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Acta Materialia 60 (2012) 831–843

Microstructure and crystallography of borides and secondaryprecipitation in 18 wt.% Cr–4 wt.% Ni–1 wt.%

Mo–3.5 wt.% B–0.27 wt.% C steel

Shengqiang Ma a,⇑, Jiandong Xing a,⇑, Hanguang Fu b, Yimin Gao a, Jianjun Zhang a

a State Key Laboratory for Mechanical Behaviour of Materials, School of Materials Science and Engineering, Xi’an Jiaotong University,

Xi’an, Shaanxi Province 710049, PR Chinab Research Institute of Advanced Materials Processing Technology, School of Materials Science and Engineering, Beijing University of Technology,

Beijing 100124, PR China

Received 16 September 2011; received in revised form 26 October 2011; accepted 1 November 2011Available online 15 December 2011

Abstract

The microstructure and crystallography of eutectic borides and secondary precipitations in 18 wt.% Cr–4 wt.% Ni–1 wt.% Mo–3.5 wt.% B–0.27 wt.% C steel have been investigated extensively. The results show that the as-cast microstructure of Cr–Ni–Mo-contain-ing Fe–B steel is composed of a dendritic martensite with large interdendritic eutectic borides. Transmission electron microscopy (TEM)results confirm that the borides are indexed to Cr- and Mo-rich M2B-type borides with the chemical formulas of Fe(1.35–1.36)Cr(0.92–

1.05)B0.96 and Fe0.73Cr0.45Mo0.78B, respectively. The cluster-like boride possesses a possible orientation relationship between body-centredorthorhombic Cr-rich M2B and martensite with h1 �1 0iM2B//h110ia growth direction. After destabilization, M23(C, B)6 secondary boro-carbide with a specific orientation relationship precipitates first and thereafter coarsens following the appearance of M7(C, B)3 precip-itation with the increasing destabilization temperature at the same soaking time, thus leading to a large decrease of Cr content in themartensite. However, no M6(C, B) secondary borocarbide is found in as-destabilized Fe–B steel. Destabilization treatment has no effecton the morphology of eutectic borides. The secondary borocarbides have the stoichiometry of (Fe18.26Cr4.74)(B, C)6 and (Fe3.86Cr3.14)(B,C)3 respectively. The high-resolution TEM results indicate that the nucleation and precipitation of M23(C, B)6 occur at the grain/sub-grain boundaries as well as partial within martensite, and a subsequent transformation from M23(C, B)6 to M7(C, B)3 takes place in situ,which is probably owing to the crystalline defects of dislocations and stacking faults in the structures caused by lattice distortion.� 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Fe–B martensitic steel; Microstructure; Crystallography; Precipitation; Transmission electron microscopy

1. Introduction

A considerable amount of economic loss occurs as aresult of the abrasion and corrosion in mechanical partsof machinery and equipment. A number of researches onsuperior wear-resistant and corrosion-resistant materials,therefore, have been given a high priority in the field of

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

doi:10.1016/j.actamat.2011.11.004

⇑ Corresponding authors. Tel.: +86 29 82665636; fax: +86 29 82663453.E-mail addresses: shengqiang012@163.com (S. Ma), jdxing@mail.xjtu.

edu.cn (J. Xing).

materials science recently so as to reduce this loss [1,2]. Itis very important to develop wear- and corrosion-resistantmaterials to meet rigorous environmental requirements.

Recently, the invention of Fe–Cr–B steels, which arewidely applied in abrasive and corrosion-resistant compo-nents (for example, mineral processing, slurry pumpingand hot-dip galvanizing), is considered a breakthrough,largely because of their high hardness, excellent impacttoughness and corrosion resistance, and chemical andmechanical stability in comparison with chromium caststeels and irons [3–6]. Due to the unique characteristics

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832 S. Ma et al. / Acta Materialia 60 (2012) 831–843

of boron in their microstructures, i.e. non-equilibriumboron segregation to grain boundaries, Fe–Cr–B alloyspossess a three-dimensional (3-D) network as a skeletonto enhance their wear and corrosion resistance in manyapplications [7–10]. The investigations on spatial distribu-tion of eutectic M2B borides in Fe–Cr–B alloy indicate thatthe continuous hard phases of Cr- and some Mo-rich bor-ides can coexist in the structures [9–13]. Similarly, carbide-forming elements, such as chromium, vanadium andmolybdenum, can also promote the formation of variousborides and remarkably enhance boron solubility in Fe–Bsteel, thus leading to a precipitation in air-quenched Fe–B steel to improve its hardness from 54.5 HRC in as-castto 59.5 HRC in heat treatment and toughness [11–15].Moreover, proper chromium addition can toughen boridesand also result in the precipitations and a good hardenabil-ity in Fe–B steel [11,14–16]. Kayser et al. revealed that thecompositions and lattice parameters of M2B borides canchange with the chromium concentration by using powderX-ray diffraction, such as (Cr1.04±0.08Fe0.96±0.08)B andCr1.0Fe1.0B phases [17–19], as found in previous work[5,11]. Christodoulou et al. suggested that the formulaand their crystal structures of M2B boride depended onthe ratio of chromium to iron plus chromium for Fe2Band Cr2B from zero to one respectively, such as 0.48 forthe ratio of chromium to iron plus chromium in CrFeBphase [20]. Investigations indicated that chromium addi-tion and boride morphology remarkably affected thewear-resistant properties [15,16].

