Surface & Coatings Technology 207 (2012) 72–78

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Surface & Coatings Technology

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Nanosized precipitates in H13 tool steel low temperature plasma nitriding

L.F. Zagonel a,b,⁎, J. Bettini a, R.L.O. Basso b, P. Paredez b, H. Pinto c, C.M. Lepienski d, F. Alvarez b

a Laboratório Nacional de Nanotecnologia, Centro Nacional de Pesquisa em Energia e Materiais, 13083-970, Campinas, Brazilb Instituto de Física “Gleb Wataghin”, Universidade Estadual de Campinas, Unicamp, P.O. Box 6165 Campinas, SP, 13083-970, Brazilc Escola de Engenharia de São Carlos, Universidade de São Paulo, 13566-590, São Carlos/SP, Brazild Departamento de Física, Universidade Federal do Paraná, Brazil

⁎ Corresponding author at: Laboratório Nacional de Nde Pesquisa em Energia e Materiais, 13083-970, Campi3579.

E-mail address: luiz.zagonel@lnnano.org.br (L.F. Zag

0257-8972/$ – see front matter © 2012 Elsevier B.V. Alldoi:10.1016/j.surfcoat.2012.05.081

a b s t r a c t

a r t i c l e i n f o

Article history:Received 10 October 2011Accepted in revised form 22 May 2012Available online 30 May 2012

Keywords:Plasma nitridingNanosized precipitatesNitrogen diffusionHR-TEMTool steelSurface hardness

A comprehensive study of pulsed nitriding in AISI H13 tool steel at low temperature (400 °C) is reported forseveral durations. X-ray diffraction results reveal that a nitrogen enriched compound (ε-Fe2-3N, iron nitride)builds up on the surface within the first process hour despite the low process temperature. Beneath the sur-face, X-ray Wavelength Dispersive Spectroscopy (WDS) in a Scanning Electron Microscope (SEM) indicatesrelatively higher nitrogen concentrations (up to 12 at.%) within the diffusion layer while microscopic nitridesare not formed and existing carbides are not dissolved. Moreover, in the diffusion layer, nitrogen is found tobe dispersed in the matrix and forming nanosized precipitates. The small coherent precipitates are observedby High-Resolution Transmission Electron Microscopy (HR-TEM) while the presence of nitrogen is confirmedby electron energy loss spectroscopy (EELS). Hardness tests show that the material hardness increaseslinearly with the nitrogen concentration, reaching up to 14.5 GPa in the surface while the Young Modulusremains essentially unaffected. Indeed, the original steel microstructure is well preserved even in thenitrogen diffusion layer. Nitrogen profiles show a case depth of about ~43 μm after nine hours of nitridingprocess. These results indicate that pulsed plasma nitriding is highly efficient even at such low temperaturesand that at this process temperature it is possible to form thick and hard nitrided layers with satisfactorymechanical properties. This process can be particularly interesting to enhance the surface hardness of toolsteels without exposing the workpiece to high temperatures and altering its bulk microstructure.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

The surface hardness, corrosion resistance, and wear resistance ofsteels can be improved by nitriding processes [1–5]. These thermo-chemical surface treatments promote the inward diffusion of nitro-gen and its subsequent reaction with the alloying elements [6–8].Depending on the nitriding temperature and nitrogen chemicalpotential at the material surface, the process may cause the formationof a continuous iron nitride surface layer. Underneath this layer, adiffusion zone will form containing nitrogen in solution and, in caseof strong chemical affinity of nitrogen with the alloying elements,nitride compounds can be formed. In order to improve as much aspossible the work piece surface properties a better understandingon the growing mechanism of nitrided surface layers is stronglyjustified.

Industrial nitriding processes are usually carried out by gaseous,plasma-assisted, or salt bath techniques. Among these, plasma nitrid-ing is particularly attractive because of its relative high efficiency

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even at low temperatures. This is relevant to materials, such as stain-less steels or tool steels, for which this condition is mandatory inorder to minimize microstructural alterations and, consequently,maintain the corrosion resistance in the bulk material [9–12].Moreovera review of relevant literature shows that a precise controlof the surface characteristics and composition during plasma nitrid-ing is a challenge, owing to the complexity of this non-equilibriumprocess. The proper setting of the process parameters (gas composi-tion, gas pressure, ion current density, voltage, and temperature)and surface state (roughness, crystalline orientation, oxidation level,etc.) are therefore indispensable to assure reproducibility of the pro-cess results. Furthermore, the bombardment of the work piece by ion-ized hydrogen and nitrogen species present in the plasma allows foran additional surface cleaning by physical and chemical sputtering(“etching”) of the oxide layers usually present on the treated workpieces, thus adding one more variable to this complex process [13,14].

