Grain size effects on the austenitization process in a nanostructured ferritic steel

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Acta Materialia 59 (2011) 3710–3719 www.elsevier.com/locate/actamat

Grain size effects on the austenitization process in 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 mechanical attrition treatment (SMAT). The austenitization processes of ferrite and carbides in the surface layers with different average grain sizes were investigated by means of in situ X-ray diffraction, differential scanning calorimetry and transmission electron microscopy. Experimental results showed that the onset temperature of the austenitization process decreases gradually with decreasing sizes of ferrite grains and carbide particles, being 120 K lower in the top SMAT surface layer compared with that in the original sample. In addition, the twostep 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 and kinetic. Ó 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 metallurgy because of a combination of fundamental scientific interests and technological importance [1–4]. Among phase transformations in steels, the on-heating formation of facecentered cubic (fcc) austenite (c) from body-centered cubic (bcc) ferrite (a) matrix has been studied extensively and some investigations have dealt with the effects of alloying elements and initial microstructures on the austenitization process [5–13]. Experimental works on low-alloy steels have indicated that cementite particles provide nucleation sites for austenite, and austenite formation is very rapid at high temperatures [6,7,12]. The case is more complicated in high-alloy steels because the carbides in the initial ⇑ Corresponding authors. Tel.: +86 24 2397 1508; fax: +86 24 2399 8660.

E-mail addresses: [email protected] (Z.B. Wang), [email protected] (K. Lu).

structure are thermodynamically more stable than cementite [14]. For example, Lenel and Honeycombe [10] observed that 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 but not in the ferrite. By studying the austenite formation and the carbide dissolution in Fe–8.2Cr–C alloys with different C concentrations, Shtansky et al. [11] noticed that the mechanisms of austenite nucleation and growth depend on the composition, starting microstructure and austenitizing temperature. Nanostructured materials have attracted intensive interest for several decades, due to their novel properties originating 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 resultant mechanical properties, have been studied in nanostructured or ultrafine-grained ferritic steels [20–22]. However, to the authors’ knowledge, there is no study on the austenitization process of nanostructured ferrite

1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.03.006

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

matrix. This might be related to the fact that significant growth of nanosized grains may occur upon heating prior to the phase transformation temperature being reached. At ambient temperature, c-Fe has been experimentally observed in nanostructured Fe when the grain size is small enough [23,24], and it was suggested to be thermodynamically stable by calculating the Gibbs free energies of interfaces in nanostructured c and a grains [25]. Therefore, notable grain size effects on the austenitization behaviors in ferritic steels might be expected. By means of surface mechanical attrition treatment (SMAT), surface layers with a gradient grain size distribution (ranging from nanometers, submicrons to microns) have been synthesized on various metallic materials [26–34] as a result of gradient variations of applied strains and strain rates with depth from the treated surface. Previous studies showed that the nanostructures with enhanced Cr diffusivity in the SMAT surface layers of low-carbon steel [35] and H13 steel [36] are effectively stabilized by the formation of fine dispersive Cr compound particles during chromizing at 873 K, resulting in the growth of much thicker chromized surface layers than that on the coarse-grained samples after chromizing treatments at temperatures as high as 1323 K. This means that the gradient microstructure of the SMAT surface layer on steels is stable at elevated temperatures with dispersive precipitates. Such a kind of gradient nanostructured surface layer provides a unique opportunity to study the grain size effect on austenitization behavior on the nanometer scale in one sample. In this work, a nanostructured surface layer with a depth-dependent microstructure is synthesized on a commercial ferritic steel plate by means of SMAT. Thermal stability and austenitization process in the surface layer are characterized with respect to the microstructure.