Fe–B martensitic steels containing chromium additionusually have to be heat-treated in order to obtain the opti-mum combination of strength and toughness [2,3,6,11,15].The heat treatment process normally leads to a fine dispersedprecipitation of iron carbides within the matrix because ofthe supersaturated alloying elements in the matrix, whichoften results in a depletion of chromium, boron and carbonin the matrix [2,6,11,12]. Destabilization treatment (800–1000 �C for 1–6 h) is often designed in the steels containingstrong carbide-forming elements such as molybdenum,vanadium and chromium because of the diffusion of substi-tutional alloying elements (such as Cr, which ensures that asupersaturated solid solution matrix is obtained andincreases the interstitial atom solubility in the matrix) duringprolonged annealing to precipitate finely dispersed alloy car-bides [21–24]. Powell and Laird revealed the precipitationsequence of secondary carbides and their morphologies dur-ing destabilization heat treatment [25,26]. Wiengmoon et al.investigated the effects of destabilization heat treatment onhigh chromium cast irons and pointed out that destabiliza-tion heat treatment played an important role on the micro-structures and their precipitations, especially for the

Table 1Chemical compositions analyzed via spark emission spectrometer (wt.%).

Element C B Cr Si

Specimen 0.27 3.48 17.97 0.58

secondary phase transformations that control the micro-structures and properties [21,27]. Powell and Bee reportedthe combined formation of M23C6 and M7C3 during pro-longed destabilization periods at the same temperature[24]. Wang et al. discovered a gradual transformation ofM23C6 to M7C3 rods after prolonged destabilization [28–30]. Bedolla-Jacuinde et al. extensively investigated decreas-ing secondary precipitations with the decreasing tempera-ture and a maximum of precipitations reached atdestabilization temperature of 1000 �C [31]. Many investiga-tions revealed the precipitation and transformation of sec-ondary carbides in high chromium cast irons in order tounderstand secondary transformation behaviour as well asthe relationships between microstructures and mechanicalproperties [24–27,32–34]. Interestingly and particularly,Fe–B steel has similar microstructural characteristics andimproved properties compared to some chromium-contain-ing martensitic steels and high chromium cast irons. Addi-tionally, molybdenum is used to stabilize carbide phases innickel–chromium-containing irons (Ni-hard) and high chro-mium cast irons and also enhances their hardenability. How-ever, there is little research on chromium–nickel- and somemolybdenum-containing Fe–B steel to investigate its micro-structures and secondary precipitations in as-cast and as-destabilized conditions systemically.

In the present work, scanning electron microscopy(SEM) and transmission electron microscopy (TEM) havebeen used to characterize the microstructures and second-ary precipitations in as-cast and heat-treated low-carbonFe–B martensitic steel and further understand the relation-ships between microstructural changes and properties ofthe steel. In addition, the effect of molybdenum additionon microstructures both in as-cast and heat-treated steelis also discussed.

2. Materials and methods

2.1. Material preparation and heat treatment

The chemical composition of the investigated alloy isgiven in Table 1. The steel was melted in a medium-fre-quency induction furnace. Initial charge materials wereclean low silicon pig iron and steel scrap. Sheet nickeland ferroalloys such as Fe–63 wt.% Cr, Fe–55 wt.% Moand Fe–18.36 wt.% B were added into a slag-free moltenalloy so as to minimize the oxidation loss and slag forma-tion. When all the alloys were melted in the furnace,0.10 wt.% pure aluminium was added into the molten alloyto deoxidize at 1520–1540 �C. The melt was subsequentlysuper-heated to 1560 �C and transferred into a pre-heatedteapot ladle. After removal of any dross and slag, the melt

Mo Ni S P Fe

1.02 3.95 0.014 0.023 Bal.

S. Ma et al. / Acta Materialia 60 (2012) 831–843 833

was poured at 1450 �C into the sodium silicate–CO2

bonded sand moulds, obtaining Y-block ingots followingASTM A781/A 781-M95. The test specimens were cutfrom the lower part of the Y-block and surface groundto remove 3 mm from the surface and eliminate any oxi-dized layer. Some of the as-cast specimens were heat-trea-ted by destabilization treatment in a conventional furnaceat 900–1100 �C for 2–10 h, followed by cooling to roomtemperature directly in air.

2.2. Sample preparation for microstructural investigation

The microanalysis examination of the specimens wascarried out by using an SEM, a back-scattered electron(BSE) image, X-ray diffraction (XRD), an electron probemicro-analyzer (EPMA), a TEM equipped with an energydispersive X-ray (EDX) spectroscope to identify micro-structures. The SEM used was a JEOL JSM-6360LVmicroscope (JEOL, Tokyo, Japan) and a Noran SystemSIX microscope (Noran Instruments, Middleton, WI,USA). XRD was performed on a MXP21VAHF diffrac-tometer with copper Ka radiation coupling continuousscanning at 40 kV and 200 mA as an X-ray source. Thespecimen was scanned in the angular 2h range from 20�to 85� with a step size of 0.02� and a collection time of 10 s.

The chemical composition of the steel was analysed byEPMA equipped with a wavelength dispersive X-ray(WDX) analysis, performed using an electron acceleratingvoltage of 20 kV and a beam current of 50 nA. Quantita-tive composition analysis of the borides was carried outby comparison of the measured X-ray line intensity ofthe detected metals to pre-measured standards to obtainestimates of the relative amounts of each metal elementpresent compared to its pure standard. The method ofEPMA is as follows: a beam of highly accelerated electronsstrikes a small surface (�1 lm2) of the sample, then theemerging X-rays are selected on the basis of their wave-length, utilizing the Bragg diffraction condition on anadopted crystal, and then the concentration of elements isquantified by comparing the intensities of characteristicX-rays from each element present in the sample (Is) withthe intensities of the same radiation emitted by a standard(Istd). The difference between the sum of the relativeamounts of all the metals and a total of 100% was takento be the relative amount of boron and/or carbon. Then,assuming the expected stoichiometry of M2B, M23(C, B)6,or M7(C, B)3, the atomic formula of the measured boridesand secondary borocarbides could be estimated.

