Abnormal Nitride Morphologies upon NitridingIron-Based Substrates

SAI RAMUDU MEKA1,3 and ERIC JAN MITTEMEIJER1,2

1.—Max Planck Institute for Intelligent Systems (Formerly Max Planck Institute for MetalsResearch), Heisenbergstrasse 3, 70569 Stuttgart, Germany. 2.—Institute for Materials Science,University of Stuttgart, Heisenbergstrasse 3, 70569 Stuttgart, Germany. 3.—e-mail: s.meka@is.mpg.de

Nitriding of iron-based components is a very well-known surface engineeringmethod for bringing about great improvement of the mechanical and chemicalproperties. An overview is presented of the strikingly different nitride mor-phologies developing upon nitriding iron-based alloy substrates. Observedabnormal morphologies are the result of intricate interplay of the thermody-namic and kinetic constraints for the nucleation and growth of both alloyingelement nitride particles in the matrix and iron nitrides at the surface of thesubstrate. Alloying elements having strong Me-N interaction, such as Cr, V,and Ti, precipitate instantaneously as internal Me-nitrides, thus allowing thesubsequent nucleation and growth of ‘‘normal’’ layer-type iron nitride. Alloy-ing elements having weak Me-N interaction, such as Al, Si, and Mo, andsimultaneously having low solubility in iron nitride, obstruct/delay thenucleation and growth of iron nitrides at the surface, thus leading to very highnitrogen supersaturation over an extended depth range from the surface.Eventually, the nucleation and growth of ‘‘abnormal’’ plate-type iron nitrideoccurs across the depth range of high nitrogen supersaturation. On this basis,strategies can be devised for tuned development of specific nitride morpholo-gies at the surface of nitrided components.

INTRODUCTION

The interplay of kinetic and thermodynamic con-straints can lead to highly unusual, i.e., unexpected,precipitation sequences and precipitate morpholo-gies in solids.1 As will be demonstrated in thisarticle, a show case is provided by nitrogen super-saturated iron-based alloys containing alloying ele-ments with an affinity for nitrogen. The importanceof this system not only derives from its contributionto fundamental understanding of phase transfor-mations in solids but also reflects the ever-wideningfield of applications of the most versatile thermo-chemical surface engineering method of presentday: nitriding.

Nitriding implies the introduction of nitrogen,typically but not always, from a gaseous (e.g., NH3

containing) atmosphere, into the surface region of ausually ferritic iron-based component.2–4 As a result,large increases of the fatigue, wear, and corrosionresistances can occur.

Nitriding is applied as a process in technologysince its advent: a patent in the United States byMachlet in 19135 and a paper in Germany by Fry in1923.6 Despite its age, a fundamental understand-ing of many important aspects of this process, suchas the development of metastable and stablenitrides, has only fragmentarily been obtained,which obstructs optimization of properties by tunedapplication of nitriding: Industrial exploitation isstill largely based on phenomenological knowledge/experience.

Alloying elements, Me, as, e.g., Al and Cr, areoften added to the material to be nitrided (thus,such special ‘‘nitriding steels’’ have been developed)so that their precipitation as alloying element ni-trides in the nitrided zone has a beneficial effect onthe mechanical properties. Extensive research inthe last decades has been devoted to the precipita-tion of Me-nitrides in binary iron-based alloys (withMe = Cr,7–9 Al,10–13 V,14–18 Ti,19–21 Si,22–24 andMo25–27) and of mixed nitrides [Me1,xMe2,1–x]N

JOM

DOI: 10.1007/s11837-013-0603-6� 2013 TMS

(Me1 = Cr and Me2 = Al28–30 and Me1 = Cr andMe2 = Ti31,32) in ternary iron-based alloys.

Nitriding involves not only the precipitationof alloying elements, as indicated above, in thenitrided zone, but also the (simultaneous) develop-ment of an iron-nitride layer at the surface of thecomponent.2 This iron-nitride surface layer isthought to improve especially the tribological prop-erties. The interaction of the precipitation of alloy-ing element nitrides in the surface region and thedevelopment of the iron-nitride layer at the surface,i.e., the competition between the formations of ironnitrides and alloying element nitrides, is the subjectof this article. The roles of the Me-N affinity, theMe-nitride-precipitate/matrix misfit, the solubilityof Me in iron nitride, and the defect density of thematrix have been exposed. The understandingattained allows the development of strategies tocontrol the greatly variable microstructure of theiron-nitride compound layer developing at the sur-face of nitrided components.

