Surface & Coatings Technology 202 (2008) 4957–4963

Contents lists available at ScienceDirect

Surface & Coatings Technology

j ourna l homepage: www.e lsev ie r.com/ locate /sur fcoat

Low-temperature nitriding of 38CrMoAl steel with a nanostructured surface layerinduced by surface mechanical attrition treatment

W.P. Tong a,b,⁎, Z. Han a, L.M. Wang b, J. Lu c, K. Lu a

a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR Chinab Key Laboratory of Electromagnetic Processing of Materials, Ministry of Education, Northeastern University, Shenyang 110004, PR Chinac Department of Mechanical Engineering, the Hong Kong Polytechnic University, Hong Kong, China

⁎ Corresponding author. Key Laboratory of NationalNortheastern University, Shenyang 110004, PR China. Tel83681758.

E-mail address: wptong@mail.neu.edu.cn (W.P. Tong

0257-8972/$ – see front matter © 2008 Elsevier B.V. Aldoi:10.1016/j.surfcoat.2008.04.085

A B S T R A C T

A R T I C L E I N F O

Article history:

A nanocrystalline surface la Received 25 November 2007Accepted in revised form 24 April 2008Available online 1 May 2008

Keywords:NitridingNanocrystalline materials38CrMoAl steelSurface mechanical attrition treatmentWear resistance

yer was fabricated on a 38CrMoAl steel plate by means of a surface mechanicalattrition treatment (SMAT). The average grain size in the top surface layer (10 μm thick) is about 10 nm, andthe grain size stability can be maintained up to 450 °C. The effect of the surface nanocrystalline layer on thegas nitriding process at a lower temperature was investigated by using structural analysis and wear propertymeasurements. The surface nanocrystallization evidently enhances nitriding kinetics and promotes theformation of an ultrafine polycrystalline compound layer. The results of the investigation showed that thisnew gas nitriding technique can effectively increase the hardness and wear resistance of the resulting surfacelayer in comparison with conventional nitriding, demonstrating a significant advancement for materialsprocessing.

© 2008 Elsevier B.V. All rights reserved.

1. Introduction

Gas nitriding is a thermochemical treatment widely used to formsurface nitrides. This technique is of great industrial interest as itforms a unique composite structure with a hard surface (nitridecompound layer) and a tough interior, so that the global mechanicalperformance and wear/corrosion resistance of alloys and steels aregreatly improved [1,2]. In recent years, new and innovative surfaceengineering technologies have been developed rapidly to meet ever-increasing demands arising from some extreme applications; but gasnitriding is one of themost widely used surface engineering techniquein industry today [3,4]. However, conventional nitriding processesperformed at high temperatures (above 500 °C) for a long duration(60–80 h) are one of the main energy consumption sources inmanufacturing industries [1, 2]. In addition, relatively high processingtemperatures promote the occurrence of porosity, distortion and thereduction of material properties. Therefore, reducing treatmenttemperatures and enhancing gas nitriding processes present a majorchallenge to surface engineers and researchers, serving as a drivingforce for the development of nitriding technologies.

The thickness of the nitrided layer developed in thermochemicaltreatment processes depends on nitriding conditions [5,6] and on theproperties of thematerial itself, notably on the reactions occurring at thesurface and the nitrogen diffusivity [7,8]. Both factors are strongly

Education Ministry for EPM,.: +86 24 83682376; fax: +86 24

).

l rights reserved.

affected by grain boundaries and defect densities. Nanocrystallinematerials possess ultrafine grains with a large number of grainboundaries that may act as fast atomic diffusion channels [9,10]. Manyenhanced atomic diffusivities in nanocrystalline materials (relative totheir conventional coarse-grained counterparts) have been experimen-tally observed [11,12]. On the other hand, a large number of grainboundaries with various kinds of non-equilibrium defects also con-stitute a high excess stored energy that may further facilitate theirchemical reactivity. It has been experimentally demonstrated thatchemical reaction (or phase transformation) kinetics are greatlyenhanced during the mechanical attrition of solids in which the grainsize is significantly reduced into the nm scale and structural defects arecreated due to severe plastic deformation [13–15]. Hence, we believe achange in the surfacemicrostructure bymeans of grain refinement is analternative option to accelerate the gas nitriding process.