Thin foil specimens for examination in a JEOL JEM-2100F field emission transmission electron microscopeequipped with EDX analysis were prepared by punching3 mm discs from wafers cut from large blocks, mechanicallygrinding by hand to 50 lm. These specimens were twin-jetelectropolished to electron transparency at 50 V in a mix-ture of 5% perchloric acid, 20% glycerol and 75% methanol,with the solution being cooled to 0 �C with liquid nitrogen,following the standard techniques of thin foil used for TEM.

The image analysis of precipitations was performed on20 randomly chosen fields of view using the Leica QWinimage analysis with microscope and high-resolutionDC500 system. Each image of 2000 magnification (an areaof 2730 lm2) was statistically examined to obtain the data.

3. Results and discussion

3.1. Solidification microstructure of the steel

3.1.1. General microstructure

The microstructures and XRD pattern of low-carbonFe–18 wt.% Cr–B steel containing 1 wt.% Mo and 4 wt.%Ni elements in as-cast condition are shown in Figs. 1 and2. From Fig. 1a (colourful mild etched morphology), theas-cast Fe–18 wt.% Cr–B steel consists of a metal matrixand a great many continuous interdendritic eutectic mor-phologies as well as a little pearlite during solidification[2,6,15]. According to the XRD result (as shown inFig. 2), the matrix is identified as a-(Fe, Cr) solid solution(JCPDS 34-0396), whereas the eutectic borides mainlycomprise M2B-type borides (M represents Cr, Fe and Moatoms), such as CrFeB (JCPDS 51-1410), Fe1.1Cr0.9B0.9

(JCPDS 72-1073) and Cr1.65Fe0.35B0.96 (JCPDS 35-1180)etc., as found in previous investigations [2,3,6,15–18].From Fig. 1b (BSE image), it can be seen that the eutecticborides mainly display rod-like, bunchy or cluster-like andgrainy structures, as illustrated in Fig. 1b. Moreover, thegrainy boride appears in some light regions which are adja-cent to the grey regions of rod-like boride (as shown inFig. 1b), and it indicates that the grainy borides facilitateformation close to rod-like boride [3,11,13]. Fig. 1c and dshows some deep etched morphologies. It is clearly seenthat the grainy boride in the light region forms near oralong the length direction of rod-like boride, while the bun-chy or cluster-like boride may favour forming separately.The eutectic M2B type boride is a 3-D network or an inter-connected skeleton through the matrix [2,6,11,13]. Further-more, the matrix of the steel is a typical martensiticstructure with supersaturated alloying elements and inter-stitial atoms (namely chromium, molybdenum, nickel, sili-con, boron and carbon atoms) [2,3,6,10–12].

Fig. 3 shows the BSE image and EDX line analysis ofvarious M2B borides. From Fig. 3a, the grainy boride inthe light region (borides containing higher-atomic-numberelements in these light regions) grows at the edge of greyrod-like boride (borides containing lower-atomic-numberelements in these grey regions) (Fig. 3a). EDX analysisalong line segment AB through various borides (Fig. 3b)indicates that the grainy boride has much higher Mo con-centration than that of rod-like boride, while the Cr andFe elements decrease in these areas. Instead, the greyregion of boride shows a large amount of Cr element[2,3,6,10–13]. Therefore, grainy boride belongs to Mo-richboride while the rod-like boride belongs to Cr-rich boride.Moreover, some Cr and Ni alloying elements dissolve intothe a-Fe matrix during solidification, which results in a

Fig. 2. XRD pattern of as-cast Fe–B steel.

Fig. 3. EDX analysis of low-carbon Fe–B steel quenched in air: (a) BSEimage; (b) EDX line analysis.

Fig. 1. Solidification microstructure of Fe–B steel: (a) optical microstruc-ture (OM, mild etched morphology, etching using the mixture of 24 gsodium thiosulfate, 2.5 g cadmium chloride, 3.5 g citric acid and 100 gdistilled water for 40 s); (b) BSE image with mild etching (10 vol.% initialsolution etching for 15 s); (c) BSE image with deep etching (10 vol.% initialsolution etching for 90 s); (d) SEM image.

834 S. Ma et al. / Acta Materialia 60 (2012) 831–843

supersaturated a-(Fe, Cr) solid solution (as shown inFig. 2) [3,11–13]. As a matter of fact, this supersaturateda-(Fe, Cr) solid solution is a lath martensite matrix.

Fig. 4 shows the deep-etched morphologies of variousborides and their EDX spectra (here analysis using energy

dispersive spectroscopy attachment to the TEM) in low-carbon Fe–Cr–B steel. From Fig. 4a, there are three mor-phological M2B-type borides in the steel. Obviously, thevarious morphological borides have formed irregular inter-connected plates in the inter-dendrites, thus leading to con-tinuous 3-D networks through the martensitic matrix (asshown in Fig. 4b and c), as reported in previous works[3,11,13]. Actually, according to the Fe–Cr–B–C diagram[4,20], the Cr-rich boride may be solidified firstly as substi-tutional solid solution during solidification owing to thesimilar crystal structural parameters of Fe and Cr[2,3,10–13]. In contrast, the Mo-rich boride may facilitate

Fig. 4. Deep-etched morphologies and EDX spectra of various eutecticborides in low-carbon Fe–B steel (10 vol.% nital solution etching for30 min): (a) BSE image at low magnification; (b) SEM image at highmagnification; (c) SEM image of local magnification; (d) EDX spectraattachment to TEM of various borides.

