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Acta Materialia 55 (2007) 3193–3201

Plasma nitridation of a novel Ni–10.8 wt.%Cr nanocomposite

J. Zhao, X. Peng *, Y. Wang, F. Wang

State Key Laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, 62 Wencui Road, Shenyang 110016, China

Received 16 November 2006; received in revised form 6 January 2007; accepted 9 January 2007Available online 2 March 2007

Abstract

A Ni–10.8Cr nanocomposite (by wt.%), consisting of nanocrystalline Ni matrix (mean grain size: 60 nm) and dispersed Cr nanopar-ticles (mean particle size: 42 nm), has been synthesized by nanocomposite electrodeposition. The unique structure causes the nanocom-posite to form a double-layered nitrided zone during plasma nitridation at 560 �C for 10 h. The outer layer (�50 lm thick) precipitatesnanometer-sized CrN (<100 nm), which increased in size but decreased in number with increasing nitridation depth (following Bohm–Kahlweit’s mode). The inner layer (�5 lm thick) exhibits larger-coarsened nitride precipitates (100–200 nm) which almost link together.The greatly enhanced nitriding kinetics in the nanocomposite compared to a compositionally similar but microstructurally differentNi–10Cr alloy (mean grain size: 30 lm) is mainly associated with the fact that the numerous grain boundaries dramatically increasethe nitrogen permeability, according to the treatment using a classical Wagner’s approach. The nanohardness profile in relation tothe microstructure of the nitrided zone in the nanocomposite has also been investigated.� 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Ni–Cr nanocomposite; Nanocrystalline; Electrodeposition; Plasma nitriding; Nanohardness

1. Introduction

Nickel is widely used as a base metal in industrial pro-cesses because of its good resistance to general corrosionand high-temperature oxidation. However, nickel-basealloys with a face-centered cubic (fcc) structure have, ingeneral, a low hardness, which limits their practical appli-cations to industrial fields where high load, erosion orabrasive wear are significant factors causing metal degrada-tion. Hence, it is essential to explore methods for increasingthe surface hardness of nickel-base alloys. Nitriding is oneof the conventional techniques used to improve the hard-ness of ferritic steels. The nitriding processes have beenreviewed elsewhere [1]. The nitrided steels are hardenedthrough solid-solution strengthening due to the introduc-tion of interstitial nitrogen atoms as a result of the highsolid solubility of nitrogen in steels, and in particularthrough hard-phase dispersion strengthening due to the

1359-6454/$30.00 � 2007 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2007.01.018

* Corresponding author. Tel.: +86 24 23893753; fax: +86 24 23893624.E-mail address: xpeng@imr.ac.cn (X. Peng).

precipitation of fine and coherent internal nitrides (suchas c 0-Fe4N and e-Fe3�2N) when the nitrogen solubilitylimit is reached [1,2]. Nickel cannot form hard nitride pre-cipitates due to a positive value of the free energy of forma-tion of its nitride [3]. One possible way to harden thenickel-base alloys is to alloy some active elements such aschromium, aluminum or vanadium, which can react withinward-diffusing nitrogen during nitriding to form internalnitrides with good thermal stability, high hardness, and lowfriction coefficient [3–7].

Because CrN has a high hardness, good thermal stabilityand corrosion resistance, plasma nitriding of nickel–chro-mium alloys has been attracting attention in recent years[5–7]. Plasma nitriding is an increasingly popular methodwhich uses a glow discharge technology to introduce ion-ized nitrogen into a metal surface [1]. Unfortunately,nitriding nickel-base alloys to a sufficient depth has notbeen achieved so far in laboratory studies [5–7]. The limitednitriding depth originates from the insurmountable, intrin-sic drawback of the conventional coarse-grained (CG)nickel-base alloys where the diffusion of nitrogen is extre-mely low at the temperatures normally used for plasma

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3194 J. Zhao et al. / Acta Materialia 55 (2007) 3193–3201

nitriding. To obtain a thick nitrided layer, it is crucial totailor an alloy structure which can greatly enhance thenitriding kinetics.

