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Grain size effects on the austenitization processin a nanostructured ferritic steel

L.M. Wang, Z.B. Wang ⇑, K. Lu ⇑

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China

Received 23 December 2010; received in revised form 25 February 2011; accepted 5 March 2011

Abstract

A surface layer with a depth-dependent microstructure was produced on a ferritic steel (P92) plate by means of surface mechanicalattrition treatment (SMAT). The austenitization processes of ferrite and carbides in the surface layers with different average grain sizeswere investigated by means of in situ X-ray diffraction, differential scanning calorimetry and transmission electron microscopy. Exper-imental results showed that the onset temperature of the austenitization process decreases gradually with decreasing sizes of ferrite grainsand carbide particles, being �120 K lower in the top SMAT surface layer compared with that in the original sample. In addition, the two-step austenitization process in the surface layers becomes a one-step one when the mean size of carbide particles is smaller than 20 nm.The effects of microstructure refinement on the accelerated austenitization processes were discussed in terms of thermodynamic andkinetic.� 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Nanostructured; Surface mechanical attrition treatment; Ferritic steel; Grain size effects; Austenitization process

1. Introduction

Solid-state phase transformations in metallic materials,especially in steels, are a central topic in physical metal-lurgy because of a combination of fundamental scientificinterests and technological importance [1–4]. Among phasetransformations in steels, the on-heating formation of face-centered cubic (fcc) austenite (c) from body-centered cubic(bcc) ferrite (a) matrix has been studied extensively andsome investigations have dealt with the effects of alloyingelements and initial microstructures on the austenitizationprocess [5–13]. Experimental works on low-alloy steelshave indicated that cementite particles provide nucleationsites for austenite, and austenite formation is very rapidat high temperatures [6,7,12]. The case is more complicatedin high-alloy steels because the carbides in the initial

structure are thermodynamically more stable than cement-ite [14]. For example, Lenel and Honeycombe [10] observedthat nucleation of austenite in an Fe–10Cr–0.2C (wt.%)steel is relatively sluggish while growth of austenite is rapid,and the dissolution of carbides occurs in the austenite butnot in the ferrite. By studying the austenite formationand the carbide dissolution in Fe–8.2Cr–C alloys with dif-ferent C concentrations, Shtansky et al. [11] noticed thatthe mechanisms of austenite nucleation and growth dependon the composition, starting microstructure and austenitiz-ing temperature.

Nanostructured materials have attracted intensive inter-est for several decades, due to their novel properties origi-nating from a large volume fraction of interfaces [15–19].The on-heating evolutions, such as dislocation recovery,grain growth and carbide precipitation, as well as the resul-tant mechanical properties, have been studied in nano-structured or ultrafine-grained ferritic steels [20–22].However, to the authors’ knowledge, there is no study onthe austenitization process of nanostructured ferrite

1359-6454/$36.00 � 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

doi:10.1016/j.actamat.2011.03.006

⇑ Corresponding authors. Tel.: +86 24 2397 1508; fax: +86 24 2399 8660.E-mail addresses: zbwang@imr.ac.cn (Z.B. Wang), lu@imr.ac.cn (K.

Lu).

www.elsevier.com/locate/actamat

Available online at www.sciencedirect.com

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matrix. This might be related to the fact that significantgrowth of nanosized grains may occur upon heating priorto the phase transformation temperature being reached.At ambient temperature, c-Fe has been experimentallyobserved in nanostructured Fe when the grain size is smallenough [23,24], and it was suggested to be thermodynami-cally stable by calculating the Gibbs free energies of inter-faces in nanostructured c and a grains [25]. Therefore,notable grain size effects on the austenitization behaviorsin ferritic steels might be expected.

By means of surface mechanical attrition treatment(SMAT), surface layers with a gradient grain size distribu-tion (ranging from nanometers, submicrons to microns)have been synthesized on various metallic materials[26–34] as a result of gradient variations of applied strainsand strain rates with depth from the treated surface. Previ-ous studies showed that the nanostructures with enhancedCr diffusivity in the SMAT surface layers of low-carbonsteel [35] and H13 steel [36] are effectively stabilized bythe formation of fine dispersive Cr compound particlesduring chromizing at 873 K, resulting in the growth ofmuch thicker chromized surface layers than that on thecoarse-grained samples after chromizing treatments at tem-peratures as high as 1323 K. This means that the gradientmicrostructure of the SMAT surface layer on steels is sta-ble at elevated temperatures with dispersive precipitates.Such a kind of gradient nanostructured surface layer pro-vides a unique opportunity to study the grain size effecton austenitization behavior on the nanometer scale inone sample.

In this work, a nanostructured surface layer with adepth-dependent microstructure is synthesized on a com-mercial ferritic steel plate by means of SMAT. Thermal sta-bility and austenitization process in the surface layer arecharacterized with respect to the microstructure.