As noted above, underneath the outmost nitride layer, nitrogendiffuses and (may) react(s) with alloying elements. These phenome-na have been the focus of numerous studies attempting to point outkey parameters affecting the whole process and the final materialproperties [15,16]. Recent studies reported by Mittemeijer andcoworkers have shown the importance of the so called “excess”nitrogen on the diffusion process [17–20]. Such increased nitrogen

73L.F. Zagonel et al. / Surface & Coatings Technology 207 (2012) 72–78

concentrations represent the atomic fraction surpassing the quantitythat should be expected in case that: (i) the alloying elements arefully consumed to form nitrides and (ii) the nitrogen equilibrium sol-ubility in an unstrained iron matrix is achieved. The nitriding of AISIH13 hot work tool steel has also been intensively studied leading toenhanced surface hardness, thermal fatigue, and corrosion resistance[6,7,21–24]. In recent works, the mechanism of N diffusion at bothlow temperatures and short nitriding times (up to 5 hours) werereported by our group. Those studies were focused on the formationof the compound layer and of the diffusion zone as well as on theeffect of the nitriding gas mixture (N2 to H2 ratio), nitriding temper-ature, and ion current density [25–28]. In those studies, we showthat for the H13 steel, at 500 °C, considerable decarburizing takesplace at the surface including the dissolution of chromium carbides[26]. Even if no bulk hardness loss could be noticed after nitriding atthis temperature, the observed microstructure changes may indicatethat nitriding could have negative side effects at this temperaturefor this steel, particularly if the nitriding processes need to be repeat-ed several times during the tool lifetime, as it is the case for some cut-ting tools.

In this paper we deal with a similar issue, i.e. the nitrogen diffu-sion mechanism at low temperature (400 °C), but considering longerprocessing times (up to 36 hours). Since temperature does not affectsignificantly the plasma structure itself, this temperature regimemaintains the effectiveness of plasma nitriding but modifies the ki-netics of the processes inside the material, such as nitrogen diffusionas well as formation and dissolution of precipitates formed with thealloying elements. For instance, at 400 °C, carbide dissolution in thediffusion zone is prevented [27,28]. Moreover, this temperature re-gime could be particularly interesting to avoid embrittlement, distor-tions or corrosion susceptibility which may result from excessivenitride precipitation at the higher temperatures normally used inthe gaseous nitriding process [29,30]. The aim the present work isto determine the properties of thick (~50 μm) nitrided layers pre-pared at relatively low temperatures (~400 °C) in the attempt to un-derstand how precipitation, nitrogen diffusion, phase transformationand hardening takes place in this process regime. In this study, weshow the effectiveness of low temperature plasma nitriding appliedto H13 hot work tool steel with sufficient processing time to reachcase depths with technological relevance.

2. Experimental procedures

The nitriding experiments were carried out using a commercialhot wall furnace plasma nitriding system operated with a pulsed cur-rent power supply working in a 50 μs pulse on and 150 μs pulse off[31]. A gaseous mixture of 20% N2+80%H2 was fed into the furnacecontrolled by mass flowmeter controllers. A 400 Pa working pressurewas used along the processing. As a compromise between plasma sta-bility and current, the negative pulse was set to−380 V. The temper-ature was controlled by three independent heating elements and theworking temperature was fixed at 400±1 °C. The studied sampleswere nitrided under identical conditions but different treatmenttimes, namely 1, 9, 16 and 36 hours.

The test samples were cut out from the interior of a singlequenched and tempered AISI H13 steel block with bulk hardness of7.5 GPa (quenched in air from 1025 °C and tempered at 580 °C for2 h). The chemical composition of this steel is C, 2.5%; Mn, 0.4%; Si,2.1%; Cr, 5.9%; Mo, 0.8%; V, 1%; Fe, balance (measured by inductive-coupled plasma-optic emission spectroscopy). The sample surfaceswere mirror polished (final step was using 1 μm diamond paste)and cleaned in an acetone ultrasonic bath before loaded into the plas-ma nitriding furnace. The final surface roughness was about 40 nm(Ra), as measured by profilometry.