S. Ma et al. / Acta Materialia 60 (2012) 831–843 835

nucleation and grow subsequently as grainy structuresadjacent to the Cr-rich boride [3,13], which may be fol-lowed by the depletion of Cr and B atoms in the surround-ing melt beneficial to the formation of molybdenum boride.Also, the Cr- and Mo-rich borides quite often precipitatealong grain boundaries (as shown in Fig. 4c). EDX spectra

of various morphological borides (as shown in Fig. 4d)indicate that Mo-rich boride dissolves a little Cr atom(green spectrum) while cluster-like boride has higher Crconcentration than grainy Cr–boride (red and black spec-tra) and it has little Mo atom.

Based on the experimental examinations, therefore, thelow-carbon Cr- and Mo-containing Fe–B steel has a differ-ent solidification process. From the Fe–Cr–B–C phase dia-gram [20], the primary crystal (c phase) separates out fromthe melt firstly. Since the partition coefficients of Cr, Mo,Ni and other alloying elements are less than one in the cphase, these alloying elements will be discharged into themelt when the c phase grows up, which facilitates the for-mation of M2B boride at the grain boundary, dependingon the surrounding alloying element concentration in themelt and their crystal structural parameters (such as atomradius, lattice constant and electronegativity, etc.)[2,5,11,20]. Finally, the austenite (c phase) naturally trans-forms into martensite that dissolves numerous alloying ele-ments (i.e. most of Ni, Si and some C, B, Cr and Mo)owing to the thermodynamics and kinetics of the systemin as-cast condition [10,11,20].

3.1.2. TEM and EPMA examination

Fig. 5 illustrates the bright-field TEM micrographs andcorresponding selected area diffraction patterns (SADPs)of borides and matrix in Fe–Cr–B martensitic steel. Thecluster-like Cr-rich M2B boride is clearly distinguishedand indexed to a body-centred orthorhombic (bco) struc-ture with the maximum iron solubility of 72 wt.% [20] (suchas Cr1.1Fe0.9B0.96, CrFeB or Cr1.65Fe0.35B0.96 as detected byXRD, i.e. (Fe, Cr)2B)) [2–4,11,13], and it belongs to Fddd

space group and Mn4B-type structure, where the latticeparameters are: a = 1.4583 nm, b = 0.7379 nm, c = 0.4245nm, respectively. The Cr-rich boride is different from sin-gle-phase Fe2B or Fe-rich M2B (dissolving the maximumchromium atom of 17 wt.%) boride that still has a body-centred tetragonal (bct) structure with the lattice constantof a = b = 0.51093 nm and c = 0.42486 nm (C16, CuAl2-type structure and c/a = 0.83) [11,20]. The SADP resultsuggests that cluster-like Cr-rich boride possesses a possibleorientation relationship between bco Cr-rich M2B andmartensite with h1 �1 0iM2B//h110ia growth direction(Fig. 5a). The TEM analysis on selected area diffractionpatterns in Fig. 5b indicates that grainy Mo-rich M2B bor-ide can be indexed to a body-centred tetragonal (bct) struc-ture (a = b = 0.5547 nm, c = 0.4739 nm, c/a = 0.85). Thestrong reflections are {51 1} (space, d = 0.227 nm) forCr-rich boride, and {002} (space, d = 0.237 nm) for Mo-rich boride. The bright-field image (Fig. 5c) taken fromthe matrix and SADP result reflects a typical martensiticmatrix in the steel [3,10–13]. The SADP taken from arod-like Cr-rich M2B-type boride reveals that the growthplane is (131), which is in good agreement with our previ-ous works [4,10,11].

The chemical compositions of various borides in the as-cast steel are analyzed via EPMA and given in Table 2.

Fig. 5. Bright-field TEM micrographs and corresponding selected areadiffraction patterns (SADPs) of borides and matrix: (a) cluster-like shapedCr-rich boride; (b) grainy Mo-rich boride; (c) martensitic matrix; (d) rod-like shaped Cr-rich boride.

836 S. Ma et al. / Acta Materialia 60 (2012) 831–843

This quantitative elemental analysis confirms that all bor-ides with various morphologies in the steel are M2B type.According to the accurate measurement using WDX at20 different areas per boride, the calculated analysis sug-gests that cluster-like Cr-rich boride has a stoichiometryof Fe1.35Cr1.05B0.96, which is close to a formula of CrFeBphase. In contrast, the rod-like Cr-rich boride has a stoichi-ometry of Fe1.36Cr0.92Mo0.06B0.96, which is more close to aformula of Fe1.1Cr0.9B0.96 phase previously found in

Fe–Cr–B alloys [15–18]. The atom ratios of Cr to Fe inthese two Cr-rich borides are 0.78 and 0.68 respectively.Experimental results indicate that the Cr content in thesteel, especially the Cr/Fe ratio in borides, plays the mainroles in the spatial structural changes of eutectic borides[3,11,13,20]. The higher the Cr/Fe ratio in boride, the moreis the formation of the cluster-like boride (bco structure),which is attributed to the replacement of the Fe atom bythe Cr atom in boride [3,11,20]. This is helpful for improv-ing the fracture toughness of boride and more Cr elementdissolving into boride can change the morphology andstructure, thus toughening borides [2,11,13,14,20]. How-ever, it can be seen in Table 2, that Mo-rich boride has aformula of (Fe0.73Cr0.45Mo0.78)B boride, which may bemuch closer to the structure of Mo2B boride according tothe corresponding SADP from TEM. All the WDX analy-sis results on the structures of various borides are in goodagreement with TEM examinations confirmed by the corre-sponding SADPs. The measurement results of the volumefraction on various morphological borides in the as-caststeel indicate that the average volume fractions of cluster-like, rod-like and grainy borides are 19.26 vol.%,15.47 vol.% and 4.72 vol.%, respectively.

