ISIJ International, Vol. 57 (2017), No. 4

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ISIJ International, Vol. 57 (2017), No. 4, pp. 737–745

* Corresponding author: E-mail: antschip@usp.brDOI: http://dx.doi.org/10.2355/isijinternational.ISIJINT-2016-553

1. Introduction

Plasma nitriding has being widely used in diverse indus-trial applications. In particular, plasma nitriding treatment is especially well suited for preparing the surface of tool steels before depositing hard ceramic coatings.1–8) This two-stage process of nitriding plus hard coating deposition is normally known as either duplex or hybrid process.1–8)

Deposition of hard ceramic coatings onto both nitrided and non-nitrided steels is a potentially effective treat-ment for improving abrasive wear resistance. Although wear resistance near the surface of the coated part does not depend on substrate mechanical properties, tool’s life is barely increased when the substrate is coated in the unhardened condition and high contact loads are applied. Tool steels that are not hardened previously to coating are prone to in-service early breaking of the coating in the metal-ceramic interface, even though abrasive wear is still negligible. In contrast, the nitride layer significantly increases load bearing capacity of the whole tools submitted to plasma nitriding prior to hard coating deposition. Load-driven cracks are avoided and the whole part can withstand severe in-service conditions.9,10)

The role of substrate properties on endurance of tool

Thermodynamic Analysis of M7C3 Carbide Dissolution during Plasma Nitriding of an AISI D2 Tool Steel

Carlos Mario GARZÓN,1) Adonias Ribeiro FRANCO Jr.2) and André Paulo TSCHIPTSCHIN3,4)*

1) Physics Department, National University of Colombia, Av. Cra. 30 No 45-03 – Ciudad Universitaria, Ed. 404, Bogotá D.C., Colombia. 2) Federal Institute of Education, Science and Technology of Espírito Santo - IFES, Av. Vitória, 1729, CEP 29050–066, Vitória - ES, Brazil. 3) Metallurgical and Materials Engineering Department, University of São Paulo - USP, Av. Prof. Mello Moraes 2463, CEP 05508–900, São Paulo - SP, Brazil. 4) LNNano – Brazilian Nanotechnology National Laboratory – CNPEM - Brazilian Center for Research in Energy and Materials, R. Giuseppe Máximo Scolfaro, 10000 - Polo II de Alta Tecnologia, Campinas - SP, 13083–970 Brazil.

(Received on October 27, 2016; accepted on December 7, 2016; J-STAGE Advance published date: March 7, 2017)

This contribution addresses the dissolution of primary M7C3 carbides and the simultaneous precipitation of nitrides during plasma nitriding of an AISI D2 tool steel. Samples of quenched and tempered D2 tool steel were plasma nitrided at 793 K (520°C) for times varying between 2.52 and 39.6 ks (between 42 min and 11 h). Microstructure and chemical composition of nitrided samples were then characterized by scan-ning electron microscopy and wavelength dispersive spectrometry. It was observed that nitrogen absorp-tion induces M7C3 dissolution and simultaneous in situ precipitation of MN nitrides. Calphad computational thermodynamic simulations were carried out in order to evaluate phase stability and atomic partitioning in the microstructure. Atomic partitioning between MN and the martensitic matrix was associated with a thermodynamic metastable state. At 793 K, the negligible metallic-element diffusion inside both M7C3 and MN phases makes the metastable state kinetically favored instead of the thermodynamic stable state.

KEY WORDS: plasma nitriding; tool steels; Calphad; diffusion; phase transformations.

steels which are hybrid processed should be emphasized.8) Besides steel hardness, in-service integrity of coated parts depends significantly not only on the phases present at regions adjacent to steel surface but also on the surface roughness.8–11)