The phase composition of the nitrided layers was analyzed byX-ray diffraction (XRD) using monochromatic Cu Kα radiation and

Bragg-Brentano (BB) θ-2θ geometry. The outmost surface zone wascharacterized using small angle XRD (referred to as SA geometry) at3º incidence angle (effective penetration depth of ~0.1 μm) [32].

X-ray photoemission electron spectroscopy (XPS) was employedto gather information on the chemical state of nitrogen and alloyingelements on the most outer material surface. The underneath layerswere also scrutinized by XPS in samples previously ex situ grindedand polished (last step with 0.25 μm diamond paste) and in situcleaned with Ar+ ion beam bombardment procedure. The XPS systemis a VG-CLAM2 electron energy analyzer with a non-monochromaticAl Kα radiation source having a total resolution 0.85 eV (X-ray linewidth plus analyzer).

The composition depth profiles were analyzed on mirror polishedsamples cross-sections by electron probe microanalysis (EPMA)employing a Cameca SX100 apparatus operated at 15 kV using theWavelength Dispersive Spectrometer (WDS). For cross-section prep-aration, the samples received an electrodeposited nickel protectivelayer after being cut, and were grinded and polished (last step with0.25 μm diamond paste). The chemical composition was determinedfrom the emitted characteristic X-ray intensities by their comparisonwith the corresponding intensities recorded from appropriated stan-dards such as pure metals and compounds (Fe3C, Fe4N). The final el-ement content was calculated applying the Φ(θz) method [33]. Theestimated error from the EPMA measurements is lower than 0.5 at% for all reported results. The sample cross-sections were etched toreveal the material microstructure and observed by scanning electronmicroscopy (SEM) using a FEI Inspect F50. The etchant was a 5% Nitalsolution containing 0.1% concentrated hydrochloric acid. Transmis-sion Electron Microscopy (TEM) analysis were performed in a JEOL2100 F microscope. Suitable thin samples were prepared applyingthe standard procedure, i.e., by gridding, polishing and dimplingdown to 20 μm. Electron transparent areas were then obtained bylow energy ion polishing.

Hardness tests were performed using a computer controllednanoindenter equipped with a Berkovich diamond tip (Nanotech).The Oliver and Pharr method was used to calculate the hardnessvalues from loading/unloading curves [34]. The precise position ofeach indentation was determined a posteriori by SEM. The maximumload was set to 400 mNwith a dwell time of 15 s. Measurements wereperformed on polished cross-sections and on the surface of the sam-ples. No piling up effect was considered in the analysis.

3. Results and Discussion

Fig. 1 shows the XRD diffractograms obtained in the BB geometryfor all the employed nitriding times. For comparison purposes, the re-sult from SA XRD for the 1 h and 36 hour(s) nitrided samples are alsodisplayed. The diffractogram recorded after 1 h nitriding using the BBgeometry reveals only the martensitic matrix phase, i.e., α-Fe. Simi-larly, the 9 hours nitrided sample shows a line pattern compatiblewith the matrix and, also, a tiny diffraction line at about 38° associat-ed with the (110) planes of the ε-Fe2-3N phase [35–37]. The relativevolume fraction of the ε-Fe2-3N phase increases significantly onlyafter nitriding for 16 and 36 hs. On the other hand, SA XRD revealsthat the ε-Fe2-3N phase is already formed during the first hour. Nev-ertheless, it exists only within the outmost sample surface and it isnot observed when using the BB geometry. These results are in agree-ment with earlier observations showing that, upon plasma nitriding,the low nitrogen diffusion coefficient at 400 °C favors the formationof nitrogen-rich iron nitrides, such as the ε-Fe2-3N phase. Indeed,the reactivity of the nitriding plasma (i.e., its efficiency) does not de-pend directly on the processes temperature [26,38]. Thus, owing tothe high reactivity of the ion plasma process, the nitrogen surfaceconcentration reaches a steady state stoichiometry characterized bythe ε-Fe2-3N phase, during the first one hour of nitriding [27,29,39].

