M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

ava i l ab l e a t www.sc i enced i r ec t . com

www.e l sev i e r . com/ loca te /matcha r

Microstructural characterization of as-cast biocompatibleCo–Cr–Mo alloys

J.V. Giacchia,c,⁎, C.N. Morandoa,c, O. Fornaroa,c, H.A. Palaciob,c

aConsejo Nacional de Investigaciones Científicas y Técnicas (CONICET), Av. Rivadavia 1917, (C1033AAJ) Buenos Aires, ArgentinabComisión de Investigaciones Científicas de la Provincia de Buenos Aires (CICPBA), Calle 526 e/10 y 11 (B1096APP) La Plata, ArgentinacInstituto de Física de Materiales Tandil (IFIMAT–FCE–CICPBA) Facultad de Ciencias Exactas, Universidad Nacional del Centro de la Provinciade Buenos Aires, Pinto 399 (B7000GHG) Tandil, Argentina

A R T I C L E D A T A

⁎ Corresponding author. Instituto de Física de M(B7000GHG) Tandil, Argentina.

E-mail address: jgiacchi@exa.unicen.edu.

1044-5803/$ – see front matter © 2010 Elsevidoi:10.1016/j.matchar.2010.10.011

A B S T R A C T

Article history:Received 4 November 2009Received in revised form12 October 2010Accepted 14 October 2010

The microstructure of a cobalt-base alloy (Co–Cr–Mo) obtained by the investment castingprocess was studied. This alloy complies with the ASTM F75 standard and is widely used inthe manufacturing of orthopedic implants because of its high strength, good corrosionresistance and excellent biocompatibility properties. This work focuses on the resultingmicrostructures arising from samples poured under industrial environment conditions, ofthree different Co–Cr–Mo alloys. For this purpose, we used: 1) an alloy built up fromcommercial purity constituents, 2) a remelted alloy and 3) a certified alloy for comparison.The characterization of the samples was achieved by using optical microscopy (OM) with acolorant etchant to identify the present phases and scanning electron microscopy (SE-SEM)and energy dispersion spectrometry (EDS) techniques for a better identification. In generalthe as-cast microstructure is a Co-fcc dendritic matrix with the presence of a secondaryphase, such as the M23C6 carbides precipitated at grain boundaries and interdendritic zones.These precipitates are the main strengthening mechanism in this type of alloys. Otherminority phases were also reported and their presence could be linked to the cooling rateand the manufacturing process variables and environment.

© 2010 Elsevier Inc. All rights reserved.

Keywords:Biocompatible alloysMicrostructureSolidificationCo–Cr–MoCarbides

1. Introduction

The Co–Cr–Mo as-cast alloys conforming to the ASTM F75standard are widely used in the manufacturing of orthopedicimplants made with investment casting techniques. Becauseof these alloys' hard workability and the shape complexity ofthe prostheses, this process reduces the high cost of machin-ing operations by producing pieces whose dimensions areclose to final ones. However, this method of fabrication leadsto poor mechanical properties compared to other manufac-turing processes such as powder metallurgy or forging [1].

It is well known that themain defects present in the as-caststate are: porosity, chemical inhomogeneity, large grain size

ateriales Tandil, Universi

ar (J.V. Giacchi).

er Inc. All rights reserved

and a microstructure with hard precipitates in interdendriticzones [2,3]. Also due to the casting process, inhomogeneitiesin carbide morphology and their size and distribution maylead to low ductility and low fatigue strength.

Nevertheless, the mechanical properties can be improvedwith ulterior heat treatments by dissolving the large carbidenetwork and produce a more homogeneous structure. It wasdemonstrated that by dissolution of the dendritic structure anisotropic behavior is obtained, particularly with respect toplastic deformation [4,5].

Pure cobalt has a hexagonal close-packed (hcp) structure atroom temperature, and it shows an allotropic transformationto face centered cubic (fcc) structure at 400 °C. The addition of

dad Nacional del Centro de la Provincia de Buenos Aires, Pinto 399 –

.

Table 1 – Chemical composition (wt.%) of the as-cast alloys.

Sample Co Cr Mo C W Si

C1 60.75 30.60 6.14 0.13 0.36 1.29C2 62.99 28.93 5.83 <0.021 0.38 1.05C3 64.13 27.01 6.30 <0.021 0.40 1.08F75 59.0/69.0 27.0/30.0 5.0/7.0 Max 0.35 Max 0.2 Max 1

Fig. 2 – OM for C1, color detected phases are shown.