Both formation of nitrides and decomposition of car-bides due to nitrogen absorption during plasma nitriding of alloyed steels have been reported.1,11–17) In particular, formation of iron nitrides, i.e., the compound layer, dur-ing plasma nitriding of diverse steels has been studied in depth.11,12) Sun and Bell1,9,11) pointed out that formation of compound layer should be inhibited in order to attain optimized coating-to-steel adherence. On the other hand, knowledge about primary-carbide dissolution during steel plasma nitriding is still incipient.14) Leroy et al.14) reported primary carbide dissolution and simultaneous nitride pre-cipitation during plasma nitriding of a high-chromium steel (wt.-%, 12% Cr, 1.6% C, 0.8% Mo, 0.4% V). They stated that ‘carbides undergo a total transformation into substitu-tional (Cr, M) N-type nitrides with a simultaneous release of α-iron from the very nitride phase’.14) They suggested that the transformation from primary carbide to MN nitride “proceeds according to an interface displacement mecha-nism”,14) i.e., the new MN nitride and the primary carbide are separated by a sharp interface which moves towards the primary carbide in a diffusion-driven process; an abrupt change in composition and structural properties takes place

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at the almost zero-thickness interface from primary carbide and the new MN nitride regions.

The aim of this work is to study the dissolution of large primary M7C3 carbides during plasma nitriding of an AISI D2 tool steel. A stability analysis of M7C3 dissolution is carried out from the results obtained by scanning electron microscopy (SEM), wavelength dispersive spectrometry (WDX), and computational thermodynamic simulations. Diverse nitriding treatments were carried out at increasing times (between 42 min and 11 h) with careful control of nitriding conditions to avoid formation of compound layer.

2. Experimental

Table 1 gives the chemical composition of the AISI D2 tool steel studied.

The starting material was a hot worked bar which diam-eter was 31 mm. Samples of 5 mm width were cut from the initial bar, and then were oil quenched and tempered twice. It is a common practice to double temper high carbon and highly alloyed tool steels to reduce the amount of retained austenite in the microstructure to a minimum, guaranteeing high hardness and adequate toughness to the final product. Table 2 gives the parameters used for heat treatments.

Samples were plasma nitrided at 793 K (520°C) using a H2 + N2 gas mixture, which contained 5%-vol of N2. The nitriding gas mixture was carefully selected with the aim of inhibiting formation of compound layer. Nitriding times varied between 2.52 and 39.6 ks (between 42 min and 11 h). Prior to nitriding treatment, a sputtering stage in hydro-gen plasma was carried out in the same treatment chamber. Table 3 gives the parameters used during plasma nitriding.

Specimen’s microstructure was examined by SEM using a JEOL JSM 5900LV microscope, operating between 20 and 25 kV. Metallographic preparation consisted of emery paper

grinding followed by diamond paste polishing and a final 1.8 ks stage of 0.05 μm colloidal silica polishing. Chemical etching was performed by using a Beraha reagent (100 ml distilled water, 24 g sodium thiosulphate, 3 g citric acid, and 2 g cadmium chloride).

Chemical characterization of individual components in the microstructure was made by X-ray WDS microanaly-sis using an Oxford WDX600 spectrometer coupled to a Cambridge Stereoscan 440 scanning electron microscope operating at 15 kV, following a procedure reported by Toro et al.18)

Hardness (H) values were determined by using depth-sensing indentation tests at different distances from the sample surface (using a cross sectioned sample). These hardness tests were carried out in a Fischerscope H100 apparatus, using a Vickers indenter tip. The maximum load was 50 mN, the loading and unloading times were 90 s, and the dwell time at peak load was 20 s. Force-depth (P-h) indentation data was analyzed by using conventional procedure.19)

Thermocalc software was used to predict thermodynamic phase stability and equilibrium chemical composition of dif-ferent structure components. The TCFE5 database was used. Equilibrium conditions are estimated in Thermocalc by using a classical thermodynamics framework, namely, the Calphad method,20) which phase stability criterion [Eq. (1)] corresponds to minimum values of Gibbs free energy func-tions. These free energy functions were constructed for each possible phase, using chemical composition, temperature, and pressure as the basis. Numerical computations allowed estimating Gibbs free energy (compositional) of different microstructure aggregates, phase fractions, and atomic par-titioning in microstructure as a function of temperature and nitrogen absorption:

G n Gii

p

i= ==∑

1

ϕ minimum, ..................... (1)

where G is the molar Gibbs free energy of the system, ni is the mole fraction of each phase, p is the number of phases, and Gi

ϕ is the Gibbs free energy of each phase.

