Review on the Oxidation of Metallic ThermalSprayed Coatings: A Case Study

with Reference to Rare-Earth PermanentMagnetic Coatings

Jo Ann Gan and Christopher C. Berndt

(Submitted December 6, 2012; in revised form May 29, 2013)

Thermal spray fabrication of rare-earth permanent magnetic coatings (PMCs) presents potential man-ufacturing routes for micro-magnetic devices. Despite this potential, thermal spray of PMCs is still notwidely explored due to oxidation concerns. It was established that oxidation leads to the loss of ferro-magnetic phases in these materials and results in deterioration of magnetic performance. Although thisreview focuses on a specific class of material, i.e., magnetic materials, there is significant technicalcrossover to all classes of feedstocks that are employed in thermal spray processing. The oxidationmechanisms and the associated influencing factors are explored in this work to implement effectiveprocessing techniques during the deposition process. This paper reviews the various stages and mecha-nisms of oxidation in thermal spray processes. The factors that influence the extent of oxidation dependon the type of oxidation that is dominant and rely on the type of spray system, powder injection position,and the particle size of feedstock. Among the aspects that are reviewed include the oxygen-fuel ratio forhigh velocity oxygen-fuel (HVOF), current intensity, gas flow rate, particle size, spray distance, andsubstrate temperature. Protection strategies to minimize oxidation in thermal spray processes, such asgas shrouding and shielding, are presented.

Keywords in-flight oxidation, oxidation, oxidation protec-tion, permanent magnet, post-impact oxidation,rare-earth, thermal spray coatings

1. Introduction

The oxidation of thermal spray coatings has become asubject of great interest because it influences the perfor-mance of coatings. Of great concern is the fact that oxi-dation deteriorates the functioning of thermal barriercoatings (TBCs) and other high temperature corrosion-resistant coatings. In TBCs, coatings are subjected toirregular degradation during thermal cycling due to thecoefficient of thermal expansion mismatch between metaland oxides (Ref 1). Chemical uniformity of surfacesexposed to corrosion is disrupted by oxidation and hencedepreciates the corrosion resistance capability (Ref 2).Coatings that are formed with oxides are also more brittle,which makes them difficult to machine for certain appli-cations since the mechanical properties of the coatings are

affected. Also, spalling of coatings has become a commonphenomenon in the presence of oxides.

However, a certain degree of oxidation is desirable inseveral applications. It has been demonstrated that oxidesimprove the wear resistance of a coating (Ref 3) andincrease the coating hardness. The strength of a coatingunder compressive loading is also enhanced in the pres-ence of oxides (Ref 4, 5). Metal oxides that reside alongthe splat boundaries can alter the intersplat cohesion,while oxide dispersion in the splat itself or supersaturatedoxygen strengthens the splat (Ref 4, 6).

Most of the studies on oxidation of thermal spraycoatings were conducted using steel or steel alloys becausesteels are (i) sensitive to oxidation and (ii) widely used ascorrosion and wear-resistant coatings. The study on ther-mal sprayed permanent magnetic coatings (PMCs) hasbeen limited due to the lack of understanding in the oxi-dation behavior of the permanent magnets, especially forrare-earth magnets.

Rare-earth transition-metal magnets exhibit the stron-gest magnetic behavior among all the other hard magneticmaterials, such as alnico and ceramic magnets. Neodym-ium iron boron and samarium cobalt are the two mostcommonly used rare-earth hard magnets that demonstratehigh overall magnetic performance. Neodymium ironboron is the stronger among the two, having bulk rema-nence, coercivity, and maximum energy product in therange of 1.0-1.4 T, 9.5-25 kOe, and 25-55 MGOe, respec-tively. Rare-earth permanent magnets are widely used inmany industries, from aerospace and automotive to

Jo Ann Gan and Christopher C. Berndt, Industrial ResearchInstitute Swinburne, Faculty of Engineering and IndustrialSciences, Swinburne University of Technology, Hawthorn,VIC 3122, Australia; and Christopher C. Berndt, Departmentof Materials Science and Engineering, Stony Brook University,Stony Brook, NY 11794. Contact e-mail: jgan@swin.edu.au.

JTTEE5 22:1069–1091

DOI: 10.1007/s11666-013-9955-2

1059-9630/$19.00 � ASM International

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telecommunication and electronic devices (Ref 7). Thickfilms or coatings of these materials can be applied inmicroelectromechanical systems (MEMS) (Ref 8, 9) forapplications such as micro-actuators, micro-motors, andmicro-generators. One downfall of the rare-earth magnet,however, is the strong affinity of rare-earth elements tooxygen, which makes them prone to oxidation duringprocessing and operation.

Oxidation leads to brittle coatings and deterioration ofmagnetic properties. However, studies to find alternativesto control the oxidation process have been limited. Moststudies on the fabrication of permanent magnet films havefocused on low temperature processes such as sputtering: athin film method that has technical limitations with regardto processing time, restrictions concerning the geometryand size of the component, and a cost penalty. In caseswhere thermal spray has been employed, spraying hasbeen mostly performed within a vacuum chamber tominimize oxidation (Ref 10-12). Recent studies haveexamined the low temperature cold spray process(Ref 13). Although the coating obtained did not exhibitoxidation or deterioration of magnetic properties, thecoating was too brittle to be of practical use.

Whether oxidation is desired or otherwise, the abilityto control the degree of oxidation in a coating is importantfor the coating to achieve optimum performance. Themechanism of oxidation must be understood to manipu-late the oxidation process to advantage. This paper isdevoted to reviewing the oxidation considerations indepositing PMCs. The stages and mechanisms of oxidationin the thermal spray process, factors that influence oxi-dation, and protection measures against oxidation duringthermal spray processing are investigated. The thermaloxidation of rare-earth permanent magnets is highlightedsince they are prone to high temperature oxidation and aprime reason regarding the poor acceptance on depositingmagnetic coatings via thermal spray.

2. Oxidation of Permanent MagneticCoatings

Oxidation has been a concern for thermal spray PMCs.PMCs are conventionally fabricated by lower temperatureprocessing methods such as sputtering (Ref 14, 15), screenprinting (Ref 16, 17), and tape casting (Ref 18, 19). Theresearch on thermal spray PMCs has been confined to theuse of vacuum plasma spray (Ref 10, 11, 20) and coldspray (Ref 13). Much research has examined the oxidationof PMCs throughout the aging process, but few havestudied oxidation during the fabrication process. There-fore, the understanding of PMC oxidation is limited. ThePMCs in this section refer to rare-earth transition metals,specifically neodymium iron boron (Nd2Fe14B) andsamarium cobalt (SmCO5 and Sm2Co17), since they aremore prone to oxidation than magnetoplumbite-type fer-rites, which are already in an oxidized state.

Earlier studies indicated that oxidation products ofdifferent phases depended on temperature and humidity

(Ref 21, 22). Generally, corrosion of rare-earth magnetscan occur under three conditions: (i) warm humid air orsteam (Ref 23-25), (ii) electrochemical reactions (Ref 26-28), and (iii) exposure to high temperature (>250 �C) forlong time periods (from several hours up to days) (Ref 29,30). Of these three adverse conditions, permanent magnetparticles are most prone to be subjected to high temper-ature oxidation during thermal spray. Although the ther-mal spray process involves very short processing time[between 0.1 and a few milliseconds (Ref 31, 32)], oxida-tion still occurred due to the high temperature involved inthermal spray processes (Ref 10, 33, 34). Therefore, theoxidation product of thermal spray permanent magnetswas similar to those of bulk magnets. A prime differencewould be that the oxidation process of thermal spraypermanent magnets occurred at a much higher rate andany intermediate oxidation products might not beobserved.

