Tribology International 61 (2013) 109–115

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Small-scale structural and mechanical characterization of the nitrided layerin martensitic steel

Masoud Asgari a, Afrooz Barnoush a,b,n, Roy Johnsen a, Rune Hoel c

a NTNU, Norwegian University of Science and Technology, Faculty of Engineering Science and Technology, Department of Engineering Design and Materials,

NO-7465 Trondheim, Norwayb Saarland University, Department of Materials Science, Building. D22, P.O. Box 151150, D-66041 Saarbruecken, Germanyc MOTecH Plasma, Oslo, Norway

a r t i c l e i n f o

Article history:

Received 18 September 2012

Received in revised form

5 December 2012

Accepted 6 December 2012Available online 20 December 2012

Keywords:

Nitriding

Nanoindentation

EBSD

XRD

9X/$ - see front matter & 2012 Elsevier Ltd. A

x.doi.org/10.1016/j.triboint.2012.12.004

esponding author at: NTNU, Norwegian Univ

ogy, Faculty of Engineering Science and

ring Design and Materials, NO-7465 Trondhe

7 735 93807; fax: þ47 735 94129.

ail address: afrooz.barnoush@ntnu.no (A. Barn

a b s t r a c t

Pulsed Plasma Nitriding (PPN) of high-strength low-alloy steels used for offshore applications is a

promising approach for controlling erosion, corrosion and hydrogen embrittlement under service

conditions. In this work, the microstructure, composition and hardness of the nitride layer produced by

an optimized PPN process on 2.25Cr–1Mo steel were examined. The nanomechanical properties of the

nitride layer were investigated via nanoindentation along the depth of the nitride layer to understand

the interconnected effect of the existing microstructure with the one developed after the nitriding

process and the nitrogen concentration. The results showed that the nitride layer is composed of a

compound layer and diffusion layer with hardness four times higher than the untreated material, which

gradually decreases across the diffusion layer.

& 2012 Elsevier Ltd. All rights reserved.

1. Introduction

The petroleum industry is currently using the forging processto build heavy components, such as blow-out preventer bodies,fluid ends for mud pumps and hydraulic fracturing pumps, andvarious other parts of well heads [1]. Usually, low-alloy steels,such as AISI 4130 or 2.25Cr–1Mo (F22) grades, are used for theseforgings in the hardened condition to achieve very high mechan-ical properties, typically on the order of 75 ksi (520 MPa) for theminimum yield strength in accordance with API 16A1 and NACEMR0175/ISO 15156-2.2 standards. The harsh service environmentof these parts in addition to their high strength and toughness,demands a very high resistances against wear, corrosion andhydrogen embrittlement. In particular, in offshore applicationsdue to the cathodic corrosion protection on the outside andcontact of H2S containing oil with the interior of these parts, therisk of hydrogen uptake and hydrogen-induced cracking in thesehigh-strength steels is very high. The advantageous effect ofnitriding on the wear resistance of steels is well known [2,3].Recently, by the application of in-situ electrochemical nanoin-dentation [4], it has been shown that nitriding can also improve

ll rights reserved.

ersity of Science and

Technology, Department of

im, Norway.

oush).

the resistance against hydrogen embrittlement [5]. Therefore,nitriding of these parts could be a very promising way to increaseservice life and to reduce the significant costs of repair andreplacement, as well as to reduce the risk of environmentalpollution as a result of an unexpected failure of these parts.However, the complicated microstructure of high-strength low-alloy steels such as 2.25Cr–1Mo and the complex interaction ofthe dissolved nitrogen with the microstructure and the alloyingelements make the nitriding of these alloys very complicated.Hence, a thorough characterization of the resultant microstruc-ture after nitriding is necessary before evaluation of the effec-tiveness of the treatment for protecting against hydrogenembrittlement. In this paper, we use different methods for themechanical and microstructural characterization of 2.25Cr–1Mosteel nitrided according to an optimized pulsed plasma nitriding(PPN) procedure developed by the MOTecH Plasma Company inNorway.

2. Experimental

2.1. Materials

Samples with a thickness of 1 mm and a diameter of 10.6 mmwere cut from 2.25Cr–1Mo steel in the quenched-temperedcondition [6] with a composition given in Table 1. An optimizedPPN process developed by the MOTecH Plasma was used fornitriding of the samples. Prior to PPN, the sample surfaces were

Table 1Chemical composition of 2.25Cr–1Mo used in this study.