A recently developed technique, the surface mechanical attritiontreatment (SMAT) [16,17], seems to be appropriate for loweringnitriding temperatures (or shortening the duration) efficiently. UsingSMAT, different microstructures can be obtained within the deformedsurface layer through the depth from the treated surface to the strain-free matrix, i.e., from nano-sized grains to submicro-sized and micro-sized crystallites on a bulk metal. The surface nanocrystalline layer isfree of contamination and porosity because the nanocrystallizationprocess is induced by severe plastic deformation at very high strainrates. Previous investigations showed that the gaseous nitriding of apure iron plate [18] and the plasma nitriding of an AISI 304 stainlesssteel specimen [19] were recently achieved at 300 °C and 400 °C,respectively, with nanostructured surface layers induced by theSMAT. In this paper, a typical nitriding of a 38CrMoAl steel plate was

Fig. 1. A cross-sectional SEM observation of the SMAT 38CrMoAl sample.

Fig. 2. A dark-field TEM image and the corresponding electron diffraction pattern (inset)of the top surface layer for the SMAT sample.

4958 W.P. Tong et al. / Surface & Coatings Technology 202 (2008) 4957–4963

employed to investigate the low temperature nitriding process with ananostructured surface layer produced by means of SMAT. Themicrostructural characterization and property measurements of thenitrided SMAT samplewere examined and comparedwith those in theoriginal coarse-grained sample. The experimental observations clearlyshow that surface nanocrystallization (SNC) greatly facilitates thenitriding process and provides an alternate approach to the surfacemodification of metallic materials.

2. Experimental

2.1. Sample preparation

The rawmaterial used in the present investigationwas a 38CrMoAltypical nitriding steel plate of 7 mm×100 mm×100 mm, withchemical compositions of (wt.%) 0.37 C, 0.32 Mn, 1.80 Cr, 0.20 Mo,0.98 Al and bal. Fe. The sample was subjected to SMAT processing inorder to achieve a nanostructured surface layer. Prior to treatment, the38CrMoAl steel sample was hardened (oil quenched from 940 °C) andtempered (550 °C for 2 h) to produce the desired mechanicalproperties (with a hardness of 290 Hv). The SMAT procedure wasperformed on an apparatus in which steel shots (8 mm in diameter)vibrated by a generator with a frequency of 50 Hz were peening ontothe sample surface. The experimental set-up and the details of theSMAT procedure were previously described in the literature [16,17].The sample surface was protected by a high-purity argon atmosphereduring SMAT processing to avoid oxidation. Repeated attritionsinduced severe plastic deformation in the surface layer and eventuallyresulted in the development of nc grains. The nanocrystallizationmechanism in the top surface layer is analogous to that during severeplastic deformation of bulk materials where the microstructure canalso be refined into nanometer size [20,21]. After the SMAT treatmentfor 60 min, the surface roughness is comparable to that of the originalsample.

Both untreated and treated samples were cut to 7 mm×10 mm×10 mm plates, washed carefully using acetone and absolutealcohol, then immediately subjected to gaseous nitriding in a home-made device. Nitriding was performed in a flowing high-purityammonia gas (NH3, 99.95%). Further nitriding parameters were asfollows: nitrogen pressure 1 atm, temperature 400 °C, treatmentduration 30 h. Both specimens were placed perpendicular to the worktable of themachine using special holders to ensure uniform nitriding.

2.2. Microstructural characterization

X-ray diffraction (XRD) analysis was used to characterize theaverage crystallite size and the mean microstrain of the surface layerof the SMAT sample and to identify phases formed during nitridingtreatments. The XRD analysis was carried out on a Rigaku D/max 2400

X-ray diffractometer (12 kW), with Cu Kα radiation. Small angularsteps of 2θ=0.02° were taken to measure the intensity of each Braggdiffraction peak. The counting time of 20 s was used to exactlymeasure the width of the diffraction peak in the step-scanning mode.The average grain size and meanmicrostrainwere calculated from theline broadening of bcc Fe (110), (200), (211), (220), (310) and (222)Bragg diffraction peaks using the Scherrer and Wilson method [22].

Cross-sectional observations of the treated 38CrMoAl steel samplewere performed on a JSM-6301F scanning electron microscope (SEM).Microstructural features in the surface layer were characterized byusing a Philip EM-420 transmission electron microscope (TEM),operating at a voltage of 120 kV. Thin foil samples for TEMobservations were prepared by cutting, grinding and dimpling witha final ion thinning at low temperatures.

