Influence of substrate pre-treatments by Xe ion ...

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Influence of substrate pre-treatments by Xe+ ion bombardment and plasma nitriding on

the behavior of TiN coatings deposited by plasma reactive sputtering on 100Cr6 steel

S. Vales1a, P. Brito2b, F.A.G. Pineda1c, E.A. Ochoa3d, R. Droppa Jr4f, J. Garcia5g, M.

Morales3h, F. Alvarez 3i, H. Pinto1j*

1Universidade de São Paulo (USP), Escola de Engenharia de São Carlos, Av. Trabalhador

São Carlense 400, São Carlos – SP, CEP 13566-590, Brazil

2Pontifícia Universidade Católica de Minas Gerais (PUC-MG), Av. Dom José Gaspar 500,

30535-901 Belo Horizonte, MG, Brazil

3Universidade Estadual de Campinas (UNICAMP), Campus Universitário Zeferino Vaz,

Barão Geraldo, Campinas - SP, CEP 13083-970, Brazil

4Universidade Federal do ABC (UFABC), Av. dos Estados, 5001, Santo André –SP, CEP

09210-580, Brazil

5Sandvik Machining Solutions, R&D Material & Processes, Lerkrogsvägen 19, SE 12160

Stockholm, Sweden

Email address: [email protected], [email protected], [email protected]; [email protected], [email protected], [email protected], [email protected], [email protected], [email protected]* *Corresponding author

Highlights:

100Cr6 steel substrates were Xe bombarded with 1.0 KeV

The XRD patters confirms that the previous Xe ion bombardment increase nitrogen

implantation in plasma nitriding process

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The AFM shows that the impact energy in combination with plasma nitriding cause a

significantly increase roughness

Abstract

In this paper the influence of pre-treating a 100Cr6 steel surface by Xe+ ion

bombardment and plasma nitriding at low temperature (380°C) on the roughness, wear

resistance and residual stresses of thin TiN coatings deposited by reactive IBAD was

investigated. The Xe+ ion bombardment was carried out using a 1.0 keV kinetic energy by a

broad ion beam assistance deposition (IBAD, Kaufman cell). The results showed that in the

studied experimental conditions the ion bombardment intensifies nitrogen diffusion by creating

lattice imperfections, stress, and increasing roughness. In case of the combined pre-treatment

with Xe+ ion bombardment and subsequent plasma nitriding of samples with relatively high

average roughness, the wear volume increases in comparison to the substrates exposed to only

nitriding or ion bombardment.

Keywords: Ion implantation, nitrides, ion beam-assisted deposition, thin films, wear.

1. Introduction

The performance of cutting tools and dies can be largely improved by surface

modification techniques, such as plasma nitriding and hard coating deposition, which are

widely applied to steels leading to increase hardness, oxidation and corrosion resistance, wear

resistance and overall superior tribological properties [1-5]. Plasma nitriding allows for

excellent parameters control and high efficiency nitrogen diffusion at relatively low

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temperatures (300-450°C), which results in a more chemically uniform compound layer (“white

layer”) and maintaining the material original hardness in comparison to conventional gas

nitriding processes [6].

Typically, plasma nitrided steel work-pieces have a compound layer (“white layer”)

formed at the surface which is composed by nitrides. This layer can be formed by ε-Fe2-3N and

γ’-Fe4N or only a monolayer with one type of nitrides. Bellow of the compound layer there are

a diffusion zone, which is composed by the α-phase saturated of nitrogen and the fine

precipitates nitrides [7]. The nitrides ε-Fe2-3N and γ’-Fe4N can exhibit very high hardness,

lower friction coefficient and good corrosion resistance, thus improving mechanical, chemical

and tribological properties [8-11].

Deposition of hard coatings such as TiN, TiC, TiCN and TiAlN, which exhibit very high

hardness, good wear resistance and low friction coefficient is also commonly utilized for

enhancing surface characteristics of steel parts. The hard ceramic coating can be applied both

directly on the metallic substrate as well as on a previously nitrided material [12-14]. In the

latter case a duplex scale is formed with superior fatigue and wear resistance due to increased

compatibility between the hard external coating and the strengthened underlying substrate [15,

16], e.g. if the substrate is much softer than the hard coating, it might undergo plastic

deformation in service causing the outer layer to collapse and fail prematurely [17]. Ion Beam

Assisted Deposition (IBAD) is a Physical Vapor Deposition (PVD) process for hard coating

deposition that may be combined with ion bombardment. The additional energy of the ion beam

allows for coating deposition at lower temperatures, avoiding some of the difficulties associated

with conventional PVD processes on nitrided substrates, such as bad adhesion, creation of

porosities or decomposition of the nitride layer leaving a relatively soft iron layer behind [18,

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19]. The IBAD technique also allows for better densification of the hard coating in comparison

to PVD processes, as well as independent control of several parameters [20, 21].