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microstructure was consequently generated. In this work, the sample was treated for 60 min in vacuum at ambient temperature at a vibrating frequency of 50 Hz. No detectable contamination was introduced into the surface layer during the SMAT process. 2.2. Microstructure characterization The microstructure evolutions with depth in both the asSMAT and the annealed-SMAT samples were characterized using a JEM-2010 transmission electron microscope (TEM) operated at a voltage of 200 kV. TEM foils of the topmost layers were cut by the electro-spark discharge technique, then mechanically polished, dimpled and finally ion-milled from the untreated side. In addition, TEM foils of different subsurface layers were cut, mechanically polished, dimpled and finally electropolished from the untreated side of the SMAT samples after the removal of surface layers of different thicknesses. The electropolishing was carried out at 253 K with an electrolyte of 5 vol.% perchloric acid and 95 vol.% alcohol. A short-period milling process at a low angle (at 4–5° for 10 min) from both sides was applied to clean the foil surfaces before the TEM observation. The grain/cell sizes were averaged from a few hundred grains/cells selected randomly from TEM images. In the case that not enough grains could be counted in a TEM image at the greater depths, an FEI Nova-nano scanning electron microscope (SEM) with a lower magnification was applied to observe the microstructure at the corresponding region of a cross-sectional sample. The phase constitution of the original sample and the SMAT surface layer at room temperature was identified by X-ray diffraction (XRD) analysis using a Rigaku D/max 2400 X-ray diffractometer (7.5 kW) with Cu Ka radiation, with a step size of 0.02°.

2. Experimental 2.3. Phase transformation measurements 2.1. Sample preparation The studied ferritic steel (P92) was supplied by Wyman Gordon Forgings Inc., with the chemical composition (in wt.%) of 0.11 C, 8.75 Cr, 0.40 Mo, 1.75 W, 0.18 V, 0.05 Nb, 0.05 N and balance Fe. The initial material was in an austenitized and tempered condition (1 h at 1343 K followed by 2.5 h at 1048 K). The plate sample (100  50  4.0 mm3 in size) of the as-received steel was submitted to SMAT, the set-up and procedure of which have been described previously [26–28]. In brief, a large number of hardened steel balls of 8 mm diameter were placed at the bottom of a cylinder-shaped chamber and vibrated at a high frequency by a generator. The sample to be treated was fixed at the upper side of the chamber and impacted by flying balls repeatedly and multidirectionally. Because the sample surface was plastically deformed with high strains and high strain rates, grains in the surface layer were effectively refined, and a depth-dependent

The surface layers at different depths of the SMAT sample were cut by electro-spark discharge technique and then mechanically polished from the untreated side to 10 lm in thickness. Thermal analyses of the prepared foil samples (15 mg in weight) were conducted on a Netzsch differential scanning calorimeter (DSC 404C). The experiments were carried out from room temperature to 1373 K at a heating rate of 20 K min1, in a flowing Ar atmosphere with a gas flow rate of 50 ml min1. The temperature was calibrated by the melting points of pure In, Au and Ni. In addition, the baseline in the region of each peak was constructed by a polynomial (or an apparent transformation function) through the tangents at the left and right sides of each peak at the evaluation limits [37]. The DSC curves shown in the present work are the subtraction results of the measured curves to the baselines. According to the standards of the International Confederation for Thermal Analysis, the onset temperature of each reaction