The microhardness was measured using a MVK-H300 microhard-ness testing machine with a load of 50 g and a time of 10 s.Microhardness values were obtained by means of the dependence ofthe residual penetration depth of the indenter (calculated from thegeometry of the Vickers pyramid import) on the applied load. Weartests were evaluated using a SRV reciprocating sliding machine(Optimal, Germany) under dry pure slide conditions. The tests wereperformed at room temperature and atmospheric pressure using aWCballing of 10 mm diameter at a sliding speed 0.1 m/s under a normalload of 50 N.

3. Results and discussion

3.1. Microstructure of the SMAT sample

Fig. 1 shows a cross-sectional SEM observation of the SMAT38CrMoAl steel sample. Evidence of severe plastic deformation (ofabout 60–70 μm thick) is present in the treated surface layer, wheremicrostructural morphology differs from that in the matrix. Fig. 2shows a typical TEM observation of the top surface layer (about 1 μmdeep) in the SMAT 38CrMoAl steel sample. The microstructure ischaracterized by ultrafine equiaxed grains (bcc Fe) with randomcrystallographic orientations, as indicated by the selected areaelectron diffraction (SAED) pattern. The mean grain size in the surfacelayer is approximately 10 nm.

By removal of the surface layer-by-layer, structure evolution alongthe depth can be determined. The measured average grain size as afunction of depth in the SMAT sample using different analysistechniques (XRD and TEM) is summarized in Fig. 3. At a depth ofabout 30–60 μm from the treated surface, the average size of theequiaxed grains is about 100–600 nm. These grains possessed ran-dom crystallographic orientations and were separated by sharp

Fig. 3. Variations in the grain size and the mean microstrain with the depth from thetreated surface of the SMAT sample determined by means of XRD analysis, TEM andSEM observations.

Fig. 5. Cross-sectional observations of an original coarse-grained sample (A) and theSMAT 38CrMoAl sample (B) after nitriding at 400 °C for 30 h.

4959W.P. Tong et al. / Surface & Coatings Technology 202 (2008) 4957–4963

conventional boundaries. With a further increase in the depth, muchlarger subgrains and dislocation cells were observed. At a depth of80 μm, the size of the equiaxed subgrains is about 1 μm. We noticedevidence of dislocation activities induced by the attrition treatment inan even deeper matrix (up to 200–300 μm deep), where no change ingrain size and atomic-level microstrain is detected. The microstraindetermined from the XRD analysis decreases significantly along thedepth in the surface layer and gradually drops to zero at about 80 μm.Detailed microstructural characterization of the surface layer in theas-treated 38CrMoAl steel sample, together with an underlyingnanocrystallization mechanism during the SMAT treatment in termsof the severe plastic deformation mode of bulk materials, is analogousto that in [17].

In order to determine the experimental conditions (temperatureand time) for gas nitriding in the nanocrystalline 38CrMoAl steelsample, the thermal stability of the nanostructured surface layeragainst grain growth was investigated. An isothermal annealingtreatment of the SMAT samples was carried out in an Ar atmosphere.The isothermal annealing procedure is accomplished by heating theSMAT samples to 450 °C, at a heating rate of 10 K/min and holding atthis temperature for 30 h before rapidly cooling to room temperature.After isothermal annealing at 450 °C for 30 h, a slight grain growthoccurs in the surface layer, as shown in Fig. 4, where grain sizes of 20–

Fig. 4. A dark-field TEM image and the corresponding electron diffraction patterns(inset) of the top surface layer for the SMAT sample annealed at 450 °C for 30 h.

40 nm are detected. The XRD analysis indicated that the average grainsize of the annealed sample is approximately 40 nm.

3.2. Microstructure of nitrided layer

After nitriding, the cross-sections of the microstructures for bothuntreated and treated samples were observed by SEM. The nitridedlayers produced at 400 °C were characterized by two differentmorphologies (see Fig. 5). These morphologies include a graycompound layer (upper part) and a diffusion zone (lower part)induced by the precipitation of nitrides. The observations for bothsamples after nitriding show that SMAT processes had a significantinfluence on the kinetics of the compound layer formation. Theoriginal 38CrMoAl steel sample (A) after nitriding has no continuouscompound layer, and only a partial nitride layer can be observed.However, for the SMAT sample (B) under the same conditions, acontinuous gray surface compound layer of about 20–30 μm thick wasobserved, showing that the formation of nitrides in the surface nclayer is faster than that in the coarse-grained sample. It also should benoted that the thickness of the compound layer in the SMAT sample isnot uniform, indicating preferential diffusion of nitrogen into thematerial along grain boundaries and dislocations.