The oxidation of neodymium iron boron and samariumcobalt is a two-stage process: (i) formation of externalscales and (ii) internal oxidation (Ref 35, 36). The externalscales are free of rare-earth elements and consist of Co3O4

and CoFe2O4 for samarium cobalt (Ref 37), and a-Fe2O3

(hematite) accompanied by Fe3O4 (magnetite) for neo-dymium iron boron (Ref 36). Internal oxidation can beprevented if the external oxide layers can restrict theaccess of oxygen. However, these oxide layers are notprotective at elevated temperatures (Ref 36) and result inan internal oxidation zone (IOZ). Internal oxidation is theprimary degradation mechanism for rare-earth transition-metal materials and is more detrimental than externaloxides since the IOZ is much thicker than the surfaceoxides and phase transformations in this region destroythe hard magnetic properties of the alloy (Ref 36-38). TheIOZ in samarium cobalt is composed of a-Co, Sm2O3, andSmCoO3 (Ref 35, 39), while the IOZ in neodymium ironboron is constituted of a-Fe, Nd2O3, NdFeO3, and B2O3

(Ref 29, 36, 40). The structure of IOZ in samarium cobaltwas described as parallel samarium oxide platelets andfibers surrounded by cobalt (Ref 41), whereas the IOZstructure of neodymium iron boron was described asneodymium oxide particles embedded in an a-Fe matrix(Ref 36). Schematics of the IOZ of neodymium iron boronand samarium cobalt are shown in Fig. 1.

Since oxidation at elevated temperature is the primarycause for IOZ formation that is detrimental to the mate-rial properties compared to external oxides, the followingsections are focused on thermal oxidation of neodymiumiron boron and samarium cobalt. However, it should benoted that the oxidation mechanisms discussed in thissection concern solid phase oxidation, while oxidation inthermal spray processes involves both solid and liquidphase oxidation. Solid phase oxidation represents theclosest approximation of oxidation reactions at a hightemperature available in the literature since there havebeen no reports in the open literature on liquid phaseoxidation for these materials. Most studies for thesematerials were carried out to investigate oxidation duringthe materials� performance as a magnet rather than duringprocessing or deposition. The significance of this section is

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that it presents an overview of the potential oxidationmechanisms for rare-earth magnetic materials in thermalspray processing and phases that will be present in thecoatings.

2.1 Neodymium Iron Boron (Nd-Fe-B)

The thermal oxidation of Nd-Fe-B occurs through theprocess of dissociation followed by oxidation. TheNd2Fe14B matrix dissociates into a-Fe nanocrystals thatconsist of small hexagonal neodymium oxide, h-Nd2O3,precipitates and form a gray surface layer that growstransgranularly into the magnet. Further oxidization led tothe formation of a-Fe2O3 and at approximately 600 �C, thea-Fe2O3 reacted with the small precipitates of h-Nd2O3

formed earlier to give NdFeO3. Another form of neodym-ium oxide exists in cubic structure, c-Nd2O3, which resultsfrom the oxidation of NdOx (Nd-rich intergranular regions)(Ref 40). The reactions were proposed to occur according toEq 1-4. The results of Li et al. (Ref 36) agreed with those ofEdgley et al. (Ref 40). The proposed overall reaction for theoxidation of Nd-Fe-B varies throughout the literature andthey are summarized in Table 1. Despite the variation in theproposed overall equation, the literature agrees on the sameoxidation products, i.e., a-Fe, Fe2O3, Fe3O4, Nd2O3, andNdFeO3, although some of these oxidation products werenot reflected in the overall equations.

Nd2Fe14Bþ 3

2O2 ! h-Nd2O3 þ 14Fe þ B ðEq 1Þ

2NdOx þx

2O2 ! c-Nd2O3 ðEq 2Þ

2Feþ 3

2O2 ! a-Fe2O3 ðEq 3Þ

a-Fe2O3 þNd2O3 ! 2NdFeO3 ðEq 4Þ

Two types of Nd2O3 were observed by Edgley et al.(Ref 40), namely, cubic and hexagonal Nd2O3. The cubicstructure, c-Nd2O3, formed at lower temperature(<275 �C), while hexagonal, h-Nd2O3, was observed at240-350 �C. In between 240 and 275 �C, both cubic andhexagonal Nd2O3 coexisted. Another Nd-Fe-B phase that

was present other than Nd2Fe14B would be sub-stoichi-ometric Nd1.1Fe4B4 (Ref 29, 36, 40). However, the oxi-dation of this phase occurred at a higher temperature(>611 �C) and this has not been investigated further. Theoxidation of Nd1.1Fe4B4 is noteworthy since it has poten-tial significance with regard to the oxidation of thermalspray PMCs. Therefore, there are generally three mainphases observed in Nd-Fe-B magnets: (i) Nd2Fe14B, (ii)Nd-rich (NdOx), and (iii) Nd1.1Fe4B4 (Ref 36, 42). Table 2outlines the oxidation products for these three phases attemperatures below and above 500 �C. Note that thetemperature discussed here is much lower than the oper-ating temperature of thermal spray processes, where thedeposition mechanism involves particle melting and rapidsolidification.

Castaldi et al. (Ref 43) and others (Ref 44, 45) men-tioned the formation of a-Fe as a result of Nd2Fe14Boxidation at process temperatures between 600 and800 �C. Other oxidation phases were not discussed indetail. Lileev et al. (Ref 45) found that there was as muchas 18 wt.% Nd-O present in a Nd-Fe-B film sputtered inan inert atmosphere. The presence of Nd-O in sputteredNd-Fe-B was explained by Chen et al. (Ref 46) as a resultof one of the following mechanisms:

� The base pressure of the chamber was too high.

� The films were deposited at high temperatures so thatmore oxygen was degassed from the chamber system.

� The target of Nd-Fe-B alloy contained a large pro-portion of neodymium that was readily oxidized.

2.2 Samarium Cobalt (Sm-Co)

Studies concerning the oxidation of samarium cobaltare relatively scarce and are not fully explained. Most ofthe studies focused on the oxidation products and kineticsand little attention have been given to the mechanism ofits thermal oxidation. The oxidation of samarium cobalthas received less attention than neodymium iron boronsince samarium cobalt demonstrated a higher Curie tem-perature, i.e., the temperature above which a ferromag-netic material becomes paramagnetic, than neodymiumiron boron (Ref 47). However, samarium cobalt is still

Fig. 1 Schematics of the IOZ of (a) neodymium iron boron and (b) samarium cobalt

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subjected to oxidation since rare-earths are prone to oxi-dation, and the oxidation of magnetic particles is morecritical than for bulk sintered magnets (Ref 48).

The primary oxidation product of samarium cobalt isSm2O3 (Ref 39, 41). Bartlett and Jorgensen (Ref 41) reportedthat thermal oxidation of SmCo5 resulted in the formation ofmonoclinic sesquioxide Sm2O3 at all temperatures, alongwith the polymorph of hexagonal a-Co and cubic b-Co.Another study by Yang et al. (Ref 39) reported that thermaloxidation of samarium cobalt commenced at 500 �C, whereconsumption of samarium by oxidation resulted in thetransformation of rhombohedral Sm2Co17 and hexagonalSmCo5 into the body-centered cubic (bcc) Fe(Co) solidsolution and face-centered cubic (fcc) Co(Fe), respectively.In other words, the Fe(Co) phase is the degraded magneticphase of Sm2Co17, while Co(Fe) is the degraded magneticphase of SmCo5. While several researchers (Ref 35, 39, 48-50) agreed with Bartlett and Jorgensen (Ref 41) on the for-mation of Sm2O3 due to oxidation, Yang et al. (Ref 39) dis-covered another oxide phase, SmFeO3, in addition to Sm2O3.The proposed overall oxidation reaction of samarium cobaltalso varies throughout the literature and these reactions aresummarized in Table 3.

Although the oxidation products of bulk magnetsfound in these studies have been observed in thermalspray PMCs (Ref 10, 51), all these studies involved alower temperature range in relation to the process tem-perature of thermal spray. There is a need for the oxi-dation process and products at higher temperature to beexplored further if thermal spray of PMCs is to beimplemented widely, since these oxidation productsinfluence the magnetic and mechanical properties of thePMCs significantly. Oxidation of neodymium iron boronand samarium cobalt resulted in the loss of ferromagneticphases to form paramagnetic oxides (Ref 42, 52), thusleading to a detrimental loss of coercivity (Ref 53, 54).For instance, Matsuura et al. (Ref 55) found that thecoercivity of a Nd-Fe-B thin film (100 nm) decreased by60% after oxidation.

The deterioration of magnetic properties as a conse-quence of oxidation is a collective action of (i) superpo-sition of soft and hard magnetic phases, (ii) oxidation ofsurface layers, and (iii) spin reorientation that changes themagnetic c-axis structure (Ref 56). Dickens and Mazany(Ref 57) determined that there is a direct correlationbetween weight gain from oxidation and magnetic prop-erties� deterioration, where a weight gain of 4.5% for Nd-Fe-B causes total magnetic loss since neodymium in allphases has been totally oxidized at this point. However,this study did not take into consideration the weight gainfrom other oxidation products such as Fe2O3 and B2O3.While the oxide phases may be helpful for the growth of ac-axis texture of Nd2Fe14B grains during deposition, itspresence during post-deposition annealing deterioratedthe magnetic properties (Ref 46). These results suggestthat some oxidation protection must be implementedduring the deposition process of PMCs to preserve theirmagnetic properties.