2. Experimental

2.1. Sample preparation

The studied ferritic steel (P92) was supplied by WymanGordon Forgings Inc., with the chemical composition (inwt.%) of 0.11 C, 8.75 Cr, 0.40 Mo, 1.75 W, 0.18 V, 0.05Nb, 0.05 N and balance Fe. The initial material was inan austenitized and tempered condition (1 h at 1343 K fol-lowed by 2.5 h at 1048 K). The plate sample(100 � 50 � 4.0 mm3 in size) of the as-received steel wassubmitted to SMAT, the set-up and procedure of whichhave been described previously [26–28]. In brief, a largenumber of hardened steel balls of 8 mm diameter wereplaced at the bottom of a cylinder-shaped chamber andvibrated at a high frequency by a generator. The sampleto be treated was fixed at the upper side of the chamberand impacted by flying balls repeatedly and multidirection-ally. Because the sample surface was plastically deformedwith high strains and high strain rates, grains in the surfacelayer were effectively refined, and a depth-dependent

microstructure was consequently generated. In this work,the sample was treated for 60 min in vacuum at ambienttemperature at a vibrating frequency of 50 Hz. No detect-able contamination was introduced into the surface layerduring the SMAT process.

2.2. Microstructure characterization

The microstructure evolutions with depth in both the as-SMAT and the annealed-SMAT samples were character-ized using a JEM-2010 transmission electron microscope(TEM) operated at a voltage of 200 kV. TEM foils of thetopmost layers were cut by the electro-spark dischargetechnique, then mechanically polished, dimpled and finallyion-milled from the untreated side. In addition, TEM foilsof different subsurface layers were cut, mechanically pol-ished, dimpled and finally electropolished from theuntreated side of the SMAT samples after the removal ofsurface layers of different thicknesses. The electropolishingwas carried out at 253 K with an electrolyte of 5 vol.% per-chloric acid and 95 vol.% alcohol. A short-period millingprocess at a low angle (at 4–5� for �10 min) from bothsides was applied to clean the foil surfaces before theTEM observation. The grain/cell sizes were averaged froma few hundred grains/cells selected randomly from TEMimages. In the case that not enough grains could becounted in a TEM image at the greater depths, an FEINova-nano scanning electron microscope (SEM) with alower magnification was applied to observe the microstruc-ture at the corresponding region of a cross-sectionalsample.

The phase constitution of the original sample and theSMAT surface layer at room temperature was identifiedby X-ray diffraction (XRD) analysis using a RigakuD/max 2400 X-ray diffractometer (7.5 kW) with Cu Karadiation, with a step size of 0.02�.

2.3. Phase transformation measurements

The surface layers at different depths of the SMAT sam-ple were cut by electro-spark discharge technique and thenmechanically polished from the untreated side to �10 lmin thickness. Thermal analyses of the prepared foil samples(�15 mg in weight) were conducted on a Netzsch differen-tial scanning calorimeter (DSC 404C). The experimentswere carried out from room temperature to 1373 K at aheating rate of 20 K min�1, in a flowing Ar atmospherewith a gas flow rate of 50 ml min�1. The temperature wascalibrated by the melting points of pure In, Au and Ni.In addition, the baseline in the region of each peak wasconstructed by a polynomial (or an apparent transforma-tion function) through the tangents at the left and rightsides of each peak at the evaluation limits [37]. The DSCcurves shown in the present work are the subtractionresults of the measured curves to the baselines. Accordingto the standards of the International Confederation forThermal Analysis, the onset temperature of each reaction

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(To) was determined as the intersecting point of the base-line with the tangential line from the point with the maxi-mum slope of the DSC curve at the left side of the peak.

In order to understand the phase transformation processof the SMAT sample, an in situ XRD experiment was con-ducted using a Brueker D8 Discover XRD (12 kW)equipped with a high-temperature attachment. The temper-ature was calibrated by the melting point of pure Al. Themeasured sample was heated to preset temperatures at arate of 60 K min�1 and held for 1 min before collectingthe XRD profiles from 41.6 to 46�, with a step size of0.02� and a scanning rate of 4� min�1. The sample temper-ature was monitored by using a PtRh–Pt thermocouple, ofwhich the accuracy is ±2 K. All analyses were carried outin vacuum (2 � 10�3 Pa).

3. Results and discussion

3.1. Microstructure of the SMAT sample

The initial microstructure of the sample before SMAT isshown in Fig. 1a. A martensitic structure is observed withnumerous rod-like precipitates at lath/subgrain boundariesin the austenitized and tempered ferritic steel. The width ofthe laths is typically �540 nm and the length is �13 lm.Subgrains are formed within laths during tempering toreduce the density of dislocations formed by austenitizing[38]. The precipitates are determined to be mostly of thetype M23C6 (M = Cr, Fe) according to the correspondingselected area electron diffraction (SAED) pattern (see theinset in Fig. 1a), with average dimensions of �80 nm alongthe short axis and �160 nm along the long axis. The totalamount of precipitates is estimated to be �3 vol.% in theoriginal sample. In addition to M23C6 precipitates, variousmuch finer M(C, N) precipitates (�16 nm) were alsodetected in the same material with the same heat treatmentby observing extraction double replicas [38,39].

Clear evidence of microstructure refining induced byplastic deformation is observed in the SMAT surface layerof �100 lm thickness. As shown in Fig. 1b, a high densityof dislocations are formed within the martensite (or ferrite;the same in this work) laths and various dense dislocationwalls (DDWs) develop mostly parallel or perpendicular tothe lath directions. These are subdivided into smallergrains/cells at a depth of �60 lm from the treated surface.The grain/cell size is �300 nm along the short axis and�600 nm along the long axis, respectively. Meanwhile, pre-cipitates are also slightly refined to sizes of �60 nm alongthe short axis and �120 nm along the long axis.