3. Results

3.1. Characterization of Heat-treated SamplesFigure 1 shows SEM micrographs of a sample in the

as-tempered condition. It can be seen a dispersion of coarse carbides (roughly between 1.0 and 5.0 μm width) in a martensitic matrix. Those coarse carbides were identi-fied by chemical microanalysis as M7C3, where M denotes the metallic atomic species Fe, Cr, V, Mo, and Mn. The martensitic matrix is actually a tempered martensite, i.e., a dispersion of fine tempering carbides in a body-centered tetragonal matrix phase.

Table 4 gives the chemical composition of both M7C3 coarse carbides and martensitic matrix as determined by WDS and as calculated by Thermocalc® calculated. Ther-modynamic simulations were carried out considering 1 353 K, which is the actual austenitizing temperature in heat treatments.

Experimental data show that the atomic partitioning between coarse primary carbides and the martensitic matrix

Table 1. Chemical composition of the AISI D2 steel (wt-%).

Cr C V Mo Mn Si P S Fe

11.7 1.53 0.77 0.89 0.53 0.07 <0.003 <0.002 Bal.

Table 2. Heat treatment conditions.

Quenching Double tempering

Temperature 1 353 K 813 K

Time 1.8 ks 7.2 ks

Cooling Oil Air

Table 3. Plasma nitriding conditions.

Time 2.52, 10.8, 21.6, and 39.6 ks

Gas mixture (95 H2−5 N2)%-vol.

Temperature 793 K

Pressure 250 Pa

Arc voltage 500 V

Pulse duration 200 μs

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is in equilibrium; M7C3 carbides are very rich in Cr, V, and C when compared with the martensitic matrix.

3.2. Characterization of Nitrided SamplesFigure 2 shows SEM micrographs taken in the transverse

section of nitrided samples. Decomposition of M7C3 car-bides at the surface-adjacent region and in situ precipitation of a different phase can be observed.

The product phase precipitated in situ on M7C3 carbides was identified by chemical microanalysis as a MN nitride. Table 5 gives WDS-measured chemical compositions of MN nitrides and M7C3 carbides at the region adjacent to steel surface.

Table 6 gives the relative fractions of metallic species inside both MN and M7C3. The ratio of metal atoms in MN nitrides is virtually the same as in M7C3 carbides (Table 6), but their carbon and nitrogen contents are quite different (Table 5).

From the metallographic structure presented in Fig. 2 and the chemical compositions in Tables 4–6 it is possible to state that MN growth is associated to diffusion of the non-metallic species inside MN, M7C3 and the martensitic

matrix, while negligible diffusion of the metallic species happens inside all of them. Figure 3 shows some metal-lographic structures which stress the role of carbon and nitrogen diffusion on M7C3 dissolution and MN growth. The dissolution reaction of M7C3 can be schematically expressed as

M C MN7 3 0 1+ ′ → + ′[ ]α αin situ

, .................. (2)

where α0′ and α1′ represent the martensitic matrix region near the precipitate before and after the in situ transforma-

Fig. 1. Scanning electron micrograph of a heat-treated sample (quenched and twice tempered). Etching: Beraha. A dis-persion of large carbides (M7C3) in a tempered martensitic matrix (α′) can be seen.

Table 4. Chemical analysis (WDS) and thermodynamic calcula-tions (TC) of main elements in M7C3 carbides and in martensitic matrix after thermal treatment (wt-%).