3. Stages and Mechanism of Oxidation inThermal Spray Process

3.1 In-Flight and Post-Impact Oxidation

Many researchers have examined the stages of oxida-tion in thermal spray processes. There have been numer-ous models of coating oxidation and there are manysimilarities among these models. In general, oxidationoccurred in two stages: (i) in-flight oxidation and (ii) post-impact oxidation.

Hackett and Settles (Ref 58) suggested that the HVOFflow field could be divided into three regions, Fig. 2, andtherefore oxidation occurred at three stages. This idea wasreinforced by Dobler et al. (Ref 59). Region I encom-passed the point of particle injection within the nozzle upto the end of the visible plume, also termed as the jet core.In this area, particles were subjected to oxidation by the

Table 1 Proposed overall oxidation reactions for Nd-Fe-B from the literature

Proposed overall reaction Reference

3Nd2Fe14B + 23O2 fi 6NdFeO3 + 10Fe3O4 + 3Fe2Ba Turek et al. (Ref 52)Nd2Fe14B + 3O2 fi 13a-Fe + c-Nd2O3 + FeBO3 Edgley et al. (Ref 40)Nd2Fe14B + 3/2O2 fi 14a-Fe + c-Nd2O3 (+ interstitial B)Nd2Fe14B + 3/2O2 fi 14a-Fe + h-Nd2O3 (+ interstitial B)4Nd2Fe14B + 9O2 fi 4Nd2O3 + 2B2O3 + 56Fe Li et al. (Ref 36)a The overall equation provided in the original literature was not balanced.The balanced equation should have 29 O2 on the left-hand side rather than the stated 23 O2.

Table 2 Oxidation reactions for Nd2Fe14B, Nd-rich, and Nd1.1Fe4B4 phases at temperatures below and above 500 �C(Ref 40)

T < 500 �C T > 500 �C

Nd2Fe14B! a-Fe! Fe2O3 Nd2Fe14B! a-Fe! Fe2O3

Nd2Fe14B! Nd2O3 ! Nd2O3 Nd2Fe14B! Nd2O3 ! NdFeO3

Nd-rich NdOxð Þ ! c-Nd2O3 ! c-Nd2O3 Nd-rich NdOxð Þ ! NdOx ! c-Nd2O3

Nd1:1Fe4B4 ! Nd1:1Fe4B4 ! Nd1:1Fe4B4 Nd1:1Fe4B4 ! Nd1:1Fe4B4 ! Unknown

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free oxygen in the products of combustion in HVOF.Although the amount of free oxygen in this zone could becontrolled by selecting the oxygen/fuel ratio of the com-bustion parameters, the flame still contained differentoxidizing combustion products and the particle oxidationcannot be completely avoided. Spray particle dwell time inRegion I was very short, approximately one millisecond,and the particle temperatures were relatively high.

Region II referred to the domain where the localatmosphere was entrained to the centreline of the flame,toward the end of the jet core, until the particle impingedonto the substrate. In-flight oxidation may have occurredin this region because the surrounding atmosphere con-sisted of up to 20% oxygen. The particle residence time inthis region was also short, while the particle temperaturewas slightly lower than Region I due to the cooling effectof the entrained air. Although the particle residence timealso depends on other parameters, such as the particle sizedistribution, gas flow rate, and injection position, theparticle residence times in Region I and II were still withinthe order of milliseconds (Ref 31, 60).

Lastly, Region III specified the splat formed on thesubstrate immediately after impact and just before thenext layer of splats was deposited. The exposure time wasapproximately ten times longer than the particle flighttime. The initial splat temperature was high due to(i) transformation of kinetic energy of the particles intothermal energy and (ii) exposure to the flame. However,the splats cooled rapidly since the substrate temperaturewas comparatively low. Seventy-five percent of total oxi-dation takes place in either Region II or III (Ref 58).Table 4 summarizes the conditions that affect the oxida-tion process at respective regions.

Vardelle et al. (Ref 61) revealed that oxidation in aplasma spray process also occurred over three stages:(i) gas-solid phase oxidation during initial preheating ofparticles to their melting point, (ii) gas-liquid oxidation ofthe molten particles in-flight, and (iii) gas-solid oxidationof the splats after deposition and before the formation ofsubsequent splat layers. These three stages of oxidationalso occur during the HVOF process. Zhang et al. (Ref 62)established that most oxidation occurs in the liquid phase,while the particle temperature and velocity have no sig-nificant effect on the thickness of the surface oxide layer(~0.1 lm) before particle melting.

On the other hand, Volenık et al. (Ref 63) classifiedoxidation processes according to stages in the plasmaspray process, i.e., (i) the in-flight stage and (ii) thedeposition, solidification, and cooling stages. In the firststage, powder was injected into the plasma plume andmelted. The molten particles then interacted with the

plasma gas that has been altered by the entrainment ofsurrounding atmosphere. The first stage was usually veryshort (several milliseconds), but the reaction temperaturewas high. The second stage of oxidation involved theinteraction of solidified deposits with a gaseous mixturethat consisted mainly of the ambient atmosphere. Thisstage was typically much longer (several seconds orlonger) than the first stage and the temperature decreasedgradually. As can be seen, Volenık�s model was similar tothat of Hackett and Settles (Ref 58) where the first stagewas equivalent to Region I and II in Hackett and Settles�model, while the second stage corresponded to Region III.

There are distinct similarities and differences in theoxidation stages between HVOF and plasma spray. Bothprocesses involve the two main stages of oxidation men-tioned earlier in this section, i.e., in-flight and post-impactoxidation. There are two regions where in-flight oxidationoccurs in both processes: (i) a visible plasma plume or jetcore consisting of the plasma gas or combustion gases,where the particles are injected, and (ii) a turbulent flowregion where surrounding atmosphere is entrained intothe system. While region (i) in the plasma spray process isoxygen free, the same region in HVOF could be oxidizingor reducing depending on the oxygen-to-fuel (O/F) ratio.The length of each region in plasma spray depends on thetype of plasma gas and flow rates, current intensity, andtorch design, while the combustion gas flow rates,O/F ratio, and torch design determine this factor in HVOFsystems. Post-impact oxidation for both processes is thesame since this stage depends on the interaction betweenthe deposited splat and the surrounding atmosphere. Thismechanism is essentially the same for both processesunless a vacuum or controlled atmosphere environment isused for either process.

3.2 Mechanisms of Oxidation

A detailed mechanism and schematic models of oxi-dation during the thermal spray process was presented byDeshpande et al. (Ref 2), where oxidation may take placein-flight, at the splat surface, and/or between passes.In-flight oxidation depended strongly on the amount ofentrained air, the particle temperature, and the nature andturbulence of the flame. In-flight particles oxidation in thejet core and in the plume was controlled by a differentmechanism, as pointed out by Syed et al. (Ref 64, 65).

Sobolev et al. (Ref 66) mentioned that there were twomass transfer processes that occurred during the interactionof liquid particles and oxygen. The mass transfer processeswere the development of oxides due to chemical reactionsbetween the surface of the liquid phase and oxygen as wellas the diffusion of oxygen in the liquid. However, laterstudies found that the amount of oxides present was too highto be explained by diffusion models (Ref 67).

In the jet core, particles were melted and the liquidphase reacted with oxygen present in the combustionproducts and entrained air. For instance, low carbon steelparticles were melted in HVOF (Ref 68) and plasma spray(Ref 69), and oxidized to form FeO, Fe2O3, and Fe3O4. Atthis stage, the motion of the liquid was induced by

Table 3 Proposed overall oxidation reaction of Sm-Cofrom the literature

Proposed overall reaction Reference

SmCo5 + 3/4O2 fi 1/2Sm2O3 + 5Co 414Sm2Co17 + 57O2 fi 4Sm2O3 + 34Co2O3 49Sm2Co17 + 3/2O2 fi 17Co + Sm2O3 50

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convective movements, forming a spherical Hill vortexthat resulted from the difference of kinematic viscositiesbetween the liquid droplet and surrounding media(Ref 70). The spherical Hill vortex is a model that de-scribes fluid motion or vortex distribution within a sphere(Ref 71). A model of oxide penetration by convectivemovement within an iron particle is depicted in Fig. 3. Thegenerated vortex destroyed the surface oxides and dis-tributed them uniformly through the bulk volume of theparticle, forming nodular oxides, Fig. 4(a). The oxideswere hurtled into the interior of the particle, leaving freshmetal on the particle surface, which enhanced the particlereactivity. The difference in surface tension between the

liquid oxides and metal encouraged the formation ofisolated spherical nodules of oxide within the metal(Ref 69, 70), Fig. 5(b) and 6(a).