Further TEM observations indicate that more and moredislocations and DDWs develop in the martensite laths andthe resulted grain/cell size decreases gradually withdecreasing depth from the treated surface, due to increas-ing strain and strain rate. Moreover, the refined ferritegrains/cells appear to be equiaxed when their size is below�35 nm (at a depth of <30 lm). As for the precipitates(M23C6), the sizes along both axes also decrease gradually

with decreasing depth, and equiaxed particles are formedwhen the size is below �20 nm (at a depth of <40 lm).

Due to the very high strain and strain rate (102–103 s�1

[27,28]), extremely fine equiaxed ferrite grains with randomcrystallographic orientations are formed in the top surfacelayer, as revealed by TEM observations and the corre-sponding SAED pattern in Fig. 1c and d. The average sizeof ferrite grains is estimated to be �8 nm from a number ofdark-field images taken from the (1 1 0) diffraction of a-Fe.In addition, significant decreases in both the size and vol-ume fraction of precipitates are observed in the top surfacelayer. The mean size of M23C6 particles is refined to be�4 nm and the volume fraction is reduced to be �1%,according to the estimated results from a number ofdark-field images taken from the (3 1 1) diffraction ofM23C6 phase (see Fig. 1e).

Microstructure observations of the SMAT surface layerindicate that both ferrite and precipitate grains in the pres-ent ferritic steel are refined by dislocation activities, whichare similar to those observed previously in AISI 52100 steel[29] and AISI H13 steel [40] during SMAT. In brief, therefinement mechanism of ferrite grains involves formationof DDWs and dislocation tangles in both the originalgrains and the refined cells (under further straining), trans-formation of these microstructures into subboundarieswith small misorientations, and evolution of subboundariesto highly misoriented grain boundaries [28,29]. When fer-rite grains are refined to a certain size, plastic deformationoccurs in precipitate particles due to grain refinementstrengthening of the ferrite matrix. Therefore, the particlesare also progressively refined into smaller particles and/ordissolved into the ferrite matrix, as suggested by significantdecrements in both the size and volume fraction of M23C6

particles in the SMAT surface layer.The measurement results of average sizes of ferrite

grains/cells and M23C6 particles are summarized in Fig. 2as a function of depth from the SMAT surface. It is clearthat the top 40 lm surface layer is nanostructured (withgrain sizes below 100 nm) and the average sizes of bothphases increase with increasing depth in the top 100 lmsurface layer. Both ferrite and M23C6 grains appear to beequiaxed in the top surface layer, but anisotropic morphol-ogies are observed in the subsurface layer at a depth largerthan �30 lm.

3.2. Austenitization process of the SMAT ferritic steel

3.2.1. Phase transformation in the top surface layer during

heating

A typical DSC curve of the top surface layer (�10 lm inthickness) of the SMAT sample is given in Fig. 3, in com-parison with the DSC curve of the original sample withoutSMAT. An exothermic peak (CP) between 800 and 870 Kis observed on the DSC curve of the SMAT sample, whileno such a peak appears on the curve of the original sample.A series of TEM observations and XRD analyses of micro-structure evolutions of ferrite and precipitates across this

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Fig. 1. Typical bright-field TEM images of (a) the original P92 steel sample and (b) at a depth of �60 lm of the SMAT sample. The inset in (a) shows thecorresponding SAED pattern of M23C6. (c) A typical bright-field TEM image of the top SMAT layer. (d and e) Dark-field TEM images of ferrite grainsand M23C6 particles, taken from diffractions of (1 1 0)a and (3 1 1)M23C6, respectively, as indicated on the SAED pattern (inset in (c)).

0 50 100 1501

10

100

1000

Depth from surface (μm)

Siz

e (n

m)

Ferrite grain/cell (short axis) Ferrite grain/cell (long axis) M

23C

6 particle (short axis)

M23

C6 particle (long axis)

Fig. 2. Variations of average grain/cell sizes of ferrite and M23C6 withdepth from the treated surface of the SMAT sample.

800 900 1000 1100 1200 1300

1019 K

AI

AIFM

FM

CP

Hea

t Flo

w

Temperature (K)

Original0.1 W/g

1122 K

AII

SMAT

Fig. 3. DSC curves of the original sample and the top surface layer of theSMAT sample, at a heating rate of 20 K min�1.

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temperature range in an ongoing study (Wang et al., inpreparation) confirmed that this peak is mostly inducedby the significantly coarsening process of ferrite grains, aswell as the reforming and/or coarsening processes of pre-cipitates from the broken and/or dissolved precursors dur-ing SMAT. The endothermic peaks FM on the DSC curvesof both the SMAT sample and the original are confirmedto be induced by the magnetic transition from a ferromag-netic state to a paramagnetic state. No phase transition isdetected across the peaks and they are completely revers-ible during DSC scans when the samples are heated to tem-peratures below the second (AI) peaks and then cooleddown. Moreover, the peak positions of FM on the mea-sured DSC curves agree very well with the reported Curriepoint (�1018 K) of a reduced activation ferritic–martens-itic steel with a nominal composition (wt.%) of 9Cr–0.09C–1W [8].