Fig. 1. BB geometry X-ray diffractograms obtained in samples nitrided by 1, 9, 16 and36 hours. SA results are shown for the sample nitrided for 1 and 36 hours. The arrowindicates the peak for (200) α-Fe or the (300) ε-Fe2-3N for a reference of the most rel-evant phase on the surface.

Table 1Full width at half maximum (FWHM) of the α-Fe diffraction lines.

hkl h²+k²+l² Untreatedmaterial

After 1 hourtreatment

After 9 hourtreatment

110 2 0.473 0.508 1.085200 4 0.903 0.901 2.653211 6 0.791 0.828 2.376220 8 1.036 1.07 1.225310 10 1.884 1.907 2.367

74 L.F. Zagonel et al. / Surface & Coatings Technology 207 (2012) 72–78

Moreover, nitriding for 1 and 9 hours leads the diffraction lines as-sociated to the α-Fe matrix to be shifted towards lower 2θ-angles, asshown in Fig. 1. For instance, the (110) reflection shifts from 44.87±0.05° in the untreated sample to 44.61±0.05° after 1 h nitriding,whereas the (200) diffraction line shifts from 65.08±0.05° to 64.83±0.05° in the same cases. This shift is due to an increased lattice parame-ter of the α-Fe phase after nitriding. This increase in lattice parameterhas two causes [40]: i) the nitrogen dissolved in the octahedral intersti-tial sites of the α-Fe lattice and adsorbed at nitride/matrix interfaces;ii) the positive volume misfit between the coherent nitrides and themartensitic matrix, which provokes a dilation of the matrix latticeand, consequently, an overall hydrostatic tensile stress in the α-Fematrix. We note, however, that coherent nitride precipitates with fewatomic layers thick are not detectable as separate diffraction lines byXRDmeasurements, since they diffract coherentlywith the surroundingmartensitic matrix [41–44]. Thus, they only contribute to the broaden-ing of the α-Fe or ε-Fe2-3N diffraction lines.

As observed in the diffractograms, all the diffraction lines associat-ed with the H13 matrix also become significantly broadened for lon-ger nitriding times (see for instance the 220 reflection for the samplenitrided during 9 hours, Fig. 1). Nitrogen dissolved in the martensiticsolid solution represents a first contribution to increasing line broad-ening. However, broadening purely caused by nitrogen dissolutionwithin the α-Fe unit cell would cause a monotonous line widthincrease with the reflection order, i.e. h2+k2+l2, without any de-pendence on particular crystallographic directions [45]. Nevertheless,the diffractograms show that the broadening of the α-Fe reflections isanisotropic with β(200) > β(211) > β(110), where β is the diffractionline breadth (see Table 1). The same trend for the α-Fe line widthswas encountered in [27]. The abnormal increase of line width associ-ated with the (200) α-Fe planes reflections suggests once again that

nanosized coherent particles have precipitated within the steel ma-trix. As discussed in the literature coherent plate-like nitrides developat the initial nitriding stages in Fe-Cr-V alloys [24,41,42,46,47]. Thesenanosized platelets (cubic CrN, VN) growth follows the Bain orienta-tion relationship with respect to the α-Fe matrix [47,48]:

100ð Þ Cr;Vð ÞN== 100ð Þα−Fe; 110½ � Cr;Vð ÞN== 001½ �α−Fe:

Because of this growth orientation of the coherent precipitates,line broadening is more significant for the (h00)-type planes of thesteel matrix, which are directly connected to the lattice of thenanosized precipitates and therefore mostly distorted [24].

The microstructure of the sample nitrided by 36 hours is revealedby SEM (Fig. 2(A) and (B)). The cross section image displays the tem-pered martensitic microstructure of the H13 steel, where some glob-ular precipitates, identified as vanadium carbides by EDX analyses,are embedded (Fig. 2(B)). The complexity of the microstructurewith small grain size is also observed. Significant microstructural dif-ferences are not apparent by comparing the diffusion zone and theoriginal matrix (Fig. 2(A)). Similarly, the bulk region of this samplehas the same microsctructure as untreated samples (not shown). Asshown later on, the nitrogen concentration varies between ~16 at.%to ~3 at.% moving from the top to the bottom of the image withoutnoticeable microstructure changes (Fig. 4). In contrast to previousreports also involving low temperature (370 °C and 400 °C) H13steel nitriding experiments, increasing nitrogen concentration doesnot seem to occur along the grain boundaries of the diffusion zone[26,27]. This effect has also been observed in both ion- and gas nitrid-ed materials [49,50]. In the case of strong preferential nitrogen diffu-sion through grain boundaries, it is observed a peak in the nitrogenprofile when a grain boundary is crossed [51]. Also, these regionsappear dark on SEM-BSE images [26,28].