Nanocrystalline (NC) materials consist of a large vol-ume fraction of grain boundaries (GBs), which have areduced atomic density with reference to the perfect crystallattice and normally function as ‘‘short-circuit’’ diffusionpaths. Thus, metallic diffusion in nanograined polycrystal-line metals would be enhanced by orders of magnitude incomparison to the diffusion in their CG counterparts withthe same chemical compositions at a given temperature,even when GB diffusion is dominant in the latter materials[8–10]. It has been reported that the surface nanocrystalli-zation of iron using the so-called surface-mechanical-attri-tion treatment (SMAT) can enable a nitrided layer with adramatically increased thickness to be obtained [11]. There-fore, if a nickel-base alloy is nanostructured, it is promisingto obtain a nitrided layer with an increased thickness.However, the relative research on this field is still lacking.

Recently, novel Ni–Cr nanocomposites consisting ofNC Ni matrix and Cr nanoparticles have been fabricatedusing co-electrodeposition by Peng et al. [12–14]. The com-posite’s structure favors the diffusion of chromium fromthe bulk to the reaction front, leading to excellent resis-tances to high-temperature oxidation [12], hot corrosion[13] and electrochemical corrosion [14]. The Ni-base nano-composites may also have expected high nitriding kinetics.In the present work, we first report plasma nitridation ofsuch a Ni–Cr nanocomposite, by comparing the nitridationof a conventional CG Ni–Cr alloy (solid solution) with asimilar Cr content. The objective of this experiment is tounderstand the nitriding process of the nanocomposite,which will be helpful to fabricate a desired nitrided layerby further tailoring the composite’s microstructure.

2. Experimental

Pure nickel (99.9%) plates of size of 15 · 10 · 2 mm wereabraded to 800 grit SiC paper. After being ultrasonicallycleaned in acetone, the Ni plate was electrodeposited witha �100 lm-thick film of a Ni–Cr nanocomposite from abath containing 150 g l�1 NiSO4 Æ 7H2O, 12 g l�1 NaCl,35 g l�1 H3BO3, 120 g l�1 sodium citrate using appropriateamounts of pure Cr nanoparticles (>99%). The nanoparti-cles were prepared using a direct-current plasma process.They were gathered from the condensed downstream gasof Cr from the chamber where pure Cr was used as ananode and evaporated in an Ar atmosphere at the voltageof 80 V and the current of 900 A. During electrodeposition,a reciprocating perforated plate maintained the suspensionof the nanoparticles in the bath [15]. The current densitywas 3 A dm�2, the bath temperature was 35 �C and thepH 5.5–6.0. The composite’s microstructure was investi-gated by scanning electron microscopy (SEM) and trans-mission electron microscopy (TEM). The Cr content inthe nanocomposites was analyzed using energy-dispersiveX-ray analysis (EDAX). In this work, the nanocomposite

with the Cr content of �10.8 (by wt.%) was used. A Ni–Cr alloy with a similar Cr content (10 wt.%) was preparedfor comparison. The alloy specimens with similar dimen-sions were cut from the cold-rolled alloy buttons, whichwere prepared by electro-arc melting pure Ni (>99.9%)and Cr (>99.9%) in vacuum with magnetic stirring at tem-perature over 1600 �C and subsequently by 2 h annealing at900 �C. The arc-melted alloy was abraded and finally pol-ished using 1.5 lm diamond paste before nitridation. Thesurface treatment was not conducted on the nanocompos-ite-coated sample for nitriding.

The nitriding chamber was first evacuated and thenfilled with NH3 gas at a stable flow rate of 900 ml min�1.Plasma nitridation was conducted at 560 �C under the envi-ronmental pressure of 800 Pa for 10 h. X-ray diffraction(XRD) and electron probe microanalysis (EPMA) wereused to determine the phase compositions and elementaldistributions of the nitrided samples, respectively. The finestructure of the nitrided zone was further investigated usingTEM. The compositional analysis was performed using aFEI Tecnai F30 TEM equipped with nanobeam energy-dis-persive X-ray spectroscopy (EDS) and high-angle annulardetector dark-field scanning transmission electron micros-copy (HAADF-STEM). The accelerated voltage used was300 keV, and the beam diameter for nanobeam electrondiffraction (NBED) was about 1.5 nm. The energy-filteredimaging was carried out using Gatan image filter in a Tec-nai F30.

The nanohardness depth profiles of the nitrided layer onboth samples were acquired using an instrumented nano-indenter (Hysitron TriboIndenter, USA) by a Berkovich-type diamond tip, with an applied load of up to 9 mN.Besides, XRD was used to characterize the mean micro-strain of the nitrided samples. The XRD line broadeningdue to the instrumental effects was calibrated using theannealed micron-sized SiO2.