Comparing the DSC curves of both samples after theferromagnetic transition, one can see that the onset temper-ature of the endothermic peak AI on the curve of theSMAT sample (1019þ26

�2 K1) is significantly decreased rela-tive to that of the original sample (1122 ± 2 K). In addi-tion, the broad endothermic peak AII observed in theoriginal sample disappears in the SMAT sample.

In situ XRD analyses were carried out on the SMATand the original samples from room temperature to1373 K, as shown in Fig. 4a and b, respectively. A decreasein the diffraction intensity of ferrite accompanied by anincrease in the diffraction intensity of austenite is observedacross the temperature intervals of peak AI for the SMATsample and of peaks AI and AII for the original sample.While a quantitative determination of the volume fractions

of a and c phases during the in situ XRD measurements isdifficult, a factor F is defined to indicate the variation ofrelative amount of c with temperature (see Fig. 4c), i.e.

F ¼ Icð111Þ

Icð111Þ þ Iað110Þð1Þ

where Ic(111) and Ia(110) are the diffraction intensities of c(1 1 1) and a (1 1 0) peaks, respectively. Austenite is firstdetected in the SMAT sample at the temperature of973 K, its volume fraction increases with increasing tem-perature and the sample is completely composed of austen-ite at 1163 K. A gradual decrease in the fraction of aaccompanies the increase in c within this temperaturerange. That is to say, the austenitization process in theSMAT ferritic steel starts at �973 K and finishes at�1163 K. In comparison, the onset and finish temperaturesof the austenitization process in the SMAT sample aremore than 100 K lower than those in the original sample,being �1093 and �1273 K, respectively.

A good agreement between the determined onset tem-perature of austenitization from in situ XRD analysis ofthe original sample and the referential data reported inthe Fe–C–Cr phase diagram (�1068 K [41]) verifies thepresent analysis. Furthermore, a difference of less than30 K might be expected between the onset temperature ofaustenitization measured by DSC and the value determinedby in situ XRD, owing to different heating modes. A con-tinuous-heating mode at a heating rate of 20 K min�1 isused during DSC measurements while a step-heating mode(isothermally at each step for �3 min) is used during in situXRD measurements. In addition, the surface layer thick-ness measured by DSC (�10 lm) is different from that inXRD experiments (�5 lm).

In general, both peaks AI and AII on the DSC curve ofthe original sample, as well as the peak AI on the curve ofthe SMAT sample (see Fig. 3), should be induced by theaustenitization process of the ferritic steel. The fact that

42 44 46 48 42 44 46 48

(b)

Inte

nsity

(a.

u.) SMAT

α(110)

γ(111)

1073 K

1023 K

1153 K

963 K

973 K

2 Theta (deg.)

1163 K

(a)

Original

γ(111)

α(110)

1173 K

1123 K

1263 K

1083 K

1093 K

1273 K

900 1000 1100 1200 1300

0.0

0.2

0.4

0.6

0.8

1.0

Original SMAT

F

Temperature (K)

(c)

Fig. 4. In situ XRD profiles of (a) the topmost layer of the SMAT sample and (b) the original sample at different temperatures. (c) Variations of intensityratio F (derived by Eq. (1)) of the top SMAT layer and the original sample with temperature.

1 Due to the overlap of FM and AI peaks, the baseline within thistemperature range could not be accurately determined. The uncertainty ofthe onset temperature of AI in the SMAT sample is determined by usingbaselines before the peak FM and after the peak AI, respectively.

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two distinct endothermic peaks appear on the DSC curveof the original sample suggests it is a two-step process tocomplete the whole austenitization during the employedheating route. The first peak (AI) is expected to be inducedby the transformation from a to c around carbide particles,following the results of in situ XRD analyses of phase tran-sition across the concerned temperature range and TEMobservations of the phase compositions of the original sam-ple. Such a transformation is also confirmed by the Fe–C–Cr phase diagram [41] and several experimental studies inferritic steels with similar compositions and starting micro-structure [8–10,42].

In situ XRD analyses show that a certain fraction of fer-rite remains at the end of AI (at �1175 K) and its amountdecreases gradually with further increasing temperature upto the end of AII (�1295 K). That is to say, the process AII

is still related to the phase transformation from a to c. Itmight correspond to the evolution of a homogeneousaustenitic microstructure with the dissolution of retainedcarbides and/or the redistribution of alloying elements(mostly Cr) from their enriched regions after AI [8,9,13].In the case of high chromium steels, it was reported thatthe incomplete realization of the a! c transformationmight result and a certain amount of carbides might be leftat the finish temperature of allotropic transformation dueto the sluggish dissolution and diffusion of carbides andchromium in austenite [8–10]. Therefore, a higher temper-ature process appears to be necessary to achieve a homoge-neous austenitic microstructure. The transformation rateduring AII is much slower than the rate during AI, so itsendothermic peak is much shallower.

The disappearance of peak AII on the DSC curve of theSMAT sample indicates that the austenitization process iscompleted within one step, i.e. during the peak AI.