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(To) was determined as the intersecting point of the baseline with the tangential line from the point with the maximum slope of the DSC curve at the left side of the peak. In order to understand the phase transformation process of the SMAT sample, an in situ XRD experiment was conducted using a Brueker D8 Discover XRD (12 kW) equipped with a high-temperature attachment. The temperature was calibrated by the melting point of pure Al. The measured sample was heated to preset temperatures at a rate of 60 K min1 and held for 1 min before collecting the XRD profiles from 41.6 to 46°, with a step size of 0.02° and a scanning rate of 4° min1. The sample temperature was monitored by using a PtRh–Pt thermocouple, of which the accuracy is ±2 K. All analyses were carried out in vacuum (2  103 Pa). 3. Results and discussion 3.1. Microstructure of the SMAT sample The initial microstructure of the sample before SMAT is shown in Fig. 1a. A martensitic structure is observed with numerous rod-like precipitates at lath/subgrain boundaries in the austenitized and tempered ferritic steel. The width of the laths is typically 540 nm and the length is 13 lm. Subgrains are formed within laths during tempering to reduce the density of dislocations formed by austenitizing [38]. The precipitates are determined to be mostly of the type M23C6 (M = Cr, Fe) according to the corresponding selected area electron diffraction (SAED) pattern (see the inset in Fig. 1a), with average dimensions of 80 nm along the short axis and 160 nm along the long axis. The total amount of precipitates is estimated to be 3 vol.% in the original sample. In addition to M23C6 precipitates, various much finer M(C, N) precipitates (16 nm) were also detected in the same material with the same heat treatment by observing extraction double replicas [38,39]. Clear evidence of microstructure refining induced by plastic deformation is observed in the SMAT surface layer of 100 lm thickness. As shown in Fig. 1b, a high density of dislocations are formed within the martensite (or ferrite; the same in this work) laths and various dense dislocation walls (DDWs) develop mostly parallel or perpendicular to the lath directions. These are subdivided into smaller grains/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, precipitates are also slightly refined to sizes of 60 nm along the short axis and 120 nm along the long axis. Further TEM observations indicate that more and more dislocations and DDWs develop in the martensite laths and the resulted grain/cell size decreases gradually with decreasing depth from the treated surface, due to increasing strain and strain rate. Moreover, the refined ferrite grains/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 formed when the size is below 20 nm (at a depth of <40 lm). Due to the very high strain and strain rate (102–103 s1 [27,28]), extremely fine equiaxed ferrite grains with random crystallographic orientations are formed in the top surface layer, as revealed by TEM observations and the corresponding SAED pattern in Fig. 1c and d. The average size of ferrite grains is estimated to be 8 nm from a number of dark-field images taken from the (1 1 0) diffraction of a-Fe. In addition, significant decreases in both the size and volume fraction of precipitates are observed in the top surface layer. 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 of dark-field images taken from the (3 1 1) diffraction of M23C6 phase (see Fig. 1e). Microstructure observations of the SMAT surface layer indicate that both ferrite and precipitate grains in the present ferritic steel are refined by dislocation activities, which are similar to those observed previously in AISI 52100 steel [29] and AISI H13 steel [40] during SMAT. In brief, the refinement mechanism of ferrite grains involves formation of DDWs and dislocation tangles in both the original grains and the refined cells (under further straining), transformation of these microstructures into subboundaries with small misorientations, and evolution of subboundaries to highly misoriented grain boundaries [28,29]. When ferrite grains are refined to a certain size, plastic deformation occurs in precipitate particles due to grain refinement strengthening of the ferrite matrix. Therefore, the particles are also progressively refined into smaller particles and/or dissolved into the ferrite matrix, as suggested by significant decrements 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. 2 as a function of depth from the SMAT surface. It is clear that the top 40 lm surface layer is nanostructured (with grain sizes below 100 nm) and the average sizes of both phases increase with increasing depth in the top 100 lm surface layer. Both ferrite and M23C6 grains appear to be equiaxed in the top surface layer, but anisotropic morphologies are observed in the subsurface layer at a depth larger than 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 in thickness) of the SMAT sample is given in Fig. 3, in comparison with the DSC curve of the original sample without SMAT. An exothermic peak (CP) between 800 and 870 K is observed on the DSC curve of the SMAT sample, while no such a peak appears on the curve of the original sample. A series of TEM observations and XRD analyses of microstructure 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 the corresponding 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 grains and 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)).