3.2. Heat-treated alloy after destabilization heat treatment

3.2.1. Structural natural features of as-destabilized steel

Fig. 6 shows the microstructures of Fe–B steel in heat-treated condition. After destabilization heat treatment,the microstructure of the steel comprises precipitated sec-ondary borocarbides within a martensite (a0) matrix,together with eutectic M2B borides whose morphologyhas changed little during destabilization owing to the highthermodynamic stability (i.e. �1.4749 eV per formular unit(eV f.u. �1), �1.4847 eV/f.u. �1 and �1.3353 eV/f.u. �1 forthe formation enthalpy (DHr) of Fe2B, Cr2B and Mo2Bphases, respectively) [35] (Fig. 6a–d). It can be seen thatthe secondary borocarbides do not precipitate on the eutec-tic borides but nucleate and grow within the dendriticmatrix near eutectic borides. Most of them precipitate atsubgrain/phase boundaries heavily [1,23–27]. In the as-castcondition, no secondary borocarbides precipitate, sincethere is insufficient time for their nucleation from austeniteduring cooling in the sand mould. The occurrence of sec-ondary borocarbides requires allowing the processes ofnucleation and growth, and such a solidification conditionin the sand mould cannot ensure that the austenite obtainssecondary precipitations during common solidificationbecause their nucleation needs some time and certain con-ditions. Instead, most of the alloying elements can dissolveinto austenite, and later a supersaturated solid solution (i.e.a-(Fe, Cr)) is achieved during solidification. However, dur-ing the destabilization heat treatment, the supersaturatedsolid solution (i.e. a-(Fe, Cr)) decomposes and has suffi-cient time for secondary precipitations to nucleate andgrow completely during heating and holding process. Addi-tionally, destabilization heat treatment in such a wide

Table 2Electron microprobe analysis of M2B-type eutectic borides in as-cast Fe–B steel.

Eutectic boride Element (wt.%) Element (at.%) Cr/Fe ratio Calculated formula (M2B-type)

Fe Cr Mo B Fe Cr Mo B

Rod-like 54.08 34.27 4.13 7.52 41.07 28.03 1.83 29.07 0.68 Fe1.36Cr0.92Mo0.06B0.96

Cluster-like 53.46 38.74 0.32 7.48 40.06 31.26 0.14 28.54 0.78 Fe1.35Cr1.05B0.96

Grainy 27.34 15.57 49.75 7.34 24.74 15.18 26.26 33.82 0.61 Fe0.73Cr0.45Mo0.78B

%B is given by difference. Calculated formula assumes MxBy-type borides. The values are based on an average of 20 different areas per samples.

S. Ma et al. / Acta Materialia 60 (2012) 831–843 837

temperature range of 900–1100 �C can change the contentsof alloying elements (such as Cr and Mo, etc.) in thematrix, especially for the contents of boron and carbon,thus affecting martensitic starting temperature and forma-tion condition, and in these cases, the entire martensiticmatrix is naturally obtained in the system of the alloy.These indicate that a solid-state decomposition of thesupersaturated solid solution depending on temperatureand duration occurs after destabilization [24–27,36]. FromFig. 6, the secondary borocarbide particles in the regionclose to eutectic borides appear to be smaller and denserthan in the centre region of original austenite dendrites.This characteristic may be associated with the segregationsof chromium, molybdenum, carbon and boron to the outerregions of the dendrites during solidification [2,11,14–16].Higher destabilization temperatures result in coarser sec-ondary particles and precipitation coalescence phenomena(Fig. 6a and b) [6,15,24,34,37]. Nevertheless, a much highersoaking temperature seems to decrease the amount and sizeof secondary particles (such as 1025 �C and 1100 �C shownin Fig. 6c and d), which may be attributed to the furtherprecipitation of secondary borocarbides. Moreover, a sig-nificant proportion of the ultra-fine particles with diameterof 0.1–0.2 lm precipitates adjacent to some larger particlesand distributes over the matrix, especially for destabiliza-tion at 1025 �C for 4 h. It indicates that some more smalland uniform secondary borocarbides may nucleate andprecipitate randomly and grow gradually during the higherdestabilization temperature [3,6,11,15].

The volume fraction and the number per area of the sec-ondary borocarbides are 12.5–20.4 vol.% and 70–260 parti-cles per 100 lm2, respectively, as shown in Fig. 7. Thesecondary precipitations can increase the buck hardnessof the steel, i.e. from �55 HRC in as-cast condition to 64HRC (applied load of 150 kgf for 15 s) in destabilizationconditions (soaking at 1025 �C for 4 h), and the microhard-ness of the matrix increases from �568 HV in as-cast con-dition to 827 HV (applied load of 50 gf for 15 s) indestabilization at 1025 �C for 4 h. These results canaccount for the precipitation hardening mechanism [23,36].