3. Results

3.1. Morphology and microstructure of as-deposited Ni–Cr

nanocomposite

Fig. 1(a) shows the bright-field TEM image of the as-received Cr nanoparticles. Most of the nanoparticles arespherical but some appeared to be cubic, while their sizeranges from 10 to 180 nm, with a mean value of42 ± 35 nm. After co-deposition, the Cr nanoparticles arerandomly distributed over the electrodeposited Ni matrix,as seen in a bright-field TEM image of the as-depositednanocomposite presented in Fig. 1(b). The Ni matrix hasa grain size ranging from 20 to 120 nm, with a mean valueof �60 nm; besides, the nanocomposite is compact, show-ing no pores or cracks between the nanoparticles andnickel. For comparison, the microstructure of the arc-melted Ni–10Cr alloy was also investigated after etchingusing SEM. Its mean grain size is around 30 lm, as seenin Fig. 1(c).

Fig. 1. (a) TEM image of Cr nanoparticles used; (b) TEM image of an as-deposited nanocomposite showing the distribution of Cr nanoparticles (someindicated by frames) over the electrodeposited nanocrystalline Ni matrix; (c) microstructure of the arc-melted Ni–10Cr alloy (etched).

J. Zhao et al. / Acta Materialia 55 (2007) 3193–3201 3195

Fig. 2(a) shows the SEM top-view of the as-depositedNi–10.8Cr nanocomposite. Its surface is generally flat.From the corresponding etched cross-section given inFig. 2(b) (where the pitting areas (dark) are enriched in

Fig. 2. (a) Surface morphology of the as-deposited Ni–10.8Cr nanocom-posite and (b) the corresponding cross-sectional morphology (etched).

chromium due to the faster etching of Cr than Ni in theetching solution used), the dispersion of Cr nanoparticlesin the electrodeposited Ni matrix is homogeneous.

3.2. Characterization of the nitrided zone in the

nanocomposite

After 10 h plasma nitridation at 560 �C, the polishedsurface of the Ni–10Cr alloy exhibits numerous nodules(Fig. 3(a)). They consist of pure Ni per se according toEDAX analysis, demonstrating that the nitrided alloy hasbeen covered by a pure Ni surface layer. This layer is gen-erally below 2 lm, except for a few locations where largerNi particles appear on a basis of cross-sectional EPMAinvestigation (see the arrow-indicated areas in the back-scattered electron image (BEI) and Ni X-ray mapping ofFig. 5 below). The surface Ni nodules are, as will be dis-cussed later, ‘‘extruded’’ from the bulk alloy during nitrid-ing, due to the nitride-precipitation-induced compressivestresses generated in the nitrided zone. The surface of theNi–10.8Cr nanocomposite also gets more roughened afternitriding (see Fig. 3(b) in reference to Fig. 2(a)). EDAXanalysis indicates that the surface Cr content is reducedfrom �10.8 wt.% to �7.5 wt.%, suggesting the occurrenceof a pure Ni layer on the surface. This Ni layer is thinnerthan that formed on the nitrided alloy according tocross-sectional investigation. Fig. 4(a) and (b) shows theXRD patterns of the nanocomposite before and afternitridation, respectively. As is evident, the Cr nanoparticlesin the probed volume have been completely converted into

Fig. 3. Surface morphologies of the Ni–10Cr alloy (a) and Ni–10.8Crnanocomposite; (b) after 10 h plasma nitridation at 560 �C for 10 h.

40 50 60 70 80

(c)CrN

(111

)

CrN

(222

)

(b)

Ni(1

11)

Ni(2

00)

Ni(2

20)

CrN

(222

)

CrN

(220

)

CrN

(111

)

Cr(

211)

Cr(

220)

Cr(

520)In

tens

ity (

a.u.

)

2 theta

Cr(

420)

(a)

Fig. 4. Comparison of the XRD patterns of the as-deposited (a) andnitrided Ni–10.8Cr nanocomposite (b), showing the conversion of Crnanoparticles into chromium nitrides after 10 h nitridation at 560 �C; (c)shows the XRD pattern of the nitrided Ni–10Cr alloy.