1000 AI

Original

FM AII

100

Ferrite grain/cell (short axis) Ferrite grain/cell (long axis) M23C6 particle (short axis)

10

Heat Flow

Size (nm)

0.1 W/g

1122 K FM

AI

SMAT

CP

M23C6 particle (long axis)

1019 K

1 0

50

100

150

Depth from surface (μ m) Fig. 2. Variations of average grain/cell sizes of ferrite and M23C6 with depth from the treated surface of the SMAT sample.

800

900

1000

1100

1200

1300

Temperature (K) Fig. 3. DSC curves of the original sample and the top surface layer of the SMAT sample, at a heating rate of 20 K min1.

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(a)

γ(111)

(b)

γ(111)

Intensity (a.u.)

(c)

0.8

Original

SMAT α (110)

1.0

0.6

F

α (110)

42

44

1163 K

1273 K

1153 K 1073 K 1023 K

1263 K 1173 K 1123 K

973 K 963 K 46 48 42

44

1093 K 1083 K 46 48

0.4 0.2 Original SMAT

0.0 900

1000

2 Theta (deg.)

1100

1200

1300

Temperature (K)

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 intensity ratio F (derived by Eq. (1)) of the top SMAT layer and the original sample with temperature.

temperature range in an ongoing study (Wang et al., in preparation) confirmed that this peak is mostly induced by the significantly coarsening process of ferrite grains, as well as the reforming and/or coarsening processes of precipitates from the broken and/or dissolved precursors during SMAT. The endothermic peaks FM on the DSC curves of both the SMAT sample and the original are confirmed to be induced by the magnetic transition from a ferromagnetic state to a paramagnetic state. No phase transition is detected across the peaks and they are completely reversible during DSC scans when the samples are heated to temperatures below the second (AI) peaks and then cooled down. Moreover, the peak positions of FM on the measured DSC curves agree very well with the reported Currie point (1018 K) of a reduced activation ferritic–martensitic steel with a nominal composition (wt.%) of 9Cr– 0.09C–1W [8]. Comparing the DSC curves of both samples after the ferromagnetic transition, one can see that the onset temperature of the endothermic peak AI on the curve of the 1 SMAT sample (1019þ26 2 K ) is significantly decreased relative to that of the original sample (1122 ± 2 K). In addition, the broad endothermic peak AII observed in the original sample disappears in the SMAT sample. In situ XRD analyses were carried out on the SMAT and the original samples from room temperature to 1373 K, as shown in Fig. 4a and b, respectively. A decrease in the diffraction intensity of ferrite accompanied by an increase in the diffraction intensity of austenite is observed across the temperature intervals of peak AI for the SMAT sample and of peaks AI and AII for the original sample. While a quantitative determination of the volume fractions 1 Due to the overlap of FM and AI peaks, the baseline within this temperature range could not be accurately determined. The uncertainty of the onset temperature of AI in the SMAT sample is determined by using baselines before the peak FM and after the peak AI, respectively.

of a and c phases during the in situ XRD measurements is difficult, a factor F is defined to indicate the variation of relative amount of c with temperature (see Fig. 4c), i.e. F ¼

I cð111Þ I cð111Þ þ I að110Þ

ð1Þ

where I c(111) and I a(110) are the diffraction intensities of c (1 1 1) and a (1 1 0) peaks, respectively. Austenite is first detected in the SMAT sample at the temperature of 973 K, its volume fraction increases with increasing temperature and the sample is completely composed of austenite at 1163 K. A gradual decrease in the fraction of a accompanies the increase in c within this temperature range. That is to say, the austenitization process in the SMAT ferritic steel starts at 973 K and finishes at 1163 K. In comparison, the onset and finish temperatures of the austenitization process in the SMAT sample are more than 100 K lower than those in the original sample, being 1093 and 1273 K, respectively. A good agreement between the determined onset temperature of austenitization from in situ XRD analysis of the original sample and the referential data reported in the Fe–C–Cr phase diagram (1068 K [41]) verifies the present analysis. Furthermore, a difference of less than 30 K might be expected between the onset temperature of austenitization measured by DSC and the value determined by in situ XRD, owing to different heating modes. A continuous-heating mode at a heating rate of 20 K min1 is used during DSC measurements while a step-heating mode (isothermally at each step for 3 min) is used during in situ XRD measurements. In addition, the surface layer thickness measured by DSC (10 lm) is different from that in XRD experiments (5 lm). In general, both peaks AI and AII on the DSC curve of the original sample, as well as the peak AI on the curve of the SMAT sample (see Fig. 3), should be induced by the austenitization process of the ferritic steel. The fact that

