a bioactive layer.

The results of this study are summarized as follows:

  • The bio-tribocorrosion resistance of cobalt boride layers immersed in calf serum solution was estimated in this work following the ASTM G119 procedure. The boride layers were formed through the powder-pack boriding process (C1) and boriding followed by a diffusion annealing process (C2) on the surface of CoCrMo alloy. The obtained bio-tribocorrosion results were then compared with those derived from the non-treated CoCrMo alloy (RM).

  • Powder-pack boriding on the surface of CoCrMo alloy, at 1000 °C for 5 h, resulted in the development of a CoB–Co2B layer with a total thickness of approximately 33 μm. Subsequent diffusion annealing completely dissolved the CoB layer on the borided CoCrMo alloy surface, resulting in the formation of a monophasic Co2B layer with an approximate thickness of 29 μm layer.

  • The results from corrosion in the absence of wear (C0 tests) indicated that the formation of B2S3 and CrPO4 on the surfaces of C1 and C2 increased the material loss rate compared with the RM. Additionally, the increase in jcorr values during corrosion in the presence of wear (CW tests) for C1, C2, and RM was associated with the ongoing passive film (and biofilm), induced by the mechanical contact effect.

  • During wear in the absence of corrosion (W0 tests), the susceptibility to failure through cracking was heightened in C1, primarily due to the presence of the CoB layer on its outer surface. This susceptibility is further exacerbated by hydrogen embrittlement, leading to increased material removal compared to C2 and RM.

  • The low CW values estimated for C2 and RM significantly contribute to the synergistic interaction, implying that wear is the dominant mechanism ( C w / W c < 0.1 ) due to the total or partial removal of passive film (and biofilm).

  • A corrosion-wear regime ( 1 ( C w / W c ) < 10 ) was observed in C1. This behavior can be explained by the chemical reaction products formed through the interaction of the calf serum with the surface of C1. In this case, the chemical composition at the borided surface was modified, leading to increased corrosion, and likely exerting additional stresses beyond those imposed by tribological contact.

  • Finally, the CoB–Co2B and Co2B layers, formed through powder-pack boriding and powder-pack boriding exposed to diffusion annealing, significantly improved the bio-tribocorrosion resistance of the CoCrMo alloy by approximately 2.4 times and 1.3 times, respectively, compared to the RM. Furthermore, the differences in the bio-tribocorrosion behavior of CoB–Co2B and Co2B layers can be attributed to the chemical characteristics of their crystal structures.

This study was supported by the Secretaria de Investigacion y Posgrado of the Instituto Politecnico Nacional under Research Grant Nos. 20230167 and 20227018.

The authors have no conflicts to disclose.

Delgado-Brito A. M.: Formal analysis (equal); Investigation (equal); Methodology (equal); Validation (equal). I. Mejía-Caballero: Formal analysis (equal); Investigation (equal); Methodology (equal). A. D. Contla-Pacheco: Formal analysis (equal); Investigation (equal); Validation (equal). R. Pérez Pasten-Borja: Formal analysis (equal); Investigation (equal); Resources (equal). V. H. Castrejón-Sánchez: Investigation (equal); Resources (equal). E. J. Hernández-Ramírez: Formal analysis (equal); Investigation (equal); Resources (equal). I. Campos-Silva: Conceptualization (equal); Investigation (equal); Methodology (equal); Resources (equal); Supervision (equal); Writing – original draft (equal).

The data that support the findings of this study are available from the corresponding author upon reasonable request.

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