Fig. 5. BEI and elemental X-ray maps of cross-section of the nitrided Ni–10Cr alloy. The arrow-indicated areas correspond to larger pure Ninodules.

3196 J. Zhao et al. / Acta Materialia 55 (2007) 3193–3201

CrN during nitriding. CrN also occurs in the nitrided Ni–10Cr alloy (Fig. 4(c)). Part of CrN precipitates and the Nibase matrix for both samples have an orientational rela-tionship of Æ1 11æCrNiÆ111æNi. Additionally, the CrN pre-cipitates in the nitrided alloy exhibits a Æ222æ texture,while the nitride in the nanocomposite is randomly precip-itated along the directions of Æ111æ, Æ220æ and Æ22 2æ.

The thickness of the nitrided zone in the alloy is below5 lm, as seen by the EPMA analysis given in Fig. 5. Thefine structure of the nitrided zone was investigated using

HAADF-STEM. A Z-contrast image (parallel) of thenitrided zone in an area adjacent to the original alloy sur-face is presented on the left of Fig. 6(a), where numerousparticulates (dark areas) are visible. Most particulates arearound 10 nm in size. The right images in Fig. 6(a) showthe variation of the intensity of N, Cr and Ni along a lineacross a particulate indicated in the Z-contrast image.Apparently, the precipitates are enriched with Cr and N.Combining this and the XRD characterization, it isassumed that the nanometer-sized particulates are CrNprecipitates. No Cr has been detected in the other areaexcluding the particulates, according to the EDS analysisin the HAADF-STEM system. The particulates were fur-ther inspected using TEM at higher magnification. Theyare projected with Moire fringes (see Fig. 6(b)). The Moirefringes appear as a result of the superimposition of CrNprecipitates and Ni matrix. Similar observation has beenreported elsewhere [16].

Fig. 7 shows a BEI and the corresponding elemental X-ray mappings of a cross-section of the nitrided nanocom-posite. This suggests the two following observations. First,the nitrided layer is quite thick (�55 lm). Thus, comparedto the CG Ni–Cr alloy (Fig. 5), the nitriding kinetics hasbeen greatly enhanced in the nanocomposite. Second, thenitrided layer is composed of two layers. The fine structureof the outer layer, �50 lm thick, was characterized usingTEM. The TEM dark-field image of a region with thenitriding depth of �5 lm, shown in Fig. 8(a), indicates thatthe size of the CrN precipitates ranges from �10 to �70 nmwith a mean value of �30 nm (chromium was hardly

Fig. 6. (a) A Z-contrast image (left) of the near-surface zone in thenitrided Ni– 10Cr alloy and the compositional analysis (right) of N, Cr, Nialong a line across a particulate indicated in the Z-contrast image; (b) abright-field TEM image of the nitrided zone at a higher magnification,showing the nanometer-sized particulates reflected with Moire fringes.

Fig. 7. BE image and elemental X-ray maps of the cross-section of thenitrided Ni–10.8Cr nanocomposite.

J. Zhao et al. / Acta Materialia 55 (2007) 3193–3201 3197

detected by EDS in the spacings among the precipitates).The number of the internal nitride precipitates per unitarea, or the area density, is �280 particles lm�2. Assuminga simple cubic structure as the unit cell for the space distri-bution of the Cr nanoparticles in the nanocomposite, thenumber (N) of the 42-nm-sized Cr particles per lm2onthe {100} planes is calculated using the followingequation:

N ¼ 106=ð2r þ sÞ2 ð1Þwhere r is the radius of the Cr nanoparticles, and s is theinterparticle spacing and can be expressed by

s ¼ffiffiffiffiffiffiffiffiffiffi4p

3DV3

r� 2

!r ð2Þ

where DV is the volume fraction of Cr nanoparticles in thenanocomposite (�13%). Therefore, the calculated numberis �60 particles lm�2, which is much lower than that ofthe precipitates. This implies that the abundant GBs in

the nanograined nanocomposite, together with the nano-particles themselves, are the preferred sites for nitridenucleation. In contrast to the ‘‘immobile’’ nucleation sitesof CrN at the nanoparticles, the nucleation site at GB is‘‘movable’’ depending on the relative magnitude of nitro-gen and chromium concentration there. It means that, inthe nanocomposite underneath the internal nitridationfront, solid dissolution of Cr nanoparticles would occurand release Cr atoms, which outward diffuses mainly alongthe nickel GBs. The latter has been indirectly but crediblyconfirmed by the EDS analysis in the TEM system whichreveals �2.0 wt.% Cr in the Ni matrix in the middle ofthe outer layer. Fig. 8(b) shows the TEM bright-field imageof a region there. The size of the CrN precipitates becomeslarger, ranging from �70 nm to �100 nm.