3.2.2. Austenitization processes at different depths

To study the effects of the gradient microstructure onthe austenitization process, a series of DSC samples(�10 lm in thickness, as labeled in Table 1) were preparedat different depths from the SMAT surface. DSC curves of

these samples are plotted within the temperature range of750–1350 K in Fig. 5, in comparison with those of thetop SMAT surface layer and the original sample (fromFig. 3). The onset and peak temperatures of the transitionsAI and AII determined from the respective DSC curves aresummarized in Table 1. It is noticed that the exothermicpeak CP, which is mostly induced by ferrite grain coarsen-ing and carbide precipitation, is weakened gradually withincreasing depth and disappears at a depth greater than30 lm. The magnetic transition peak FM seems also tobe slightly depth-dependent.

As shown in Fig. 5 and Table 1, the onset temperatureof the endothermic peak AI, representing formation of aus-tenite from ferrite, decreases gradually with decreasingdepth in the subsurface layer, from �1122 K in the originalsample to �1101 K at a depth of 20–30 lm. It decreasessharply in the top 20 lm surface layer, from �1095 K ata depth of 10–20 lm to �1019 K in the top 10 lm surfacelayer. In comparison, the decrement amplitude in the peaktemperature of the austenite formation is smaller, from

Table 1The onset (To) and the peak (Tp) temperatures of AI and AII peaks on theDSC curves of the surface layers at different depths and of the originalsample, at a heating rate of 20 K min�1.

SampleID

Depth fromSMAT surface(lm)

AI AII

To1 (K) Tp1 (K) To2 (K) Tp2 (K)

D05 0–10 1019þ26�2 1106 ± 1 – –

D15 10–20 1095 ± 2 1114 ± 1 – –D25 20–30 1101 ± 2 1120 ± 1 – –D35 30–40 1101 ± 2 1120 ± 1 – –D45 40–50 1103 ± 2 1121 ± 1 – –D55 50–60 1105 ± 2 1121 ± 1 1192 ± 4 1232 ± 2D65 60–70 1105 ± 2 1121 ± 1 1185 ± 4 1236 ± 2D75 70–80 1108 ± 2 1122 ± 1 1197 ± 3 1242 ± 1D85 80–90 1109 ± 2 1123 ± 1 1187 ± 3 1246 ± 1Original – 1122 ± 2 1138 ± 1 1203 ± 3 1258 ± 1

800 900 1000 1100 1200 1300Temperature (K)

Hea

t Flo

w

0.1 W/g Tp2

Tp1

To2T

o1

AIFM

CP

AII

D15

D25

D55

D65

D75

D85

Original

D05

Fig. 5. DSC curves of the surface layers at different depths from theSMAT surface, at a heating rate of 20 K min�1. The characteristictemperatures on the curves are listed in Table 1.

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�1122 K at 80–90 lm to �1106 K in the top 10 lm surfacelayer. The endothermic peak AII representing the homoge-nization process of austenite is detectable on the curves ofthe samples with depth greater than 50 lm and its heightdecreases gradually with decreasing depth. Its onset andpeak temperatures also decrease gradually with decreasingdepth. This peak disappears when the depth is less than50 lm.

In principle, the sum of integrated peak areas of AI andAII on the DSC curves might correspond to the enthalpychange of the austenitization processes. However, thisquantity is difficult to determine accurately due to the inev-itable oxidation of the sample surface at high temperatureduring the DSC measurement, even under the protection offlowing Ar.

3.3. Microstructure effects on austenitization behaviors

It is notable that no detectable contamination is intro-duced into the surface layer of the treated sample bySMAT, and that the austenitization process does not varysignificantly with slight chemical composition deviations ofalloying elements (such as C and Cr) from the nominationvalues of the present steel according to phase diagrams [41].Therefore, the obvious varied austenitization behaviors, i.e.the decreased transition temperature from a to c (AI) andthe weakened homogenization process (AII), are expectedto result from the microstructure refinement in the SMATsurface layer.

3.3.1. Microstructure of the SMAT sample before

austenitization

The microstructure of the SMAT surface layer waschecked after the annealing treatment at 923 K, a temper-ature below the starting temperature of austenitization, ata heating rate and a cooling rate of 20 K min�1. Fig. 6

shows the typical microstructures of ferrite and M23C6

precipitate in the topmost surface layer of the annealedsample. It is clear that the average grain sizes of both ferriteand M23C6 have increased slightly in comparison with theirrespective sizes in the as-SMAT sample (as shown inFig. 1c–e). According to the dark-field images taken fromthe (1 1 0) diffraction of ferrite (see Fig. 6b) and (3 1 1) dif-fraction of M23C6 (see Fig. 6c), the mean grain sizes of fer-rite and M23C6 are determined to be �18 and �9 nm,respectively. In addition, the volume fraction of M23C6 pre-cipitates increases from �1% to �3% in the annealedSMAT surface layer.

The microstructure evolutions of both ferrite grains andM23C6 precipitates with depth in the surface layer of theSMAT sample annealed at 923 K are summarized inFig. 7. It is clear that the grain/cell sizes of both ferriteand precipitates also increase slightly in the subsurfacelayer of the SMAT sample after the annealing treatment.The growth kinetics decreases with increasing depth, andno microstructure change can be observed in the temperedmatrix due to the gradually decreasing stored energy withdecreasing strain during SMAT. Gradient microstructurecharacteristics as in the as-SMAT surface layer remainafter the annealing treatment, i.e. the grain/cell sizes ofboth ferrite and precipitates increase gradually withincreasing depth in the surface layer of �100 lm in thick-ness. And the mean grain size of ferrite is below 100 nmin the top surface layer of �30 lm in thickness. Further-more, both ferrite grains and precipitates remain equiaxedat smaller depths (<30 lm) and appear to be elongated atlarger depths.