The ion bombardment of metallic surfaces produces a variety of microstructural

changes, such as the creation of lattice defects, phase transformations, material removal

(sputtering) and amorphization on the annealed 100Cr6 steel [22]. Recent studies have

demonstrated that the atomic bombardment of steel substrates with Xe+ ions improves the

diffusion of nitrogen in the material by prompting stress, grain refinement, defects creation and

diffusion channels in the 4140 (quenched and tempered) steel [23, 25-27]. This process can also

be used to modify the substrate surface on the atomic level by increasing the surface roughness

and thus improving the adherence of hard coatings to the treated areas [7, 28]. Thus, the

development of processes which allow improving the performance of tools and dies can benefit

from the understanding of pre-treating techniques for metallic substrates aiming at the

subsequent deposition of ceramic coatings with enhanced hardness and thermal stability. In this

work we report on individual and combined surface processing involving the bombardment

with Xe+ ions and plasma nitriding at low temperature, i.e. 380°C, of annealed 100Cr6 steel

substrates, as pre-treatment alternatives for the deposition of TiN hard coatings.

2. Materials and Methods

2.1. Substrate material

The substrate material was an annealed 100Cr6 bearing steel. Its chemical composition

is given Table 1. The steel specimens were employed in the annealed instead of quenched and

tempered condition in order to avoid phase transformations during processing by ion plasma

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nitriding. The samples used in all experiments were in the form of 20 mm diameter discs with

2 mm thickness. Prior to application of all surface modification techniques, the steel samples

were ground with SiC grinding paper, polished in 6, 3 and 1 µm diamond paste and finally

polished in a 0.25µm colloidal silica suspension.

Table I: Chemical composition of the 100Cr6 steel.

Element C Si Mn Cr

Composition [wt.%] 1.00 0.25 0.35 1.5

2.2. Surface modification and TiN deposition

Samples from the 100Cr6 steel were submitted to preliminary Xe+ ion bombardment

and plasma nitriding prior to TiN film deposition. The Xe+ ion bombardment/implantation and

TiN deposition processes were performed using an ion beam assisted deposition (IBAD) system

located at the Ionic Implantation and Surface Treatment Laboratory of the Campinas State

University (Unicamp, Brazil). The system consists mainly of a deposition chamber equipped

with two Kaufman ion sources of 3 cm diameter, a support for four targets and a temperature

controlled (<1000 °C) sample holder. More details regarding this system can be found in the

work of Hammer et al. [29].

For ion bombardment a xenon gas flux was introduced into the ~3 cm diameter ion beam

Kaufman sources (9.9995%) and ions produced were accelerated to 1.0 keV. The samples were

bombarded perpendicularly to the polished surface with a ~2.8 mA/cm2 current density for

30min. During Xe+ bombardment, the substrate temperature was kept below 260 °C and the

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work pressure in this case was 0.15 Pa. Plasma nitriding was performed using a 0.2 keV N+

beam with ~2.7 mA/cm2 current density focused on the 100Cr6 sample surface for 30min. The

temperature of the process was set to 380 °C and the working pressure was 0.015 Pa for all

nitrided samples. The deposition of TiN coatings was accomplished by sputtering a titanium

target with an argon ion beam of 1.45 keV (80 mA) in a N2 atmosphere. The deposition N2

pressure was 0.055 Pa and the substrate temperature was set to 400 °C. The deposition time

was 120 min for all the samples.

All those processes were carried out independently in the following manner: initial

sputter cleaning of the substrates using a 600 eV Ar+ beam for 5 minutes, followed by Xe+ ion

bombardment and/or nitrogen implantation when applicable, and finally, TiN coating

deposition. Thus, a series of four samples was produced by combining the processes described

above yielding the samples described in Table 2.

Table II: Sample identification and structure.