The X-ray diffraction analysis in Fig. 6 provides information onabout 5 μm thick phase composition from the nitrided surface layer,The surface layer in the original 38CrMoAl nitrided sample containsstrong diffraction peaks for the α-Fe phase, as well as weak diffractionpeaks of ɛ-Fe2–3N and γ′-Fe4N compounds, whereas that of the SMATsample nitrided under the same conditions consists of strong ε-Fe2–3Nand γ'-Fe4N. The diffraction peaks of the ε-Fe2–3N and γ'-Fe4N phasesobviously increase and those of the α-Fe phase disappear, indicating asubstantial enhancement of the volume fraction of nitrides in thesurface layer. XRD analysis simultaneously confirmed the presence ofCr2N in the SMAT sample after nitriding, which differs from theoriginal Fe sample where no Cr2N was detected.

TEM observations combined with electron diffraction analysesshow that the compound layer for the nitrided SMAT sample iscomposed of ultrafine polycrystalline iron nitrides and alloy nitrides(see Fig. 7). Those nitrides include ε-Fe2–3N and γ'-Fe4N phases, aswell as a small amount of Cr2N, which are in agreement with the XRDanalysis results. Compared to the conventional nitriding of 38CrMoAlsteel, some nitrides were not formed due to the low mobility ofnitrogen and alloy elements in the material at 400 °C. The compoundlayer can produce amuchmore complex structurewith various nitride

Fig. 8. A TEM image of the microstructure beneath the compound layer (about 40 μmdeep from the top surface). Indicated by arrows are γ′ phase particles, formed at grainboundaries and junctions ofα-Fe phase, as confirmed by the electron diffraction pattern(inset).

Fig. 6. X-ray diffraction patterns of a SMAT sample and an original coarse-grainedsample nitrided at 400 °C for 30 h.

4960 W.P. Tong et al. / Surface & Coatings Technology 202 (2008) 4957–4963

phases only by increasing the nitriding temperature to 500 °C andhigher [23]. The temperature for Cr2N formation reaches even higherthan 550 °C. A detailed description of the nitrided structure for theoriginal 38CrMoAl steel sample in this temperature range has beengiven in the literature [24]. Therefore, the constitution of thiscompound layer for the SMAT sample nitrided at 400 °C is similar tothat formed during conventional nitriding at high temperatures(about 550 °C). The formation of Cr2N for the nitrided SMAT samplewas observed in the electron diffraction pattern and XRD analysis,suggesting that the diffusion ability of Cr at 400 °C in the grain-refinedmaterial is sufficient for Cr2N formation. No Aluminium nitrideparticles were detected in the nitrided SMAT sample. This is becausethat the chemical composition of Al was smaller than that of Cr, and itis known that AlN with a more complex crystal structure can onlynucleate under a high driving force [25]. Probably, the driving force ofnitriding at 400 °C was not enough for AlN forming.

The diffusion zone of the nitrided SMAT sample (Fig. 8) wasinvestigated by TEM. Underneath the compound layer, a large amountof rod-like and particle-like nitrides (about 100–300 nm) wereprecipitated. The electron diffraction pattern from these nitrides wasfound to correspond to the γ'-Fe4N (cubic) phase. Clearly, thesenitrides (γ'-Fe4N) formed largely at grain boundaries (or junctions) ofthe ultrafine-grained α-Fe phase, dislocations and other defects, andthe lower nitriding temperature limits the growth of precipitates [18].With the increase in depth, the volume fraction of these nitrides in thediffusion zone decreases gradually.

The above experimental results clearly illustrate that the SMATtechnique can significantly speed up the lower temperature gasnitriding process. This phenomenon has also been observed by sev-eral other investigations, in both gas nitriding and plasma nitriding

Fig. 7. A dark-field TEM image and the corresponding electron diffraction pattern for thecompound layer of a SMAT sample nitrided at 400 °C for 30 h.