Significantly enhanced thermal stability of nanocrystal-lites was observed in the present ferritic steel with respectto that in other steels after SMAT processing[29,35,36,43]. For example, a mean grain size of >100 nmwas observed by TEM in the SMAT Fe annealed at

Fig. 6. (a) Typical bright-field TEM image of the topmost layer of the SMAT sample annealed at 923 K. (b and c) Dark-field TEM images of ferrite grainsand M23C6 particles, taken from the diffractions of (1 1 0)a and (3 1 1)M23C6, respectively, as circled on the SAED pattern (the inset in (a)).

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823 K [43], and the grain sizes of ferrite in the nanostruc-tured AISI H13 steel increase to �150 nm after theannealing treatment at 923 K [36]. It has been confirmedthat the grain growth of the ferrite matrix is markedlyretarded by the presence of a large number of dispersiveCr-enriched carbides and M(C, N)-type precipitates, whichexert a strong pinning influence on the growth of ferritegrains and thereby yield fairly stable grain sizes at highertemperatures in ferritic steels [8,9]. This effect might bemore distinct in the nanostructured surface layer, in whichprecipitates have been intensively refined by SMAT (Wanget al., in preparation).

3.3.2. Size effects on kinetic and thermodynamic of

austenitization process

As discussed previously, the austenitization process inthe original sample without SMAT is composed of a trans-formation from a to c around carbide particles (AI) and ahomogenization process with the dissolution of retainedcarbides and/or redistribution of Cr from Cr-enrichedregions (AII) upon heating. In the Fe–Cr–C ferritic steelswith spheroidized alloy carbides [11], it was observed thatthe austenite first nucleates at a/a grain boundary triple

points in the vicinity of Cr-enriched carbides or in contactwith carbides located on the a/a grain boundaries, and thekinetics of austenite growth is controlled by Cr diffusionwhile the diffusion of C is much faster. With increasingtemperature above the onset temperature of AI, moreand more austenite phase will be formed to enclose the car-bides. Suppressed transformation might result if the aus-tenite phase enveloping the carbides is thick enough, dueto the slower diffusion kinetics of alloy elements in the aus-tenite than in the ferrite [12,44]. The transformation willalso restart at a higher temperature, when both the diffu-sion kinetics and the driving force to form more austeniteare enhanced.

With decreasing depth in the SMAT surface layer, sizesof both ferrite grains and carbide particles decrease gradu-ally (see Fig. 7) and the defect density induced by deforma-tion (such as various interfaces formed by dislocationactivities) increases. While austenite nucleates at interfacesbetween ferrite and carbide and its growth rate is also con-trolled by the distribution of carbide particles [45], thedependences of both the onset temperature (To1) andthe peak temperature (Tp1) of the transformation AI onthe mean particle size of M23C6 are plotted in Fig. 8. Themean particle sizes (Dm) of M23C6 are determined directlyby TEM observations for the equiaxed particles at depthsless than 40 lm or derived as

Dm ¼ffiffiffiffiffiffiffiffiffiffiffiffiffiDs � Dl

pð2Þ

for elongated particles at larger depths in the SMAT sur-face layer annealed at 923 K (see Fig. 7b). Here Ds andDl are the sizes along the short and long axes, respectively.It is clear that both the onset and the peak temperatures de-crease with decreasing particle size of M23C6. This might bewell understood according to the increasing nucleation rate

0 50 100 1501

10

100

1000

1

10

100

1000

M23

C6 (SMAT, short axis)

M23

C6 (SMAT, long axis)

M23

C6 (SMAT923, short axis)

M23

C6 (SMAT923, long axis)

Siz

e (n

m)

Depth from surface (μm)

Ferrite (SMAT, short axis) Ferrite (SMAT, long axis) Ferrite (SMAT923, short axis) Ferrite (SMAT923, long axis)

Fig. 7. Variations of average grain/cell sizes of ferrite and M23C6 withdepth from the treated surface of the SMAT sample annealed at 923 K(SMAT923), in comparison with the size–depth dependences of ferrite andM23C6 in the as-SMAT sample (see also Fig. 2).

0 40 80 120 160960

1000

1040

1080

1120

To1

(DSC)

Tp1

(DSC)

To1

(in-situ XRD)

Tem

pera

ture

(K

)

Mean M23

C6 particle size (nm)

Fig. 8. Variations of the onset and the peak temperatures of transforma-tion AI (determined from DSC curves) with the mean size of M23C6

particles. The onset temperatures of austenitization process determined byin situ XRD analyses on the top SMAT surface layer and on the originalsample are also included for comparison.

L.M. Wang et al. / Acta Materialia 59 (2011) 3710–3719 3717

Author's personal copy

of austenite at interfaces between ferrite and carbide [46],i.e.