Sample Sample structure

S Steel + TiN

SN Steel + nitrided layer + TiN

S1000 Steel + Xe+ bombarded layer + TiN

S1000N Steel + Xe+ bombarded and nitrided layer + TiN

2.3. Phase analyses and Microstructure

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Angle-dispersive X-ray diffraction (XRD) using synchrotron radiation was carried out

at the experimental station XPD of the Brazilian Synchrotron Radiation Laboratory (LNLS,

Campinas) in order to assess the phase composition within the TiN coated samples subjected

to the different surface pre-treatment procedures. The radiation energy was set to 10.5 keV (

= 1.1823 Å) and the beam was 4 mm wide and 1 mm high. Diffractograms were obtained both

under a grazing incidence angle () of 1°, which corresponds to an average penetration depth

of 500 nm, and also using the symmetric -2 mode.

The microstructure of the 100Cr6 steel substrate and the TiN coatings was investigated

by scanning electron microscopy (SEM) using a Philips XL-30 FEG microscope. The sample

cross-sections were etched to reveal the material microstructure with 3% Nital solution.

Transmission Electron Microscopy (TEM) of sample cross-sections was used to characterize

the size and distribution of the iron nitrides formed during plasma nitriding and to visualize the

impact of atomic peening on the microstructure of the steel surface by using a JEOL 2100

microscope. Suitable thin samples were prepared applying a wedge shaped thin slice, i.e.,

polishing by gentle abrasive rubbing of a specimen with a slight tilt (0.3-0.7°) to obtain an

electron transparent wedge with smooth and highly polished surfaces (20 µm). Finally, electron

transparent areas were then obtained by low energy ion polishing.

2.4. Topography, wear and coating adhesion

The influence of each pre-treatment procedure on the surface topography of the coated

samples was evaluated by atomic force microscopy (AFM) using the tapping mode in a

FlexAFM 3 – Nanosurf equipment. To evaluate the adhesion of the coatings, an Anton Paar

Nanoscratch-Tester (NST) was employed using a linear increase of the applied normal force up

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to 500 mN. The wear resistance of the TiN coatings deposited on the different substrate surfaces

was evaluated by means of ball on disc tests. A ball of cemented carbide (WC-Co) with 3.0 mm

radius and 5 N normal load were applied to the tests. The total sliding distance was set to 1000

m.

2.5. Residual Stress Analysis

The effect of Xe bombardment and plasma nitriding on the average residual stresses of

the TiN films was determined by applying the sin2 method at a fixed penetration depth of 500

nm using a four-circle diffractometer located at the experimental station XPD of the

Synchrotron facilities (LNLS, Campinas, Brazil) [30, 31]. The radiation energy was 10.5 keV

( = 1.1823 Å) and beam dimensions were 1 x 1 mm. A fixed penetration depth was chosen in

order to suppress the influence of stress gradients, which typically occur in thin surface films,

on the stress evaluation procedure. Owing to the weak macroscopic texture of the IBAD

deposited TiN films, unrestricted lattice spacing measurements could be performed for all

sample directions defined by the inclination angle and the azimuth .

The penetration depth of the synchrotron X-rays depends on the incident () and exit

() angles of the incoming and diffracted X-rays, respectively. This can be also expressed as a

function of the diffractometer angles according to Eq.(1):

sin sin sin sin(2 ) cos

[sin sin ] [sin sin(2 )]

(1)

where is the linear absorption coefficient of TiN and , 2 and represent the

instrumental angles of a four-circle diffractometer. The azimuthal and the inclination

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angles, which define the sampled specimen directions, can be also expressed in terms of the

diffractometer angles according to Eq. (2) and (3):

arccos cos cos( ) (2)

sinarctan

tan( )

(3)

Considering Eq. (1) to (3), a sin2-based stress analysis can be carried out at a constant

penetration depth by applying the following procedure:

(I) the penetration depth of interest is chosen taking into account e.g. the film

thickness;

(II) a diffraction line (hkl) of the material under consideration is chosen, thus fixing the

diffraction angle 2 = 2hkl;

(III) a list of the azimuths and the inclination -angles has to be defined, taking

into account the symmetry of the stress state, the crystallographic texture of the material and

the required measurement accuracy;

(IV) a list of diffractometer settings (, , ) has to be determined for each specimen

direction (, ) using Eq. (1) to (3).