EXPERIMENTAL PROCEDURES

The investigated alloy casts were produced byinduction melting of the pure elemental metals in aprotective argon atmosphere. Such produced castswere cold rolled to obtain sheets of about 1 mmthickness. Rectangular specimens were cut from therolled sheets and the surfaces of specimens wereprepared by grinding and polishing steps followedby ultrasonic cleaning in ethanol. These specimenswere subjected to the (thermodynamically) con-trolled gaseous nitriding treatment using a flowingmixture of high-purity ammonia/hydrogen gases.Important nitriding parameters are the tempera-ture and the nitriding potential. The nitriding po-tential rN is given by:

PNH3

P3=2H2

where P represents partial pressure; the nitridingpotential is proportional with the nitrogen activity inthe solid at the surface provided local equilibriumoccurs at the gas–solid interface.33 Temperature andnitriding potential were chosen such that thec¢-Fe4N1�x (x< 0.04; face-centered cubic sublatticeof iron atoms) iron nitride can develop on the surfaceof pure iron.34 The resulting microstructures werecharacterized by light microscopy (LM), scanningelectron microscopy, transmission electron micros-copy, electron probe microanalysis (EPMA), andx-ray diffraction (XRD) (for details, see Refs. 35 and 36).

IRON-NITRIDE DEVELOPMENT IN CASEOF INSTANTANEOUS Me-NITRIDE

PRECIPITATION

Consider the case of nitriding pure iron at condi-tions where c¢-Fe4N1�x iron nitride develops at the

surface. Due to the competition between the finiterates of surface reactions (ammonia dissociation andthe association of surface-adsorbed nitrogen atomsand their subsequent desorption as N2 molecules)and the nitrogen diffusion into the solid, uponnitriding the surface-nitrogen content increaseswith time.37,38 Once the nitrogen solubility limit ofthe ferrite matrix gets surpassed, c¢-Fe4N1�x nucleidevelop at the surface which by lateral growthconstitute a closed layer; continued growth of suchlayer occurs by diffusion of nitrogen through thelayer (Fig. 1).

For a comparison of the energetics of the precip-itation of different Me-nitrides in the matrix, a so-called Me-N interaction parameter39 can be defined,which is the ratio (per unit volume of precipitate) ofthe (chemical) Gibbs energy of formation of precip-itate and the (mechanical) misfit-strain energy in-duced in the precipitate/matrix system. Based onthe thus characterized strength of the Me-N inter-action, alloying elements can be classified as ele-ments of a ‘‘strong,’’ ‘‘medium,’’ or ‘‘weak’’ Me-Ninteraction.

Alloying elements such as Cr, V, and Ti can moreor less instantaneously precipitate as (rock-saltstructured) Me-nitrides (CrN, VN, and TiN) due to‘‘strong interaction’’ as a result of strong/moderateMe-N affinity and a moderate volume misfit inassociation with the establishment of a(semi)coherent interface with the matrix. Uponnitriding (at conditions where for pure iron c¢develops at the surface) iron-based substrates al-loyed with elements of strong Me-N interaction, theinwardly diffusing nitrogen atoms get consumed bymore or less instantaneous precipitation of Me-nitride particles. This process does not allow thenitrogen solubility limit of the matrix to be sur-passed; i.e., a nitrogen supersaturated matrix can-not develop effectively. Only after all Me hasprecipitated as Me-nitride in the surface adjacentregion, the nitrogen solubility limit of the matrix atthe surface gets surpassed leading to nucleation ofc¢ at the surface. The resulting c¢ layer grows intothe substrate under incorporation of the Me-nitrideparticles already developed in the matrix; i.e., thec¢ layer ‘‘overruns’’ the existing Me-nitride particles.The experimental results from nitrided Fe-Cr andFe-V substrates support the above interpretation(Fig. 2): Incorporation of Me-nitride particles (CrNand VN) into the growing c¢ layer was proven byEPMA and XRD analysis.40,41