C Si Mn Cr Ni Mo Cu V

wt% 0.14 0.18 0.54 2.45 0.28 1.1 0.16 0.012

at% 0.651 0.359 0.548 2.628 0.267 0.639 0.141 0.013

0 100 200 300 4000

5

10

15

Distance (m)

20

Fig. 1. Cross-sectional OM micrograph obtained perpendicular to the direction of

nanohardness testing from the PPN 2.25Cr–1Mo sample with a E20 mm com-

pound layer thickness and E400 mm diffusion layer, the change in the nitrogen

atomic ratio from the edge, is overlaid on the micrograph 200� magnification,

etched in 2% nital.

M. Asgari et al. / Tribology International 61 (2013) 109–115110

ground, then mechanically polished up to a 0.1 mm diamondslurry and finally electropolished at room temperature in a 1 Mmethanolic H2SO4 solution [7].

PPN was carried out in a vacuum furnace with pressuresranging from 45 Pa to 300 Pa. During heating and prior to surfacehardening, a negative potential of 600 V was applied to the testparts to create a hydrogen glow discharge within the vacuumretort. Under such conditions, the surfaces are cleaned andactivated prior to the actual surface hardening. The PPN processwas carried out at a temperature above 500 1C in dilutednitrogen-rich plasma containing more than 66 wt% nitrogen.

2.2. Characterization methods

2.2.1. Microstructural characterization

A nitrided sample was cut from the middle and prepared as across-section by subsequent embedding, grinding and polishingdown to 0.04 mm colloidal silica. The cross-section was examinedvia light optical microscopy using an Olympus BX51 optical micro-scope after etching in 2% nital solution (2% Nitric acid in Methanol).The same etched cross-section as well as samples that have beenpolished down to a certain depth revealing the nitride layer atdifferent depths were studied by a Zeiss Supra 55VP scanningelectron microscope (SEM). These subsurface regions of the nitridelayer were made by mechanical polishing of three identical nitridedsamples for 4, 8 and 20 min. The corresponding revealed depthswere estimated by observation of the change in a Vickers indentmade prior to mechanical polishing. The depths were estimated tobe approximately 5 mm, 10 mm and 17 mm below the surface for 4,8 and 20 min of mechanical polishing, respectively. The resultingsamples are called throughout the text according to the estimateddepth below the nitride layer, i.e., ‘‘5 mm below PPN’’, ‘‘10 mm belowPPN’’ and ‘‘17 mm below PPN’’.

To determine the composition of the nitrided layer, electronprobe micro-analysis (EPMA) was performed, employing a JEOLJXA-8500 EPMA, on the cross-section at intervals of 1 mm fromthe surface. At each depth at least four measurements were made,and the mean values of these measurements are reported.

Additionally, the ‘‘5 mm below PPN’’, ‘‘10 mm below PPN’’ and‘‘17 mm below PPN’’ samples were characterized by high resolu-tion electron backscatter diffraction (EBSD) using the Zeiss SupraSEM and a NORDIF UF750 camera. Data were recorded andanalyzed using the TSL OIM Analysis software. Tempered mar-tensite was indexed as body-centered cubic ferrite. To index theFe3N (e) and Fe4N (g0) phases, these phases are defined using thelattice constants and phase space groups reported in literature.The hexagonal close-packed e phase belongs to the P63/mmc

space group and has the nominal lattice parameters ofa¼2.529 A and c¼4.107 A [8]. The face-centered cubic g0 phasebelongs to the Pm3m space group and has a lattice parameter ofa¼3.798 A [8].

To determine which phases were present after nitriding, XRDwas performed using a Siemens D5005D X-ray diffractometer.Measurements were made using Cu Ka (l¼0.154 nm) radiation at40 kV and employing the Bragg–Brentano geometry with agraphite monochromator in the diffracted beam. The range ofthe diffraction angle (2y) was from 351 to 751, with a step size of0.041. X-ray diffractograms were recorded from the surface of the

‘‘5 mm below PPN’’, ‘‘10 mm below PPN’’ and ‘‘17 mm below PPN’’samples. To identify the phases from the positions of the diffrac-tion peaks, data from the EVA-Bruker-AXS database were used.