Outside the jet core, the velocities and temperatures ofthe plasma and particles converged, thus bringing downthe Reynolds number (Re = qvD/l where q, v, D, and lare the density, velocity, diameter, and viscosity of thefluid, respectively) from over 25 at spray distance 35 mmto about 7.5 at 75 mm (Ref 69). As the kinematic viscositydecreased rapidly, the convective movement and oxidemass transfer from particle to surface also ceased.Therefore, formation of nodules only occurred within thefirst 40-50 mm downstream of the nozzle exit under

Fig. 2 Schematic of the potential oxidation regions within a generic thermal spray process

Table 4 Conditions affecting coating oxidation in Regions I, II, and III in HVOF jet (Ref 58)

Region I Region II Region III

Estimated fraction of oxygen (%) ~2 ~20 ~20Exposure time (s) 0.001 0.001 0.01-0.02Relative temperature High Medium High, then lowParticle shape Sphere Sphere Splat

Fig. 3 A model of convective movement by the spherical Hill vortex mechanism within an iron particle in the core of an Ar-H2 plasma(Adapted from Ref 91)

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standard experimental conditions (Ref 70). At distancesfurther than this, the dominating mechanism is diffusion-based surface oxidation.

The oxides started to solidify and formed a thinoxide shell around the droplet as the particle tempera-ture decreased, Fig. 4(b) and 5(b). Deshpande et al.(Ref 2) also proposed for the plasma spray process thatoxides segregated toward the front end of the particle asthe particle velocity increased due to the difference inrelative velocity between the particle and plasma. Thehigh radiation from the segregated oxides caused cool-ing at the front end of the particle. As a result, the

metal oxides in the droplets were pushed from the hotregion toward the cooler front end by capillary flow toform an oxide cap, i.e., segregation of oxide at one endof the particle, Fig. 6. Syed et al. (Ref 70) also estab-lished that smaller particles promoted the entrainmentof oxide toward the tail of the particle. Figures 5 and 6highlight the formation of nodular oxides and surfaceoxide in collected sprayed particles. The presence ofthese oxides, in shell and nodular morphologies, hasbeen verified by other researchers including Espie et al.(Ref 69), Sobolev and Guilemany (Ref 72), and Li andLi (Ref 73).

Fig. 4 Schematics that illustrate in-flight oxidation. Molten particles and formation of the first splat with either (a) internal nodularoxides or (b) a surface oxide shell (Adapted from Ref 2)

Fig. 5 Micrographs of collected particles: (a) surface of particles collected at 55 mm and (b) cross section of particles collected at120 mm. Points 1 and 2 represent the oxide shell and nodular oxide, respectively (Adapted from Ref 111)

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In plasma spray, the in-flight oxidation of metallicparticles is also influenced by the rate of particle evapo-ration, which occurs due to the high heat flux from theplasma gas (Ref 62, 74). The injected particles penetratethe plasma jet core and are rapidly oxidized and melted.The particle temperature continues to increase since theparticles are within the hottest zone of the plasma. Theparticles would start to vaporize once they reach the boil-ing point of the material. The evaporated particles areeither oxidized to the corresponding metal oxide or con-densed in the form of metallic fume. In the case where theboiling point of the oxides is much lower than that of thepure material, such as tungsten, the oxides are evaporatedat such a high rate that the coatings will be almost oxide-free. The rate of particle evaporation is determined by thepartial pressure of oxygen in the boundary layer of theparticles, which is restricted by the formation of an oxidemist that surrounds the particle. When the particles travelwithin the jet core, this oxide mist can be swept away bythe gas flow around the particles, and the particles con-tinue to evaporate. However, a less optimum particletrajectory will cause the partial pressure of oxygen to in-crease due to the buildup of counter-diffusing oxygen andmetallic vapor at the particle surface (Ref 62). Conse-quently, the particle evaporation diminishes, and surfaceand nodular oxides continue to form as previouslydescribed. Therefore, the rate of oxidation is limited bythe extent of particle evaporation.

Upon reaching the substrate, splats are formed on top ofeach other that are separated by alternating oxide layers asshown in Fig. 7. These oxide layers originated from the oxideshell formed during in-flight oxidation. The surface of thesolidifying splat was exposed to a surrounding oxygen-richatmosphere and, hence, can be further oxidized by theimpinging jet until the splat was covered by the next arrivingmolten droplet, i.e., post-impact oxidation. According toEspie et al. (Ref 69), splat oxidation can be neglected since

the time lag for each splat to be covered by the next splat, andhence the time for splat exposure to the oxygen in the jet, ismuch shorter than the time lag between two successive pas-ses. Each splat was covered by the next one within 10-20 ls,while a few seconds occur between two successive passes(Ref 31, 69). The splat layers are exposed to hot gases in theboundary layer flow that contains more than 90 vol.% of airbetween two successive passes (Ref 31, 61). On the otherhand, Hackett and Settles (Ref 58, 75) believed that signifi-cant oxidation occurred after particle impact on the sub-strate.

The powder feed rate, torch traverse speed, depositionefficiency, and time between passes, which affect the time lagbetween the formation of two successive splats (Ref 2), playan important role in these opposing views. In addition, splatmorphology influences the extent of post-impact oxidation.Splats with a higher surface area, such as fragmented splats,are oxidized to a greater extent than disk-shaped splats of thesame volume under the same processing conditions.

In studying oxidation of a Cr3C2-NiCr coating, Mat-thews et al. (Ref 76) found that the splats were oxidizedindependently. That is, (i) carbide oxidation occurred bydiffusion of oxygen through the oxide layer to the oxide-carbide interface, while (ii) oxidation of NiCr occurred byoutward diffusion of chromium to the oxide-gas interface.Nickel oxide nucleated rapidly before being encapsulatedfrom below by the slower growth of chromium oxide. Thedifferent oxidation mechanisms are attributed to the dif-ferent structures of Cr3C2 and NiCr. Unlike NiCr, thechromium atoms in Cr3C2 are locked into the compoundstructure and are not free to diffuse out to the gas-oxideinterface, hence resulting in the diffusion of oxygen into acarbide phase instead (Ref 76).

Zeng et al. (Ref 77) reiterated that spray distancegreatly affected in-flight oxidation, while post-impactoxidation was influenced by particle size and substratetemperature. There is controversy concerning the

Fig. 6 (a) Micrograph of the polished cross sections of a collected particle and image treatment analysis of (b) nodular oxides, and (c)oxide cap (Adapted from Ref 70)

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oxidation stage that dominates the thermal spray process.The oxidation process depends on the spray techniqueinvolved, spray parameters, and feedstock material. Thereis no single explanation that rules out the oxidation pro-cess. It is a complex process that involves many factors.Table 5 compares the various aspects of in-flight and post-impact oxidation in the thermal spray process.

4. Factors that Influence Oxidation

4.1 Types of Spray System

The factors affecting oxidation are governed by the typeof oxidation process that is dominant. The oxidation processin control must be determined before the parameters can beadjusted to achieve the desired outcome. Oxidation isfavored in both HVOF and plasma spray due to the pres-ence of combustion product containing excessive oxygen inHVOF, the high temperatures involved in the plasma pro-cess (Ref 66), and mixing and entrainment with the ambientenvironment. The prevailing oxide varies for different spraytechniques. Deshpande et al. (Ref 2) conducted a study onthe oxidation of coatings sprayed under different tech-niques. It was found that APS and wire arc spray producedcoatings with more porosity and oxides compared to

HVOF. This was due to the fact that in-flight oxidation wasmore extensive in APS and wire arc spray because of thelonger residence time and greater degree of melting com-pared to HVOF. Fukushima and Kuroda (Ref 78) agreedwith Deshpande et al. (Ref 2) that HVOF suppressed oxi-dation of metallic materials better than APS, whereasKnight and Smith (Ref 79) argued otherwise.