Upon prolonged nitriding, the following micro-structural abnormalities occur:

1. Underneath the surface c¢ layer, c¢ can developalong the grain boundaries of the matrix (Fig. 3).This is attributed to Me (e.g., Cr) segregated atgrain boundaries, already in the unnitrided con-dition, leading to preferred precipitation of MeN(CrN) at the grain boundaries upon nitriding. As aconsequence, the prevailing supersaturation of

Meka and Mittemeijer

nitrogen in the grain-boundary surroundingregions then, in the absence of dissolved Me, canlead to the development of c¢ adjacent to such grainboundaries.

2. The initially developed nanosized, coherent MeNprecipitates, in the so-called continuous precipita-tion (CP) region, may coarsen via a discontinuouscoarsening reaction that results in the develop-ment of a lamellar structured, discontinuous coars-ened (DC) microstructure. As the ferrite matrixsurrounding the continuous MeN precipitates canpossibly contain significantly more excess nitro-gen* than the ferrite matrix surrounding thediscontinuously coarsened MeN precipitates,18

during the discontinuous coarsening a hugeamount of excess nitrogen gets released whicheither locally enhances the nitrogen supersatura-tion of ferrite leading to the development of c¢ alonggrain boundaries and also at the CP/DC interface,or associates at the grain boundaries leading to theformation of N2 gas filled pores, coalescence ofwhich leads to the development of open grainboundaries/cracks (Fig. 4). The penetration of thenitriding atmosphere (NH3/H2 gas mixture)through the cracks opened to the surface thenleads to the development of c¢ along the crack faces.

IRON-NITRIDE DEVELOPMENT IN CASE OFDELAYED Me-NITRIDE PRECIPITATION

Concerning the iron-based alloy substrates, withAl, Si, and Mo as alloying elements, the followingdistinct characteristics of these elements should berecognized for the case of nitriding at conditionswhere for pure iron c¢ develops at the surface:

1. Al, Si, and Mo have strong to moderate Me-Naffinity. However, their Me-N interaction (asdescribed in the beginning of the Iron-NitrideDevelopment in Case of Instantaneous Me-Nitride Precipitation section) is weak because ofthe large volume misfit with the ferrite matrix,and thus, the (internal) precipitation of these Me-nitrides is difficult.

2. Al, Si, and Mo have low solubility in c¢-Fe4N1–x

iron nitride. This hinders the development of c¢as this requires either Me partitioning in thematrix to form Me free c¢ or forced c¢ precipitationunder incorporation of Me. This process in anycase has a low driving force.1

As a consequence of (1) and (2), initially a highnitrogen supersaturation over a large depth rangecan develop in the nitrided substrate. If this nitro-gen supersaturation has become large enough, iteventually forces the formation of c¢ plates over thislarge depth range under incorporation (dissolution)of the alloying elements (Al, Si, and Mo) in thenitride (Fig. 5) (upon prolonged nitriding the lowsolubility of alloying elements in c¢ causes thedevelopment of some e-Fe3N1+x (hexagonal close-packed sublattice of iron) phase within the iron-

Fig. 1. Light optical micrograph recorded from the cross-section of a nitrided (580�C, rN = 0.8 atm�1/2, 2 h) pure iron specimen. ‘‘Normal’’ layer-type c¢ has developed.

Fig. 2. LMs recorded from the cross-sections of nitrided (580�C, rN = 0.8 atm�1/2, 2 h) Fe-4 at.% Cr (a) and Fe-4 at.% V specimens. ‘‘Normal’’layer-type c¢ has developed; the growing layer incorporates the Me-nitride particles, which had precipitated in the matrix before being ‘‘overrun’’by the growing layer.

*The N content in surplus of the sum of the N associated with theMexNy precipitates ð N½ �MexNy

Þ and the equilibrium N solubility ofthe unstrained ferrite matrix N½ �oa�Fe

� �is called ‘‘excess nitro-

gen,’’39 which is taken up as (I) additionally dissolved N in theferrite matrix due to the hydrostatic component of the misfit-strain field around the nitride precipitates and (II) adsorbed N atthe nitride-precipitate/ferrite-matrix interface.