Cr C V Mo N Fe

Tempered martensite

WDS 8.0 ±1.0 0.67±0.5 0.25 ±0.2 0.7±0.3 0.4 ±0.2 89.8 ±4.0

TC 7.3 0.73 0.3 0.8 0.53 90.3

Coarse M7C3 carbides

WDS 43.6 ±2.5 8.3 ±1.0 5.4 ±1.0 2.2 ±0.5 0.5 ±0.3 36.9 ±3.5

TC 51.2 8.7 4.9 1.6 0.5 33.0

Fig. 2. SEM micrographs of the transverse section of the nitrided samples. Etching: Beraha. Decomposition of M7C3 carbides at the surface-adjacent region can be noticed. (a) 2.52 ks; (b) 10.8 ks; (c) 21.6 ks and (d) 39.6 ks.

Table 5. WDS chemical analysis (wt-%) of MN and M7C3 phases near the surface of a sample nitrided for 39.6 ks.

Phase Cr C V Mo Mn N Fe

MN 38.8 ±2.5 0.43 ±0.5 3.9 ±1.0 1.8 ±0.5 0.3 ±0.2 17.5 ±3.5 35.3 ±3.5

M7C3 44.1 ±2.5 8.0 ±1.5 5.2 ±1.0 2.1 ±0.5 0.5 ±0.3 N.D.* 38.0 ±3.5

*Not-detected

Table 6. Relative fraction (mole-%§) of metallic elements in the MN and M7C3 phases near the surface of a sample nitrided for 39.6 ks.*

Cr V Mo Mn Fe

MN 49 5.0 1.2 0.2 42

M7C3 50 6.0 1.2 0.5 42§The molar character of reported chemical composition should be noticed.*The metallic fraction inside each phase was calculated from the data in Table 5. The data uncertainties are also presented in Table 5.

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tion, respectively. The role of the martensitic matrix as sink for carbon atoms arising from in situ decomposed M7C3

should be emphasized.

3.3. Growth Kinetics of in situ Precipitated MN Nitrides

Figure 4 shows the in situ MN-case depth, the hardened-case depth, and the ratio between those case depths as a function of the nitriding time.

The in situ MN-case depth was defined as the deepest microregion where in situ dissolution of primary carbides was observed by metallographic analysis. The hardened–case depth was defined as the mean depth where the hard-ness is 0.5 GPa higher than the hardness of the non-nitrided core. The relationship observed between hardened-case depth and nitriding time is typical for a diffusion-controlled layer thickening: the rate of layer thickening decreases while elapse time is increased. On the contrary, no decrease of the growth rate of in situ MN-case depth with increasing elapsed time was observed. The role of limited duration of nitriding treatments (maximum nitriding time of 39.6 ks) on the layer thickening should be emphasized.

The ratio between hardened-case depth and in situ MN-case depth varies significantly as nitriding time is increased (Fig. 4(b)), due to the very dissimilar dependence of them with elapsed time. At the earliest nitriding stages the in situ MN-case depth is a small fraction of the hardened-case depth (~2%) but this ratio increases with elapsed time, attaining a value of ~14% after 39.6 ks elapsed time.

Figure 5 shows carbon and nitrogen content profiles

inside the martensitic matrix as a function of distance from the surface, for a sample nitrided at 39.6 ks. Chemical com-position profiles of metallic species are not reported because no variation in the content of those atoms was observed.

Figure 6 shows carbon and nitrogen activity profiles inside the martensitic matrix as a function of distance from the surface calculated for a sample nitrided at 39.6 ks. These activity profiles were assessed with chemical microanalysis as a basis. The activity (dimensionless) reported for each component was computed relative to the stable element reference (SER), at 25°C (298.15 K). It was assumed arbi-trarily that inside every microregion of tempered martensite the thermodynamic equilibrium conditions were attained, where microstructure was a mixture of BCC iron and fine carbides and nitrides. The atomic partitioning among the phases inside each microregion of tempered martensite was the one ruled by the equilibrium conditions. In Fig. 6 it can be seen that nitrogen absorption from the plasma induces an activity gradient of both N and C, where N and C atoms have their activities increased at regions near the steel sur-face with regard to sample inner regions, i.e. a net diffusion flux of both C and N is expected towards sample core.