N I ¼ xCI expDGm

I

kT

� �� exp

DG�IkT

� �nuclei m�3 s�1 ð3Þ

where x is a factor including the vibration frequency of theatoms and the surface area of the critical nucleus, T is tem-perature, and DGm

I and DG�I are the activation energies foratomic migration and the barrier against nucleation atinterfaces, respectively. Assuming a constant volume frac-tion (�0.03) of carbide particles at different depths, theconcentration of nucleation sites at interfaces (CI) can begiven by

CI ’0:72d

D2m

C0 ð4Þ

where d is the interface thickness and C0 denotes the num-ber of atom sites per unit volume. In addition, the nucle-ation rate of austenite at junctions between carbides andferrite grain boundaries increases much faster with decreas-ing carbide size, due to the facts that the concentration ofnucleation sites (CJ) depends on (1=D3

m) and the activationenergies for atomic migration (DGm

J ) and the barrieragainst nucleation (DG�J) at junctions are much lower thanthose at interfaces [45,46]. When M23C6 particles and fer-rite grains are refined to a mean size of �10 nm, the CJ be-comes considerable [17] and their enhancing effects onaustenitization kinetics might not be negligible.

In addition to the increased nucleation rate of austenite,its apparent growth rate is also expected to be significantlyenhanced by the more dispersive distribution of carbideand more “short-circuit” diffusion channels (interfaces)with decreasing sizes of M23C6 particles and ferrite grains.Therefore, the first step (AI) of austenitization might beaccelerated and less carbides/ferrite might remain goinginto the second step (AII). Complete austenitization mightbe achieved during the AI transformation for the samplesat depths less than 50 lm.

From a thermodynamic point of view, microstructurerefinements of ferrite and carbides in the SMAT surfacelayer enhance the Gibbs–Thompson effect [47]. In addition,mechanically induced nanostructures usually possess inter-faces with higher stored/excess energy than in the well-annealed states [35,48]. For example, the upper limit forenergy change induced by cold deformation was estimatedto be �0.2 kJ mol�1, which is not negligible with respect tothe energy change induced by solid-state transformations(0.5–3 kJ mol�1) [47]. This value is expected to be muchhigher in the nanostructured materials produced bySMAT, in which very high strains and strain rates areapplied. The enhanced free energy might provide an addi-tional driving force for the transformation from ferriteand carbides to austenite. For example, a stored energyof 10–100 meV per atom was suggested for Cu with grainsizes on the order of nanometers, resulting in possiblesize-induced structural transformations [49]. Moreover,some experimental observations or calculations have even

claimed that fcc Fe is more stable than bcc Fe at room tem-perature when the grain size is small enough [24,25].

The much larger decrement in the onset temperaturethan in the peak temperature might be induced by the factthat the former denotes the transformation around the car-bides with the smallest size while the latter denotes thetransformation around the carbides with the mean size.As discussed previously, a small decrease in the particle sizewill bring large differences in not only the nucleation ratebut also the driving force of austenite formation whenthe size is small enough (�10 nm in this case). In addition,the strong microstructure gradient in the first 10 lm sur-face layer might result in extra uncertainty in deriving theonset temperature of the autenitization process, by signifi-cantly broadening the transformation peak. Fortunately,such uncertainty is negligible in other samples, consideringthe fact that the difference between the peak temperaturesof neighboring slabs is much smaller than the transforma-tion spreading along the temperature axis (see Table 1 andFig. 5).

4. Summary

A gradient microstructure has been produced in the sur-face layer of a ferritic steel plate by means of SMAT. Themean sizes of ferrite grains and M23C6 particles at the topsurface are about 8 and 4 nm, respectively. They increasegradually with increasing depth and reach the respectivesizes in the original sample at a depth of �100 lm. Mean-while, the measured volume fraction of carbides is reducedfrom �3% in the substrate to �1% in the top surface layer.Such gradient characteristics remain after being annealedat 923 K, though with a slight grain growth and/or repre-cipitation of dissolved carbides.

The onset temperature of the austenitization processdecreases gradually with the decreasing grain sizes of ferriteand carbides in the SMAT sample. It is �120 K lower inthe top surface layer than in the original sample. In addi-tion, at a heating rate of 20 K min�1, complete austeniteis achieved during transformation AI (formation of caround carbides) in the surface layer with a mean carbidesize below 20 nm, while an additional transformation,AII (homogenization process), occurs to achieve completeaustenite in the surface layer with coarser carbides.

The refined microstructure accelerates the austenitiza-tion process in the SMAT sample in two ways: it promotesnucleation and growth rates of austenite grains by provid-ing more nucleation sites and fast-diffusion channels, and itincreases the driving force for the transformation from a toc with a higher stored energy. The latter might be more sig-nificant in the top surface layer with extremely fine carbideparticles.

Acknowledgements

Financial support from the National Natural ScienceFoundation of China (50701044 and 50890171) and the

3718 L.M. Wang et al. / Acta Materialia 59 (2011) 3710–3719

Author's personal copy

MOST International S&T Cooperation Project of China(2010DFB54010) is acknowledged. The authors thankMr. S.C. Wang for in situ XRD analyses and Prof. J.D.Embury for suggestive discussions and comments.