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

3.2.2. Austenitization processes at different depths To study the effects of the gradient microstructure on the austenitization process, a series of DSC samples (10 lm in thickness, as labeled in Table 1) were prepared at different depths from the SMAT surface. DSC curves of

these samples are plotted within the temperature range of 750–1350 K in Fig. 5, in comparison with those of the top SMAT surface layer and the original sample (from Fig. 3). The onset and peak temperatures of the transitions AI and AII determined from the respective DSC curves are summarized in Table 1. It is noticed that the exothermic peak CP, which is mostly induced by ferrite grain coarsening and carbide precipitation, is weakened gradually with increasing depth and disappears at a depth greater than 30 lm. The magnetic transition peak FM seems also to be slightly depth-dependent. As shown in Fig. 5 and Table 1, the onset temperature of the endothermic peak AI, representing formation of austenite from ferrite, decreases gradually with decreasing depth in the subsurface layer, from 1122 K in the original sample to 1101 K at a depth of 20–30 lm. It decreases sharply in the top 20 lm surface layer, from 1095 K at a depth of 10–20 lm to 1019 K in the top 10 lm surface layer. In comparison, the decrement amplitude in the peak temperature of the austenite formation is smaller, from

Tp1 0.1 W/g

Tp2 Original

To2

To1

D85

D75

D65

Heat Flow

two distinct endothermic peaks appear on the DSC curve of the original sample suggests it is a two-step process to complete the whole austenitization during the employed heating route. The first peak (AI) is expected to be induced by the transformation from a to c around carbide particles, following the results of in situ XRD analyses of phase transition across the concerned temperature range and TEM observations of the phase compositions of the original sample. Such a transformation is also confirmed by the Fe–C– Cr phase diagram [41] and several experimental studies in ferritic steels with similar compositions and starting microstructure [8–10,42]. In situ XRD analyses show that a certain fraction of ferrite remains at the end of AI (at 1175 K) and its amount decreases gradually with further increasing temperature up to the end of AII (1295 K). That is to say, the process AII is still related to the phase transformation from a to c. It might correspond to the evolution of a homogeneous austenitic microstructure with the dissolution of retained carbides 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 that the incomplete realization of the a ! c transformation might result and a certain amount of carbides might be left at the finish temperature of allotropic transformation due to the sluggish dissolution and diffusion of carbides and chromium in austenite [8–10]. Therefore, a higher temperature process appears to be necessary to achieve a homogeneous austenitic microstructure. The transformation rate during AII is much slower than the rate during AI, so its endothermic peak is much shallower. The disappearance of peak AII on the DSC curve of the SMAT sample indicates that the austenitization process is completed within one step, i.e. during the peak AI.

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D55

AII D25

Table 1 The onset (To) and the peak (Tp) temperatures of AI and AII peaks on the DSC curves of the surface layers at different depths and of the original sample, at a heating rate of 20 K min1. Sample ID D05 D15 D25 D35 D45 D55 D65 D75 D85 Original

Depth from SMAT surface (lm)

AI To1 (K)

Tp1 (K)

To2 (K)

Tp2 (K)

0–10 10–20 20–30 30–40 40–50 50–60 60–70 70–80 80–90 –

1019þ26 2

1106 ± 1 1114 ± 1 1120 ± 1 1120 ± 1 1121 ± 1 1121 ± 1 1121 ± 1 1122 ± 1 1123 ± 1 1138 ± 1

