EBSD Analysis of Microstructure Evolution of Pure Iron Subjectedto Sliding Wear and Related Change in Vickers Microhardness

Yoshihisa Kaneko+ and Taiyo Sugimoto

Department of Mechanical Engineering, Osaka City University, Osaka 558-8585, Japan

Sliding wear tests were carried on pure iron to investigate evolution of microstructure below worn surface. After the wear tests, grainboundary formation and lattice rotation were analyzed with electron backscatter diffraction (EBSD) method. In the vicinity of the worn surface,submicron grains separated by high-angle grain boundaries were generated. Below the submicron grain region, dominant microstructures weretwo kinds of low-angle grain boundaries which were horizontal and inclined to the worn surface, respectively. At deeper area from the wornsurface, continuous lattice rotation was detected. To correlate the microstructure and strength, Vickers microhardness was measured over a crosssection of the wear-affected zone. In the submicron grain and the low-angle grain boundary regions, the microhardness was proportional to thereciprocal square root of boundary spacing. In the lattice rotation region, we calculated geometrically-necessary (GN) dislocation density fromgradient of lattice rotation. The microhardness value in the lattice rotation region showed good correlation with the square root of the GNdislocation density. [doi:10.2320/matertrans.MA201301]

(Received August 1, 2013; Accepted October 29, 2013; Published December 6, 2013)

Keywords: wear, iron, ultrafine grained material, electron backscatter diffraction, Hall­Petch relation, hardness

1. Introduction

A conventional severe plastic deformation (SPD) tech-nique such as the equal channel angular pressing (ECAP)1,2)

can achieve relatively uniform plastic straining in a bulkmaterial. On the other hand, plastic strain induced byfrictional force is non-uniform and localized near the wornsurface. When a ductile metallic material is subjected tofrictional force repeatedly, fine grains are generated nearsurface.3­5) The grain refinement during the frictional processcan be attributed to the SPD which occurs locally nearsurface.

Observations with transmission electron microscopy(TEM) have shown that shape and size of the generatedgrains depended on distance from surface.3­5) At area distantfrom the surface, only dislocation cell structures have beenobserved.3) In addition to the grain boundary formations,large lattice rotation has been detected in the worn copper.3)

Hence, it seems reasonable to investigate both the grainboundary formation and the lattice rotation simultaneouslyto understand the microstructural evolution process belowworn surface. The electron backscatter diffraction (EBSD) isthe most suitable method for this purpose.

In our previous EBSD analyses on the worn copper singlecrystals,6­8) the evolutions of low- and high-angle boundariesand their orientation dependence were argued in terms of thelattice rotation. However, the EBSD analysis cannot evaluateplastic strain components directly, and thus it was difficultto consider the grain boundary formation with reference toa local shear mode. In the present study, by measuring aninclination angle of a preexisting boundary, we investigatedthe relationship between the local shear mode and the low-angle boundary formation which corresponds to the initialstage of the grain refinement.

In accordance with the Hall­Petch relation strength ofmetallic materials increases with decreasing grain size, and

hence the SPD is a promising process to obtain high-strengthmaterials.9­11) Also for the surface friction, improvementof surface strength is expected. In recent studies on thefriction stir processing (FSP), it has been reported that thehardness12­14) and strength15,16) certainly changed as afunction of the distance from a rotation center. However,since the FSP involves quite large surface deformation andheat production, complicated microstructures such as ther-momechanically-affected zone (TMAZ)17) are generated.Such complicated microstructural evolution would eliminatesignatures produced during an initial stage of grain refine-ment, and thus detailed mechanisms for the grain refinementand related strengthening are still unclear. In the presentstudy, we imposed relatively moderate friction on ironplates to simplify microstructures. Then, microhardness wasmeasured at the various microstructures formed below theworn surface. Particularly, we paid attention to the effect ofgrain size and dislocation density on the microhardness.