Abnormal Nitride Morphologies upon Nitriding Iron-Based Substrates

nitride layer that apparently has a larger solubility forMe than c¢.35,36,42). Subsequently, the Me still dis-solved in the matrix adjacent to the c¢ plates, precipi-tates as alloying element nitrides (AlN, Mo2N, andSi3N4) using the dislocations generated to accommo-date the misfit of c¢ with ferrite.12,27 Experimentalresults support this interpretation: The developmentof Me-nitrides in the ferrite matrix takes place, after c¢has nucleated at the surface and had penetrated dee-ply, as plates, the substrate: An EPMA line scancrossing the c¢ plates developed at the initial stages ofnitriding shows a nitrogen content in the ferrite ma-trix surrounding the c¢ plates incompatible with thehigh amount of nitrogen expected there if MeNn pre-cipitationhad occurred (Fig. 6) (in the exampleshown,the nitrogen content of the matrix should have beenabout 5 at.%). Only at later stages of nitriding Me-nitrides developed in the matrix (Fig. 7).

The precipitation kinetics of the Me-nitrides canbe enhanced by (I) increasing the nitriding tem-perature or (II) introducing defects (dislocations)into the substrate on which easy development ofprecipitates occurs. Then, the barrier for c¢ nucle-ation and growth has been taken away, and indeed

for these two cases, a layer-type growth of c¢ isobserved (Fig. 8a and b).

The above interpretation is also supported by thefollowing experiments: Fe-Me substrates were firstnitrided such that only MeNn development is pos-sible (i.e., employing temperature and nitridingpotential corresponding to the a-iron region of theLehrer diagram34) and then nitrided at highernitriding potential allowing c¢ formation. Indeed, alayer-type growth of c¢ was observed upon such two-stage nitriding treatment.35,36

CONCLUSION

Nitriding of Fe-Me alloys leads to (high) super-saturation with nitrogen and competition betweeniron-nitride and alloying-element-nitride precipita-tion.

Delayed/obstructed precipitation of Me-nitride (asholds for AlN, Mo2N, and Si3N4 owing to high-vol-ume misfit; so-called ‘‘weak Me-N interaction’’), incombination with low solubility of Me in iron ni-tride, leads to high nitrogen supersaturation overan extended depth range. Eventually, c¢ iron-nitride

Fig. 3. Light optical micrograph recorded from the cross-section of a (longer than in Fig. 2a) nitrided (580�C, rN = 0.8 atm�1/2, 4 h) Fe-4 at.% Crspecimen. c¢ has also developed along the grain boundaries of the matrix.

Fig. 4. Light optical micrograph recorded from the cross-section of a (longer than in Fig. 2b) nitrided (580�C, rN = 0.8 atm�1/2, 4 h) Fe-4 at.% Vspecimen. c¢ has also developed along the grain boundaries of the matrix and along the CP/DC interface.

Meka and Mittemeijer

plates, nucleated at the surface, develop across thedepth range of high nitrogen supersaturation underincorporation (dissolution) of the alloying elements Me.

‘‘Normal’’ closed iron-nitride layers at the surfaceoccur if Me-nitride precipitation occurs fast/instanta-neously upon nitriding, i.e. (I) at higher temperature,

(II) in a matrix of high defect density, or (III) foralloying elements with so-called ‘‘strong Me-N inter-action’’ (as holds for Ti, V, and Cr). In the latter case,underneath the iron-nitride layer, often iron nitridedevelops along grain boundaries of the matrix. Thisphenomenon can be explained as follows:

Fig. 5. Light optical micrographs recorded from the cross-sections of nitrided iron-based alloys. (a) Fe-4.7 at.% Al alloy nitrided at 500�C for10 min using rN of 1.73 atm�1/2. (b) Fe-1 at.% Mo alloy nitrided at 480�C for 2 h using rN of 0.7 atm�1/2. (c) Fe-4.5 at.% Si alloy nitrided at 550�Cfor 2 h using rN of 0.82 atm�1/2. ‘‘Abnormal’’ plate-type c¢ has developed.