2.2.2. Nanomechanical characterization

The nanomechanical properties of the PPN layers were exam-ined by means of a Hysitron TI 750 Ubi

TM

scanning nanoindenta-tion system with a performech

TM

control unit. Tests wereperformed on the cross-section of the sample and the surface ofthe samples after PPN and consecutive mechanical polishing ofthe surface. All indents were made by a Berkovich diamond tip,which was used to image the surface of the sample before andafter indentation with the in-situ imaging option of the nanoin-dentation system. The hardness and elastic modulus were calcu-lated from the load–displacement (L–D) curves according to theOliver–Pharr method [9]. Indents on fused quartz were used tocalibrate the tip area function.

3. Results and discussion

3.1. Microstructural and chemical analyses

3.1.1. LM, SEM and EPMA

A detailed micro-structural analysis of this steel grade hasbeen performed previously and is available in the literature[6,10,11]. Fig. 1 shows an optical micrograph of the PPN samplecross-section with an overlay of the nitrogen concentration fromthe surface into the bulk as measured by EPMA. There is a brightregion within the first 20–30 mm of Fig. 1 that corresponds to ahigh N concentration between 22 and 17 at%. This region isfollowed by a gradual change in the color from dark to brightwithin 300–400 mm of the diffusion layer where the N concentra-tion gradually changes from approximately 5 at% to zero. Fig. 2shows an SEM micrograph of the first 40 mm of the nitride layer athigher magnification. The nitrogen concentration from the surfaceinto the bulk as measured by EPMA is also overlain on the SEMmicrograph in Fig. 2.

Considering the presence of Cr and Mo with relatively highaffinity for N, thermodynamically it is plausible to assume theformation of MoN and CrN precipitates through the dissolution ofMoC and CrC precipitates. In Fig. 3, the difference between themeasured N concentration and the sum of the Cr and Moconcentrations is presented. It can be concluded from Fig. 3 thatwithin the first �250 mm of the diffusion layer all the N is

10 15 20 25

Nitr

ogen

Ato

mic

Rat

io

0

5

10

15

20

25

0 5Distance from surface (µm)

Fig. 2. Cross-sectional SEM micrograph obtained perpendicular to the nitrogen-

diffusion direction from the PPN 2.25Cr–1Mo sample with a 20 mm compound

layer thickness and the change in the nitrogen atomic ratio from the edge, 3k�

magnification. The 5 mm below PPN, 10 mm below PPN and 17 mm below PPN

sample regions obtained after 4, 8 and 20 min of mechanical polishing are also

colored (red, blue and green, respectively). (For interpretation of the references to

color in this figure legend, the reader is referred to the web version of this article.)

Fig. 3. The difference between the N concentration and the sum of the Cr and Mo

concentrations.

M. Asgari et al. / Tribology International 61 (2013) 109–115 111

precipitated in the form of nitrides of the alloying elements andall the carbides of the alloying elements are dissolved and thatcarbon is released in the matrix. After �250 mm, the total Navailable in the matrix is again in the form of nitrides of thealloying elements but the N concentration is not enough tocombine with all the available alloying elements. Therefore, amixture of Mo and Cr nitrides and carbides are present in thediffusion layer after �250 mm of depth.

3.1.2. EBSD analysis

EBSD analysis was performed in the marked region of the SEMimage shown in Fig. 4a. This SEM image was taken with back-scattered electron (BE) contrast from the nitrided layer. The phasemap from the EBSD analysis of the selected area in Fig. 4bcorresponds perfectly with the BE image of the SEM but with abetter contrast showing the grain boundaries and the change inthe microstructure from the surface into the bulk material.The map clearly shows that the compound layer formed on thesurface mainly consists of two regions, i.e., an outer region thatmainly consists of columnar grains and an inner part that consistsof equiaxed grains.

Fig. 4b reveals that the outer region with columnar grainsmainly consists of the e phase while the inner region partlycontains the g0 phase. This g0 phase is distributed within thee grains with an slight increase in their distribution close to thesubstrate.

The EBSD map from the cross-section PPN sample shows thatduring the PPN treatment two different layers were formed in thecompound layer. From the surface down to 10–15 mm, thecompound layer consists of columnar e crystals. After this region,

the rest of the compound layer consists of equiaxed e grains withg0 islands.

The reliability of the indexed Kikuchi patterns can be judgedon the basis of the confidence index (CI) resulting from the fit ofthe measured pattern with the calculated one from the materialdata. As mentioned before, we defined the material data for boththe g0 and e files. The resulting CI maps for each phase show arelatively large portion of the points in which each phase isindexed with a high CI. This result confirms that both themeasurement and analysis of the phases are reliable and alsoshows how EBSD analysis in SEM can successfully be applied tostudy the microstructure of the nitride layer in steels and providehigh-resolution information on the phases formed, which waspreviously attainable mainly through cumbersome TEM analysis[12,13].