2. Experimental Procedure

2.1 Sliding wear testPure iron plates of 99.99mass% purity were cut into

30 © 30 © 2mm3 shape, and then annealed in vacuum at1223K for 1 h. A resultant average grain size wasapproximately 100 µm. Specimen surfaces were polishedmechanically and electrolytically. The electrolytic polishingwas conducted at 288K at 21V, in the solution consisting of60% perchloric acid and acetic acid, mixed at the ratio of1 : 9.

Sliding wear tests were carried out in a pin-on-disk typeapparatus (Fig. 1) at room temperature. A SUJ2 steel pin of5mm diameter was set at the specimen surface. Static loadof 10N was applied to the pin along axial direction. Thespecimen was fastened at the rotating table. By rotating thespecimen, we could form a circular wear track on the surface.Wear velocity was set at 10mm/s. The sliding wear testswere interrupted at 50 rotation cycles.+Corresponding author, E-mail: kaneko@imat.eng.osaka-cu.ac.jp

Materials Transactions, Vol. 55, No. 1 (2014) pp. 85 to 92Special Issue on Strength of Fine Grained Materials ® 60 Years of Hall­Petch®©2013 The Japan Institute of Metals and Materials

2.2 Microstructural observationAfter the sliding wear tests, the specimens were cut along a

tangential line of the circular wear track (i.e., along a dottedline in Fig. 1(a)) to obtain a cross section. Before the crosssectioning, the worn surface was electrodeposited with ironto protect microstructures generated near the worn surface.The cross sections were elecrolytically polished. Wear-affected zone below the worn surface was observed withthe JEOL JSM-6500F scanning electron microscope (SEM).The grain refinement and lattice rotation in the wear-affectedzone were analyzed by the EBSD technique with the EDAX-OIM system.

2.3 Vickers microhardness measurementTo correlate strength with the microstructures produced

below the worn surface, Vickers microhardness wasmeasured over the cross section. The Vickers microhardness

test was carried out with Shimadzu DUH-201 at theindentation force of 4.9mN. A shape of the resultantindentation was measured with the SEM, and the microhard-ness value was calculated from diagonal line lengths of eachindentation.

3. Results and Discussion

3.1 MicrostructureFigure 2 is a backscattered electron (BSE) image showing

wear-affected zone of the worn iron. In this photograph,the sliding wear direction is from left to right-hand side. Inaddition to subsurface microstructure produced by the slidingwear, some large grains are visible. These are the grainswhich preexisted at the as-annealed samples. The preexistinggrain boundaries were basically straight before wear test,and hence the grain boundary inclination seen in Fig. 2 isattributed to the sliding wear.

It is confirmed that the microstructure produced by thesliding wear depends on distance from the worn surface.Figure 3 shows enlarged images of typical microstructures atdifferent depths. In the vicinity of the worn surface, elongatedgrains with submicron widths are recognized. The elongationdirections are nearly parallel to the worn surface. Below thesubmicron grain region, the BSE images of microstructurescan be grouped roughly into three types, i.e., inclined layeredstructure, horizontal layered structure, and lattice rotationstructure. In Fig. 3(b), the layered microstructure parallel tothe preexisting grain boundary is seen. The inclined layeredmicrostructure is not always observed at cross sections of theworn irons. Figure 3(c) shows the other kind of layeredstructure which is horizontal to the worn surface. This layeredstructure is always observed below the submicron grainregion. Below these layered structures, the BSE intensitychanges gradually as shown in Fig. 3(d). In this region, noappreciable boundary is found. The change in the BSE

Pin

wear trackiron plate

rotation12mmφ

(a) sliding wear test (b) cross sectioning

vertical load

wear-affected zone

xy

z

x

yz

EBSD analysis

30mm

30m

m2m

m

Fig. 1 Schematic illustrations of (a) an iron specimen with circular weartrack and (b) the position of cross sectioning.

20μm

electrodeposited iron

wear direction

worn surface

pre-existing grain boundary

xy

z

Fig. 2 A BSE image of microstructure developed below surface of the worn iron.

Y. Kaneko and T. Sugimoto86

intensity would be due to continuous lattice rotation. Belowthe lattice rotation region, microstructure evolution is nolonger observed.