For 39.6 ks elapsed time, nitrided case depth was around 35–45 μm while the region where carbon diffusion occurred was around 75–100 μm in depth (Fig. 4). On the other hand, the hardened-case depth was around 70–90 μm for that nitriding time. Thus, the hardened-case depth and the depth of the region where carbon diffusion occurred are closely related. Consequently, the fact that the hardened-case depth for 39.6 ks elapsed time was almost twice the nitrided-

Fig. 3. Metallographic evidences of the role of carbon and nitrogen diffusion on the dissolution of M7C3 and growth of MN phases. Optical micrographs. Etching: Beraha. The role of bulk diffusion on M7C3, MN and martensitic matrix and interface diffusion on phase interfaces is emphasized. (Online version in color.)

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case depth can be attributed to carbon diffusion against the concentration gradient (Fig. 5). Carbon diffusion from low-carbon to high-carbon content regions occurs because nitrogen absorption increases the activity of carbon atoms at the region adjacent to the steel surface, with regard to atoms at deeper regions from the surface (Fig. 6). Actually, nitrogen absorption affects the activity of all atomic species in the martensitic matrix, but the temperature is very low for significantly allowing diffusion of metallic species.

4. Analysis and Discussion

Results of in situ growth of MN nitrides onto M7C3 carbides were presented and depicted. It was formulated that this phase transformation is driven by atomic nitrogen absorption from plasma to the solid phases into the steel. MN formation with low dissolution of carbon as well as negligible accumulation of nitrogen inside M7C3 carbides was observed. Figure 7 shows a schematic model to describe evolution of atomic partitioning during this phase transformation.

Figure 8 shows several two-dimensional projections of M7C3 and MN crystal structures along diverse zone axis (depicted on the figure). Orthorhombic Fe7C3 and cubic FeN crystal structures were assumed to construct in Fig. 8, however it should be noticed that quite small differences in lattice parameter should exist between Fe7C3 and Cr7C3 as well as between CrN and FeN, because of the close posi-tion in the periodic table of Fe and Cr atoms. Thus, it is a

Fig. 4. (a) Mean depths of the in situ MN-case and hardened-case as a function of elapsed time and (b) ratio between MN-case depth and hardened-case depth as a function of elapsed time.

Fig. 5. WDS measured (a) nitrogen and (b) carbon contents in the martensitic matrix (martensite plus fine tempering precipi-tates) as a function of distance from the surface for a sam-ple nitrided for 39.6 ks.

Fig. 6. Activity of (a) nitrogen and (b) carbon in the martensitic matrix (martensite plus fine tempering precipitates) as a function of distance from the surface for a sample nitrided for 39.6 ks. Activity profiles were computed with chemical microanalysis as a basis.

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reasonable to represent M7C3 and MN crystal structures in Fig. 8 with Fe7C3 and FeN as a basis. Crystallographic data used for constructing Fig. 8 are reported in the figure. Crys-tallographic simulations were carried out using the Vesta 3 software. Downloading information and use of Vesta 3 are depicted in depth elsewhere.21)

Figure 8 was constructed for descriptive purposes and no experimental study of crystallographic features of either M7C3 or MN phases was undertaken. It is relevant to point out, as observed in Fig. 8, that M7C3 to MN phase trans-formation could not take place with a simple one-by-one substitution of carbon atoms by nitrogen atoms. Instead, a long-range carbon and nitrogen diffusion is mandatory. In addition, no epitaxy between M7C3 and MN can be predicted from the results shown in Fig. 8. Short range rearrangement of metallic atoms, at lengths lower than interatomic distance, is a must in such phase transformation type. It implies changes in the average distance between

Fig. 7. Schematic model of in situ precipitation of MN nitrides on M7C3 carbides during plasma nitriding of the D2 tool steel.