– – – – – 1192 ± 4 1185 ± 4 1197 ± 3 1187 ± 3 1203 ± 3

– – – – – 1232 ± 2 1236 ± 2 1242 ± 1 1246 ± 1 1258 ± 1

1095 ± 2 1101 ± 2 1101 ± 2 1103 ± 2 1105 ± 2 1105 ± 2 1108 ± 2 1109 ± 2 1122 ± 2

D15

AII

FM

AI D05

CP 800

900

1000

1100

1200

1300

Temperature (K) Fig. 5. DSC curves of the surface layers at different depths from the SMAT surface, at a heating rate of 20 K min1. The characteristic temperatures 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 surface layer. The endothermic peak AII representing the homogenization process of austenite is detectable on the curves of the samples with depth greater than 50 lm and its height decreases gradually with decreasing depth. Its onset and peak temperatures also decrease gradually with decreasing depth. This peak disappears when the depth is less than 50 lm. In principle, the sum of integrated peak areas of AI and AII on the DSC curves might correspond to the enthalpy change of the austenitization processes. However, this quantity is difficult to determine accurately due to the inevitable oxidation of the sample surface at high temperature during the DSC measurement, even under the protection of flowing Ar. 3.3. Microstructure effects on austenitization behaviors It is notable that no detectable contamination is introduced into the surface layer of the treated sample by SMAT, and that the austenitization process does not vary significantly with slight chemical composition deviations of alloying elements (such as C and Cr) from the nomination values 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) and the weakened homogenization process (AII), are expected to result from the microstructure refinement in the SMAT surface layer. 3.3.1. Microstructure of the SMAT sample before austenitization The microstructure of the SMAT surface layer was checked after the annealing treatment at 923 K, a temperature below the starting temperature of austenitization, at a heating rate and a cooling rate of 20 K min1. Fig. 6

shows the typical microstructures of ferrite and M23C6 precipitate in the topmost surface layer of the annealed sample. It is clear that the average grain sizes of both ferrite and M23C6 have increased slightly in comparison with their respective sizes in the as-SMAT sample (as shown in Fig. 1c–e). According to the dark-field images taken from the (1 1 0) diffraction of ferrite (see Fig. 6b) and (3 1 1) diffraction of M23C6 (see Fig. 6c), the mean grain sizes of ferrite and M23C6 are determined to be 18 and 9 nm, respectively. In addition, the volume fraction of M23C6 precipitates increases from 1% to 3% in the annealed SMAT surface layer. The microstructure evolutions of both ferrite grains and M23C6 precipitates with depth in the surface layer of the SMAT sample annealed at 923 K are summarized in Fig. 7. It is clear that the grain/cell sizes of both ferrite and precipitates also increase slightly in the subsurface layer of the SMAT sample after the annealing treatment. The growth kinetics decreases with increasing depth, and no microstructure change can be observed in the tempered matrix due to the gradually decreasing stored energy with decreasing strain during SMAT. Gradient microstructure characteristics as in the as-SMAT surface layer remain after the annealing treatment, i.e. the grain/cell sizes of both ferrite and precipitates increase gradually with increasing depth in the surface layer of 100 lm in thickness. And the mean grain size of ferrite is below 100 nm in the top surface layer of 30 lm in thickness. Furthermore, both ferrite grains and precipitates remain equiaxed at smaller depths (<30 lm) and appear to be elongated at larger depths. Significantly enhanced thermal stability of nanocrystallites was observed in the present ferritic steel with respect to that in other steels after SMAT processing [29,35,36,43]. For example, a mean grain size of >100 nm was 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 grains and 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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1000

100

Size (nm)

10

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

1

Temperature (K)

1120

1080

1040 To1 (DSC) Tp1 (DSC)

1000

To1 (in-situ XRD) 960 0

1000

40

80

120

160

Mean M23C6 particle size (nm) Fig. 8. Variations of the onset and the peak temperatures of transformation AI (determined from DSC curves) with the mean size of M23C6 particles. The onset temperatures of austenitization process determined by in situ XRD analyses on the top SMAT surface layer and on the original sample are also included for comparison.