Figure 4 shows typical inverse pole figure (IPF) maps ofcross sections of the worn irons. Both high-angle and low-angle boundaries are indicated in the maps. The drawn low-angle boundaries are limited to misorientation angles from 2to 15°. Because we paid attention to the boundary inclinationin the present study, the analyzed areas were selected so as tocontain a single preexisting boundary which was perpendic-ular to surface. Because of the local SPD, the quality of theEBSD analysis was relatively poor at the submicron grainregion. But, it is still recognized that the submicron grainregions contain a lot of high-angle grain boundaries. Fractionof the high-angle boundaries amounts to approximately 70%of all the grain boundaries.

Below the submicron grain regions, there are low-anglegrain boundaries which are arranged parallel to each other.

Collectives of the parallel low-angle boundaries can beclassified into two types. One is the boundaries formedhorizontal to the worn surface (Region I in Figs. 4(a) and4(b)). The other type is the boundaries almost parallel to thepreexisting boundary inclined (Region II in Fig. 4(b)). TheBSE images of Figs. 3(b) and 3(c) correspond to these tworegions, respectively. Below the low-angle boundary region,the IPF maps reveal continuous lattice rotations. From thestacked microstructures seen in Figs. 2 and 4, it is deducedthat the evolution of the submicron grains should have beenpreceded by the stage of the low-angle boundary formationand the lattice rotation.

To correlate the low-angle boundary formation with localshear deformation, we investigated the lattice rotation inconjunction with the grain boundary inclination. Figure 5shows the lattice rotation angles º and the inclination angles¡ of the preexisting grain boundaries, which are measuredfrom Fig. 4. The rotation angle º is a misorientation anglefrom a deeper position of corresponding matrix grain. Therotation angles º are calculated in the Grains A and B inFig. 4 along curved lines which are 0.5 µm distant fromthe respective preexisting boundaries. At deeper areas, therotation angles º are zero. With decreasing distance fromworn surface, the rotation angles º of both grains increaserapidly. The increases in rotation angle continue up to x = 9and 8 µm in Grains A and B, respectively. These positionscoincide roughly with the uppermost part of Region I.Beyond these positions, the rotation angles no longerincreased in such a rapid rate.

An inclination angle ¡ shown in Fig. 5 is measured as anangular difference between an inclined boundary and anoriginal line of a preexisting boundary which is assumedstraight before a wear test. The inclination angles ¡ are zeroat the deeper areas and increase with decreasing depth fromsurface in both grains. It should be noted that the grain

wear direction

10μm

Region I

(a)

Grain A

High-angle GBsLow-angle GBs

(b)

Region I

10μm

Region II

Grain B

High-angle GBsLow-angle GBs

xy

z

Fig. 4 IPF maps of wear-affected zones, where a vertical grain boundaryexisted before each wear test. Regions I and II contain low-angleboundaries which are parallel and inclined to the worn surface,respectively.

5μm

(a)

(b)

(c)

(d)

5μm

5μm

2μm

pre-existing GB

xy

z

Fig. 3 BSE images showing typical worn microstructures which contain(a) submicron grains, (b) layered structure parallel to a preexistence GB,(c) layered structure parallel to worn surface and (d) gradual latticerotation.

EBSD Analysis of Microstructure Evolution of Pure Iron Subjected to Sliding Wear and Related Change in Vickers Microhardness 87

boundary inclination angles ¡ are almost equal to the latticerotation angle º, up to the uppermost part of Region I in bothgrains. However, in Region II of Grain B, the inclinationangle ¡ was significantly deviated from the rotation angle º.

IPF maps of another cross section and (100) pole figure ofRegions I and II of Grain C are shown in Fig. 6. In additionto the orientations viewed from the cross-section normal(i.e., z-axis), the orientations viewed from y-axis are mapped.Like Grain B of Fig. 4(b), Grain C includes the low-angleboundary regions consisting of Regions I and II. From the(100) pole figures, the lattice rotation occurred around the

z-axis in both regions. In the y-axis IPF map (Fig. 6(b)), twolarge orientation changes are detectable at the horizontalboundaries in Region I. From this orientation changes, it isfound that the subgrains in Region I were rotated around onerotational direction. In contrast, the orientation rotations inRegion II were identical between right- and left-hand sidesat both the z-axis and y-axis IPF maps, although severalinclined boundaries were certainly present. It is reasonable toconsider that Region II contained two kinds of the inclinedlow-angle boundaries which can induce positive and negativerotation around the z-axis.