As pointed out earlier, oxides formed during the plas-ma spray process segregated toward the front end of theparticle due to relative velocity effects. On the contrary,the lower temperature of HVOF led to higher viscosityand a lower relative velocity difference between the flameand the particle, thus resulting in internal oxides and a thinoxide shell, but no oxide segregation. Post-impact oxida-tion was also more pronounced for an APS coating, whileonly thin oxide bands were observed for a HVOF coatingdue to the reduced exposure time of the previous splat.

The level of oxidation observed for HVOF coatingsvaries with respect to the different HVOF systems. TheJP-5000 system (Praxair Surface Technologies, IN, USA)reduced oxygen content using over-stoichiometric condi-tions, while an under-stoichiometric condition was neededto attain low oxygen contents when using the Diamond Jet(DJ) Hybrid systems (Sulzer Metco, Wohlen, Switzerland)(Ref 59). This variation was attributed to the different fuelsused in each system and hence different combustion prod-ucts and temperature. The two systems showed contrastingeffects on oxidation due to their distinctive powder injectionposition. Feedstock powder was injected directly into thecombustion chamber, before the de Laval nozzle, in a DJHybrid system. On the other hand, in the JP-5000 system,the feedstock was injected after the combustion chamber,leading to lower temperatures in the system. The position ofpowder injection was important since it determined the typeof oxidation that dominated. In-flight oxidation dominatedfor traversing particles, while post-impact oxidation pre-vailed for external powder injection techniques (Ref 75).The lower temperature system resulted in lower meltingthat reduced the degree of splashing during the splatquenching process. In this manner, the surface area exposedto the atmosphere decreased, and hence there was lessoxidation. Nonetheless, the oxidation level was still influ-enced by the effects of other parameters such as the oxygen/fuel (O/F) ratio, chamber pressure, gas flow rate, particlesize, and spray distance.

Another example of differing spray system and injec-tion position effects on the dominating oxidation processis the coatings produced by gas-stabilized plasma (GSP)and water-stabilized plasma (WSP). Although the coatingssprayed by both techniques exhibited similar oxygencontent, the oxide morphologies of the coatings vary sig-nificantly. The GSP coatings contain mainly nodular oxi-des with a very thin oxide shell (less than 100 nm), whilethe WSP coatings contain a thick oxide shell with nonodular oxide (Ref 69). The oxide shell in WSP low car-bon steel coatings was two to four times thicker than thatof GSP coatings. Such difference in oxide morphologieswas attributed to variation of injection position. The par-ticles have to be injected further downstream (65 mmfrom exit nozzle) for WSP compared to GSP (3 mm from

Fig. 7 Schematics of consecutive splat formation on oxidizedand non-oxidized regions of pre-deposited splat for (a) theatmospheric plasma spray (APS) process and (b) the highvelocity oxy-fuel (HVOF) process (Adapted from Ref 2)

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exit nozzle) to reduce vaporization of particles (Ref 69).Consequently, the Reynolds number for particles in GSPwas more than four times that for particles in WSP, andhence there was more convective movement within theparticles that caused the formation of nodular oxide. Theoxide shell in WSP coatings was much thicker since theparticles travel a greater distance and at a lower velocitycompared to the GSP process, which allows more time forthe growth of the oxide shell.

4.2 Oxygen-to-Fuel Ratio and Current Intensity

The O/F ratio was considered the most influentialparameter that affected oxygen content for a HVOF-sprayed coating (Ref 59). As mentioned earlier, differentspray systems have varying O/F ratio effects on oxidation,and the under-stoichiometric condition does not neces-sarily favor oxidation. There are opposing opinions con-cerning the influence of the O/F ratio on oxidation. On theone hand, it was proposed that increasing the O/F ratiodecreased oxidation because combustion temperaturedecreased (Ref 75). Voggenreiter et al. (Ref 80) estab-lished that the particles must be sufficiently cold so thatfragmentation does not occur upon impact since moltenparticles would tend to fragment upon impact and createan increased area that would be liable to oxidize. On theother hand, decreasing the O/F ratio decreased oxidationsince excess oxygen was minimized at a low O/F ratio(Ref 81). The reduction of excess oxygen directly dimin-ished in-flight oxidation.

Syed et al. (Ref 64, 65) investigated the effect of cur-rent intensity and the hydrogen ratio of plasma spray onin-flight oxidation. Figure 8 shows the effects of currentintensity and hydrogen flow rate on the oxygen content aswell as formation of nodular oxides and oxide cap. It wasfound that higher current intensity favored in-flight oxi-dation and the formation of nodular oxides. An increaseof current intensity led to a rise of temperature as well asthe particle velocity, Fig. 8(a). Shorter dwell time and highvelocity led to a decrease of surface oxide formation, whilethe particle Reynolds number and convective reactivity ofparticles escalated so that more nodular oxides wereformed. Since surface oxides were dominant in this case,the total oxide content decreased even though the amountof nodular oxides increased, Fig. 8(c). The rise in hydro-gen ratio also enhanced the formation of nodular oxides,while formation of an oxide shell increased sharply andthen deteriorated gradually, Fig. 8(d). A higher hydrogenflow rate encouraged the surface oxides of molten parti-cles to penetrate the powder in-flight and led to manynodular oxides inside the powder after solidification(Ref 82).

4.3 Gas Flow Rate

Generally, an increase of gas flow rate is expected tocause an increase in the flow momentum of either aplasma plume or a HVOF flame and hence an increase inparticle velocity and a decline in oxidation. However, theinteraction between the different gases involved in bothT

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plasma spray and HVOF systems, i.e., primary and sec-ondary plasma-forming gases, combustion gases, and car-rier gas, is complex. The effects of flow rates of these gaseson the particle characteristics and degree of oxidation areinterdependent.

The total gas flow rate has little effect on the gasvelocity and temperature within the nozzle since they arelimited by the nozzle and are considered to be at a chokingstate (Ref 83). However, a high total gas flow rate willincrease the mass flow rate of the free jet at the nozzleexit. The increased mass flow rate causes an increase oftotal thermal and momentum inertia, where the flame gascan be kept at high temperature and velocity for a longerdistance. The effects of total gas flow rate are valid forboth plasma spray and HVOF systems where the increaseof total gas flow rate of (i) the primary and secondarygases in plasma spray, and (ii) the oxygen and fuel gases inHVOF, increases the ratio of heat-generating gas to non-heat-generating gases such as the carrier gas. Since theparticle temperature and velocity both influence oxidationthrough the effects of degree of melting and in-flight time,respectively, the extent of oxidation will depend on thetype of dominating oxides as discussed in section 4.2.

In a plasma spray system, an increase in the primary gasflow rate imparts a greater momentum to the in-flightparticle, hence increasing the particle velocity anddecreasing the in-flight time. Also, the increase of plasmajet velocity constricts the penetration of particles fedradially into the jet core, therefore lowering the particletemperature. Increasing the primary gas flow rate, whilemaintaining the secondary gas flow rate, also decreases theplasma energy density and particle temperature since the

fraction of secondary gas that generates greater enthalpyis lowered (Ref 84). The shorter in-flight time and lowerparticle temperature induced by the increase of primarygas flow rate results in reduced oxidation.

On the other hand, the secondary gas flow rate did notalter greatly the plasma temperature and velocity. How-ever, the thermal conductivity and plasma-to-particle heattransfer increased distinctly (Ref 64). Planche et al. (Ref 85)also suggested the same mechanism, but claimed that theincrease in hydrogen flow rate led to higher temperatureand better melting. Greater heat transfer magnified particletemperature until a point where an asymptotic limit wasreached where further changes in average particle temper-ature became negligible (Ref 64). Once this temperaturewas reached, hydrogen served as a reducing species andreacted with the highly oxidizing species in the plasma jet,hence decreasing oxidation. While each powder exhibited adifferent melting state, Kim et al. (Ref 82) also agreed withPlanche et al. (Ref 85) that an increase in hydrogen flow rateboosted the fraction of molten feedstock in the plasma, i.e.,the lamellar structure of splats increased with hydrogen flowrate. Hydrogen flow rate had almost no influence on particlevelocity independent of the particle size, while the rise inparticle temperature became more significant as the particlediameter increased.