In the inner layer the CrN precipitates are more coars-ened (generally in a range of 100–200 nm) and most ofthem link together, as seen in Fig. 9(a). Below this layerthe un-nitrided nanocomposite exists, where Cr has beenhardly probed using EDS. This result is consistent withFig. 7 where a Cr-depleted band appears underneath theinner nitrided layer. Intriguingly, EDS analysis of someprecipitates, as one example given in Fig. 9(b) with the cor-responding content (at.%) of 63.5Cr, 31.3N, 3.5Ni, 1.7Cu(containment from the TEM foil holder), suggests the con-version of partial CrN into Cr2N.

Here, it should be pointed out that the ‘‘isolated’’ nano-particles do not appear in the double-layered nitridationzone, indicating that they have been, if not nitrided, com-pletely dissolved.

Fig. 8. (a) TEM dark-field image and the corresponding diffraction pattern of the near-surface zone of the outer nitrided layer showing numerous nano-sized CrN precipitates. (b) TEM bright-field image and the corresponding selected area diffraction pattern of the middle zone of the outer nitrided layershowing the CrN precipitates (in brighter contrast) which increase in size but decrease in number.

3198 J. Zhao et al. / Acta Materialia 55 (2007) 3193–3201

3.3. Nanohardness measurement

Fig. 10 shows the nanohardness profiles of the nitridedsamples. The nitrided Ni–10.8Cr nanocomposite has a highhardness value of �7.5 GPa on the measured spot closestto the surface. Thereafter, the hardness gradually decreasesto �6.0 GPa on the spot furthest from the surface in theouter nitrided layer. In the inner ‘‘continuous’’ CrN layer,the maximum hardness of �12 GPa appears. In contrast,the nitrided Ni–10Cr alloy has the hardness of 7.1 GPa inthe near-surface zone, the value of which quickly reducesto the hardness of the alloy substrate (�4 GPa) becauseof the shallow nitriding depth (Fig. 2).

4. Discussion

4.1. Nitriding kinetics of the nanocomposite

The nitriding kinetics of alloys has been mostly treatedusing the classical Wagner’s ‘‘internal oxidation’’ model,as having been reported elsewhere [16,17]. In this case, itis generally assumed that the nitriding kinetics is domi-nated by the inward permeation of nitrogen (that is,DNN S

N � DCrN Cr, where N SN is the mole fraction of nitrogen

at the sample surface (equal to the lattice solubility ofnitrogen), DN and DCr are the diffusion coefficients ofnitrogen and chromium in the nitriding sample, and NCr

is the Cr mole fraction of the bulk metal, respectively)

and that the concentration of nitrogen decreases linearlyfrom the alloy surface to the internal nitridation front,and is zero in the bulk metal, the nitriding kinetics isexpressed by [18,19]

n ¼ 2N SNDN

NCr

t� �1=2

ð3Þ

where n represents the depth (thickness) of the internallynitrided zone. The values of NS

N, DN and DCr were not mea-sured in the present work, and the literature on the mea-surement of these parameters in Ni–Cr alloys andnanocomposites are scanty. If we assume that NS

N andDN (equal to the lattice diffusivity of nitrogen) in the Ni–10Cr alloy are close to those measured in the Ni–2V alloyin Ref. [20], where N S

N ¼ 1:5� 10�6 exp½7457=T � (molefraction) and DN = 2.01exp[ � 20,688/T] (cm2 s�1) are gi-ven respectively, the nitriding depth of the alloy can be ex-pected using Eq. (3) for 10 h nitridation at 833 K (560 �C).Its value is �5.2 lm. It is very close to the nitriding depthmeasured, indicating that the nitriding kinetics of the Ni–10Cr alloy follows a parabolic law.