5 10 15 20 25 300 5 10 15 20 25 300°

20°

40°

60°

80°Rotation angle, GB inclination,

Depth from Surface, x /μm

Ang

le,φ

α ,

(a) Grain A (b) Grain B

Region I Region II

Region I

φα

Fig. 5 Rotation angle º from a matrix grain and grain boundary inclination angle ¡ in the Grains A and B of Fig. 4.

(a) z-axis orientation (b) y-axis orientation

High-angle GBsLow-angle GBs

10μm 10μm xy

z

Region I

Region II

Grain CGrain C

(c) (100) PF of Regions I

(d) (100) PF of Regions II

x

z y

x

z y

Fig. 6 IPF maps showing the same cross section, where the drawn colors indicate the orientations viewed from (a) z-axis and (b) y-axis,and (100) pole figures at (c) Region I and (d) Region II.

Y. Kaneko and T. Sugimoto88

To consider the low-angle boundary formations in termsof the lattice rotation and the grain boundary inclination,we proposed two simple deformation models differing inshear direction, as illustrated in Fig. 7. In the vertical shearmodel, it is assumed that the wear process induces sheardeformation along slip planes vertical to the worn surface.Edge dislocations are emitted from the worn surface, andeventually they give rise to lattice rotation below the surface.The stored dislocations can behave as the GN dislocations.18)

Hence, the observed lattice rotations at the deeper partcould be understood by the vertical shear model. In thismodel, the preexisting boundary would be inclined as aneighboring lattice is sheared towards the vertical direction:in ideal case the inclination angle ¡ is equal to the rotationangle º. Indeed, as seen in Fig. 5, the angles ¡ and º areroughly identical at Region I and the lattice rotation region.The introduced edge dislocations can rearrange in the formof the polygonization. Since the Burgers vector is parallelto the worn surface normal, the low-angle boundarieshorizontal to the worn surface should appear. This formationof the low-angle boundaries horizontal to the surface hasbeen discussed also in the copper single crystal subjected tosliding wear.7)

In the horizontal shear model, glides of dislocations occuralong slip planes which are horizontal to the worn surface.The inclination of the preexisting boundary is caused bysimple shear towards the wear direction. If dislocation pileup at the preexisting boundary is ignored, no lattice rotationis anticipated. Hence, the discrepancy between rotation angleº and the inclination angle ¡ must exist. Let us considerthe deformation characteristics in Region II of Grain B(Fig. 4(b)). In this region, the rotation angle º was roughlyconstant and apparently smaller than the inclination angle ¡,as shown in Fig. 5(b). This result is consistent with thehorizontal shear model.

It is assumed that the edge dislocations of opposite signsare multiplied from dislocation sources in the horizontalshear model. One can expect that polygonization of thesedislocations possibly occur when dislocation density in-creased sufficiently. Because the dislocations of oppositeBurgers vector signs are emitted, the polygonization canproduce two kinds of low-angle boundaries where thedislocations with the positive and negative signs aredominantly included. The low-angle boundaries give rise tolattice rotation at a neighboring subgrain. The lattice rotationdepends on the sign of the included dislocations: the rotationdirection of the positive dislocation boundary is opposite tothat of the negative dislocation boundary. Hence, net latticerotation would be reduced at the collective of the positiveand negative dislocation boundaries. This idea is compatiblewith the fact that the orientation of Region II in Grain Cwas unchanged as shown in Fig. 6(b). As a consequence, theinclined low-angle boundaries observed at Region II couldbe understood by the horizontal shear deformation.