3.1.3. XRD analysis

The X-ray diffractograms recorded from the surfaces of theuntreated sample, the 5 mm below PPN, 10 mm below PPN and17 mm below PPN samples are shown in Fig. 5. The XRD pattern ofthe untreated sample showing peaks of a (110) and a (200)proves that the untreated sample is mainly composed of ferrite, asexpected for the tempered martensitic microstructure of thespecimen. In contrast, in the X-ray diffractograms recorded fromthe samples at various depths of PPN, different peaks belonging tothe e and g0 phases are observable. In the 5 mm below PPNsample, peaks with relatively low intensity were detected thatare attributed to MoN, CrN and CrO. As Lepienski et al. [14] haveshown, due to the presence of residual oxygen in the plasmachamber the formation of oxide during the nitriding process ispossible. In addition, TEM studies by Salas et al. [15] haveconfirmed the fine distribution of CrN precipitates near thesurface. In the samples at 10 and 17 mm below PPN, no otherphases were detected, except for e and g0. This finding is in goodagreement with the results of the EBSD analysis shown in Fig. 4b.There is also a gradual increase in the intensity of the g0 peak fromthe 5 mm below PPN sample in comparison to the 17 mm belowPPN sample. This result again correlates perfectly with the EBSDphase map shown in Fig. 4b, in which the increase in the presenceof the g0 grains at lower depths can be observed. The correlationbetween the EBSD phase map and the XRD results should beperformed very carefully and with consideration for the probedvolume by the XRD and the EBSD measurements. The area fromwhich an EBSD pattern is acquired with an electron beam focusedon a 701 tilted sample is approximately elliptical, and the majoraxis, which is perpendicular to the tilt axis, is approximately threetimes the length of the minor axis. It is a function of the material,beam accelerating voltage, specimen tilt and probe size, and theresolution parallel to the tilt axis. For a-iron with a field emissiongun SEM the small axis of the ellipsoid is not more than 10 nmand the penetration depth could also be considered equal to thisvalue. The step size in the phase map shown in Fig. 4b is 100 nm,i.e., if any precipitates smaller than 100 nm exists in the micro-structure it is not possible to resolve them. However, the depth ofpenetration of the XRD and, therefore, the volume probed duringXRD depends on various parameters, including the incident angle,the density of the material analyzed and the energy of theincident X-rays. Luo and Liu [16] have shown that the densityof iron nitride and iron are the same. We used a density of iron of7.87 g/cm3 and 8.048 keV (X-ray energy) to calculate the pene-tration depth. These depths are approximately 1.586 mm for2y¼401 and 3.279 mm for 2y¼901. This finding in fact showshow in a complementary manner XRD and EBSD analysis can beused to analyze globally and locally the presence of the phases inthe nitride layer.

Fig. 4. (a) Cross-sectional SEM image from the EBSD area, (b) grain boundry and phase map from the cross-section PPN sample, with the phase-relevant color-code map,

qFe (a) phase (red), e phase (yellow) and g0 phase (green). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this

article.)

M. Asgari et al. / Tribology International 61 (2013) 109–115112

3.1.4. Nanoindentation

To measure the variations in the mechanical properties as afunction of the microstructure and nitrogen concentration, nanoin-dentation tests were performed on untreated samples, cross-sections of the PPN samples and the samples 5, 10 and 17 mmbelow PPN. Nanoindentation is an extremely surface sensitivetechnique, and the data analysis is based on the presence of aperfectly flat surface. Therefore, the surface preparation is a criticalstep prior to nanoindentation. The imaging capability of thenanoindentation system was used to examine the surface rough-ness. The arithmetical mean roughness (Ra) and root mean squareroughness (RMS) over an area of 64 mm2 were measured usingGwyddion [17], open-source atomic force microscopy (AFM) ana-lysis software, and the results are reported in Table 2.

The resulting L–D curves after nanoindentation of theuntreated sample with different maximum loads are shown inFig. 6. The nanomechanical behavior of the 5, 10 and 17 mm belowPPN samples in the form of L–D curves are shown in Fig. 7.