Fig. 8. Two-dimensional projections of (a) M7C3 and (b) MN crystal structures along diverse zone axis. Orthorhombic Fe7C3 and cubic FeN crystal structures were assumed.

metallic atoms.The growth of MN nitrides can be associated with: (i)

bulk nitrogen diffusion inside MN and bulk carbon dif-fusion inside M7C3; (ii) interface diffusion of C and N in both matrix-MN and matrix-M7C3 phase interfaces; (iii) nitrogen and carbon diffusion inside the martensitic matrix and (iv) negligible bulk diffusion of metallic species inside the stoichiometric phases as well as the martensitic matrix (metallographic evidences of the role of carbon and nitrogen diffusion on the dissolution of M7C3 and growth of MN phases were already presented in Figs. 2 and 3).

Despite the diffusion-controlled character of MN growth, the in situ MN-case depth varies with elapsed time in a non-parabolic manner (see Fig. 4). This can be explained as follows. At the earliest stages of nitriding treatment, growth of MN nitrides is mainly driven by nitrogen diffusion inside MN, while at later stages nitrogen diffusion inside the mar-tensitic matrix plays an increased role in supplying nitrogen to the MN phase just ahead of the migrating phase interface M7C3–MN.

Table 7 gives the molar Gibbs energy (G) computed for M7C3, M2N, and MN phases at 793 K, i.e., the nitriding temperature. Atomic molar fractions used for computing G

Table 7. Molar Gibbs energy (G) computed for M7C3, M2N, and MN phases at 793 K.

Phase G J/mol

Chemical composition* (mole-fraction)§

Cr Mo V C N Fe

M7C3 −3.6×104 0.408 7.01×10 −3 3.94×10 −2 0.3 0.00 bal.

M2N −6.2×104 0.3887 6.68×10 −3 3.75×10 −2 0.0 0.333 bal.

MN −7.2×104 0.2916 5.01×10 −3 2.82×10 −2 0.0 0.50 bal.

*Metallic fraction inside each phase was the same reported in Table 4 for primary carbides, whereas the C or N fractions were varied according to the metallic/nonmetallic exact stoichiometry for each phase.§The molar character of reported chemical composition should be noticed.

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Table 8. Equilibrium phase fractions and their chemical compositions as a function of nitrogen content absorbed in the steel at 793 K.*

Average N-content (wt-%) Phases Phase fractions

(vol-%)Phase chemical composition (wt-%)

Cr C V Mo Mn N Fe

0.5 N

BCC 80 0.54 9×10 − 5 4×10 − 4 0.038 0.24 1.0×10 − 6 Bal.

M7C3 12.9 66.5 8.8 0.14 1.82 2.74 0.0 Bal.

M23C6 1.6 41.5 5.1 4×10 − 5 17.1 0.47 0.0 Bal.

M(C,N) Cr rich 4.5 68.4 10.1 4×10 − 5 9.55 0.03 8.7 Bal.

MN V rich 1.0 0.01 0.72 78.8 0.003 0.001 20.18 Bal.

2.0 N

BCC 76.8 0.02 0.001 1×10 − 5 0.002 0.21 3.0×10 − 5 Bal.

M(N,C) Cr rich 19.0 72 8.1 1.4 5.5 0.60 11.51 Bal.

MN V rich 0.9 0.10 0.38 77 0.002 0.016 21.04 Bal.

Cementite 3.3 3.08 6.71 0.002 0.017 7.9 3.0×10 − 6 Bal.

6.0 N

BCC 40.1 2×10 − 5 0.002 1×10 − 9 0.001 0.02 0.04 Bal.

MN Fe rich 32.9 31.9 0.09 2.6 0.43 0.02 17.0 Bal.

MN Cr rich 5.5 46.6 0.22 0.36 15.1 2.06 19.2 Bal.

Cementite 21.1 0.002 6.68 3×10 −7 0.01 0.73 0.004 Bal.

Fe2(N,C) 0.4 0.024 2.15 2×10 − 4 0.83 9.03 7.42 Bal.

*Data reported in this table correspond to selected points in the phase diagram Fig. 6.

values are also given in this table.These molar fractions were estimated assuming that

all the phases have an even stoichiometric ratio between metallic and non-metallic species and that the relative frac-tion of metallic species in each phase is the same as the relative metallic fractions in M7C3 phase stable at 1 353 K (i.e., the austenitizing temperature during heat treatment). According to the data in this table, the thermodynamic sta-bility hierarchy from the most stable to the least stable is: MN–M2N–M7C3.