3.2 Vickers microhardnessIn the present study, the Vickers microhardness was

measured at very low indentation force of 4.9mN, at which aresultant indentation size is of a few micrometer. This isbecause the thickness of the wear-affected zone was typicallyless than 50 µm, within which the microstructure dependenceof hardness is unable to be estimated under a conventionalindentation force. It should be noted that hardness valueusually depends on the indentation force. This is knownas the indentation size effect (ISE).19,20) This dependenceusually follows a power-law relation, which is sometimereferred to as Meyer’s law. In order to determine the propertreatment of a hardness value at 4.9mN, we first investigatedindentation-force dependence of the microhardness, as shownin Fig. 8. The sample used was an as-annealed iron plate. It is

wear directionworn surface

slip plane preexisting GB

wear direction

α

φ =0

(b) Horizontal Shear Model

αdislocation φ

(a) Vertical Shear Model

s

x

yz

Fig. 7 Deformation models during the sliding wear test, which assume (a) vertical and (b) horizontal shear deformation.

EBSD Analysis of Microstructure Evolution of Pure Iron Subjected to Sliding Wear and Related Change in Vickers Microhardness 89

confirmed that the microhardness value decreases as anegative power function of the indentation force. Themicrohardness at the indentation force P = 4.9mN wasHv144. In a pure iron with 100 µm average grain size, thelower-yield stress was approximately 100MPa at strainrate of 10¹3 s¹1.21) Hence, the microhardness of Hv144 isapparently overestimated, assuming that the yield stress isthree times higher than the Hv value. As long as the relationbetween microhardness (Hv144) and the yield stress(100MPa) holds at P = 4.9mN, it is deduced that amicrohardness value (Hv) measured at P = 4.9mN is higherthan yield stress value (MPa) of corresponding position by afactor of 1.5.

The microhardness values of the worn iron were measuredover a polished cross section, inside a single preexistinggrain. Figure 9 shows typical shapes of indentations whichwere put at the different microstructures under the sameapplied force (P = 4.9mN). The indentation size at thesubmicron grain region was certainly small, in comparisonwith the others. Figure 10 is the relationship between theVickers microhardness and the depth from surface. Themicrohardness increases with decreasing depth from theworn surface.

For the aluminum processed by accumulated roll bonding(ARB), Kamikawa et al. has suggested that low-angledislocation boundaries with misorientation angles above 2­3° act as conventional grain boundaries in terms of strengthcontribution.22) Hence, the role of the low-angle boundarieson microhardness should be argued also in the present worniron. To show grain-size dependence of the microhardness,boundary spacing in both the submicron grain and the low-angle boundary regions were measured from the correspond-ing IPF map, which was obtained prior to the microhardnesstests. In the submicron grain region, a grain width vertical toeach elongated direction is taken as the boundary spacing.For the low-angle boundary region, a distance betweenadjoining parallel boundaries is simply taken as the boundaryspacing. For each indentation point, we averaged theboundary spacings existing within 2 µm distance. It shouldbe noted that both the high- and low-angle boundaries arecounted for the averaging. Figure 11 is the Vickers micro-hardness plotted against the boundary spacing. The micro-hardness increases linearly with increasing d¹1/2 value (i.e.,the reciprocal square root of the boundary spacing) for both

the submicron grain and the low-angle boundary regions: themicrohardness was in accordance with the Hall­Petch typerelation (· = ·0 + kd¹1/2). It should be emphasized thatdata points of both the submicron grain and the low-angleboundary regions are on the single line. Hence, it can besaid that the low-angle boundaries ­­­whose misorientationangles exceed 2° ­­­ have the same impact on strengtheningas the high-angle boundaries. This agrees well with theabove-mentioned low-angle dislocation boundary strengthen-ing in the ARB aluminum.22)

The occurrence of the lattice rotation shown in Figs. 4­6can be explained by uniform distribution of the GNdislocations. By utilizing the EBSD technique, we can

Indentation Force, P /mN10 100

Vic

kers

Mic

roha

rdne

ss, H

v

0

20

40

60

80

100

120

140

160

Fig. 8 Dependence of Vickers microhardness on indentation forcemeasured at an as-annealed iron plate.