54 M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

alloying elements such as chromium (Cr) and tungsten (W)increases this transformation temperature. Co-base alloysexhibit a dendritic α-fcc metastable matrix due to the sluggishnature of the fcc→hcp transformation [6], and a precipitateformedmainly by M23C6 carbides, an intermetallic σ compoundand a lamellar phase formed by interlayed plates of M23C6

carbide and a phase that has not been clearly identified yet,possibly a σ phase, both α and σ phases and either α or M6Ccarbide phases, as has been proposed by several authors [7–11].This lamellar structure is formedduringcoolingat temperaturesunder 990 °C, and it has been found to be the most detrimentalfeature of the microstructure [12] because once this grainboundary precipitate was dissolved, the fatigue life improves.

Carbide precipitation at grain boundaries and interdendriticregions is the major strengthening mechanism in the as-castcondition [13] but also by dislocation interactions with stackingfault intersections [14,15]. Coarse blocky carbides play animportant role acting as sources of dislocations and stackingfaults when stresses are applied [15]. Due to the intrinsic lowcooling rates of the manufacturing process, two carbidemorphologies can coexist, a “blocky type” and a “pearlitic type”as a result of the eutectoid reaction [16,17]. Other authors havereported that the formation of lamellar “eutectoid” carbidesoccurs for cooling rates varying from 8 to 16 °C/min [18] andprevious results that established that themaximumcooling ratepossible for the eutectoid phase precipitation is 35 °C/min [7].Also, theprecipitate size anddensity of both blocky and lamellarphase,was found to increasewith thecarboncontent [17],whilstthe ductility is not significantly affected.

On the other hand, the metallographic analysis showedthat the nitrogen content has a little effect on the grain size. Asthe nitrogen content increases, it promotes the formation offiner M6C carbides instead of M23C6 carbides that tend to

Fig. 1 – Optical micrograph of the as-cast microstructure of C1at the upper section.

accumulate along the grain boundaries. When present insolution, nitrogen might act as a solid solution strengthener,resulting in improved mechanical properties [1], at theexpense of other properties which often turn out insufficientto satisfy the requirements of the aforementioned standard.

Kilner et al. [19], suggested the possible precipitation of theM7C3 carbide during solidification of an ASTM F75 alloy and itsdissolution by diffusion during the subsequent cooling, since itwas not observed at room temperature. Clemow and Daniell etal. [20] stated that the reactions during solution heat treatmentsat temperatures above 1300 °C include carbide incipientmeltingand that M23C6 carbide initially present in the alloy tends totransform into M6C carbide and σ phase. It is important tomention that this type of heat treatment requires a precisecontrol of the variables involved, due to the narrow range oftemperature at which carbide dissolution takes place [21].

On the other hand, Opris et al. [6] confirmed the presence ofa complex combination of Co3Mo/W3C, M23C6 and Cr7C3

carbides embedded in the Co-rich matrix since the metalcarbides may actually have a ternary solid solution composi-tion [22]. Besides the interdendritic phases, some of whichshould be avoided because they negatively affect some

Fig. 3 – Optical micrograph of the as-cast microstructure of C1at the bottom section.

Fig. 4 – Optical micrograph of themain phases detected on C1by color metallography.

Fig. 6 – Optical micrograph of the as-cast microstructure fromsample C2 at the upper section of the cylinder.

55M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

mechanical properties, microsegregation occurs, which even-tually promotes the formation of non-equilibrium phases [16].

The as-cast microstructure could also be modified bychanging alloying elements. Replacing tungsten by an equal

Fig. 5 – a. SE-SEMmicrograph of main phases present in C1. b. EDSd. EDS analysis of the C1 Co–Cr–Mo matrix.

amount of molybdenum improves the ductility, increases thesolidification range and changes the morphology of theprecipitates by segregation [7] generating additional quanti-ties of eutectic carbides [4]. At the same time, the increase of

analysis of theM23C6 carbide. c. EDS analysis of the σ-phase.

Fig. 8 – Main present phases in C2 detected by colormetallographic techniques.

56 M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

the tungsten content decreases the stacking fault density andsimultaneously increases the volume fraction of the interme-tallic σ compound in Co–Cr–Mo alloys [23]. A similar behaviorhas been reported by Zhuang et al. [24] for nickel additions.Adding fcc-phase stabilizers, such as Fe, Ni, Mn and C,facilitates the retention of fcc at room temperature by slowingthe fcc→hcp transformation [4].