[18,19,26]. In recent works it has already been demonstrated that thenitrogen activity of nitriding atmosphere, surface roughness as well ascrystalline structure of the substrate influence the nitriding kinetics.Nanocrystalline nature [27], high surface roughness [28] beforenitriding and nitrogen potential [29,30] (the combination of hydrogenand ammonia) are very effective in improving the growth of nitridinglayer. As the surface roughness of SMAT sample was comparable tothat of the original sample, and both samples were nitrided under thesame atmosphere condition, the present study therefore supports theconclusion that the surface nanocrystallization of SMAT sampleaccelerated the nitriding process.

The growth of the nitriding layer is a diffusion-controlled process(of nitrogen). The nitrided layer thickness of 38CrMoAl steel can bedescribed by [5]

n2 ¼ 2rN½ �X½ �Dt

which is derived from diffusion theories accompanied by a phasechange which has been applied to internal oxidation, where, [N] is thesurface nitrogen concentration, [X] is the original alloy element in thenitride phase, r is the ratio of nitrogen to alloy element in the nitridephase, D is diffusivity of nitrogen, t is nitriding time.

Clearly, in the nitriding of conventional coarse-grained Fe, nitrogendiffusion in the lattice dominates, where the kinetics is related totemperature. For the nc 38CrMoAl steel sample, however, nitrogenmay preferably diffuse at low temperatures along numerous grainboundaries with a larger enhanced diffusivity compared with lat-tice diffusion because there is smaller activation energy at grainboundaries (approximately half that of the lattice diffusion) [18].Therefore, the nc structure of 38CrMoAl steel in the SMAT sampleprovides a large number of fast diffusion channels (grain boundaries)for nitrogen at low temperatures which may significantly enhance thenitriding kinetics.

The microstructural characterization of the SMAT 38CrMoAl steelsamples indicated a high concentration of the non-equilibrium defects(dislocations and subgrain boundaries) in the sub-surface layer thatare induced by plastic deformation. These defects may decrease theactivation energy of diffusion and act as fast atomic transfer channelsas well [31]. Similar kinetics enhancement can be accomplished forthe diffusion zones obtained in the nitriding process.

Furthermore, the formation of nitrides is a process of nucleationand growth. Grain boundaries are preferable heterogeneous nuclea-tion sites for the ε-Fe2–3N and γ'-Fe4N phases, for which the activationenergy barrier is much smaller than that for homogeneous volume

Fig. 10. Variations in the wear volume loss with duration for different samples asindicated.

4961W.P. Tong et al. / Surface & Coatings Technology 202 (2008) 4957–4963

nucleation in the α-phase lattice [32,33]. According to References[34,35], for nc materials with large amounts of grain boundaries,nucleation site densities will evidently increase, i.e., measurably highnucleation rates can be obtained with very small driving forces. Forexample, the heterogeneous nucleation rate at grain boundaries mayincrease by an order of 104 when the grain size is reduced from100 μmto 10 nm. Therefore, the enhanced formation ability of a compoundlayer in the SMAT sample can be attributed to a larger increase in theheterogeneous nucleation rate at numerous grain boundaries.

A very high heterogeneous nucleation rate was the primary reasonwhy the compound layer for the SMAT sample consists of nanos-tructured nitrides. A continuous compound layer formed by abundantheterogeneous nucleation will not provide a diffusion barrier fornitrogen due to a special ultrafine polycrystalline structure, whichhelps to increase the thickness of the diffusion zone.

3.3. Hardness and wear resistance

The absolute hardnesses of the surface layer in the nitrided SMATsample have been determined in this work. In the case of the SMATsample nitrided for 30 h, there is a peak hardness of 1340 Hv at thesurface, and the hardness decreases slowly with a longer distancefrom the surface (Fig. 9). The diffusion zone depth (defined by thedistance from the surface to the point where the microhardness was50 Hv higher than that in the matrix) was 200 μm. For the original38CrMoAl steel sample nitrided under the same conditions, thehardness is only 1015 Hv at the surface and decreases rapidly withlonger distances from the surface; the diffusion zone depth is about50–60 μm. The hardness values found for the SMAT sample are a factorof 30% higher than those of conventional 38CrMoAl steel afterindustrial nitriding at a high temperature (nitrided at 550 °C and ahardness of about 1000–1100 Hv). This greater hardening in the outerpart of the nitrided layer is probably due to the formation of thenanostructured compound layer and Cr2N precipitation. It is wellknown from previous work that nanocrystalline materials can greatlyenhance mechanical properties compared with their traditionalcoarse-grained polycrystalline counterparts due to the grain sizeeffect. According to the Hall–Petch relationship [36–38], the yieldstress of crystalline material and grain sizes have the followingrelationship:

rs ¼ ri þ ky � d�1=2

where σs is the yield stress, σi is the friction resistance to dislocationmotion in single crystal, ky is the coefficient of the stress concentration