Depth from Surface, x /μμm

0 10 20 30 40 50

Vic

kers

Mic

roha

rdne

ss, H

v

0

100

200

300

400

500

Fig. 10 Vickers microhardness measured over a cross section of the worniron. The indentation force was P = 4.9mN.

(a)

(b)

(c)

1μm

1μm

1μm

Fig. 9 Typical shapes of Vickers indentations at (a) the submicron grain,(b) the low-angle grain boundary and (c) the lattice rotation regions.

Y. Kaneko and T. Sugimoto90

evaluate the GN dislocation density µGN by the followingequation.18)

µGN ¼ 1

b

d£xy

dx� 1

b

dº

dxð1Þ

where b is the Burgers vector and dº/dx is the gradient ofrotation angle. Figure 12 presents the rotation angle º and itsgradient dº/dx calculated along a line which is perpendicularto the worn surface. After the EBSD analysis, the micro-hardness measurements were carried out along the line. Wecalculated the gradient dº/dx by liner curve-fitting of therotation angle data which are divided at every 5 µm depth.The curve fitting was conducted for the depths where therotation angle increased monotonically. Since the positivedirection of the x-axis is set toward the depth direction, therotation gradient dº/dx is calculated to be negative. Theabsolute value of the gradient dº/dx gradually increased withdecreasing depth as shown in Fig. 12.

Figure 13 is the Vickers microhardness in the latticerotation region, plotted against the GN dislocation densitywhich is estimated from the rotation gradient dº/dx. Themicrohardness seems to increase as a function of the squareroot of the GN dislocation density

ffiffiffiffiffiffiffiffi

µGNp

: the microhardness

in the lattice rotation region appears to follow the usualdislocation hardening rule23) given by

¸f ¼ c®bffiffiffiffi

µtp ð2Þ

where ¸f is the flow stress, c is a geometrical constant(c ³ 0.1­0.4), ® is the shear modulus and µt is the totaldislocation density.

At the lowermost part of the low-angle grain boundaryregion, the microhardness was almost identical to that atthe uppermost part of the lattice rotation region (see Figs. 9and 11). Let us consider the strength at the lowermost partof the low-angle boundary region, from the viewpoint of thedislocation hardening. A low-angle tilt boundary is expressedas an edge dislocation array whose spacing is h = b/ªLA,where ªLA is the misorientation angle of the boundary. Hence,the layered low-angle boundary structure can be regarded as akind of forest dislocations. The average density µLA of thesedislocations can be given simply by µLA = ªLA/bdLA, wheredLA is the average spacing between the low-angle boundaries.From the EBSD analysis, the average misorientation and theaverage boundary spacing were determined as ªLA = 7.5° anddLA = 2 µm. As a result, the calculated dislocation density isµLA = 2.73 © 1014m¹2. This value is almost equal to the GNdislocation density plotted at the uppermost part of the latticerotation region (µGN = 2.59 © 1014m¹2). Hence, it is reason-able to consider that the low-angle boundaries can contributeto strengthening in the same manner as forest dislocations.

Dislocations stored during a plastic work can be classifiedinto the GN dislocations and the statistically-stored (SS)dislocations.17) Because inducing no significant straingradient, the SS dislocations are normally difficult to bedetected with a conventional EBSD technique. If we assumethat the deformation in the lattice rotation region occursonly in the manner of the proposed vertical shear model(Fig. 7(a)), all the dislocations stored in the lattice rotationregion should have the character of the GN dislocation whichis supplied from the worn surface. To understand the con-tribution of the GN dislocations to the strength, we estimatedthe flow stress from the calculated GN dislocation density.By substituting ® = 84.7GPa, b = 0.2482 nm, µt = µGN =2.59 © 1014m¹2 into eq. (2), we obtain the flow stress of

(Boudary Spacing)-1/2, d -1/2 / μm-1/2

0.7 0.8 0.9 1.0 1.1 1.2

Vic

kers

Mic

roha

rdne

ss, H

v

100

200

300

400

500

Submicron grain regionLow-angle GB region

0.71.01.5

Boudary Spacing, d / μm2.0

Fig. 11 Vickers microhardness measured at the submicron grain and thelow-angle boundary regions, plotted against the reciprocal square root ofboundary spacing.