3.2.2. XRD characterization of heat-treated steel

Fig. 8a and b shows the XRD patterns of the steel afterdestabilization at 950 �C and 1025 �C for 4 h respectively.The XRD results indicate that the population of secondaryborocarbides is mainly composed of M23B6 and/or M23(C,B)6 (JCPDS 47-1332 or 12-0570), accompanied by a small

quantity of M7(C, B)6 (JCPDS 05-0720 or 22-0211) precip-itation (Fig. 8a) [3,6,11]. Little M6(C, B) (JCPDS 47-1191)precipitation is detected by XRD analysis owing to thesmall amount of dissolution of Mo in the matrix(Fig. 8a). Typical carbide-forming elements, such as Fe,Cr and Mo, dissolve into austenite matrix and form differ-ent kinds of borocarbides during destabilization. Compar-ison of the experimental determined d spaces with thosegiven in the standard X-ray powder diffraction data indi-cates that the lattice parameters of martensite (0.2886nm) in the steel are greater than the normal lattice constantof ferrite (0.2866 nm) [34]. The increase of the latticeparameter suggests that the martensite in the steel is super-saturated with carbon, boron and other alloying elements(mainly Cr, Ni and a little Mo), and a significant dissolu-tion of these elements in the matrix facilitates the forma-tion of secondary borocarbides in Fe–B steel [11].However, M7(C, B)3 borocarbides increase during theincreasing destabilization temperature at the same soakingtime (Fig. 8b). With the increase of destabilization temper-ature, the precipitation of more M23(C, B)6 and M7(C, B)3

secondary borocarbides from the matrix induces the car-bon and boron contents in austenite to decrease at thesame soaking time, which decreases the carbon and boroncontents of martensite transformed by boron-, carbon- andchromium-depleted austenite. On the other hand, the Ost-wald ripening of precipitated M23(C, B)6 secondary boro-carbides occurs (i.e. coarsening process) and results in areduction of interfacial energy theoretically treated by Lif-shitz, Slyozov and Wagner [31,33,36–40]. The coarseningrate of Ostwald ripening depends on the diffusion path con-trolling process and is accelerated by the present stress/stain produced from the lattice distortion [31,33,36,37].Thus, some M23(C, B)6 secondary borocarbides begin tocoarsen following the appearance of more M7(C, B)3 sec-ondary borocarbides (Fig. 8b). At the same time, partialM23(C, B)6 borocarbides may transform to much smallerM7(C, B)3 borocarbide during higher destabilization tem-perature. The precipitation and transformation of borocar-bides take place gradually during destabilization, therebygenerating ultra-fine precipitations extensively. Theseresults suggest that M23(C, B)6 borocarbide precipitatespreferentially, and subsequently the ultra-fine M7(C, B)3

particles generate at the subgrain boundary probably viaM23(C, B)6 offering chromium, boron and carbon atoms.Therefore, the final structures after destabilization consistof M2B-type eutectic borides and a martensite mixed with

Fig. 6. Secondary electron images of as-destabilized Fe–B steel at varioussoaking temperatures for 4 h: (a) destabilization at 900 �C; (b) destabili-zation at 950 �C; (c) destabilization at 1025 �C; (d) destabilization at1100 �C.

Fig. 7. Volume fraction and count per unit area of secondary borocar-bides of Fe–B steel after destabilization: (a) volume fraction of secondaryborocarbides; (b) count per area of secondary borocarbides.

838 S. Ma et al. / Acta Materialia 60 (2012) 831–843

M23(C, B)6 and M7(C, B)3 secondary borocarbides in it.The existence of these small secondary borocarbides,including their volume fraction, mean particle radius andthe degree of random distribution, is significantly helpfulfor the mechanical properties of Fe–B martensitic steel(such as strength, fatigue, creep, impact toughness, wearresistance, etc.) owing to the precipitation-strengtheningof the Orowan mechanism [41–43].

3.2.3. TEM and EPMA examination

Fig. 9 shows the bright-field TEM micrographs and thecorresponding SADPs of secondary borocarbides precipi-tated within the martensite in the steel. The results revealthat the secondary precipitations display spherical androd-like shaped morphologies with a grain size of 200–400 nm in diameter (Fig. 9a). Electron diffraction patternsfrom the regions of spherical and rod-like shaped particlescan be indexed to M23(C, B)6 and M7(C, B)3, respectively(Fig. 9b). For the face-centred cubic (fcc) structure, weakreflections correspond to the M23(C, B)6 secondary boro-carbides with a = 1.0682 nm, and for the body-centredtetragonal (bct) structure, strong reflections correspondto the martensite (a0) matrix with a = 0.2831 nm andc = 0.3143 nm (c/a = 1.11). The SADPs between M23(C,B)6 secondary borocarbides and the martensite (a0) matrixreveal that a possible orientation relationship can beobtained from Fig. 9b, allowing for a maximum scatteringof 5�: (511)M23(C, B)6//(110)a0. The orientation relationshipbetween M23(C, B)6 and M7(C, B)3 secondary borocarbidesmay be: (420) M23(C,B)6//h1 1 �1 0iM7ðC;BÞ3. Fig. 9c and dshows the TEM bright-field image and the correspondingSADP of rod-like shaped secondary borocarbides. Thereflection spots along the [0001] zone axis suggests that

Fig. 8. XRD patterns of Fe–B steel after destabilization heat treatment:(a) destabilization at 950 �C; (b) destabilization at 1025 �C.

Fig. 9. Bright-field TEM micrographs and the corresponding SADPs ofsecondary borocarbides precipitated within the martensite in Fe–B steel:(a) a bright-field TEM micrograph of M23(C, B)6 reflected from ½2 �1 �6�zone axis; (b) corresponding SADPs from M23(C, B)6 secondaryborocarbide; (c) a bright-field TEM micrograph of M7(C, B)3 reflectedfrom [0001] zone axis; (d) corresponding SADPs from M7(C, B)3

secondary borocarbide.