Moreover, the H13 tool steel investigated in the present paper ex-hibits a matrix free of chromium precipitates, i.e. most of the chromi-um content appears to remain dissolved within a α-Fe solid solution.This is in contrast to previous observations that the tempering ap-plied to the H13 steel after quenching leads to an extensive precipita-tion of chromium-rich particles in the initial state prior to nitriding[5,13,28]. These observations suggest that the presence of chromiumpreferentially in the solid solution could be linked to the enhancednitrogen bulk diffusion. This agrees with the predictions made byKrawitz [52] and Gouné [53] that the thermodynamic interactionbetween nitrogen and the alloying elements determines the extentof nitrogen up take in some iron-based materials. Thus, chromium,having high affinity with nitrogen, seems to induce nitrogen diffusionthrough the bulk material, increasing nitrogen trapping probability(i.e. adsorption) at dislocations and interfaces present in the tem-pered martensitic microstructure. This effect is also observed in pow-der metallurgy alloys containing vanadium [54].

To complete these observations, TEM analysis (including ElectronEnergy Loss Spectroscopy, EELS) has been carried out, allowing theobservation of fine nanometric precipitates [41,48,55]. Fig. 3 showsa bright field TEM images from the sample nitride during 36 hoursjust a fewmicrons below the sample surface. Although chemical map-ping was not possible due to the magnetic character of the sample, it

Fig. 2. SEM micrograph (secondary electrons) of the sample nitrided by 36 hours.(A) Low magnification image showing the microstructure on the diffusion layer. Thesample surface is on the top of the figure and discontinuous layer on top is a protectivenickel coating for sample preparation purposes. (B) Detail on vanadium carbideprecipitates.

Fig. 3. TEM images from the sample nitrided for 36 hours. (A) Bright field imageshowing a big feature on the top-left corner identified by EELS as a vanadium carbide.(B) Dark-gray areas are identified as nanoprecipitates of chromium nitride. (C) HR-TEMshowing that nanosized precipitates can be coherent with the steel matrix.

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was possible to perform EELS analysis in selected areas by focusingthe electron beam to a small spot (~5 nm). Therefore, as revealedby EELS, the feature observed on the top-left of the Fig. 3(A) is com-patible with vanadium carbide. The dark regions with fine structuresare identified as fine dispersed chromium nitride precipitates, withlow carbon content. Fig. 3(B) shows a closer look on this area. Dark-gray areas show mainly nitrogen, chromium and iron (no quantifica-tion was attempted since the exact probed volume could not be de-termined accurately). Further EELS analysis indicates that light-grayareas in Fig. 3(B) are similar in composition to dark-gray areas,indicating the presence of both chromiumand nitrogen in the steelma-trix possibly as very fineprecipitates. Indeed, by a careful observation ofFig. 3(B), it is possible noting very fine features, about 1–2 nm in width,which could be ascribed to chromium nano precipitates. The coherentnature of some precipitates could be confirmed by HR-TEM images.As an example, Fig. 3(C) shows atomic columns on one bigger chromi-um nitride nanosize precipitate while in its vicinities, smaller precipi-tates are observed with faint contrast over the steel matrix. Thisprecipitate has a projected area of about 4x6 nm2. Finally, in agreementwith SEM investigations, no micrometric nitride precipitate was found.