There exists much previous research about the character-ization of the ASTM F75 and similar alloys, focusing on thebehavior of the microstructure and properties achieved underdifferent heat treatment conditions. Knowledge of precipita-tion of secondary phases during solidification and subsequentcooling may provide appropriate alternatives to optimizethese thermal treatment processes [7,25].

The aim of this work was to provide more informationabout the Co–Cr–Mo alloy microstructure by detecting differ-ent phases and morphologies present, and to evaluate theirpossible transformation during solidification and cooling in anindustrial environment.

2. Experimental Procedure

Three different Co–Cr–Mo alloys were poured, all of themcomplying with the ASTM F75 chemical composition standard.Sample number one C1was obtained bymelting an appropriatequantity of alloyed materials of commercial purity. Samplenumber two (C2) was obtained from the remelted remnant of(C1), checking and adjusting the composition. For this reason itis expected that C2 could have more of metallic inclusions andmore inhomogeneities. Sample number three (C3)was obtainedusing a commercial material ASTM F75 certified alloy, forcomparison purposes.

The runs were carried out in an industrial environment, byusing a 100 kW induction furnace. The alloy was melted at1530 °C under an argon atmosphere and then poured intoceramicmoldswhichwere preheated to 1200 °C, continuing themethodology adopted in a previous work [26]. The chemicalanalysis of the resulting samples is shown in Table 1. It isworthmentioning that the tungsten content exceeded the value

Fig. 7 – Optical micrograph of the as-cast microstructure of C2at the bottom section.

suggested by the standard in all cases, therefore it might alsobe expected to contribute to carbide formation.

The ceramicmolds used included a set of three cylinders toobtain tensile specimens and three hip Charnley prostheses tostudy the solidification structure in this particular geometry.Under the presumption that an important segregation mightoccur, two cross sections of each tensile test bar were cut intoslices, keeping the remaining part of the bars for furtheranalysis. The cuts correspond to the top and bottom sectionsof the cylinders.

The preparation of samples for optical and scanning electronmicroscopy were carried out using conventional procedures forpolishing, from 600 grid SiC paper to 0.05 μm levigated alumina.Afterwards the specimens were thoroughly cleaned so as toremove any polishing residue. Following Szala et al. [27], twodifferent etching techniques were used to reveal themicrostruc-ture and to identify secondary phases, selecting a combination ofetchants that, in a first stage, revealed the dendritic matrixfeatures with a light electrolitic etchant and a second stage thatstained the secondary phases [28]. Optical micrographs (OM)from80× to 800×were used in order to obtain awide field of view.The high magnifications were used to identify different phases.

Fig. 9 – Identification of different phases present in C2.

Fig. 10 – a. SE-SEM micrograph corresponding to C2. b. EDS analysis of the M23C6 carbide phase. c. EDS analysis of the M6Ccarbide phase. d. EDS analysis of silicon-rich inclusión. e. EDS analysis of the C2 cobalt-rich matrix.

57M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

The samples were examined by scanning electron micros-copy (SE-SEM) using a JEOL JSM-6460LV microscope andenergy dispersive spectroscopy (EDS) techniques were alsoemployed to provide a more accurate characterization of thedifferent precipitated phases, even though the EDS techniqueprovides a semi-quantitative analysis.

The system used was an EDAX Genesis XM4-Sys 60,equipped with multichannel analyzer EDAX mod EDAM IV,

Sapphire Si(Li) detector and super ultra thin beryliumwindow,with a 20 kV acceleration voltage.

3. Results and Discussion

Metallographic analysis of thedifferent sampleswas carried outby analyzing the samples in different sections. A comparison

Fig. 11 – Optical micrograph of the as-cast microstructure ofC3 at the top section.

Fig. 12 – Optical micrograph of the as-cast microstructure ofC3 at the bottom section of the cylinder.

58 M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

between the microstructure of the top against bottom sectionsof the cylinders was also carried out for each sample.

3.1. Sample C1

The microstructure observed by optical microscopy consistedof a dendritic matrix with the presence of precipitates at grainboundaries and interdendritic zones, which agreed well withthose reported in the literature for these types of alloys in theas-cast condition. Also the presence of intrinsic castingmethod defects such as micropores were detected, probablydue to the high pouring temperature. Figs. 1 and 2 show the as-cast microstructure in detail for the upper section and Figs. 3and 4 show, the bottom section of the bar. Although there is nogreat difference between these micrographs, a continuouscarbide network and thicker dendrites can be observed in thebottom section indicating different cooling conditions. InFigs. 2 and 4 the phases present are also marked, themicrostructure detected included M23C6 carbides, a coarselamellar phase and a σ tetragonal phase precipitated at theinterdendritic space and grain boundaries.