S. Ma et al. / Acta Materialia 60 (2012) 831–843 839

for a hexagonal-close-packed (hcp) structure, strong reflec-tions correspond to M7(C, B)3 secondary borocarbideswith a = 1.3978 nm and b = 0.4251 nm. The TEM micro-graphs in Fig. 9 indicate that the nucleation and precipita-tion of M23(C, B)6 and M7(C, B)3 secondary borocarbidesprobably take place in an in situ manner within the mar-tensite matrix.The present work reveals that M23(C, B)6

secondary borocarbides may preferentially precipitatearound the original martensite where there exists high den-sity regions of dislocations and stacking faults resultingfrom lattice distortion owing to a mass of dissolution ofthe alloying elements [33,34,36,37,44]. The precipitationsfrom the supersaturated solid solution actually includethe nucleation of secondary borocarbide particles of thenew phase, their growth and companying the alloying ele-mental depletion of the matrix and the coarsening of theprecipitates. The evidence suggests that secondary borocar-bides grow by a displacive transformation mechanism inwhich the borocarbide lattice is yielded by a deformationof the supersaturated ferrite lattice, as reported in previouswork [1,22]. It is well known that the deformation of ferritelattice by a number of “zig-zag” displacements of ironatoms can generate the carbide lattice [1,22,24]. However,the formation of M7(C, B)3 may occur by an in situ mech-anism in some stacking faulting regions of M23(C, B)6 pre-cipitation. Fig. 10 shows bright-field and dark-field TEM

micrographs, a bright-field micrograph of internal stackingfaults and a corresponding SADP of M7(C, B)3 secondaryborocarbides from M7(C, B)3 reflection (½0 1 �1 1� zone axisin Fig. 10d). From Fig. 10a and b, it can be seen that manyfaulting streaks exist in the inside of M7(C, B)3 precipita-tion, and these internal faulting streaks with a width of25–50 nm are parallel to the ð1 �1 0 1ÞM7ðC;BÞ3 plane. Some

Fig. 10. Bright-field TEM micrographs showing the stack faults in theM7(C, B)3 secondary borocarbide and corresponding SADP: (a) a bright-field TEM micrograph; (b) a dark-field TEM micrograph; (c) a bright-fieldTEM micrograph of faults; (d) SADP of M7(C, B)3 precipitation reflectedfrom ½0 1 �1 1� zone axis.

Table 3Electron microprobe analysis of secondary borocarbides in Fe–B steel after d

Precipitation Element (at.%)

Fe Cr Mo

M23(C, B)6 53.42 13.86 4.15M7(C, B)3 36.26 29.48 2.65

%B (and/or %C) is given by difference. Calculated formula assumes MxBy-typesamples.

840 S. Ma et al. / Acta Materialia 60 (2012) 831–843

steps are observed among the stacking faults (Fig. 10c) andformed on {420}M23(C, B)6 planes based on the indexing ofthe electron diffraction pattern in Fig. 9b. Phase transitionof M7(C, B)3 rods could occur at the grain/subgrain andphase boundaries via these defects. Hence, as discussedabove, the M23(C, B)6 borocarbide first precipitates, andthen grows up accompanied with M7(C, B)3 borocarbideduring destabilization.

The secondary borocarbides in an as-destabilized condi-tion are analysed by using EPMA as shown in Table 3. Theboron and some carbon elements are calculated by balanc-ing with the other elements. According to the accuratemeasurement using WDX at 20 different areas of per sec-ondary borocarbide, the calculated analysis suggests thatM7(C, B)3 precipitation has a stoichiometry of(Fe3.86Cr3.14)(B, C)3, while M23(C, B)6 precipitation has astoichiometry of (Fe18.26Cr4.74)(B, C)6. Additionally,M23(C, B)6 precipitation has more Mo atoms than M7(C,B)3 precipitation, which is attributed to encouragement ofMo on M23(C, B)6. The results indicate that the phasetransformation of M7(C, B)3 generates by M23(C, B)6 pre-cipitation offering boron and alloying atoms which aregradually made by a diffusion-controlled growth model[33,40]. Table 4 shows the average composition of thematrix in as-cast and destabilized conditions measuredvia WDX analysis at 20 different areas (balance by otherelements). It indicates that the chromium content of thematrix in the destabilized condition is much lower thanthat of the matrix in the as-cast condition, which is attrib-uted to the secondary precipitates during the destabiliza-tion heat treatment. All the WDX analysis results on thestructures of various borocarbides are in good agreementwith TEM examinations confirmed by the correspondingSADPs.

3.2.4. High-resolution TEM analysis

Fig. 11 shows a high-resolution TEM micrograph andcorresponding SADPs of Fe–Cr–B steel after

estabilization heat treatment.

Cr/Fe ratio Calculated formula

B/C

28.57 0.26 Fe18.26Cr4.74(C, B)6

31.61 0.81 Fe3.86Cr3.14(C, B)3

borocarbides. The values are based on an average of 20 different areas per

Table 4Average composition of the matrix in as-cast and destabilized conditionsmeasured via WDX analysis at 20 different areas (balance by otherelements).

Condition Element (wt.%) Element (at.%)

Cr Si Mo Ni Cr Si Mo Ni

As-cast 9.58 1.04 0.42 4.95 12.14 1.62 0.24 4.57Destabilized 6.24 0.98 0.37 4.85 7.08 1.59 0.22 4.51

Fig. 11. High-resolution transmission electron microscopy image of Fe–B steel destabilized at 950 �C for 4 h and corresponding SADPs of secondaryborocarbides: (a) HRTEM image; (b) SADP of M23(C, B)6 precipitation; (c) SADP of M7(C, B)3 precipitation; (d) SADP of a0 matrix.