0 10 20 30 40 50 600°

10°

20°

30°

40°

50°

60°

-70

-60

-50

-40

-30

-20

-10

0

Depth from Surface, x /μm

Rot

atio

n A

ngle

fro

m M

atri

x, φ

Gra

dien

t, d

φ/d

x / r

ad m

m-1

gradient dφ /dx

rotation angle φ

Fig. 12 Rotation angle º from a matrix grain and its gradient dº/dx, whichwere measured along a line where the microhardness test had beenconducted after the EBSD analysis.

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8

Vic

kers

Mic

roha

rdne

ss, H

v

100

120

140

160

180

200

220

240

Square Root of the GN Dislocation Density, /x107 m-1GN

10 50 150100 200 250 300

GN Dislocation Density, GN /x10ρ 12 m-2

ρ

Fig. 13 Vickers microhardness measured at the lattice rotation region,plotted against the square root of the GN dislocation density which wasestimated by the lattice rotation gradient.

EBSD Analysis of Microstructure Evolution of Pure Iron Subjected to Sliding Wear and Related Change in Vickers Microhardness 91

·f = 93­372MPa, assuming that Taylor factor is M = 2.75for bcc polycrystal.24)

From the indentation force dependence of the microhard-ness (Fig. 7), it is deduced that the microhardness measuredat P = 4.9mN is approximately 1.5 times higher than a yieldstress (MPa). Thus, at the uppermost part of the latticerotation region, the yield stress estimated from the micro-hardness (Hv199) is ·y = 133MPa. This stress is in the rangeof the flow stress (·f = 93­372MPa) estimated from the GNdislocation density. Keh et al.25) has reported the relationshipbetween flow stress and dislocation density in iron.According to their relation, the dislocation density requiredfor ·f = 133MPa is estimated to be µ = 2.3 © 1014m¹2. Thisvalue is comparable with the GN dislocation density at theuppermost part of the lattice rotation region. Accordingly, itis plausible that the increase in microhardness at the latticerotation region was attributed mainly to the GN dislocationssupplied from surface.

4. Summary

(1) The wear-affected zone below surface of the ironsubjected to sliding wear was composed of the differentregions stacked. In the close vicinity of the wornsurface, the submicron grain structure with theelongated grains was generated. Below the submicrongrain region, two kinds of the layered low-angleboundaries, which were horizontal and inclined to theworn surface, were recognized. At deeper area fromsurface, the continuous lattice rotation was detected.

(2) The formations of the low-angle boundaries wereargued from the relationship between the lattice rotationangle and the inclination angle of the preexisting grainboundary. In the region where the low-angle boundariesarranged horizontal to surface, the lattice rotation anglewas almost identical to the inclination angle of apreexisting grain boundary. It was hence proposed thatthe horizontal low-angle boundaries resulted from therearrangement of dislocations coming from the wornsurface. In contrast, near the inclined low-angleboundaries, the lattice rotation angle was deviatedsignificantly from the inclination angle. The formationof the inclined low-angle boundaries was explainedfrom the shear deformation horizontal to surface.

(3) At the submicron grain and the low-angle boundaryregions, the boundary spacing d were measured bytaking account of both the high and low-angleboundaries. The Vickers microhardness at both regionsincreased with increasing d¹1/2 value. It was expectedthat the low-angle boundaries could act as the high-angle boundaries for strengthening.

(4) From the gradient of the lattice rotation angle, thechange in the GN dislocation density was investigatedfor the lattice rotation region. The microhardnessrevealed the good correlation with the square root ofthe GN dislocation density.

Acknowledgments

This study was financially supported by the Grant-in-Aidfor Scientific Research on Innovative Areas “Bulk nano-structured metals” through MEXT Japan under contractNo. 22102006, and by that on Challenging ExploratoryResearch through JSPS under contract No. 24656091.

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