Thus, in order to maintain a constant penetration depth during a sin2-based stress

analysis, the traditional or tilting, which are used individually at a certain -azimuth, has

to be replaced by a coordinated movement of the diffractometer angles , and . As a

consequence of the controlled penetration depth, only a limited region within the sin2-plot is

accessible for a certain hkl reflection. For further details regarding this technique, the reader is

referred to [30].

The lattice strains were measured for the {220} diffraction line of TiN. This reflection

was chosen based on a compromise between the ease for separating this line from neighbouring

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ones, its intensity and the highest possible 2θ-angle. The Diffraction Elastic Constants (DEC)

were calculated based on the Eshelby-Kröner model using the single crystal elastic constants

for TiN [31].

3. Results and Discussion

3.1. Microstructure

The microstructure of the 100Cr6 steel subjected to Xe+ bombardment followed by

plasma nitriding and TiN deposition (sample S1000N) is presented in Figure 1. The SEM image

initially reveals that the annealed hypereutectoid steel substrate consists of a ferritic matrix with

embedded spheroidized cementite particles. It can also be observed that the average thickness

of the TiN coatings deposited by IBAD amounts to about 400 nm. Plasma nitriding conducted

at 380 °C did not lead to the formation of a white compound layer but only to a restricted

diffusion zone, which cannot be visualized using SEM analyses.

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Figure 1: Cross section micrograph of the S1000N sample showing the TiN coating (20,000X).

The results of the phase analyses conducted by XRD on the S, SN, S1000 and S1000N

specimens are presented in Figure 2. The XRD patterns of non-nitrided samples (S and S1000)

contain reflections that correspond to ferrite and Fe3C from the substrate, as well as TiN from

the deposited hard coating. For the SN sample, which was not subjected to Xe+ bombardment,

the diffractograms reveal additionally the presence of the iron-rich γ’-Fe4N nitride. With

regards to the S1000N sample, which was in turn bombarded with Xe+ prior to nitriding, both

ε-Fe2-3N and γ’-Fe4N iron nitrides could be detected. The formation of ε-Fe2-3N nitride confirms

that the previous ion bombardment was responsible for increasing nitrogen retention [26, 27]

on the steel surface during the plasma nitriding process. In addition, it is worth noticing that the

volume fraction of γ’-Fe4N appears to be larger in the S1000N specimen than in the SN sample,

as indicated by the difference in the relative intensity of the respective diffraction lines. This

probably happens because bombardment with Xe+ ions texturized the treated surface, causing

the generation of lattice defects and stress which accelerate the nitrogen diffusion into the first

atomic surface layers [26, 27]. Overall, the TiN reflections exhibit line broadening (e.g. in

comparison with the reflections which stem from metallic iron in the steel substrate), an

indication of small crystallite size and possibly nanostructured layer (i.e., disorder). In addition,

there appears to be no strong preferential orientation in the TiN coatings deposited on the

different substrates since in all cases several TiN reflections could be detected with no marked

difference in their relative intensities, in agreement with a recent investigation [33].

The microstructure of the S1000N specimen was further investigated by TEM (Figures

3 and 4). Observation of Figure 3 indicates elevated dislocation density in the sub-surface

region of the steel substrate (approximately 0.5μm thickness), likely caused by plastic strain

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associated with the Xe+ ion bombardment. Possible sources of deformation caused by ion

bombardment include: formation of gas bubbles from within the substrate lattice, phase

transformations and the peening effect itself caused by ion collision [21, 24].

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Figure 2: XRD spectra and phase identification for the S, SN, S1000 and S1000N samples.

The presence of lattice defects accounts for the higher diffusivity necessary for the

formation of the nitrogen-rich ε-Fe2-3N nitride layer detected by XRD (Figure 2). Figure 4

displays a TEM image of the coating/substrate interface in S1000N sample. The interface

between the steel substrate and TiN coating exhibit fringes due to very fine nitride formation

(red arrows). Analyses of the electron diffraction data from the coating-steel interface indicate

two distinct superposing patterns, corresponding to the -Fe2-3N and γ’-Fe4N nitrides, as

indicated in the inset of Figure 4, thus corroborating the results of the XRD analysis presented

in Figure 2

Figure 3: TEM image from the S1000N sample showing increased defect density near the interface with

the TiN coating.

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Figure 4: TEM image and electron diffraction pattern from the sample S1000N showing the

presence of -Fe2-3N and γ’-Fe4N nitrides.