100

M23C6 (SMAT, short axis)

10

M23C6 (SMAT, long axis) M23C6 (SMAT923, short axis) M23C6 (SMAT923, long axis)

1 0

50

100

150

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

823 K [43], and the grain sizes of ferrite in the nanostructured AISI H13 steel increase to 150 nm after the annealing treatment at 923 K [36]. It has been confirmed that the grain growth of the ferrite matrix is markedly retarded by the presence of a large number of dispersive Cr-enriched carbides and M(C, N)-type precipitates, which exert a strong pinning influence on the growth of ferrite grains and thereby yield fairly stable grain sizes at higher temperatures in ferritic steels [8,9]. This effect might be more distinct in the nanostructured surface layer, in which precipitates have been intensively refined by SMAT (Wang et al., in preparation). 3.3.2. Size effects on kinetic and thermodynamic of austenitization process As discussed previously, the austenitization process in the original sample without SMAT is composed of a transformation from a to c around carbide particles (AI) and a homogenization process with the dissolution of retained carbides and/or redistribution of Cr from Cr-enriched regions (AII) upon heating. In the Fe–Cr–C ferritic steels with spheroidized alloy carbides [11], it was observed that the austenite first nucleates at a/a grain boundary triple

points in the vicinity of Cr-enriched carbides or in contact with carbides located on the a/a grain boundaries, and the kinetics of austenite growth is controlled by Cr diffusion while the diffusion of C is much faster. With increasing temperature above the onset temperature of AI, more and more austenite phase will be formed to enclose the carbides. Suppressed transformation might result if the austenite phase enveloping the carbides is thick enough, due to the slower diffusion kinetics of alloy elements in the austenite than in the ferrite [12,44]. The transformation will also restart at a higher temperature, when both the diffusion kinetics and the driving force to form more austenite are enhanced. With decreasing depth in the SMAT surface layer, sizes of both ferrite grains and carbide particles decrease gradually (see Fig. 7) and the defect density induced by deformation (such as various interfaces formed by dislocation activities) increases. While austenite nucleates at interfaces between ferrite and carbide and its growth rate is also controlled by the distribution of carbide particles [45], the dependences of both the onset temperature (To1) and the peak temperature (Tp1) of the transformation AI on the mean particle size of M23C6 are plotted in Fig. 8. The mean particle sizes (Dm) of M23C6 are determined directly by TEM observations for the equiaxed particles at depths less than 40 lm or derived as pffiffiffiffiffiffiffiffiffiffiffiffiffi Dm ¼ Ds  Dl ð2Þ for elongated particles at larger depths in the SMAT surface layer annealed at 923 K (see Fig. 7b). Here Ds and Dl are the sizes along the short and long axes, respectively. It is clear that both the onset and the peak temperatures decrease with decreasing particle size of M23C6. This might be well understood according to the increasing nucleation rate

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of austenite at interfaces between ferrite and carbide [46], i.e.  m   DGI DGI  exp nuclei m3 s1 N I ¼ xC I exp ð3Þ kT kT where x is a factor including the vibration frequency of the atoms and the surface area of the critical nucleus, T is tem perature, and DGm I and DGI are the activation energies for atomic migration and the barrier against nucleation at interfaces, respectively. Assuming a constant volume fraction (0.03) of carbide particles at different depths, the concentration of nucleation sites at interfaces (CI) can be given by CI ’