The nitrogen diffusion should be greatly enhanced in thenanocomposite containing numerous GBs. In this case, aneffective diffusivity Deff

N is proposed, which is a summationof lattice (Dl

NÞ and GB diffusivity ðDgbN Þ of Cr as given

below:

DeffN ¼ ð1� f ÞDl

N þ fDgbN ð4Þ

0 20 40 60 80 1002

4

6

8

10

12

Har

dnes

s (G

Pa)

depth (um)

Fig. 10. Comparison of the nanohardness profiles of the nitrided layers inthe nanocomposite and the alloy counterpart.

Fig. 9. TEM bright-field image of the inner nitrided layer (parallel to thenitridation direction) showing the precipitation of larger-coarsened nitride(in brighter contrast). Part of nitrided zone (on the upper right) spalledduring TEM sample preparation. The arrow-indicated area corresponds tothe nanocomposite beneath this nitrided layer.

J. Zhao et al. / Acta Materialia 55 (2007) 3193–3201 3199

where f is the GB area fraction and equals 2d/d (d is the GBwidth and d the grain size) if considering the electrodepos-ited Ni grains as cubic. The active energy (Qgb) for an atomdiffusion along GBs is normally 0.5–0.7 times smaller than

the active energy (Ql) for the atom lattice diffusion [21].Supposing Qgb= 0.6Ql, and using the aforementionednitrogen diffusivity [19], Dgb

N � 6:8� 10�7 cm2 s�1 at560 �C. Thus, Deff

N can be quantitatively evaluated usingEq. (4). By combining Eqs. (3) and (4), the calculatednitridation depth at 560 �C for 10 h is �97 lm, providedthat d = 0.5 nm and the coarsening of NC Ni grains isnot significantly (d = 60 nm). The calculated data is muchhigher than the practical nitriding depth of �55 lm(Fig. 7). The reason primarily lies in two aspects. First,the inevitable coarsening and growth of the electrodepos-ited NC Ni grains at the exposing temperature, as havingbeen reported elsewhere [22,23], would cause a gradualbut noticeable reduction of Deff

N . The TEM characterization(Figs. 8(b) and 9(a)) also indicates that after 10 h nitridingthe Ni grains are increased in size from �60 nm to gener-ally <200 nm. Second, a partial or complete dissolutionof Cr nanoparticles into the nickel matrix would occur be-fore the internal nitridation front reaches, which outward-diffuse along the nickel GBs and consequently, minimizethe inward permeability of nitrogen. Therefore, the treat-ment on the nitriding kinetics of the nanocomposite shouldbe more complicated because of its unique structure.

4.2. Microstructure of the nitrided zone in the nanocomposite

Similar to internal oxidation, internal nitridationinvolves a competition between the nucleation of freshinternal precipitates and the growth of those alreadyformed. According to the classical Bohm–Kahlweit model[24], nitride nucleation depends on the relative magnitudeof the concentrations of nitrogen and the active metal atthe internal nitridation front. This magnitude easilyreaches the critical value for nitride nucleation at manylocations when nitridation starts, because of a relativelyhigh flux of nitrogen. The nitride nucleation dominatesthe internal nitridation process, accompanying the precipi-tation of numerous nitrides. When the nitriding depthincreases, the reaction mechanism progressively shifts fromthe nitride-nucleation-dominated process into one domi-nated by nitride growth. However, the low solubility ofnitrogen in the arc-melted alloy limits the build-up of a sig-nificant nitrogen concentration gradient (driving force fornitrogen diffusion), causing the formation of the shallownitrided layer (Fig. 5). Compared to this, the nitridingkinetics is greatly enhanced in the Ni–10.8Cr nanocompos-ite, leading to the formation of the thick nitrided layer(Fig. 7).

The microstructural characterization of the outernitrided layer in the nanocomposite reveals that the near-surface zone precipitates fine and nano-sized CrN particles(Fig. 8(a)), while the middle zone produces larger nitrideparticles (Fig. 8(b)). It suggests that this nitrided layerforms following the Bohm–Kahlweit mode, which isaddressed below.

At the onset of nitridation, the Cr nanoparticles on andadjacent to the surface are directly in situ transformed into

3200 J. Zhao et al. / Acta Materialia 55 (2007) 3193–3201

CrN at the arrival of nitrogen. That is, the nanoparticlesact as ‘‘active sites’’ for the nitride nucleation. Simulta-neously, the dissolved Cr from the bulk nanocompositeoutward-diffuses along the composite’s GBs, and is precip-itated as nitride there when the magnitude of the local con-centrations of chromium and nitrogen reaches the criticalvalue for the nitride nucleation. With further advancementof the internal reaction front, the nitrogen flux, togetherwith the concentration of Cr in the un-nitrided zone ofthe nanocomposite, decreases. This makes new CrN nucle-ation difficult. At this point, the growth of the precipitatednitrides becomes the dominant process, causing the forma-tion of a smaller number of larger internal precipitates(Fig. 8(b)).