According to the Oliver–Pharr [9] method, it is possible toextract the hardness (H) and reduced elastic modulus (Er) fromthe L–D curves of Figs. 6 and 7. The Oliver–Pharr analysis is basedon elastic isotropic calculations given by the following equations:

H¼Pmax

AC, ð1Þ

Er ¼

ffiffiffiffipp

2

1

bSffiffiffiffiffiffiAC

p , ð2Þ

1

Er¼

1�n21

E1þ

1�n22

E2, ð3Þ

where Pmax is the maximum applied load, Ac is the pro-

jected contact area evaluated from the contact depth (hc)and the tip area function, S is the slope of the L–D curve atthe initial unloading and b is a correction factor dependingon the tip geometry (1.034 for a Berkovich indenter).E1 and E2 are the elastic module of the material and tip,respectively, and n1 and n2 are the Poisson ratios of the materialand tip, respectively.

The resulting hardness values for each sample at differentmaximum loads are summarized in Fig. 8. As expected, thehardness of the pristine material after the PPN treatmentincreased approximately three times, but the most importantpoint in Fig. 8 is the observed gradient in the hardness within thenitride layer. The highest value of the hardness is measuredwithin 17 mm below the PPN layer. The same trend is observedfor the measured Er as shown in Fig. 9. For the sake of comparison,both the expected Er values for iron and e [18] are shown in Fig. 9.The deviation of the measured Er value for the untreated samplefrom the reported value for iron (Fig. 9) as well as its increasingtrend with increasing indentation depth can be explained by theoverestimation of the contact area as a result of the pile-up [9].Fig. 10 shows a typical topography height normalized to theindentation depth for an untreated sample where the obviouseffect of the pile-up can be seen. The same uncertainity in the

Fig. 5. XRD results from the untreated sample and various depths of the

compound layer.

Table 2Samples surface roughness before indentation.

Sample Ra (Sa) (nm) RMS (Sq) (nm)

Untreated 0.84 1.11

5 mm below the PPN surface 1.17 1.8

10 mm below the PPN surface 1.28 1.84

17 mm below the PPN surface 1.28 1.72

Fig. 6. L–D curves of the untreated 2.25Cr–1Mo samples at different

maximum loads.

Fig. 7. L–D curves with a 8 mN load and 4 mN s�1 loading rate on 2.25Cr–1Mo

after PPN at depths of 5 mm, 10 mm and 17 mm below the surface.

Fig. 8. Hardness versus contact depth measured at depths of 5 mm, 10 mm and

17 mm below the surface and in the untreated 2.25Cr–1Mo samples.

Fig. 9. Er versus contact depth measured at depths of 5 mm, 10 mm and 17 mm

below the surface and in the untreated 2.25Cr–1Mo samples. The expected values

of the Er for iron and e [21] are also shown.

M. Asgari et al. / Tribology International 61 (2013) 109–115 113

estimation of the contact area overlaid on the indentation sizeeffect can explain the trend in the hardness values of theuntreated sample.

Fig. 10. Topography height normalized to the maximum indentation depth for

evaluation of the pile up effect in an indent made by 5 mN maximum load into the

untreated sample. Maximum indentation depth was 210 nm.

Fig. 11. Reduced modulus, hardness of the compound across the depth of the

layer after PPN treatment.

M. Asgari et al. / Tribology International 61 (2013) 109–115114

According to Fig. 5, the Er measurements at depths of 5 mm and10 mm below the surface are all performed in the region withcolumnar microstructure, while the nanoindentations on 17 mmbelow PPN sample are representative of the nanomechanicalproperties of the region with the equiaxed e grains and g0 islands.Therefore, the variations in both hardness and Er are related to themicrostructural changes, increasing presence of the g0 phase,decarburization from the surface and residual stresses [19,20].

In addition to the indentation tests performed directly on thesurface of the 5, 10 and 17 mm below PPN samples, the cross-section of the PPN sample was also characterized by nanoinden-tation for very high-resolution nanomechanical characterizationof the diffusion zone. A 50�50 mm2 area was positioned asclosely as possible to the sample surface by in-situ imaging. Then,a matrix of 26�5 indents was made across the PPN layer in thisregion. Over a length of 200 mm farther into the material, 52�5indents were made at intervals of 4 mm. Finally, the next 150 mmwere nano-mechanically tested every 6 mm by application of27�50 indents. All these indents were made with a maximumload of 5 mN.