3.2. Topography, wear and coating adhesion

The surface topography of the samples investigated by AFM is displayed in Figure 5. It

is possible to verify that the different types of pre-treatment lead to marked differences in the

surface roughness of the TiN films. The sample S exhibits the lowest levels of roughness,

compatible with the surface finishing of the metallographic preparation to which the substrates

were initially subjected. Thus, the non-treated sample represents the reference state for the

purpose of comparison. Plasma nitriding of the substrate (sample SN) leads to a small increase

in roughness, which can be explained by different mechanisms such as: volumetric expansion

of the substrate lattice due to incorporation of N in solid solution and formation of nitrides,

sputtering of the surface due to the impact of N ions during deposition, as well as redeposition

of the sputtered material [34, 35].

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Figure 5: AFM surface topography of the investigated samples obtained after wear testing.

The bombardment with Xe+ ions at 1.0 keV kinetic energy produces a more significant

increase of roughness in the sample S1000 than the nitriding process. This occurs due to a more

intense atomic attrition prompted by the heavy Xe+ ions. As previously shown reduction the

Xe+ bombardment kinetic energy diminishes surface roughness [22, 26]. Finally, the exposure

of the 100Cr6 steel surface to prolonged atomic attrition caused by the combination of ion

bombardment and plasma nitriding leads to the roughest surface profile in the sample S1000N.

The evolution of the average surface roughness of the TiN coatings as a function of the pre-

treatment technique applied to the 100Cr6 substrates is presented in Figure 6.

The average worn volumes obtained after ball on disc tests are presented in Figure 7.

The lowest wear rates were observed for the samples subjected to individual treatments of

nitriding (SN) or ion bombardment (S1000) instead of a combination of both plasma nitriding

and Xe+ ion bombardment (S1000N).

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Figure 6: Evolution of average roughness Ra as a function of surface pre-treatment in TiN coating.

This indicates that the bombardment with Xe+ ions causes strain hardening of the steel

surface that is comparable with the hardness increase promoted by nitride formation, or possibly

a combination of surface hardening with slightly increased roughness that contributes to

enhance the coating adhesion to the substrate. Both effects, i.e. increased substrate hardness

and coating adhesion, ought to be considered here because in all cases the thin TiN coatings

(~400 nm, Figure 1) from all samples appear to have been completely removed during the wear

tests. The wear behavior of the S1000N specimen was found to be comparable to the untreated

sample S. In this case, excessive surface roughness causes insufficient TiN adhesion on the

substrate. As the tribological system investigated (cemented carbide against TiN) involves

materials of elevated hardness, the main wear mechanism appears to be abrasive, thus leading

to appreciable breakup of the sharper coating asperities.

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Figure 8 exhibits the scratch scars and the critical loads for coating delamination

obtained during the nano-scratch tests. It is possible to observe chevron cracking, indicating

cohesive failure; the dashed rectangles highlight the start of chipping failure (adhesive failure).

These values indicate a transition point in the wear behavior, which can occur due to

plastic deformation, increase in temperature [36] or initial failure of the coating (microcracking,

surface flaking, loss of cohesion) [37]. In the present case, the critical load is an indicative of

adhesion of the TiN coating to the substrate material. The highest critical load was observed on

the SN sample, consistent with the minimum wear volume observed in Figure 7. The same trend

is observed for the remaining samples, within the error margins of the experiment.

Figure 7: Average wear volume of the different samples after ball-on-disc tests.

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a)

b)

Figure 8: a) Image of the scratch scars and b) critical load for TiN coatings deposited onto the different

substrates.

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3.3. Residual stress analyses

Residual stress in thin coatings produced by deposition can usually be ascribed to two

sources. On one hand, the energetic ions cause atoms from the substrate to be shifted from their

original lattice positions producing residual interstitials. This leads to an outwards expansion

of the thin film from the substrate. In the plane of the coating, however, expansion is restricted

and the entrapped atoms cause elevated compressive stress [38]. In addition to these intrinsic

stresses, additional residual stresses arise during cooling from the deposition temperature due

to incompatibility of thermal expansion between substrate and coating [39, 40]. These cooling

stresses are therefore influenced by the surface pre-treatments which can alter the phase

composition and the average thermal expansion coefficient near the interface with the hard

metal coating.