Fig. 9. Variations in the microhardness along the depth from the top surface layer for aSMAT sample and an original sample nitrided at 400 °C for 30 h.

in the grain interface, and d is the grain size; σi and ky are independentof the grain sizes. The grain sizes in the nanostructured compoundlayer for the SMAT sample were about 10–50 nm. The average grainsize in the compound layer (conventional 38CrMoAl steel afterindustrial nitriding at a high temperature) was of about micronmagnitude and normally formed a pillar structure [39]. Therefore, thenanostructured compound layer has amuch higher hardness than thatof conventional 38CrMoAl steel.

On the other hand, the greatly improved hardness of the nitridedSMAT sample can be attributed to the appearance of fine alloy nitrideCr2N. Precipitate formation is inconsistent with the diffusivity of Cratoms in Fe. Hence, there was no sign of Cr2N formation at 400 °C forthe conventional industrial nitriding of 38CrMoAl steel because thediffusivity of Cr atoms was too low at this temperature. For thenitrided SMAT sample, however, the fine Cr2N precipitates are alreadypresent due to the fast diffusion of the Cr atoms in the nanostructuredcrystalline layer. The formation of Cr nitride precipitates at grainboundaries has been reported in the literature [40,41]. The segrega-tion of Cr at grain boundaries will be enhanced in the nc material dueto a larger grain boundary area and the shorter distance a Cr atom hasto travel to diffuse to a grain boundary, even though diffusion of Cr isoccurring at 400 °C. So Cr2N precipitation could be determined by fastgrain boundary diffusion. The size of individual precipitates will bevery small due to an increased grain boundary area and a limitedvolume fraction of Cr2N which has to be distributed over the grainboundaries. Compared with 38CrMoAl steel after conventionalindustrial nitriding (at 550 °C), the particle sizes of Cr nitrides werenormally larger or formed a lamellar structure due to the enhanceddiffusion of Cr atoms at a high temperature [24]. Cr2N has very hardparticles, so the appearance of this phase in the compound layer canlead to increased hardness. The fine dispersive precipitate of Cr2N(or other alloy nitrides) in a low temperature treatment is responsiblefor the increased hardening response in the nitrided SMAT sample.

High hardness was obtained in the first 100 μm of the diffusionzone for the nitrided SMAT sample. This phenomenon can beexplained by the precipitation of nitrides at grain boundaries, dis-locations and other defects in the α phase. The low nitridingtemperature limits the growth of precipitates, resulting in finedispersions of the γ'-Fe4N phase in the diffusion zone. Togetherwith ultrafine crystallites and the high defect density in the subsurfacelayer, the precipitation produces a greater hardness in the diffusionzone.

Sliding wear test results showed that gas nitriding in theinvestigated temperature range can significantly increase the wearresistance of the materials. Fig. 10 summarizes the results for various

Fig. 11. SEM observations of the wear scars formed at the applied load of 50 N in a nitrided SMAT sample subjected to abrasion for 10 min (A) and 30 min (B) and the nitrided originalsample abraded for 10 min (C) and 30 min (D).

4962 W.P. Tong et al. / Surface & Coatings Technology 202 (2008) 4957–4963

38CrMoAl steel samples. For both nitrided samples, wear volume lossincreases with a longer duration, but the wear of the nitrided SMATsample is lower than that of the original nitrided sample during eachsliding time, especially between 30–60 min, indicating the wearresistance of the nitrided 38CrMoAl steel sample has been improvedby SMAT nitriding processes.