Kilner et al. [19] have reported that the “pearlitic” coloniesobserved, consisted of interlayed plates of M23C6 and α-fcccobalt-rich matrix. Their result is consistent with the reportedlamellar carbide formation during continuous cooling ataround 990 °C [7,18] and our results.

Fig. 5a is a secondary electron SEM image with thedifferent phases identified. Fig. 5b, c and d show the EDSanalysis for the carbide, the σ tetragonal phase and thematrixrespectively, where the M23C6 carbide type resulted in achromium-rich one.

Regarding the σ phase, Ramirez et al. [18] stated that thisphase is present in both as individual particles and as part ofa binary eutectic within the α-fcc matrix. In this work, theσ phase was found both as individual particles and precipi-tated over the M23C6 which may corroborate the transforma-tion relationship between these phases established by Sims[15].

3.2. Sample C2

Figs. 6 and 7 show themicrostructure present in sample C2 forthe upper and bottom sections of the cylinders respectively, atdifferent magnifications. Once again, a slight difference in thedendrite arm size between the upper and bottom sections ofthe cylinder was determined and related to inhomogeneouscooling conditions.

The microstructure of this remelted sample, consisted of adendritic Co-rich fcc matrix and interdendritic precipitates asin C1, but two main differences could be noticed: the absenceof the intermetallic compound (σ phase) and the coexistenceof two different kinds of carbides due to the different color thateach phase presented, as seen by optical microscopy.

Fig. 8 shows the M23C6 and M6C carbides, which wereidentified by color metallography and EDS analysis. As sampleC2 was obtained by the remelting of sample C1 remnantmaterial, one distinctive feature was a large presence ofmetallic inclusions. These weremainly detected in themiddleof the precipitated phases, but it is not clear whether thesemetallic inclusions promote the nucleation of these phases or

they are dragged to the interdendritic space. The secondoption seems more probable, as the structure of carbides is ingeneral similar to C1. Fig. 9 shows Si-rich inclusions withinprecipitated carbides in the middle of interdendritic zones.

As manganese and silicon are mainly in solution and are notsegregated toward grain boundaries, they are difficult to reveal. Itisworthmentioning that Co–Cr–Moalloyswith lowandmediumC content, as in this case, typically show very little grainboundary attack, and are only revealed by the precipitation ofsecondaryparticles there [21]. Such is the case inFig. 7whereM6Cand M23C6 carbides can be noted on a grain boundary.

The coexistence of these two carbide types by stain etchingand heat tinting was reported by Lane et al. [29], but thesetechniques are significantly dependent upon the chemicalcomposition so these results cannot be considered conclusivedue to the variability of each solute in the carbide composi-tion. Clemow and Daniell [20] by using extracting methods,proposed the M23C6→M6C transformation at 1210 °C duringheat treatments.

Fig. 13 – Main colored phases detected in C3.

59M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

Fig. 10a shows a secondary electron SEM image of C2showing an M6C carbide phase with a Si-rich inclusionprecipitated within, embedded in a α-fcc matrix. Because thestaining techniques cannot be conclusive, the phases werefurther identified by EDS analysis as two different chromium-rich carbides as Fig. 10b and c show. The EDS analysis for the

Fig. 14 – a. Main phases present in C3, shown in an electron micrc. EDS analysis of the C3 Co–Cr–Mo rich uniform matrix.

phases identified as silicon-rich inclusions and cobalt-richmatrix are shown in Fig. 10d and e respectively.

3.3. Sample C3

In Figs. 11 and 12, the as-cast microstructures correspondingto sample C3 are shown for the top and bottom sections of thecylinders respectively. The microstructure consisted of adendritic Co-rich α-fcc matrix and interdendritic precipitates,similar to C1 but with a smaller size than in C2. The carbideswere identified by color metallography as M23C6 carbides. Itwas not possible to register any other phase by opticalmicroscopy. Analogous to C1 and C2, the secondary armspacing is smaller in the top section, due to different coolingrates. Also, this sample shows a very little precipitation ingrain boundaries, making its metallographic detection diffi-cult. Some inclusions were also found, but smaller and inlesser quantity than in C1 and C2 and only detected byscanning electron microscopy. A distinctive characteristic ofthis alloy was the uniformly distributed precipitation of asmaller size than previous samples as Fig. 13 shows. Fig. 14ashows the morphology of the main phases found in C3 in asecondary electron SEM micrograph. These were also furtheridentified by EDS analysis as chromium-rich carbides embed-ded in a Co–Cr–Mo uniform matrix, as Fig. 14b and c show.

ograph (SE-SEM). b. EDS analysis of the M23C6 carbide phase.