Below the outer layer where the internal nitride was dis-continuously precipitated, an approximately continuousCrN layer occurred (Fig. 7). This result is not in agreementwith Bohm–Kahlweit’s model. This occurs undoubtedly asa consequence of the complex structure and of its evolutionwith time. It seems possible that at a critical reaction timethe amount of chromium present at the front of internalnitridation in the form of Cr particles plus that reachingthe same location by diffusion from the bulk may exceedthe corresponding amount of N, slowing down the rateof displacement of the front and increasing the local vol-ume fraction of nitride above the critical value requiredfor the transition from internal to ‘‘external’’ growth ofCrN [18,19]. After the establishment of the continuousnitride layer, continuous advancement of nitrogen becomesdifficult (see the N X-ray map in Fig. 7). Accordingly, theconstant supply of Cr to the reaction front from the bulknanocomposite may convert partial CrN precipitates bythe reaction CrN + Cr! Cr2N.

The CrN precipitation causes a considerable volumeincrease in the internally nitrided zone, resulting in a com-pressive stress gradient between the internal reaction frontand the stress-free sample surface. The stresses may berelieved to some extent by the outward diffusion of Nithrough several mechanisms, such as grain boundary slid-ing [25], Nabarro–Herring (lattice diffusion) creep [26]and dislocation-pipe-diffusion [16,27–29]. From the surfacefeature of the nitrided Ni–10Cr alloy shown in Fig. 3(a),numerous Ni nodules outgrow during nitriding. These nod-ules do not like the typical surface feature formed in rela-tion to grain boundary sliding [25]. The experimentallymeasured volume diffusivity of Ni through the N–H creepis also too slow to be responsible for the formation of sucha thick surface layer of pure Ni [16,29]. Its appearance isplausibly associated with the outward diffusion of nickelthrough dislocation-pipe-diffusion-controlled creep. Basedon the average size and the density per area of the CrN pre-cipitates given in Fig. 6, and the theoretically anticipatednumber of interface misfit between CrN and the alloymatrix (presumptively along the Æ111æ directions), the dis-location density can be roughly evaluated, the value ofwhich is in a magnitude of �1016 m�2. Such a high disloca-tion density should cause a sufficient Ni diffusion and con-

sequently, the formation of the surface Ni nodules. GBdiffusion of Ni (Coble creep [30]) and grain boundary slid-ing are proposed to operate for relieving the stress gener-ated in the nanocomposite, because the Ni grains areultra-fine (generally below 200 nm even though the graincoarsening is considered during nitriding). This may bethe reason that the pure Ni surface layer is thinner on thenitrided nanocomposite than on the nitrided alloy.

4.3. Hardening mechanism

Numerous nano-sized internal CrN are precipitated inthe near-surface zone (see Figs. 6 and 8(a)). The nitride pre-cipitates usually has a semi-coherent interface with thematrix [7,16]. If we assume the precipitation of Cr in nickelalong the Æ1 11æ directions for the two materials (see Fig. 4),the mismatch between the CrN and the Ni-base matrix canbe calculated by

Dd ¼ 2dCrN � dNi

dCrN þ dNi

ð5Þ

where d is the lattice parameters for the nitride and the ma-trix obtained from the XRD line analysis in this work. Thevalue is �16.1% for the CrN formation in the alloy and�15.8% in the nanocomposite. Such a high misfit betweenthe precipitates and the matrix can induce a considerablelattice microstrain in the nitrided zone. From the acquiredXRD lines, the microstrain (e) could be calculated using theJade 5.0 software, based on the following approximationwhich combines the Scherrer formula and the Wilson for-mula [31,32]:

w ¼ 2e tan hþ 0:9kd cos h

ð6Þ

where w is the width at half maximum intensity of XRDlines, h the Bragg angle, k the wavelength of the X-ray radi-ation and d the average crystallite size. The calculatedmicrostrain is 0.28 ± 0.03% in the near-surface zone ofthe nitrided nanocomposite. For evaluating the microstrainin the nitrided alloy, the second item in Eq. (6) (the Scher-rer formula) is negligible due to its large grain size(�1000 nm). The microstrain generated in the alloy is0.24 ± 0.12%. The calculated microstrain coincides wellwith the measured value of the nanohardness in the near-surface zone of the nitrided nanocomposite (�7.5 GPa)and its alloy counterpart (�7.1 GPa). The nitrided nano-composite has a slightly higher surface hardness, whichmay result from the synergistic effect of two contributionswhich are expected during the composite nitridation: oneis the numerous sites available for CrN nucleation leadingto the precipitation of internal nitrides with higher molarvolume fraction, the other is the ultra-fine-grainedstructure leading to the microstrain relaxation differentfrom the route of that in the nitriding alloy, as addressedabove.

With the advancement of the nitridation front in thenanocomposite, the nitride growth becomes a dominant

J. Zhao et al. / Acta Materialia 55 (2007) 3193–3201 3201

process (Fig. 8(b)). This reduces the area of the semi-coherent precipitate/matrix interface (the Gibbs energy ofthe interface), and thus the measured nanohardness of theouter nitrided layer (Fig. 10). However, the size of thenitride precipitates in the nanocomposite is generally inthe nanorange. That is, the internal nitrides are ratherclosed-spaced in dispersion as compared to those in con-ventional coarse-grained Ni-base alloys [4–7,16,29]. Hence,the hardness profile throughout the outer nitrided layer inthe nanocomposite is not as steep as that measured fromthe nitrided Ni–Cr alloys [4].

5. Conclusions and implications

During plasma nitridation at 560 �C for 10 h, the Ni–10.8Cr nanocomposite forms a nitrided zone with a dou-ble-layered structure. The outer layer with a thickness of�50 lm precipitates numerous nanometer-sized CrN,which increase in size but decrease in number with increas-ing nitriding depth. This microstructure clearly indicatesthat the outer nitrided layer forms following the classicalBohm–Kahlweit’s mode, i.e., the nitriding process shiftsfrom the dominant nucleation of fresh internal nitride tothe dominant growth of the precipitated nitride with theadvancement of the internal reaction front. This nitridedlayer having the typical structure yields a nanohardnessprofile with a decrease gradient from �7.5 GPa in thenear-surface area to �6.0 GPa in the innermost area. Theinner nitrided layer with a thickness of �5 lm has the max-imum hardness of �12.0 GPa, because the layer is almostcomposed of larger-coarsened nitride precipitates. Thislayer occurs, it is proposed, as a result of continuousgrowth of the internal precipitates there in relation to theCr content increase due to the constant Cr supply fromthe dissolved nanoparticles in the bulk nanocompositeand simultaneously the nitrogen permeability reductionwhich makes it impossible the further advancement of thenitridation front.

The nanocomposite consisting of NC Ni matrix and dis-persed Cr nanoparticles exhibits dramatically enhancednitriding kinetics, as compared to the arc-melted CG Ni–10Cr alloy which only forms a <5 lm-thick internallynitrided zone in the same condition. Besides, the nanocom-posite has abundant sites for the nitride precipitation andas a consequence precipitates numerous ultra-fine-grainedCrN (<200 nm) throughout the entire nitrided zone. Thisis rather different from the nitridation of conventionalNi–Cr alloy where the nitride precipitates are microme-ter-sized [4]. Thus, the formation of the fine-nitride-precip-itated and double-layered nitrided zone depends on thecomposition and structure of the nanocomposite. On thisbasis it is anticipated that the property of the nitridednanocomposite might be further adjusted by tailoring thesize of Ni grains and/or particles, and/or by varying thenanoparticle content.

The result offers a novel approach for achieving anitrided coating with a preferred property on conventionalNi-base alloys through microstructure tailoring, and pavesa new way for fabricating a desired nitrided layer on themetals which are not ideally used for nitridation hardening,through prior coating them with the nanocomposites byelectrodeposition.

Acknowledgements

We acknowledge financial support from the NationalScience Foundation of China (NSFC) project (Grant No.50674082), and the Chinese Academy of Sciences (CAS)through ‘‘One Hundred Talents Plan’’ program. XP alsothanks Professor F. Gesmumdo at DICheP, Universita diGenova in Italy, for his comments.

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