Figure 9 shows equilibrium phase stability domains as a function of both nitrogen content absorbed in the steel and

temperature.Table 8 gives, at 793 K, phase volumetric fractions and

the chemical compositions of stable phases for increasing nitrogen contents. Data reported in Table 8 corresponds to selected points in the phase diagram in Fig. 9.

Nitrogen absorption clearly decreases the stability of highly alloyed carbides, while increases the stability of both low-alloyed carbide cementite and nitrides. Moderate nitrogen additions induce increasing stability of M(N,C) carbonitrides, while high additions of nitrogen induce increasing stability of Fe-rich nitrides, namely, formation of a compound layer. Those numerical results are in accor-dance with results in literature.1,2,11,12) Mittemeijer et al.22) also calculated the relative stability of Fe3C cementite and CrN nitride upon increasing nitrogen content in the micro-structure, obtaining an increased stability of low-alloyed Fe3C carbide and CrN nitrides.

As it was reported in Section 2, the nitriding parameters used, namely, the nitriding temperature (793 K) and the composition of the nitriding gas (95 H2−5 N2%-vol.), were selected with the aim of inhibiting compound layer forma-tion. The average nitrogen content (measured by WDX) at the steel surface of samples nitrided for 9.6 ks, was between 4.2 and 4.8 N wt-%. This N-content and the microstructure observed in the nitrided samples are in good agreement with the phase diagram shown in Fig. 9 and with the aimed compound layer-free microstructure. However, it should be pointed out that the atomic partitioning among tempered martensitic matrix and large M(N,C) or M7C3 phases is very different from the one predicted for stable equilibrium at 793 K, as discussed in the next paragraph.

Chemical compositions of phases at 793 K, for increased nitrogen contents in the alloy, are given in Table 8. Chemi-cal composition of stable M(N,C) carbonitrides numerically predicted differs significantly from chemical compositions measured experimentally for nitrides actually precipitated

Fig. 9. Equilibrium phase diagram (isopleth) where phase stabil-ity domains are related to temperature and nitrogen addi-tions to the D2 tool steel. Chemical composition of phases at selected points (namely, 0.5, 2.0, and 6.0 N wt-%) in this diagram, for 973 K, are given in Table 8. The pseudobi-nary character of the diagram should be noticed.

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(Table 5). Formation of diverse MN nitrides rich in Cr or V or Fe is predicted, and its carbon content is not expected to be negligible. Thus, it is clear that the actual chemical composition of MN nitrides corresponds to a metastable state, which is very far from the stable atomic partition-ing in microstructure. The metastable thermodynamic state regarding MN chemical composition is kinetically favored instead of the stable one, due to the negligible atomic-element diffusion inside both M7C3 and MN at the nitriding temperature, namely, 793 K. These results are quite differ-ent from those published by Zagonel et al.,23) who observed that, during plasma nitriding at 673 K, chromium carbides present initially in the matrix were not converted to CrN, most probably due to the lower nitriding temperature used in their experiments.

The significant difference regarding chemical composi-tion between in situ precipitated MN and that for MN phase at the stable equilibrium is potentially detrimental for high temperature stability in-service of nitrided parts. Thermo-dynamic simulations predict a significant net driving force for decomposition of in situ precipitated MN nitrides into a mixture of nitrides plus ferrite.

High residual stresses inside MN nitrides, associated with the in situ phase transformation, can be expected because of high volumetric restrains. Mittemeijer et al.22) reported a specific volume of 145.4 cm3kg −1 for Cr7C3 carbide and 161.3 cm3kg −1 for CrN nitride, indicating a possible source of constraint during the in-situ dissolution of Cr7C3 carbide and the formation of CrN nitride. In fact, in some samples, mainly those long-term nitrided, some microcracking inside MN nitrides was observed (see insets in Fig. 2). Those microcracks could be formed during metallographic section-ing of samples, and they probably arise from high levels of residual stresses inside MN nitrides. Although this phenom-enon has a major technological relevance, unfortunately the results presented in this work do not allow undertaking a deep insight into these likely residual stresses inside MN and further work is then mandatory.