60 M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

4. Conclusions

Three Co-based alloys with a composition conforming to theASTM F75 standard poured in an industrial environmentwere studied. The different phases present were identifiedby using optical and electron microscopy and EDS. In all thecases, the microstructures agreed well with those reportedin the literature and were well identified by color metallog-raphy. The EDS analysis was employed for an accurateidentification, even though it is a semi-quantitative analysistechnique.

M23C6 type carbides and intermetallic σ phase weredetected at grain boundaries and interdendritic zones. Thefirst ones are present in two different morphologies, a “blockytype” and colonies of a “pearlite type”. This latter structure isformed once the solidification is complete at temperaturesunder 990 °C, when the cooling rate is lower than 35 °C/min asis this case.

In this work the σ phase was found both as individualparticles and precipitated over the M23C6 which may corrob-orate the transformation relationship between these phasesestablished by Sims [15].

A quasi-continuous carbide precipitation was noticeable atgrain boundaries. Also fine colonies of lamellar carbide formedat grain boundaries or adjacent regions. Further investigationis required in order to establish their formation mechanism,since there is no clear evidence to affirm if they have grown bya discontinuous reaction or originated from a eutectoidreaction, as has been proposed [18].

Sample C2 was characterized by the coexistence of twokinds of carbides, identified by color metallographic and EDStechniques as M23C6 and M6C carbides, dispersed in interden-dritic zones. This particular behavior could be attributable to achange in the cooling rate during solidification, as suggestedby Clemow and Daniell [20]. Also in this sample, a greaterquantity of metallic Si-rich inclusions located within thecarbides was detected, probably due to C2 being a remelt of C1.It is not clear yet, whether these metallic inclusions promotethe nucleation of the secondary phases or are dragged to theinterdendritic space.

The precipitates found in sample C3 were homogeneouslydistributed and smaller in size with respect to C1 and C2. Thiseffect could be the result of the lower carbon and chromiumpresent in solution in this alloy when compared to previoussamples.

It is noteworthy that in all the samples, the microstruc-tural defects as shrinkages, pores and segregation as well aslarge grain size were also present, as result of the pouringconditions. While these conditions were controlled, we musttake into account that typical industrial environment controlmethods are not as strict and rigorous as those in researchlabs.

Acknowledgements

This work was carried out at IFIMAT, Facultad de CienciasExactas,UNCPBA, andpartially supportedbyANPCyT,CONICET,CICPBA and SECAT-UNCPBA.

R E F E R E N C E S

[1] Escobedo J, Méndez J, Cortés D, Gómez J, Méndez M,Mancha H. Effect of nitrogen on the microstructure andmechanical properties of a CoCrMo alloy. Mater Des 1996;17:79–83.

[2] Campbell J. Castings: The new metallurgy of cast metals. 2ndedition. Oxford/GB: Elsevier Science & Technology; April 2003.

[3] Stefanescu DM. Science and engineering of castingsolidification. 2nd edition. USA: Springer US; 2009.

[4] Amigo Borrás V, Paolini A, Moreno Ballester JF, VicenteEscuder A, Romero Sanchis F. Estudio de la influencia de lostratamientos térmicos en la microdureza y microestructurade aleaciones CoCrMo. Congreso Nacional de PropiedadesMecánicas de Sólidos, Gandia; 2002. p. 487–96.

[5] Montero Ocampo C, Talavera M, Lopez H. Effect of alloypreheating on the mechanical properties of as-cast CoCrMoCalloys. Metall Mater Trans A 1999;30:611–20.

[6] Opris CD, Liu R, Yao MX, Wu XJ. Development of Stellite alloycomposites with sintering/HIPing technique forwear-resistant applications. Mater Des 2007;28:581–91.

[7] Ramírez LE, Castro M, Méndez M, Lacaze J, Herrrera M, LesoultG. Precipitation path of secondary phases duringsolidification of the CoCrMoC alloy. Scr Mater 2002;47:811–6.