References

[1] Doherty RD. In: Cahn RW, Haasen P, editors. Physical metal-lurgy. Amsterdam: Elsevier Science; 1996. p. 1363.

[2] Bhadeshia HKDH, Honeycombe RWK. Steel: microstructure andproperties. Oxford: Elsevier; 2006.

[3] Tsuji N, Maki T. Scripta Mater 2009;60:1044.[4] Offerman SE, van Dijk NH, Sietsma J, et al. Science 2002;298:1003.[5] Molinder G. Acta Metall 1956;4:565.[6] Dirnfeld SF, Korevaar BM, van’t Spijker F. Metall Trans

1974;5:1437.[7] Garcia CI, Deardo AJ. Metall Trans A 1981;12:521.[8] Raju S, Ganesh BJ, Rai AK, et al. J Nucl Mater 2009;389:385.[9] Raju S, Ganesh BJ, Banerjee A, Mohandas E. Mater Sci Eng A

2007;465:29.[10] Lenel UR, Honeycombe RWK. Met Sci 1984;18:201.[11] Shtansky DV, Nakai K, Ohmori Y. Z Metallkd 1999;90:25.[12] Speich GR, Szirmae A, Richards MJ. Trans Metall Soc AIME

1969;245:1063.[13] Wycliffe P, Purdy GR, Embury JD. J Metal 1980;32:9.[14] Coltters RG. Mater Sci Eng 1985;76:1.[15] Gleiter H. Prog Mater Sci 1989;33:223.[16] Lu K. Mater Sci Eng R 1996;16:161.[17] Suryanarayana C. Int Mater Rev 1995;40:41.[18] Wilde G. Nanostructured materials. Amsterdam: Elsevier; 2009.[19] Valiev RZ, Islamgaliev RK, Alexandrov IV. Prog Mater Sci

2000;45:103.[20] Ivanisenko Y, Wunderlich RK, Valiev RZ, Fecht HJ. Scripta Mater

2003;49:947.[21] Yang G, Huang CX, Wang C, et al. Mater Sci Eng A 2009;515:199.[22] Zhou L, Liu G, Han Z, Lu K. Scripta Mater 2008;58:445.[23] Fukano Y. Jpn J Appl Phys 1974;13:1001.

[24] Haneda K, Zhou ZX, Morrish AH, Majima T, Miyahara T. Phys RevB 1992;46:13832.

[25] Meng Q, Zhou N, Rong Y, Chen S, Hsu TY. Acta Mater2002;50:4563.

[26] Lu K, Lu J. J Mater Sci Technol 1999;15:193.[27] Lu K, Lu J. Mater Sci Eng A 2004;375–377:38.[28] Tao NR, Wang ZB, Tong WP, Sui ML, Lu J, Lu K. Acta Mater

2002;50:4603.[29] Zhou L, Liu G, Ma XL, Lu K. Acta Mater 2008;56:78.[30] Wang K, Tao NR, Liu G, Lu J, Lu K. Acta Mater 2006;54:5281.[31] Zhang HW, Hei ZK, Liu G, Lu J, Lu K. Acta Mater 2003;51:1871.[32] Zhu KY, Vassel A, Brisset F, Lu K, Lu J. Acta Mater 2004;52:4101.[33] Wu X, Tao N, Hong Y, et al. Acta Mater 2005;53:681.[34] Sun HQ, Shi YN, Zhang MX, Lu K. Acta Mater 2007;55:975.[35] Wang ZB, Lu J, Lu K. Acta Mater 2005;53:2081.[36] Lu SD, Wang ZB, Lu K. Mater Sci Eng A 2010;527:995.[37] Hohne GWH, Hemminger WF, Flammersheim HJ. Differential

scanning calorimetry. Berlin: Springer-Verlag; 2003.[38] Ennis PJ, Zielinska-Lipiec A, Czyrska-Filemonowicz A. Mater Sci

Technol 2000;16:1226.[39] Ennis PJ, Zielinska-Lipiec A, Wachter O, Czyrska-Filemonowicz A.

Acta Mater 1997;45:4901.[40] Lu SD, Wang ZB, Lu K. J Mater Sci Technol 2010;26:258.[41] Villars P, Prince A, Okamoto H. Handbook of ternary alloy phase

diagrams, vol. 5, 2nd ed.. Materials Park (OH): The MaterialsInformation Society; 1997. p. 6622.

[42] Tokunaga T, Hasegawa K, Masuyama F. Mater Sci Eng A 2009;510–511:158.

[43] Wang ZB, Tao NR, Tong WP, Lu J, Lu K. Acta Mater 2003;51:4319.[44] Huntz AM, Guiralde P, Aucoutur M, Lacombe P. Mem Sci Rev

Metall 1969;66:85.[45] Christian JW. The theory of transformations in metals and alloys. 3rd

ed. Oxford: Pergamon Press; 2002.[46] Porter DA, Easterling KE. Phase transformations in metals and

alloys. London: Chapman & Hall; 1992.[47] Embury JD, Deschamps A, Brechet Y. Scripta Mater 2003;49:927.[48] Tong WP, Tao NR, Wang ZB, Lu J, Lu K. Science 2003;299:686.[49] Spaepen F. Philos Mag 2005;85:2979.

L.M. Wang et al. / Acta Materialia 59 (2011) 3710–3719 3719