0:72d C0 D2m

ð4Þ

where d is the interface thickness and C0 denotes the number of atom sites per unit volume. In addition, the nucleation rate of austenite at junctions between carbides and ferrite grain boundaries increases much faster with decreasing carbide size, due to the facts that the concentration of nucleation sites (CJ) depends on (1=D3m ) and the activation energies for atomic migration (DGm J ) and the barrier against nucleation (DGJ ) at junctions are much lower than those at interfaces [45,46]. When M23C6 particles and ferrite grains are refined to a mean size of 10 nm, the CJ becomes considerable [17] and their enhancing effects on austenitization kinetics might not be negligible. In addition to the increased nucleation rate of austenite, its apparent growth rate is also expected to be significantly enhanced by the more dispersive distribution of carbide and more “short-circuit” diffusion channels (interfaces) with decreasing sizes of M23C6 particles and ferrite grains. Therefore, the first step (AI) of austenitization might be accelerated and less carbides/ferrite might remain going into the second step (AII). Complete austenitization might be achieved during the AI transformation for the samples at depths less than 50 lm. From a thermodynamic point of view, microstructure refinements of ferrite and carbides in the SMAT surface layer enhance the Gibbs–Thompson effect [47]. In addition, mechanically induced nanostructures usually possess interfaces with higher stored/excess energy than in the wellannealed states [35,48]. For example, the upper limit for energy change induced by cold deformation was estimated to be 0.2 kJ mol1, which is not negligible with respect to the energy change induced by solid-state transformations (0.5–3 kJ mol1) [47]. This value is expected to be much higher in the nanostructured materials produced by SMAT, in which very high strains and strain rates are applied. The enhanced free energy might provide an additional driving force for the transformation from ferrite and carbides to austenite. For example, a stored energy of 10–100 meV per atom was suggested for Cu with grain sizes on the order of nanometers, resulting in possible size-induced structural transformations [49]. Moreover, some experimental observations or calculations have even

claimed that fcc Fe is more stable than bcc Fe at room temperature when the grain size is small enough [24,25]. The much larger decrement in the onset temperature than in the peak temperature might be induced by the fact that the former denotes the transformation around the carbides with the smallest size while the latter denotes the transformation around the carbides with the mean size. As discussed previously, a small decrease in the particle size will bring large differences in not only the nucleation rate but also the driving force of austenite formation when the size is small enough (10 nm in this case). In addition, the strong microstructure gradient in the first 10 lm surface layer might result in extra uncertainty in deriving the onset temperature of the autenitization process, by significantly broadening the transformation peak. Fortunately, such uncertainty is negligible in other samples, considering the fact that the difference between the peak temperatures of neighboring slabs is much smaller than the transformation spreading along the temperature axis (see Table 1 and Fig. 5). 4. Summary A gradient microstructure has been produced in the surface layer of a ferritic steel plate by means of SMAT. The mean sizes of ferrite grains and M23C6 particles at the top surface are about 8 and 4 nm, respectively. They increase gradually with increasing depth and reach the respective sizes in the original sample at a depth of 100 lm. Meanwhile, the measured volume fraction of carbides is reduced from 3% in the substrate to 1% in the top surface layer. Such gradient characteristics remain after being annealed at 923 K, though with a slight grain growth and/or reprecipitation of dissolved carbides. The onset temperature of the austenitization process decreases gradually with the decreasing grain sizes of ferrite and carbides in the SMAT sample. It is 120 K lower in the top surface layer than in the original sample. In addition, at a heating rate of 20 K min1, complete austenite is achieved during transformation AI (formation of c around carbides) in the surface layer with a mean carbide size below 20 nm, while an additional transformation, AII (homogenization process), occurs to achieve complete austenite in the surface layer with coarser carbides. The refined microstructure accelerates the austenitization process in the SMAT sample in two ways: it promotes nucleation and growth rates of austenite grains by providing more nucleation sites and fast-diffusion channels, and it increases the driving force for the transformation from a to c with a higher stored energy. The latter might be more significant in the top surface layer with extremely fine carbide particles. Acknowledgements Financial support from the National Natural Science Foundation of China (50701044 and 50890171) and the

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