Nitrogen absorption inside the steel induces an increased stability of cementite (phase diagram Fig. 9). This increased stability of cementite is in accordance with reported pre-cipitation of cementite carbides in plasma nitrided M2 tool steels24–26) and chromium bearing steels.27) Cementite

Fig. 10. SEM micrograph of the transverse section of a nitrided sample. Etching: Beraha. Precipitation of second phases at prior austenitic grain boundaries is emphasized. That second phase precipitation mainly takes place at the end of the very nitrided region (towards sample core) just in the region were carbon enrichment is observed in carbon profiles.

precipitation at prior austenite grain boundaries have been reported.25,26) Results of transmission electron microscopy (TEM) characterization of that grain boundary cementite were reported by Tier.25) A quite similar second-phase pre-cipitation along prior austenite grain boundaries was veri-fied in this work (neither TEM or chemical characterization was carried out). Figure 10 shows an SEM micrograph where second phase precipitation along prior austenitic grain boundaries can be seen (the similarity of this micro-structure with reports of Tier et al.28) should be noticed). There is no evidence of such second-phase precipitation in the nitrided region. Instead, it was detected in the region where carbon heightening occurs. A priori it can be assumed that carbon arising from dissolution of M7C3 carbides can be an auxiliary factor inducing cementite precipitation at the surface-adjacent region.14) However, it was shown in this work that the surface-adjacent region is actually car-bon depleted, and that carbon arising from dissolution of M7C3 enriches the non-nitrided region next to the nitrided region, towards the sample core. The sharp drop of carbon activity towards steel inner regions, just ahead the region where M7C3 to MN phase transformation has already taken place (Fig. 6), induces fast carbon diffusion towards inner regions in the sample and depletes in carbon the tempered martensitic matrix, hindering precipitation of cementite carbides at regions near the surface of the steel. Rocha et al.29) report that depletion in carbon near sample surface is driven by decarburization processes at the plasma-steel interface, being observed in higher extend under nitriding conditions where no compound layer is developed. They29) argued that depletion in carbon in the very nitrided region should be associated with carbon diffusion towards sample surface, being slowed down notably when a compound layer is formed (carbon diffusivity in iron nitrides is rather low than in tempered martensite). However, the plasma-induced decarburization hypothesis is not in agreement with the carbon heightening at the end of nitrided region depicted in this work, neither is it in agreement with a continuously decreasing activity profile of carbon (from sample surface towards inner regions). Thus, results presented in this work do not support the decarburization hypothesis but are consis-tent with the hypothesis of counter gradient carbon diffusion driven by nitrogen absorption from the plasma.

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5. Conclusions

M7C3 primary carbides in an AISI D2 tool steel are dis-solved when the steel is submitted to plasma nitriding at 793 K (520°C). This dissolution takes place at regions adjacent to steel surface and is associated with in situ diffusion-controlled precipitation of MN nitrides.

MN growth is associated with diffusion of the non-metallic species, namely, C and N, and negligible diffusion of metallic species, namely, Fe, Cr, V, Mo, and Mn, inside MN, M7C3, and the martensitic matrix.

The negligible atomic mobility of metallic species kineti-cally favors the in situ precipitation of MN nitrides with metastable chemical composition very far from that cor-responding to the stable equilibrium condition.

At the earliest stages of nitriding, M7C3 dissolution is mainly controlled by N-diffusion inside the MN product phase. At latter stages, an increased relevance of nitrogen diffusion inside the martensitic matrix on M7C3 dissolution kinetics is established.

AcknowledgementsThe JSM-5900 LV microscope of the LME/LNNano

(Laboratory of Electron Microscopy of the Brazilian National Nanotechnology Laboratory - Campinas – Brazil) was used for performing part of the electron microscopy work.

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