The resulting hardness profile together with the change inelastic modulus and nitrogen concentration are shown in Fig. 11.Both the hardness H and elastic modulus Er increase within thefirst 20 mm in accordance with the measurements at depths of5 mm, 10 mm and 17 mm below the sample surfaces shown inFigs. 8 and 9. Within the first 200 mm of the diffusion layer, thehardness gradually decreases while the nitrogen concentrationshows a very small reduction (Fig. 3). Interestingly, Er shows asmall increase within this first region of the diffusion layer. Atapproximately 200 mm of depth, there is a sharp reduction in bothH and Er while only the N concentration gradually decreases.Afterwards, while Er remains relatively constant, both hardnessand N concentration decreases. Explanation of this observationrequires a consideration of nitrogen diffusion during PPN andseveral reactions that occur simultaneously in the diffusion zone.These reactions include saturation of ferrite with N, precipitationof metal nitrides, carbon redistribution and the generation of theresidual stress [22]. In the studied 2.25Cr–1Mo steel, as shown inFig. 3, the chemical analysis of N, Cr and Mo shows that we canexpect all the nitrogen atoms in the diffusion zone to combine

with chromium and molybdenum to form CrN and MoN aspredicted thermodynamically [22]. Consequently, the CrC andMoC precipitates existing in the pristine material are dissolved,and carbon is released into the ferrite matrix. As a result ofprecipitation of nitrides and the release of carbon atoms into thematrix, a compressive residual stress develops in the diffusionlayer due to the volume change, as reported by Jegou et al. [23].This compressive stress causes carbon redistribution. The carbonatoms initially in the base metal will diffuse to the stress-freeregions, i.e., towards the surface and the nitriding front, leading todecarburization of the compound layer and the formation of acarbon-rich region in the nitriding front. The hardness profileshown in Fig. 11 results from a combination of three position-dependent factors, namely: (i) the hardening effect of fine CrNand MoN precipitates, which is a function of their quantity, sizeand distribution, (ii) the change in the hardness due to theredistribution of the carbon, and (iii) error in measurements ofthe contact area and overestimation of the hardness by nanoin-dentation due to the presence of the residual stress [24–26].The error in estimation of the contact area due to the presence ofthe residual stresses also affects the Er measurements [24–26].The sharp reduction in the Er observed at approximately 200 mmdepth seems to be due to the presence of the maximum com-pressive residual stress existing at this point as measured byJegou et al. [23] for a similar steel and nitriding condition.However, further examination is required to understand thecomplicated interaction between the residual stress, N concen-tration, nitride precipitation, carbon redistribution and mechan-ical properties of the nitride layer.

4. Conclusions

PPN-treated 2.25Cr–1Mo high-strength low-alloy steel wascharacterized using different techniques, the results of whichwere linked to understand the position-dependent nanomecha-nical properties and microstructure of the nitride layer. Thefindings are as follows:

1.

The nitride layer is composed of two regions: a compoundlayer with a thickness of approximately 20–30 mm and adiffusion layer that extends approximately 400 mm into thebase material.

2.

Within the compound layer, two regions can be distinguished:an outer region mainly composed of columnar epsilon grains,which is approximately 10–15 mm and an inner region con-sisting of equiaxed e grains with g0 islands, which has the

M. Asgari et al. / Tribology International 61 (2013) 109–115 115

highest mechanical properties (HE13 GPa). From the innerregion into the outer columnar region, the hardness decreases(HE11 GPa). The reduced Young’s modulus within the com-pound layer shows the same trend. This trend in elastic andplastic properties is due to the presence of residual stressesand decarburization from the surface.

3.

The diffusion layer is also composed of two regions: the innerpart within the first 200 mm where the N concentrationremains constant and an outer region where the N concentra-tion gradually decreases. In the inner region, N mainly com-bines with the alloying elements to form sub-micrometer CrNand MoN precipitates. This precipitation introduces highresidual stresses and releases the C that had previouslycombined with the alloying elements in the form of carbideprecipitates. These effects result in a complex variation of thehardness and Young’s modulus within the diffusion layer.

Acknowledgments

The results presented in this paper were generated as part ofthe JIP ‘‘Pulsed plasma surface treatment (PPST)—effect onhydrogen embrittlement and hydrogen absorption and diffusion’’executed by the Department of Engineering Science and Materi-als, NTNU, with financial and technical support from Total E&PNorge, Aker Solutions and Motech plasma Material and Over-flateknologi (MOTecH). MOTecH is the owner of the applied PPSTprocess and delivered all the project’s treated samples.The authors would like to thank all of the participating companiesfor their fruitful technical discussions and for permission topublish the results from the JIP.

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