S. Ma et al. / Acta Materialia 60 (2012) 831–843 841

destabilization. The populations of secondary borocarbidesprecipitate at the grain boundaries and the martensiticregions with high dense dislocations and stacking faults,which is mainly attributed to the high energies, such asfaulting energy and lattice distortion energy. Because thealloying elements (Cr, Mo) and interstitial atom carbonand boron (RC = 0.077 nm, RB = 0.082 nm) dissolve intothe ferrite regions largely, the lattice structures are seriouslydistorted and deformed, thus resulting in these regions withhigh dense dislocations and faults. The secondary borocar-bides can precipitate and dissolve constantly during desta-bilization. For example, the formation of M23(C, B)6 andM7(C, B)3 occurs in situ in defect regions, which dependson the structure of internal faults and dislocations withinmartensite. However, no molybdenum borocarbides, suchM6B precipitation, are observed by TEM in the destabi-lized steel.

Fig. 12 shows the bright-field high-resolution TEMimage and corresponding Fourier-filtered HR-TEM imageof as-destabilized Fe–B steel. It is evident from Fig. 12athat the martensitic matrix has high dense dislocationsclose to the grain/subgrain and phase boundaries. TheFourier-filtered HR-TEM images of the martensitic matrix

and secondary M23(C, B)6 precipitation as well as theirgrain boundary (i.e. as shown in regions of A, B and Cin Fig. 12a) in the steel after destabilized at 1025 �C for4 h are illustrated in Fig. 12b–d, respectively. It can be seenthat a large amount of dislocations exists within the mar-tensite, as shown in Fig. 12b. Moreover, there are numer-ous stacking faults around the dislocations. Theinterplanar distance of the (110) plane for martensite is0.2038 nm, which is in good agreement with the XRDresults. The results indicate that a tensile stress takes placealong the (002) plane, as shown in Fig. 12b. Only one dis-location can be observed in the M23(C, B)6 precipitation,but the stacking faults coexist with the dislocations, asshown in Fig. 12c and d. The interplanar distance of the(511) plane for M23(C, B)6 precipitation is 0.2082 nm.The Fourier-filtered HR-TEM results reveal that disloca-tions and stacking faults may facilitate providing the nucle-ation and growth positions of secondary borocarbidesowing to the high energies in these regions. It is well knownthat the number of microstructural crystalline defects, i.e.dislocations and faults, has an effect on the precipitationof secondary borocarbides and their phase transitions[33,42,44].

Fig. 12. High-resolution transmission electron microscopy image and Fourier-filtered images of M23(C, B)6 precipitation, matrix and their boundary: (a)HR-TEM image of Fe–B steel destabilized at 950 �C for 4 h; (b) Fourier-filtered image of a0 matrix in (a); (c) Fourier-filtered image of M23(C, B)6

precipitation in a; (d) Fourier-filtered image of boundary between a0 matrix and M23(C, B)6 precipitation.

842 S. Ma et al. / Acta Materialia 60 (2012) 831–843

4. Summary

The metallurgical microstructures and precipitationamong these secondary borocarbides as well as theirmutual crystal orientation relationships after destabiliza-tion heat treatment of Fe–B martensitic steel have beeninvestigated by scanning electron microscopy and transmis-sion electron microscopy. The conclusions are as follows:

1. The as-cast microstructure of a 18 wt.% Cr–4 wt.% Ni–1 wt.% Mo–3.5 wt.% B–0.27 wt.% C steel is composedof a dendritic martensitic matrix with large interdendrit-ic eutectic M2B-type borides. The volume fractions ofcluster-like, rod-like and grainy eutectic borides are19.26 vol.%, 15.47 vol.% and 4.72 vol.%, respectively.

2. The borides are identified as Cr- and Mo-rich M2B-typeborides with the chemical formulas of Fe(1.35–1.36)Cr(0.92–

1.05)B0.96 and Fe0.73Cr0.45Mo0.78B. The cluster-like Cr-rich boride processes a possible orientation relationshipbetween body-centred orthorhombic Cr-rich M2B andmartensite with h1 �1 0iM2B//h110ia growth direction.

3. After destabilization, M23(C, B)6 secondary borocarbidewith a specific orientation relationship precipitates first,

and thereafter coarsens following the appearance ofM7(C, B)3 precipitation with the increasing destabiliza-tion temperature at the same soaking time, thus leadingto a large decrease of Cr content in the martensite. How-ever, no M6(C, B) Mo-rich secondary borocarbide isfound in as-destabilized Fe–B steel. Destabilization treat-ment has no effect on the morphology of eutectic borides.

4. The secondary borocarbides have the stoichiometry of(Fe3.86Cr3.14)(B, C)3 and (Fe18.26Cr4.74) (B, C)6, respec-tively. The M23(C, B)6 secondary borocarbides andmartensite (a0) matrix process a possible orientationrelationship with (511)M23(C, B)6//(110)a0, and the orien-tation relationship between M23(C, B)6 and M7(C, B)3

secondary borocarbides may be (420)M23(C, B)6//h1 1 �1 0iM7ðC;BÞ3.

5. The high-resolution transmission electron microscopyresults indicate that the nucleation and precipitation ofM23(C, B)6 at the grain/subgrain boundaries as well aspartial within martensite occur, and a subsequent trans-formation from M23(C, B)6 to M7(C, B)3 takes placein situ, which is probably owing to the crystalline defectsof dislocations and stacking faults in the structurescaused by lattice distortion.

S. Ma et al. / Acta Materialia 60 (2012) 831–843 843

Acknowledgements

The authors appreciate the financial support for thiswork from the Natural Science Foundation of China underGrants Nos. 50871084 & 51054008, the National High-Tech R&D program of China under Contract No.2007AA03Z510.

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