Fig. 4(A) shows the elemental profile from the sample nitrided for36 hours obtained by electron probe micro-analysis (EPMA) close tothe surface. These profiles were taken as line scans perpendicular tothe surface and contain fluctuation due to the local microstructurevariations. Fig. 4(B) shows a micrograph acquired just after the profileacquisition showing the location of the measured points. Vanadium

carbides are identified in points where vanadium and carbon concen-trations are far larger than those from the matrix average ones. Atthese locations, molybdenum concentration also increases slightly in-dicating that the precipitates contain small molybdenum quantities inaddition to vanadium. The chromium depth profile is fairly homoge-neously distributed, indicating that chromium carbides or nitridesdo not form large precipitates, in agreement with TEM analysis.Peaks associated with vanadium and carbon concentrations alongtheir depth profiles reveal that vanadium carbides are not dissolvedin the nitrided layer, also confirmed by TEM results [27]. Similar

Fig. 4. (A) Profile of principal elements from 36 hours nitrited sample. Concentrationsof V and C above average are identified as Vanadium Carbides. (B) Micrograph takenjust after the composition profile (WDS) acquisition showing the measured regionand the electron probe line scan.

Fig. 5. Nitrogen concentration profile for different studied nitriding times. The arrowsindicate the nitrogen penetration depth for 9, 16 and 36 hours.

Table 2Surface (~1 μm depth) hardness and Young's modulus for different treatment times.

Nitriding Time (hours) Young's modulus (GPa) Surface Hardness(GPa)

0 (matrix) 260±7 7.5±0.49 250±8 14.0±0.716 250±12 14.3±0.936 250±20 14.4±0.7

76 L.F. Zagonel et al. / Surface & Coatings Technology 207 (2012) 72–78

features were observed in other samples prepared by nitriding for 9and 16 hours treatments. At such low temperatures (400 °C), theprecipitation kinetics might be too slow to allow the formation ofmicroscopic nitride precipitates or the dissolution of carbides presentin the matrix [56,57].

For comparison, the nitrogen profiles of samples nitrided for 9, 16,and 36 hours are shown in Fig. 5. For these samples, the nitrogen sur-face concentration is close to 16 at%. These concentrations stem fromthe presence of (possibly dispersed) ε-Fe2-3N iron nitrides in thematerial surface convoluted by the electron excitation volume pro-bed, which is much larger than the iron nitride layer. Beneath of a~3 μm layer (i.e. beneath iron nitrides), the nitrogen concentrationmeasured on the cross-sections by EPMA is ~12%. This amount isstill far larger than the nitrogen solubility limit (~0.1 at.% at 400 °C)in un-strained iron matrix [44]. Moreover, this nitrogen concentra-tion is also larger than the quantity that would be necessary to con-sume the full Cr and V contents (~7.0 at.%) in the H13 steel to formnitrides (moreover, V is observed to be forming carbides and henceunavailable to form nitrides). Also, the additional line broadening ob-served for all α-Fe diffraction features after nitriding times longerthan 9 hours indicates that interstitial excess nitrogen also contrib-utes to strain the steel matrix [8,35]. Therefore, nitrogen appears tobe trapped at the interfaces of the nanosize coherent precipitates ob-served on the TEM [48]. Since the tempered martensitic microstruc-ture of this steel is very complex, high densities of dislocations andinterfaces could favor the presence of immobile excess nitrogen[19]. Indeed, nitrogen must be at interstitial sites, interfaces, and for-ming precipitates to explain such high nitrogen concentrations in the

steel matrix. In this complex scenario, it is observed a coincidence inthe nitrogen profiles (shown in the inset of Fig. 5) if they are scaled bynormalizing the diffusion depth by the square root of the nitridingtime. This could indicate that the mobile nitrogen has diffusedunder the effect of traps [58]. Also, even for 9 hours nitriding at400 °C, the depth of the diffusion zone reaches more than 40 μm,indicating the efficiency of this process even at such low tempera-tures [24].

As expected, the presence of the ε-Fe2-3N phase and coherent nan-oprecipitates on the surface increases the surface hardness of thestudied samples by nearly 100%, see Table 2 [8,26]. We attributedthe observed slight surface hardness differences to the increasingthickness of the compound layer on time. Moreover, the surfaceYoung's modulus is not significantly changed by the nitriding process,indicating that the modified surface maintain its original elasticity,see Table 2. We note that in a bimetallic two dimension planar struc-ture, the radius of curvature has a minimum for equal Young´s modu-lus of the materials forming the junction [59]. Therefore, it is a veryinteresting finding since the Young´s modulus matching betweenthe hard nitride case and the H13 substrate diminishes the chancesof delamination due to interface shear stress.

