Effects of Alloying Elements on Microstructure, Hardness and Growth Rate ofCompound Layer in Gaseous-Nitrided Ferritic Alloys+1

Fanhui Meng1,+2, Goro Miyamoto2 and Tadashi Furuhara2

1Hitachi Construction Machinery Co., Ltd., Kasumigaura 300-0134, Japan2Tohoku University, Sendai 980-8577, Japan

Pure iron and Fe­M (M =Mo, Si,Mn,Cr, Al,V) binary ferritic alloys were nitrided at 843K for various times. Phase constituent, hardnessdistribution and growth rate of the compound layer are investigated by means of X-ray diffraction, EBSD, EPMA, 3DAP and nanoindentation. Inpure iron, ¾ and £ A (beneath ¾) form on the surface, and voids form along the interface. The hardness of £ A is higher than ferrite matrix, while voidformation causes softening of compound layer. Although alloying effects on the hardness of compound layer are small, Si, V and Al additionssuppress the formation of voids, and therefore the softening is also contained. All the elements investigated increase growth rate of compoundlayers. Especially the Si, V and Al additions significantly accelerate their growth rate due to the presence of excess nitrogen in compound layer.However, in the case of Al-added specimen, the growth layer becomes sluggish at longer nitriding time, presumably because of the extensiveprecipitation of AlN into the diffusion layer. [doi:10.2320/matertrans.H-M2021807]

(Received November 27, 2020; Accepted February 2, 2021; Published March 19, 2021)

Keywords: nitriding, alloying elements, compound layer, electron backscatter diffraction (EBSD), three-dimensional atom probe (3DAP)

1. Introduction

Many mechanical power transmission components, in-cluding automobiles and construction machinery, have beensubjected to nitriding treatment to improve fatigue strengthand wear resistance in order to extend their product life andreliability. The nitriding treatment under a certain nitridingpotential at the temperature of ferrite region in the Fe­Nphase diagram generates a compound layer which on thesurface consists of iron nitrides ¾ and £ A, outmost of nearsurface. In addition, beneath the compound layer, a diffusionlayer containing solute nitrogen and alloy nitrides developsby diffusion of nitrogen from the surface. There are manysystematic studies on the effect of alloying elements ondiffusion layer growth and hardening.1­4) On the other hand,regarding the growth and hardening of the compound layer,the effect of the nitriding time on the thickness of thecompound layer and the void layer has been mainlyinvestigated on practical steel.6­9) The effect of alloyingelements on formation of compound layer has also beenobserved in the plasma nitrided Fe­M (M: substitutionalelement) binary alloy.5) However, the alloying effects on theformation of compound layers during gas nitriding, has notbeen systematically investigated yet.

Therefore, in this study, the phase composition, hardnessdistribution, and growth rate of the compound layer duringgas nitriding were investigated in Fe­M (M: Mo, Si, Mn,Cr, Al, V) binary alloy, by means of X-ray diffraction,electron backscatter diffraction (EBSD), electron probemicroanalyzer (EPMA), three-dimensional atom probe(3DAP) and nanoindentation.

2. Experimental Procedure

2.1 Specimen and pretreatmentPure iron and Fe­M binary alloys which contain 1mass%

of Mo, Si, Mn, Cr, Al, or V were used. Specimens withequiaxed ferrite structure obtained by homogenizationtreatment were cut into dimensions of 15mm © 8mm ©3mm. The surface for nitriding treatment was wet polishedby using emery paper, mirror-polished with a diamond buff(3 µm), and cleaned with acetone, then subjected to nitridingtreatment.

2.2 Nitriding treatment and microstructure character-ization

The gas nitriding furnace (manufactured by Toei ScientificIndustrial Co., Ltd.) was used for the nitriding treatment.After heating the furnace to 843K, a gas combination of H2/NH3 (1:1 composition) was introduced. After the temperatureand atmosphere in the furnace has reached a steady state, itwas heated in the same atmosphere and nitrided from 14.4 ksto 115.2 ks. Immediately after nitriding, the specimen wasquenched in oil kept at room temperature. As a result ofmeasuring the outlet gas from the furnace with a hydrogensensor (manufactured by STANGE ELEKTRONIK GMBH),the gas composition was determined to be PNH3 = 0.465 atm,PH2 = 0.525 atm, PN2 = 0.01 atm, and the nitriding potential(= PNH3/(PH2)3/2) was 1.22 atm¹1/2. Under this condition,the nitriding treatment of pure iron has demonstrated thegeneration of ¾ on the outermost surface.

Phase identification of outermost surface was performedby X-ray diffraction (Ultima IV, manufactured by RigakuCo., Ltd., Cu tube, tube voltage 40 kV, tube current 40mA).Microstructure observations were made by laser microscopetogether with phase constituent maps in cross section ofspecimen were evaluated by SEM-EBSD apparatus (JSM-7001, manufactured by JEOL Ltd. accelerating voltage25 kV, TSL-OIM system). The thickness of compound layerwas measured at a total of 60 points at intervals of 50 µm on

+1This Paper was Originally Published in Japanese in Netsu-shori 59 (2019)336­343.