An increase in the carrier gas flow rate increases theparticle velocity, but at the same time the gas temperatureis lowered since the carrier gas absorbs the thermal energygenerated by the plasma-forming or combustion gases(Ref 83, 86). Consequently, the heat transfer to the par-ticle and degree of melting are lowered, the in-flight timeis shortened, and the extent of oxidation is diminished.

Fig. 8 (a) Effects of current intensity on in-flight particle temperature and velocity (Ref 117), (b) effects of current intensity andhydrogen flow rate on the oxygen content (Ref 65), (c) percentage area of oxide nodules and oxide cap as a function of current intensity(Ref 65), and (d) percentage area of oxide nodules and oxide cap as a function of hydrogen flow rate (Ref 65)

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4.4 Particle Size

The studies by Gourlaouen et al. (Ref 87, 88) andFukushima and Kuroda (Ref 78) showed different degreesof oxidation, although both groups used the same sprayprocess and material. It has been established that particlesize plays a significant role in this discrepancy. Con-tradicting findings on the effects of particle size on oxi-dation of thermal spray coatings have also been noted.Effects of particle size on oxidation depend on the type ofoxidation that dominated. Generally, oxidation wasenhanced when the particle size decreased if in-flightoxidation was in control (Ref 2, 64, 72, 89, 90). The rea-soning was that the smaller particle size provided a highersurface area and hence more exposure to oxygen in-flight.In addition, Sobolev and Guilemany (Ref 72) recom-mended the use of feedstock with a narrow size distribu-tion and large mean size of particles to reduce oxidation.

Li and Li (Ref 73) found that oxygen content did notchange much for particle sizes below 45 lm. On the con-trary, oxygen content increased exponentially withdecreasing particle size by a power of approximately 3.4for particles larger than 45 lm. It was suggested that therewas a transient particle size range where the dominantmechanism changes, and this was reflected in the oxida-tion effects on the particles. The size range depended onthe spray conditions. For larger particle sizes, post-impactoxidation was responsible for the oxygen content incoatings. However, it was claimed that the contribution ofpost-impact oxidation to total oxygen content in thecoating was significant only when the total oxygen contentin the coating was low. In other words, in-flight oxidationwas almost always the major type of oxidation. This claimwas supported later by Espie et al. (Ref 69). Figure 9(a)(Ref 73) shows the effect of particle size on the oxygencontent in coatings as a function of spray distance.

4.5 Spray Distance

Spray distance or standoff distance was another impor-tant parameter that determined the extent of oxidation in acoating. As mentioned in section 3.2, the mechanism ofoxidation was proposed to change from convective to

diffusion when the spray distance increased, thus resultingin the formation of nodular oxides and an oxide shell,respectively. Figure 10 shows the type of oxides observed incollected particles sprayed at different spray distances.

Espie et al. (Ref 91) found that the oxygen content ofplasma sprayed low carbon steel particles increased withspray distance. Zeng et al. (Ref 77) presented comparableresults on plasma sprayed iron and NiCr-based alloys.Although Fukushima and Kuroda (Ref 78) reported sim-ilar results on plasma sprayed Ni-20Cr, a HVOF-sprayedstainless steel coating showed the opposite effect. Onceagain, the dominating oxidation process, i.e., in-flight orpost-impact, played an important role. The changes inoxygen content as a function of spray distance for in-flightoxidation and post-impact oxidation are shown inFig. 9(b) (Ref 73). In cases where in-flight oxidation was incontrol, a short spray distance reduced oxidation becauseparticles spent less time in the flame and were less exposedto in-flight oxidation. On the other hand, in situationswhere post-impact oxidation matters more, a longer spraydistance minimized oxidation since the temperature on thesubstrate surface was reduced with the heat source furtheraway.

4.6 Substrate Temperature and Heat Treatment

Substrate temperature did not affect in-flight oxidation,but it influenced post-impact oxidation (Ref 77). Hackettand Settles (Ref 75) demonstrated that oxidation wasenhanced when the substrate temperature was high. Ahigh substrate temperature impeded the solidificationprocess and hence exposed the molten splat to the sur-rounding atmosphere for a longer time. On the contrary,Espie et al. (Ref 69) contended that substrate temperaturecontrol did not affect oxygen content of the coating sig-nificantly, but it led to better adhesion between the coat-ing and substrate. The improved adhesion was proposed tobe caused by the formation of an appropriate oxide layeron the substrate.

Heat treatment of a coating often led to intersplatsintering and the disappearance of splat boundaries(Ref 76). Post-spray heat treatments also reduced thehardness of the coating, while metallurgically enhancing

Fig. 9 Diagram illustrating the dependency of oxygen content in a coating on the spray distance in terms of different (a) particle sizesand (b) dominating oxidation process (Adapted from Ref 73)

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the bond strength. The fine grain microstructure andenhanced bond strength helped prevent corrosive liquidsfrom penetrating the coating between particles. In otherwords, oxygen may also be prevented from entering thevoids between particles and, hence, mitigate against oxi-dation during any aging process. In terms of oxides, heattreatment was claimed to break off the continuous surfaceoxide layer and form oxide spheres (Ref 59). It was sug-gested that the oxide spheres filled the spaces betweenparticles and improved the bond strength.

In short, the oxidation process can be controlled bymanipulating the process parameters provided that themechanism and the dominating oxidation process areunderstood. Generally, the extent of oxidation is regulatedby two main factors: (i) the in-flight time or particlevelocity and (ii) the degree of melting or particle tem-perature. These two factors can be used to control oxi-dation concurrently or separately and are dictated by thevariables discussed in section 4. Firstly, oxidation can belimited by shortening the particle in-flight time orincreasing the particle velocity so that the time availablefor reaction between the particles and oxygen is limited.This can be achieved using (i) a high velocity process suchas HVOF, (ii) a high current intensity, (iii) a high total,primary, and/or carrier gas flow rate, and/or (iv) a shortspray distance.

Alternatively, a low level of oxidation can be achievedby lowering the degree of particle melting or the particletemperature, where the reactivity of the particle withoxygen is limited. This condition can be established using(i) a lower temperature process, (ii) feedstock injectionafter the combustion chamber, (iii) a low current intensity,(iv) a high primary and/or carrier gas flow rate, (iv) a lowtotal and/or secondary gas flow rate, (v) a large meanparticle size with narrow particle distribution, and/or (vi) alow substrate temperature. Particle temperature can alsobe lowered for the HVOF process using a high O/F ratio,but care is required when exercising this approach since italso creates an oxidizing environment that will enhanceoxidation. The extent that these process parametersinfluence the oxidation differs and will further depend on

the type of dominating oxidation process, i.e., in-flight orpost-impact.

5. Protection Against Oxidation DuringProcessing

Metals and alloys deposited by thermal spray processescan have some level of protection against oxidation(Ref 69). Such protection strategies include the following:

� Shrouding with a combustible gas such as acetylene.

� Effective use of a non-combustible protective gas.

� Maintaining the substrate at the lowest possible tem-perature that is compatible with good adhesion/cohesion.

� Rapid cooling.

� Use of a vacuum or controlled atmosphere chamberto evacuate air.

Alloy powders have been sprayed with the highvelocity air-fuel (HVAF) process and cold spray to pre-vent oxidation (Ref 92, 93). Oxidation was greatlyreduced, but the coating density tended to be inadequateand post-spray sealing was usually required. Sprayingunder vacuum conditions was effective, but at the expenseof an increased processing cost. Also, generating a vacuummedium was time consuming and the size of the parts tobe created was restricted by the chamber size. As a result,many researchers have employed gas shrouding, and thishas proved to be effective and relatively inexpensive.

5.1 Gas Shroud Versus Gas Shield

Protection measures against oxidation have beenimplemented as early as in the 1960s. Patents have been filedon varying designs of shields and shrouds for thermal spraypurposes (Ref 94-97). Schematics of these gas shrouddesigns are shown in Fig. 11. Both gas shield and shroud arecomponents that are flushed with inert gas during operation

Fig. 10 Micrographs of particle cross sections showing the effects of spray distance on the dominating type of oxide: (a) nodular oxidesfor particles collected at 55 mm and (b) oxide shell for particles collected at 75 mm (Adapted from Ref 111)

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for oxidation protection, which differ in their attachmentposition, i.e., a shield is attached to the substrate surface,while a shroud is attached to the thermal spray source.