Thehardnesswasmeasured on the cross-sections by nanoidentationas function of depth and the results are correlated to the nitrogen pro-files (shown in Fig. 6). Moreover, it is observed a fairly good linear cor-relation between hardness and nitrogen concentration within the errorbars, as previously reported by other studies [60–62]. The correlationcoefficient obtained from the statistical analysis is R2=0.91 and therelation is H – H0=0.73 [N] (GPa/at.%), where H0 is the bulk hardness,7.5±0.4 GPa. This behavior is consistent with the occurrence of ahomogeneous precipitation of nanosized nitrides [19,48,63].

XPS measurements can provide information on the average localchemical bonding state of the alloying elements. Fig. 7 displays thespectra associated with the N1s and Cr 2p3/2 electrons core levelsobtained at different depths from the sample nitrided by 16 hours.As observed, for these elements the binding energy and band shapeare similar for all mapped depths. Moreover, the Cr binding energy

Fig. 6. Hardness variations H vs. nitrogen concentration [N]. The best straight linefitting is indicated in the plot.

77L.F. Zagonel et al. / Surface & Coatings Technology 207 (2012) 72–78

corresponds fairly well to the value expected for the solid solutionor nanosized precipitates of the element in iron α-Fe matrix [64].Regarding nitrogen, the peaks are relatively enlarged, phenomenonsteaming from nitrogen in different chemical states, i.e. dissolved insolid solution, adsorbed at interfaces, and forming different fineprecipitates. Indeed, previous in situ XPS experiments, showedsimilar results, confirming that at such low temperatures, no clearchemical state is reached for nitrogen [27,28,65]. In the assumptionof extensive formation of nitrides, binding energy shifts are expected,as previously reported in experiments performed at higher tempera-tures [27,28,60]. We remark that the residual nitrogen concentrationat a depth of 60 μm is most likely due to the incertitude in the depthmeasurement.

Fig. 7. XPS spectra associated with the N1s and Cr 2p3/2 core level electrons. Thecurves show no change in the chemical state of the elements as function of depth(16 hr sample).

4. Conclusions

The effect of nitriding AISI H13 tool steel at a low temperature(400 °C) is reported. The high plasma reactivity even at this temper-ature favored the formation of the ε-Fe2-3N phase at the barely sur-face of the material. High nitrogen concentration was observed inthe diffusion layer of the studied samples, indicating the presence ofexcess nitrogen (concentration above expected values). Nitrogen dif-fuses forming coherent nanosized precipitates while original carbidespresent in the matrix are not dissolved during the process. Indeed,large sized precipitation was inhibited by the low temperature pro-cess even for high nitrogen concentrations (above 12 at% underneaththe surface) and, consequently, nitrogen diffuses as deep as 40 to70 μm in the material bulk. These findings can be explained by the ex-cess nitrogen in the matrix interstitial sites, the α-Fe matrix expan-sion due to coherent nanoprecipitates, and finally immobile excessnitrogen trapped at the interfaces of the complex steel structure.This scenario is supported by the linear correlation of hardnessand nitrogen concentration that suggests progressive homogeneousnanosize precipitation (contrary to coarse precipitation leading tohardness decrease). Moreover, the similar chemical bonding state ofnitrogen and alloying elements suggested by the XPS measurementsobtained at several depths of the treated sample support this inter-pretation. Finally, these results support the idea that plasma nitridingprocesses performed at lower temperatures is effective to formnanosized precipitates, thus enhancing surface hardness withoutaffecting bulk microstructure or hardness. Therefore, this processcould be indicated for tool steel applications demanding increasedsurface hardness for which high nitriding temperature processesare not acceptable due to workpiece distortion or change in bulkmechanical properties.

Acknowledgments

Part of this work was performed during a stay of LFZ in the Prof. E. J.Mittemeijer group at the Max Planck Institute (MPI) for MetalsResearch in Stuttgart, Germany. The authors are grateful to Mrs S.Haug from theMPI for assistancewith the electron probemicroanalysismeasurements. This work was partially supported by FAPESP, grant# 05/53926-1. FA and LFZ are CNPq fellows. HP thanks the financialsupport of the DFG project 444Bra-113/25/0-1.

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