The sin²ψ versus lattice spacing plots are presented in Figure 9. As can be observed for

all samples, residual stresses are compressive and the {200} lattice spacing varies linearly with

sin²ψ, in agreement with previous investigations [41, 42]. It is also important to observe that

XRD measurements for residual stress analysis on all materials could be performed over a large

range of sample tilt positions, indicating that no strong preferential orientation is present in TiN

coatings, as expected from the results presented in Figure 2.

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Figure 9: sin²ψ versus TiN (200) lattice spacing plots of the investigated substrates.

The residual stress values obtained for each coating procedure are presented in Figure

10. In all cases, stress values are highly compressive: (-6.25±0.2), (-4.55±0.05), (-6.77±0.12)

and (-5.95±0.1) GPa for the S, SN, S1000 and S1000N specimens, respectively. By comparing

the results obtained for samples S and S1000, it is possible to observe that Xe+ bombardment

alone does not have a significant effect on the residual stress state on the deposited TiN coating.

However, plasma nitriding caused a significant reduction in the residual stresses. In

duplex coatings, where a hard ceramic layer is deposited onto a nitrided substrate, a transition

zone appears between substrate and coating [37, 43]. This happens because of the high ion

energy during IBAD that promotes the transference of chemical elements across the

substrate/coating interface, leading to the formation of a Ti-rich interlayer [33]. In the nitrided

sample, this would lead to the reaction with nitrogen and an increase in adhesion and chemical

compatibility between the TiN coating and substrate. This also contributes to reduce the residual

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stresses in the TiN coating. For similar reasons, the SN sample exhibited superior wear

characteristics (Figures 7 and 8) in comparison to the other tested materials.

Figure 10: Average residual stress values in the TiN as a function of surface pre-treatment.

The application of Xe+ bombardment prior to plasma nitriding causes an increase in the

residual stresses of the TiN film to values similar to those of the untreated substrate in the S

sample. This is because of the occurrence of -Fe2-3N that has a Thermal Expansion Coefficient

(TEC) greater than ’-Fe4N: TEC(’-Fe4N) < TEC(-Fe2-3N) < TEC(Fe).

4. Conclusions

The impact of individual and combined pre-treatments based on the bombardment with

1.0 keV Xe+ ions and plasma nitriding at 380°C on the adhesion and wear of TiN coatings

deposited by IBAD onto spheroidized 100Cr6 steel was investigated. The results showed that

the bombardment with Xe+ ions causes grain refinement, texturization, and increases the density

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of lattice imperfections at the treated surface, thus enhancing nitrogen diffusion at 380°C and

probably causing strain-hardening as well.

The combined pre-treatment led therefore to a diffusion zone containing both ε-Fe2-3N

and γ’-Fe4N nitrides, whereas simple plasma nitriding promoted the formation of only γ’-Fe4N.

The presence of ε-Fe2-3N confirms an enhanced nitrogen retention in the diffusion zone and the

formation of a more expressive nitride fraction. This corroborates the positive influence of Xe+

ion bombardment on the nitrogen retention during nitriding treatments.

Superior wear behavior was however verified for the application of individual plasma

nitriding. Xe+ ion bombardment with 1.0 keV impact energy in combination with plasma

nitriding caused sputtering of substrate and significantly increased surface roughness of the TiN

film. In such condition, the wear rate of the TiN became similar to that observed in the substrate

without pre-treatment. Reduction of the Xe+ ion energy shall lead to less sputtering and surface

roughness. This might enhance the wear performance in the case of a combined pre-treatment

due to a more pronounced nitrogen retention.

All TiN coatings evolved compressive residual stresses, ranging from -4.55 and -6.77

GPa. The lowest residual stress values were observed for the sample subjected to individual

plasma nitriding due to the formation of only γ’-Fe4N with reduced TEC and the increased

chemical compatibility by a Ti-rich interlayer between substrate and coating during IBAD.

Acknowledgments

The authors HP and JG acknowledge the financial support of the NanoCom Network

(Proposal FP7 no 247524) and the CREATe-Network project (H2020-MSCA-RISE/644013),

both funded by the European Union. SV is grateful to Conselho Nacional de Desenvolvimento

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Científico e Tecnológico (CNPq) and Comissão de Aperfeiçoamento de Pessoal do Nível

Superior (CAPES) for granting the scholarship for acting in this research. FA and HP are CNPq

fellows. Part of this work was supported by FAPESP, projects 2010/11391-2 and 2012/10127-

5. The Brazilian Nanotechnology National Laboratory (LNNano) is greatly acknowledged for

support during the TEM analyses.

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