[8] Weeton JW, Clauss FJ. Effect of heat treatment uponmicrostructures microconstituents and hardness of awrought cobalt-base alloy, Stellite 21. National AdvisoryCommittee for Aeronautics, Technical Note 3107; March 1954.

[9] Weeton JW, Signorelli RA. Effect of heat treatment uponmicrostructures microconstituents and hardness of awrought cobalt-base alloy. Trans ASM 1955;47:815–45.

[10] Silverman R, Arbiter W, Hodi F. Effect of sigma phase onCo–Cr–Mo base alloys. Trans ASM 1957;49:805–22.

[11] Sims CT. Contemporary view of cobalt-base alloys. J Met1969;21:27–42.

[12] Dobbs HS, Robertson JLM. Heat treatment of cast Co–Cr–Mofor orthopaedic implant use. J Mater Sci 1983;18:391–401.

[13] Asgar K, Peyton FA. Effect of casting conditions on somemechanical properties of cobalt-case alloys. J Dent Res1961;40:73–86.

[14] Rajan K, Vander Sande JB. Room temperature strengtheningmechanisms in a Co–Cr–Mo–C alloy. J Mater Sci 1982;17:769–78.

[15] Sims CT, Hagel W, Stoloff N. The Superalloys II: Hightemperature materials for aerospace and industrial power.2nd edition. Wiley & Sons Inc.; 1987. Chapter 5.

[16] Montero Ocampo C, Salinas A. Effect of carbon content on theresistance to localized corrosion of as-cast cobalt-basedalloys in an aqueous chloride solution. J Biomed Mater Res1995;29:441–53.

[17] Herrera Trejo M, Espinoza A, Méndez J, Castro M, López J,Rendón J. Effect of C content on the mechanical properties ofsolution treated as-cast ASTM F75 alloys. J Mater Sci MaterMed 2005;16:607–11.

[18] Ramírez Vidaurri LE, Castro RománM, Herrera Trejo M, GarcíaLópez CV, Almanza Casas E. Cooling rate and carbon contenteffect on the fraction of secondary phases precipitate inas-cast microstructure of ASTM F75 alloy. J Mater ProcessTechnol 2009;209:1681–7.

[19] Kilner T, Pilliar RM, Weatherly GC, Alibert C. Phaseidentification and incipient melting in a cast CoCr surgicalimplant alloy. J Biomed Mater Res 1982;16:63–79.

[20] Clemow AJ, Daniell BL. Solution treatment behavior ofCoCrMo alloy. J Biomed Mater Res 1979;13:265–79.

[21] Mancha H, Carranza E, Escalante JI, Mendoza G, Méndez M,Cepeda F, et al. M23C6 carbide dissolution mechanisms duringheat treatment of ASTM F75 implant alloys. Metall MaterTrans A 2001;32A:979–84.

61M A T E R I A L S C H A R A C T E R I Z A T I O N 6 2 ( 2 0 1 1 ) 5 3 – 6 1

[22] VandammeNS, Que L, Topoleski LDT. Carbide surface coatingof Co–Cr–Mo implant alloys by a microwave plasma-assistedreaction. J Mater Sci 1999;34:3535-3521.

[23] Karaali A,MirouhK,Hamamda S, Guiraldenq P.Microstructuralstudy of tungsten influence on CoCr alloys. Mater Sci Eng, A2005;390:255–99.

[24] Zhuang L, Langer EW. Effects of alloy additions on the fatigueproperties of cast Co-Cr–Mo alloy used for surgical implants.J Mater Sci 1990;25:683–9.

[25] Amigó V, Vicente A, Romero F, Paolini Y. Influencia deltratamiento HIP en la distribución de los carburos en prótesisCoCrMo. Bol Soc Esp Ceram Vidrio 2004;43:573–7.

[26] Ges A, Garbellini O, Fornaro O, Palacio H. Solidificacióncontrolada de aleaciones bicompatibles enprótesis de cadera. Anales SAM/CONAMET-AAS 2001; 2001.p. 1189–96.

[27] Szala J, Szczotok A, Richter J, Cwajna J, Maciejny A. Selectionof methods for etching carbides in MAR-M509 cobalt-basesuperalloy and acquisition of their images. Mater Charact2006;56:325–35.

[28] Vander Voort G. Metallography and microstructures. ASMHandbook, vol. 9; 2004. Ohio.

[29] Lane JR, Grant NJ. Carbide reactions in high temperaturealloys. Trans ASM 1952;44 p. 113, 62.