Fig. 6. Light optical micrograph showing the location of the EPMA elemental line scan (shown with white dashed line) crossing the c¢ plates,together with the elemental concentrations for the nitrided (550�C, 10 min, rN = 0.82 atm�1/2�) Fe-4.7 at.% Al alloy. Peaks in the nitrogen contentcorrespond with the EPMA line scan crossing a c¢ precipitate (plate). The ferrite matrix present between the c¢ plates contains a very low nitrogencontent (less than 0.3 at.%), which indicates that no AlN precipitates have formed inside the ferrite matrix; i.e., c¢ has formed with the locallypresent Al dissolved in it.35

Abnormal Nitride Morphologies upon Nitriding Iron-Based Substrates

1. The presence of segregated alloying element Me,at grain boundaries, already in the unnitridedcondition, leads to preferred precipitation of Me-nitride at the grain boundaries of the matrixbelow the surface iron-nitride layer. The prevail-ing supersaturation of nitrogen in regions adja-cent to such grain boundaries then, in theabsence of Me, can lead to the development ofiron nitride adjacent to such grain boundaries.

2. A loss of the capacity for excess nitrogen uptake, asdue to a continuous-to-discontinuous coarseningof Me-nitride particles, causes this excess nitrogento precipitate, e.g., as pores filled with N2 gas atgrain boundaries. Coalescence of such pores leadsto the development of open grain boundaries/cracks, allowing penetration of the nitridingatmosphere (NH3/H2 gas mixture) along these

open grain boundaries/cracks, leading to thedevelopment of c¢ along the crack faces.

A fundamental understanding of the influence ofvarious alloying elements on the nucleation andgrowth of iron nitrides, as provided by thisarticle, leads to a new approach for designingthe chemistry of nitriding steels and for thechoice of the nitriding parameters to control themicrostructure of the nitrided layers: Havingalloying elements of strong Me-N interaction(Me = Ti, V, and Cr) in steel together withalloying elements of weak Me-N interaction andnegligible solubility in iron nitrides (Me = Al, Moand Si), allows avoidance, or modification of themorphology and microstructure, of the iron-nitride-based compound layers. In particular,

Fig. 7. Backscattered electron (BSE) image showing the location of the EPMA elemental line scan (white dotted line) crossing the c¢ plates,together with the elemental concentrations for the (much longer than in Fig. 6) nitrided (550�C, rN = 0.82 atm�1/2, 20 h�) Fe-4.7 at.% Al alloy.The EPMA line scan was made parallel to the surface at a depth of about 40 lm. In the BSE image, c¢ appears dark due to the highnitrogen content (about 20 at.%). Peaks in the nitrogen content correspond with the EPMA line scan crossing a c¢ precipitate (plate). Higher nitrogencontent of the matrix than for the specimen shown in Fig. 6 is due to the development of AlN precipitates (note that the maximum nitrogen solubilityof ferrite matrix is about 0.4 at.%). The 8.5 at.% nitrogen in the matrix is much larger than the normally expected maximum nitrogen content for theFe-4.7 at.% Al alloy, which is the sum of the nitrogen incorporated in the AlN precipitates (4.7 at.%) and the equilibrium nitrogen solubility of theferrite (0.4 at.%). The surplus nitrogen represents the excess nitrogen (see the footnote in the Iron-Nitride Development in Case of InstantaneousMe-Nitride Precipitation section) as well as nanosized c¢ precipitates, which developed in the matrix.35

Fig. 8. Light optical micrographs recorded from the cross-sections of nitrided Fe-1 at.% Mo alloy. (a) Recrystallized specimen nitrided for 2 h at a(higher than in Fig. 5b) temperature of 550�C using an rN of 0.7 atm�1/2. (b) Cold-rolled specimen nitrided for 4 h at 520�C using an rN of0.7 atm�1/2. In both cases, a ‘‘normal’’ layer-type c¢ has developed.36

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obstructing the formation of the outer, brittle,iron-nitride compound layer can be beneficial inmany industrial applications.

ACKNOWLEDGEMENTS

The authors thank Mrs. S. Haug for EPMAmeasurements and Prof. Dr. P. van Aken for pro-viding access to transmission electron microscopyfacilities.

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Abnormal Nitride Morphologies upon Nitriding Iron-Based Substrates