The morphology of the worn surface is shown in Fig. 11. Whensliding time was less than 20 min, both nitrided samples shared asimilar wear mechanism. The wear loss results mainly from plowingandmicro-cutting under the action of abrasive particles such as the tipof the WC ball and oxides derived from the sample surface [42–44].From the morphology of the wear scar under these conditions, thereare many parallel grooves formed by the plowing action and micro-cutting. But the worn surface of the original nitrided sample was veryrough with deep, broad grooves as well as spalling (Fig. 11C), whereasthe worn surface of the nitrided SMAT sample was relatively smoothwith some narrow grooves (Fig. 11A). When the sliding time isincreased (30–60 min), for the original nitrided sample, the WC ballindentsmore deeply into the surface, resulting in awearing through ofthe hardened surface layer (compound layer). Thus the repetitivework-hardening in the new surface layer will be more severe under arepetitive sliding action. Eventually, propagating and growing cracksinduced spalling of thematerial in the newsurface layer. The dominantwearmechanism causes plastic removal and surface fatigue fracture ofthe deformed layer [45,46]. Fig. 11D shows cracks and pits after thesurface spalled off (marked by arrow). This wear mechanism is similarto that of the original annealed 38CrMoAl steel sample. The apparentvariation in the wear rate for the original sample nitrided under asliding timeof about 30min (in Fig.10)might result from such a changein the wear mechanism. Compared with the nitrided SMAT sample(Fig. 11B), the wear mechanism tends to a steady state during all testprocesses due to a very thick hardened surface layer (compound layer).

This improved abrasion resistance for the nitrided SMAT samplecan be attributed mainly to the higher hardness and case depth(especially the higher effective case depth) because wear resistance isgenerally related to the hardness ratio. An ultrafine polycrystallinecompound layer (grain size of about 30–60 nm) which formed during

the gas nitriding process is also believed to contribute to the higherabrasive resistance, since wear resistance is associated with the grainsize of the material surface layer.

4. Summary

In this paper it is shown that a nanostructured surface layer on a38CrMoAl steel sample can be obtained by a surface mechanicalattrition treatment (SMAT). The subsequent gas nitriding kinetics ofthe SMAT sample with the nanostructured surface layer was found tobe greatly enhanced. Effective nitriding can be performed attemperatures as low as 400 °C, which is much lower than conven-tional nitriding temperatures. The nitrided layer in the SMAT sample iscomposed of nanostructured nitrides, resulting in relatively higherhardness and wear resistance as compared to those of the nitridedcoarse-grained sample. The greatly reduced nitriding temperature forSMAT samples can be attributed to nano-scale-sized grains with alarge amount of defects such as grain boundaries. These defects notonly provided more nitride nucleation sites, but also greatly enhancedthe atomic diffusion which facilitates the nitriding process.

The remarkable depression in the nitriding temperature by meansof the surface nanocrystallization demonstrates a significant advance-ment for materials processing and provides a new approach forselective surface reactions in solids.

Acknowledgments

Financial support from the National Science Foundation of China(50571023), the 111 Project (B07015), Foundation for the Author ofNational Excellent Doctoral Dissertation of PR China (200745),Program for New Century Excellent Talents in University and ResearchFund for the Doctoral Program of Higher Education (20060145022).

References

[1] T. Bell, Source Book on Nitriding, ASM, Metals Park, OH, 1977, p. 266.[2] J. Mongis, J.P. Peyre, C. Tournier, Heat Treat. Met. 3 (1984) 71.

4963W.P. Tong et al. / Surface & Coatings Technology 202 (2008) 4957–4963

[3] M.P. Fewell, J.M. Priest, M.J. Baldwin, G.A. Collins, K.T. Short, Surf. Coat. Technol. 131(2000) 284.

[4] A. Mucha, M. Braun, Surf. Coat. Technol. 50 (1992) 135.[5] B.J. Lightfoot, D.H. Jack, Heat Treatment '73. Book No. 163, the Metals Society,

London, 1975, p. 59.[6] E. Lehrer, Z. Electrochem. 36 (1930) 383.[7] H. Ferkel, M. Glatzer, Y. Estrin, R.Z. Valiev, Scr. Metall. 46 (2002) 623.[8] J. Baranowska, M. Wysiecki, Surf. Coat. Technol. 125 (2000) 30.[9] H. Gleiter, Prog. Mater. Sci. 33 (1989) 223 and references therein.[10] K. Lu, Mater. Sci. Eng. R 16 (1996) 161.[11] S. Schumacher, R. Birringer, R. Strauss, H. Gleiter, Acta Metall. 37 (1989) 2485.[12] L. Lu, M.L. Sui, K. Lu, Science 287 (2000) 1463.[13] H.J. Fecht, in: G.C. Hadjipanayis, R.W. Siegel (Eds.), Nanophase Materials, Kluwer,