Jackson (Ref 94) used argon and helium gases asshielding gas to reduce the oxygen concentration in coat-ings. This type of protection has been used for variousthermal spray techniques such as plasma spray, wire arc,and also HVOF. Figure 12 shows the comparison of theoxide contents in coatings sprayed with and without a gasshroud. The type of protection used, however, dependsgreatly on the oxidation process involved, i.e., whetheroxidation was dominated by in-flight or post-impactmechanisms. In cases where in-flight oxidation dominated,a gas shroud was a good option, while at times where post-impact oxidation was in control, a shield must be used.

Shield and shroud methods have been presented byFukushima and Kuroda (Ref 78), Fig. 13.

Mostaghimi and co-workers developed a cylindricalshroud that successfully minimized in-flight oxidation.However, the shrouded nozzle causes recirculation of flowwithin the nozzle and decreases the particle velocity(Ref 98). In addition, the recirculation of cold air at the exitof a conventional conical gas shroud nozzle reduces theplasma enthalpy and the particle temperature (Ref 99). TheMostaghimi group further developed a curvilinear insert tothe gas shroud nozzle that conformed to the streamlines ofthe primary gas flow and eliminated the air recirculation(Ref 99, 100). This modification impeded mixing with thesurrounding atmosphere and improved the heat andmomentum transfer to the particles.

Fig. 11 Schematic of several gas shroud patents designed for thermal spray by (a) Jackson, Union Carbide Corporation (Ref 94), (b)Rotolico and Weodarczy, The Perkin-Elmer Corporation (Ref 95), (c) Simm, Steine, and Streb, Castolin S.A. (Ref 96), and (d) Reiter,Nova-Werke AG (Ref 97)

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In the process of understanding coating oxidationcontrol, Hackett and Settles (Ref 58) used Schlierenimages to visualize the operational mechanics of gasshrouding. It was shown that the shroud gas flows coaxially

to the HVOF jet from the spray torch and was rapidlyentrained. The extent of the entrainment was indicated bylimiting the streamline between the shroud gas and theambient atmosphere. Further downstream of this point,the ambient atmosphere became entrained into the jet.The inert shroud gas became distributed throughout themixing layer by vigorous turbulent mixing. This mecha-nism served to dilute the amount of oxygen present in thejet centreline. The effects of gas properties and flow ratewill be discussed further. Hackett and Settles (Ref 58)established that shrouding controlled the formation ofcoating oxides and also reduced the oxidation levels iniron coatings by as much as 50 wt.%.

Hackett and Settles (Ref 58) demonstrated that coat-ings sprayed using a ducted controlled atmosphere spray(CAS) has lower oxygen content than the coatings sprayedusing an unconfined shroud for the same gas mass flowrate. In other words, the working area of a gas shroudaround the torch must be enclosed for better efficiency.Wang et al. (Ref 101), who implemented two types ofshrouding on the wire arc spray process, also verified thesame outcome. An unconfined shroud demonstratedhigher oxide levels due to air entrainment. Spraying in aconfined gas shroud also resulted in more uniform particlesize distributions, more focused spray patterns, higherparticle velocities, and improved coating properties(Ref 101). Figure 14 shows the gas shroud configurationused by Wang et al. (Ref 101) where Fig. 14(a) shows anunconfined shroud, while Fig. 14(b) is confined.

5.2 Selection of Inert Gas

Inert gases, such as argon, nitrogen, and helium, arenormally used as shroud gases. Hackett and Settles

Fig. 12 Cross section of plasma sprayed CoNiCrAlY coatingssprayed (a) without and (b) with nitrogen shrouding (Adaptedfrom Ref 118)

Fig. 13 Schematics of two types of gas shielding attachmentsused to control oxidation: (a) nitrogen gas shielding attachmentused on a substrate and (b) gas shroud used to control the oxi-dation of HVOF-sprayed particles during in-flight (Adaptedfrom Ref 78)

Fig. 14 Schematics of nozzle and shroud configuration for (a)straight bore nozzle with shroud inert gas protection and (b)straight bore nozzle with secondary gas injection (Adapted fromRef 101)

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(Ref 58) compared the effect of different gases for shroudingon oxidation and concluded that the choice of shroud gashad little influence on the oxygen content in the coating. Anargon-shrouded coating exhibited a similar oxygen contentto a nitrogen-shrouded coating at the same shroud gas massflow rate. A helium shroud produced the lowest oxygencontent coatings at a low shroud gas mass flow rate, takingnote that the volume flow rate was equivalent to the highestnitrogen flow rate due to the low density of helium(0.17 kg/m3 for helium compared to 1.17 kg/m3 for nitrogenand 1.66 kg/m3 for argon). Thus, coatings produced undernitrogen and helium shrouds exhibited similar oxygen levelsat the same volume mass flow rates.

An interesting study by Volenık et al. (Ref 63) usedacetylene (C2H2 with a density of 1.09 kg/m3) as a shroudgas instead of the conventional inert gases. It was foundthat nitrogen-hydrogen shrouding was less effective indecreasing oxide contents and the oxide layers werethicker with pronounced dendritic growth. A combustiblegas, acetylene, on the other hand demonstrated a twofoldeffectiveness in reducing the oxide levels. Acetyleneworked because it displaced air and reduced the amountof oxygen available in the plume by consuming it in acombustion process. The use of carbon dioxide (CO2 witha density of 1.84 kg/m3) as a gas shroud instead of theconventional inert gases has also been found effective byWang et al. (Ref 102). The oxide content of the stainlesssteel coatings decreased by 15 wt.% when produced undera carbon dioxide shroud.

Kang and Hong (Ref 103) found that a shroud gas flowrate 3.75 times that of the torch working gas was necessaryto achieve a ‘‘nearly oxygen-free plasma’’ region with anair content of approximately 10 mol.% at 100 mm down-stream of the nozzle exit. Gawne et al. (Ref 104), on theother hand, showed that the air content can be reducedfrom 82 to 46 vol.% with a 100 m/s velocity shroud gas atthe turbulent jet center, 105 mm from the torch nozzleexit. Hackett and Settles (Ref 58) found that an increase ingas shroud mass flow rate completely enveloped theHVOF jet and very little atmospheric oxygen wasentrained. Despite this, some minimum amount of oxygenwas always present in coatings due to either (i) free oxy-gen combustion in Region I (refer Fig. 1) or (ii) oxidesthat were present in the feedstock powder. It was alsofound that an increase in the inert shroud gas mass flowrate resulted in a linear decrease in oxygen content.

It has been established that the gas shroud and/orshield can affect the temperature, velocity, and air con-centration fields in turbulent plasma jets. While gasshrouding was very effective in minimizing oxidation,Dolatabadi et al. (Ref 98) claimed that the coatings weremore porous due to a reduction in particle impact velocity.On the contrary, Wang et al. (Ref 101) disagreed andfound that shrouding reduced the porosity level in coat-ings. Instead of reducing particle impact velocity, Wanget al. (Ref 101) argued that shrouding accelerated thedroplets and air entrainment caused less cooling.Hence, the coating was built up at higher velocity and

Fig. 15 Cross-sectional images of captured particles and coatings depicting the effects of SEs: (a) and (c) without SEs, (b) and (d) withsilicon and boron as SEs (Adapted from Ref 77)

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temperature that led to reduced porosity. It was estab-lished that shrouding reduced the amount of oxides andthere was more uniform contact between the metallicsplats. Planche et al. (Ref 85) also agreed that coatingsfrom shrouded spraying were less porous than conven-tionally sprayed coatings due to the increase in tempera-ture and velocity in the shroud.

The above contradicting views are attributed to thedifference in materials, nature of gas shroud used, and theflow rate. In all cases, each shroud must be designedspecifically according to the spray gun used so that oxi-dation can be reduced and, at the same time, any adverseeffects of using such a shroud can be minimized. Theoptimum geometry must be used to achieve this. Modelingand theoretical computation have been performed fordifferent shroud designs (Ref 103-106) to verify theeffectiveness of gas shrouding as well as to design the idealshroud geometry.