the Netherlands, 1994.[14] C.C. Koch, Nanostruct. Mater. 2 (1993) 1.[15] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1.[16] K. Lu, J. Lu, J. Mater. Sci. Technol. 15 (1999) 193.[17] N.R. Tao, Z.B. Wang, W.P. Tong, M.L. Sui, J. Lu, K. Lu, Acta Mater. 50 (2002) 4603.[18] W.P. Tong, N.R. Tao, Z.B. Wang, J. Lu, K. Lu, Science 299 (2003) 686.[19] H.W. Zhang, L. Wang, Z.K. Hei, G. Liu, J. Lu, K. Lu, Z. Met. kd 94 (2003) 1143.[20] N. Hansen, M. Etall, Metall. Mater. Trans. 32A (2001) 2917.[21] R.Z. Valiev, R.K. Islamgaliev, I.V. Alexandrov, Prog. Mater. Sci. 45 (2000) 103.[22] H.P. Klug, L.E. Alexander, X-ray Diffraction Procedures for Polycrystalline and

Amorphous Materials, Wiley, New York, 1974, p. 661.[23] Y. Sun, T. Bell, Mater. Sci. Eng. 140A (1991) 419.[24] K.H. Jack, nitriding, Heat Treatment '73. Book No. 163, The Metals Society, London,

1975, p. 39.[25] H.H. Podgurski, H.E. Knechtel, Trans. AIME 245 (1969) 1595.

[26] W.P. Tong, N.R. Tao, Z.B. Wang, H.W. Zhang, J. Lu, K. Lu, Scripta Mater. 50 (2004)647.

[27] M. Wohlschlögel, U. Welzel, E.J. Mittemeijer, Appl. Phys. Lett. 91 (2007) 141901.[28] G.P. Singh, J. Alphonsa, P.K. Barhai, P.A. Rayjada, P.M. Raole, S. Muherjee, Surf. Coat.

Technol. 200 (2006) 5807.[29] J.G. Buijnsters, P. Shankar, J. Sietsma, J.J. ter Meulen, Mater. Sci. Eng. A 341 (2003)

289.[30] S.S. Hosmani, R.E. Schacherl, E.J. Mittemeijer, Mater. Sci. Technol. 21 (2005) 113.[31] H. Ferkel, M. Glatzer, Y. Estrin, R.Z. Valiev, Mater. Sci. Eng. A348 (2003) 100.[32] O. Salas, U. Figueroa, J. Bernal, L. Oseguera, Surf. Coat Technol. 163 (2003) 339.[33] H. Du, N.L. Ågren, Surf. Eng. 11 (4) (1995) 301.[34] D.A. Porter, K. E. Easterling, Phase Transformations in Metals and Alloys,

Chapman & Hall, London SEI 8HN, UK, 1992, p. 277.[35] W.P. Tong, C.S. He, J.C. He, L. Zuo, N.R. Tao, Z.B. Wang, Appl. Phys. Lett. 89 (2006)

021918.[36] E.O. Hall, Proc. Phys. Soc. Lond. B 64 (1951) 747.[37] N.J. Petch, Iron Steel Inst. 174 (1953) 25.[38] Z.N. Farhat, Y.O. Ding, A.T. Northwood, Alpas. Mater. Sci. Eng. A 206 (1996) 302.[39] M.A. J Somers, Heat Treat. Met. 4 (2000) 92.[40] N.G. Chechenin, et al., Metall. Mater. Trans. 33A (2002) 3075.[41] P.C.V. Wiggen, H.C.F. Rozendaal, E.J. Mittemeijer, J. Mater. Sci. 20 (1985) 4561.[42] D.H. Jeong, F. Gonzalez, G. Palumbo, K.T. Aust, U. Erb, Scripta Mater. 44 (2001) 493.[43] K. Jia, T.E. Fischer, Wear 266 (1997) 203.[44] I. Iliuc, M. Jokl, Wear 176 (1994) 73.[45] D.K. Chaudhuri, D. Xie, A.N. Lakshmanan, Wear 209 (1997) 140.[46] K. Jia, T.E. Fischer, Wear 200 (1996) 206.