5.3 Sacrificial Elements

Besides gas shrouding, some light alloying elements,such as silicon and boron, can be added as ‘‘sacrificialelements’’ (SEs) in the feedstock. The SEs must (i) readilyT

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Table 7 Hardness values of various metals and theiroxides (Ref 129)

MetalHardness,kg/mm2

Metaloxide

Hardness,kg/mm2

Aluminum 15 Al2O3 1800Chromium (plated) 850 CrO3 1600Copper 40 Cu2O, CuO 130Iron 105 FeO, Fe2O3, Fe3O4 400Lead 5 PbO, PbO2 250Magnesium 27 MgO 400Molybdenum 220 MoO3 50Nickel 86 NiO, Ni2O3 430Tin 5 SnO2 1650

SnO 350Zinc 35 ZnO 200

Fig. 16 Variation of Vickers hardness values with oxygen con-tent in high velocity oxy-fuel-sprayed stainless steel 316L coat-ings (Adapted from Ref 59)

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react with oxygen to form oxides and (ii) exhibit greateraffinity to oxygen than the core elements. Figure 15 showsthe effects of SEs on the reduction of oxidation for col-lected sprayed particles and coatings. The reduction ofon-substrate oxidation becomes more obvious with theincrease of SE content. Zeng et al. (Ref 77, 107) foundthat the SE content was reduced in captured in-flightparticles and coatings compared to the feedstock, butrevealed no difference between the captured particles andcoatings. To look at this another way, the addition of suchSEs was more effective against in-flight oxidation than forpost-impact oxidation. Boron was found to be the mosteffective SE against the oxidation of iron. However, extraprecautions must be taken when using such SEs becausethe oxidation of such elements forms gaseous oxides.

Powders that contain SE are often referred to as self-fluxing alloys since they were intended to produce densecoatings by having similar effects as fused coatings,although post-deposition fusing is sometimes still required.The SE used in such applications typically has a lowermelting point and is more easily oxidized than the basematerial (BM). Self-fluxing alloys function by incorporatinga thermally oxidizable element, such as aluminum, that

reacts exothermically with oxygen to release a largeamount of heat to produce good bonding. For this reason,such powders are also used as a bond coat to improveadhesion of coating to metal substrates. These powders areeither manufactured as (i) SE-coated particles with ametallic core or (ii) powder mixture or agglomerated par-ticles containing SE. Examples of such powders are listed inTable 6. For powders manufactured in the form (ii), the SEparticles are typically smaller than the particles of BM sincethis will provide the SE particles with greater surface areaand hence greater tendency to oxidize in preference to theBM (Ref 108). The percentages of SE in the powders varydepending on the type of material and oxidation effectsdesired; refer to examples in Table 6.

6. Gaging the Effects of Oxidation viaHardness Testing

The extent of oxidation that occurred during processingcan be measured by various techniques, such as imageanalysis (Ref 70, 109, 110) and analysis of the oxygen/

Fig. 17 Flowchart of oxidation considerations for thermal spray processing of materials

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nitrogen content (Ref 59, 111). Hardness or indentationtests present a simple, indirect approach to measure theapproximate oxide content in a material. The occurrenceof oxidation is reflected in the hardness values of thecoatings since metal oxides generally exhibit higherhardness compared to their metallic counterparts. Acomparison of hardness values between metal and theiroxides is presented in Table 7. Oxidation of Nd-Fe-B hasalso been evidenced to result in increased hardness(Ref 29).

Hardness of thermal sprayed coatings has been estab-lished to depend on the combined effects of unmelts,porosity, residual stresses, and oxide inclusions within thecoatings. Although there have been no direct correlationsestablished between hardness and oxide content, thecoating hardness tends to increase with an increase inoxide content (Ref 59). Other work (Ref 101, 112) has alsoattributed the variation of hardness in thermal sprayedcoatings to the presence of oxides. Therefore, hardnesstesting methods can provide an estimate of oxide contentwhen a suitable calibration is carried out. Figure 16 showsan example of hardness value variations with oxygencontent in a coating.

The practicality of using a hardness test or micro-indentation test to evaluate the effects of oxidation isdemonstrated by its use in evaluating mechanical propertychanges and damage in TBCs induced by oxidation andresidual stresses (Ref 113, 114). TBCs are applied tosuperalloy materials to reduce oxidation and thermalfatigue of load-bearing alloys for high temperature appli-cations. TBCs consist of a ceramic top coat and a metallicbond coat that is applied to a superalloy substrate. Athermally grown oxide (TGO) layer, which results fromoxidation of the bond coat, develops at the interfacebetween the top coat and bond coat during thermal cy-cling. The TGO layer causes a thermal expansion mis-match and development of stresses, which result in coatingspalling and failure.

However, the evaluation of oxidation behavior in TBCsis more complex than as-sprayed coatings where oxidationoccurs during deposition, due to different oxidationmechanisms. Other factors, such as effects of sintering,oxide scale growth, and development of cracks, need to betaken into consideration since they have varying effects onthe mechanical properties. Reviews on TBCs and theirfailure mechanisms are available in (Ref 115, 116).

7. Conclusions

Oxidation of thermal spray coatings poses concerns inapplications in which the oxides interfere with the per-formance of the coatings, such as in corrosion resistance,magnetic properties, and mechanical properties. Themechanisms and influencing factors of oxidation must beunderstood in order to control the extent of their oxida-tion by implementing protection measures. The flow chartin Fig. 17 summarizes the oxidation considerationsfor thermal spray process and forms part of the

decision-making process in determining the appropriatespray parameters and protection measures to use for thedesired oxidation effects.

Application of PMCs by thermal spray processes is stilllimited due to oxidation concerns. Oxidation causes theloss of ferromagnetic phases and leads to deterioration ofmagnetic properties. Oxidation of PMCs during the fab-rication process is still not completely understood sincemost studies are centered on oxidation during aging.Thermal oxidation of neodymium iron boron and samar-ium cobalt occurs over two stages: (i) formation ofexternal scales and (ii) internal oxidation. Neodymiumiron boron undergoes dissociation followed by oxidationto form a-Fe2O3 and Fe3O3 on its external scales and a-Fe,Nd2O3, NdFeO3, and B2O3 in its IOZ. On the other hand,the oxidation mechanism for samarium cobalt is still notfully understood. However, most oxidation studies agreedon the oxidation products of Co3O4 and CoFe2O4 on theexternal scales and a-Co, Sm2O3, and SmCo3 in its IOZ.Further work is still necessary to understand the oxidationof PMCs in thermal spray process.

There are two stages of oxidation in thermal spraycoatings: (i) in-flight and (ii) post-impact oxidation.In-flight oxidation encompasses two regions: (i) the jetcore and (ii) beyond the jet core and before impact ontothe substrate, where local atmosphere is entrained into theflame. Post-impact oxidation refers to the deposition,solidification, and cooling stages. The majority of oxida-tion occurs at the in-flight stage due to the high temper-ature and oxygen entrainment, although the exposure tooxygen during post-impact is longer than the in-flight time.In-flight oxidation resulted in nodular oxides and oxideshells, while post-impact oxidation produces surface oxi-des between splat layers.

The factors that influence oxidation rely on the oxida-tion process that dominates, and this depends on the spraysystem, powder injection position, and particle size of thefeedstock. In cases where in-flight oxidation is in control,the extent of oxidation is influenced by the oxygen-fuelratio, current intensity, gas flow rate, particle size, andspray distance. Extent of oxidation is governed by thesubstrate temperature and spray distance when post-impact oxidation is dominant.

Considering the detrimental effects of oxidation forrare-earth magnetic materials, some extent of oxidationprotection is possible for these materials to be depositedvia the thermal spray process. The protection measurewould depend on the oxidation process that is in domi-nance, i.e., in-flight or post-impact oxidation. Shrouding isconsidered the most effective as well as economical andpractical protection method, where its principal is basedon minimizing air entrainment. Shrouding should beemployed where in-flight oxidation is in control, whileshielding is more appropriate where post-impact oxidationdominates. Inert gases such as argon, nitrogen, and heliumare typically used, although acetylene and carbon dioxidewere also proven to be effective shrouding agents. Theextent of air entrainment reduction depends on the massand volume flow rate of the gases due to the differences ingas densities. SEs can also be used to minimize oxidation.

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The oxidation protection measures used should bedesigned for a specific spray system to achieve the opti-mum effect since each system has unique geometricalneeds. Consideration also needs to be turned to thegeometry of the substrate since this will influence the gasflow fields during the spray process.

Finally, indentation tests present a viable, simplealternative in evaluating the effects of oxidation in ther-mal sprayed coatings. However, this approach needs to beused judiciously since hardness values depend not only onoxide inclusions but also on other microstructural artifactssuch as unmelts and porosity.

Acknowledgments

The authors would like to acknowledge the financialsupport by a Swinburne University Postgraduate ResearchAward (SUPRA) for this research and development ofthermal spray processing of magnetic coatings.

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