+2Corresponding author, E-mail: f.meng.hn@hitachi-kenki.com

Materials Transactions, Vol. 62, No. 5 (2021) pp. 596 to 602©2021 The Japan Society for Heat Treatment

a microstructure photograph taken with a laser microscopeover the entire length of the nitride surface of the sample. Thealloy nitrides in the compound layer and the diffusion layerwere analyzed using 3DAP (LEAP 4000HR produced byCAMECA). The needle-shaped samples for 3DAP measure-ment were prepared from the cross section by the FIBmethod.

Local hardness was measured using nano indenter(Triboindenter MN55344, manufactured by Hysitron, Inc.load 6mN, Berkovich indenter). Furthermore, by using FE-EPMA (JXA-8500F, manufactured by JEOL Ltd., accel-eration voltage 15 kV, measurement current 10mA), thedepth-concentration profiles of nitrogen and various alloyingelements were measured on the cross section of the nitridedspecimens.

3. Results and Discussions

3.1 Microstructure and hardness of nitrided pure ironFigure 1 shows the laser micrographs of compound layer

of the pure iron nitrided at 843K for 14.4 ks and 115.2 ks.The thickness of compound layer increases with increasingnitriding time. In the specimen nitrided for 14.4 ks, almostno voids were observed, but the voids became denser and thethickness of the void region increased simultaneously withthe nitriding time.

Figure 2(a) shows a cross-sectional laser micrograph ofpure iron nitride at 843K for 115.2 ks. A compound layer isobserved on the surface, and many voids are observed on the

surface side of the compound. Nano-hardness was measuredusing nano indentation along 3 lines at 3 µm intervals fromthe compound layer to the matrix, whose indents werepresented in Fig. 2(a). The phase map obtained by EBSDmeasurement in the same field of view as this lasermicrograph is shown in Fig. 2(b). It is evident that thephases change in the order of ¾, £ A, and ¡ from the surface toinner region. This result corresponds to the X-ray diffractionprofile measured on the specimen surface (described later inFig. 4(a)). The hardness distribution measured in the samefield of view is demonstrated in Fig. 2(c). Considering thatthe compound layer thicknesses at the three measuredlocations are slightly different, the hardness is plotted againstthe distance from the £ A/¡ boundary in Fig. 2(c). Since ¾ onthe outermost surface is thin and its hardness could not bemeasured, only the hardness of £ A at deeper region is shownhere. The maximum hardness of about 7GPa is obtainednear the £ A/¡ boundary, and the hardness decreases in thecompound layer as it approaches the surface. This is probablybecause the void density is higher near the surface. On theother hand, even though the ¡ hardness is lower than £ A, it isincreased compared to un-nitrided specimen.

Fig. 1 Micrographs of compound layer in cross-sections of the pure Fenitrided at 843K for (a) 14.4 ks, (b) 115.2 ks.

Fig. 2 Pure Fe nitrided at 843K for 115.2 ks of (a) Laser micrograph,(b) EBSD phase map and (c) nano-hardness-depth profiles.

Effects of Alloying Elements on Microstructure, Hardness and Growth Rate of Compound Layer in Gaseous-Nitrided Ferritic Alloys 597

3.2 Effect of alloying elements on the microstructureand hardness of compound layers

Figure 3 shows the laser micrographs of the Fe­M binaryalloys nitrided at 843K for 115.2 ks. A compound layer isobserved on the surface of all the specimens, but in Fe­1Mo(a), Fe­1Cr (d), and Fe­1Valloys (f ), many voids are formedas in the pure iron. The voids in Fe­1Mn alloy (c) are less innumber and coarser than pure iron. In Fe­1Si (b) and Fe­1Alalloys (e), void formation is suppressed in comparison withpure iron, and the compound layer/matrix interface is notsmooth but irregular. It is also observed that the compoundlayer becomes thicker than that of pure iron by the addition ofall the elements except Al.

Figure 4 shows the X-ray diffraction patterns measured onthe nitrided surfaces of pure iron and Fe­M binary alloysnitrided at 843K for 115.2 ks. The X-ray penetration depthunder the present measurement condition is about 5 µm. Inpure iron (Fig. 4(a)), only the ¾ peaks and a weak Fe3O4 peakare observed. In Fe­1Mo, Fe­1Mn and Fe­1Cr alloys(Figs. 4(b), (d), and (e)), only the peaks of ¾ are observedas in pure iron. Conversely, the Fe­1V alloy (Fig. 4(g)) hasstrong ¾ peaks and weak £ A peaks. Moreover, in Fe­1Si andFe­1Al alloys (Figs. 4(c) and (f )), the peak of £ A is strongerthan ¾, in contrast to the pure iron. Figure 5 shows the cross-sectional phase map of the Fe­M binary alloys nitrided at843K for 115.2 ks. In the Fe­1Mo alloy (Fig. 5(a)), ¾ isobserved on the outermost surface as well as £ A and ¡ in theinner layer as in the pure iron shown in Fig. 2(b). In the Fe­1Si alloy (Fig. 5(b)), only £ A and ¡ can be observed. Thecompound layer of the Fe­1Mn alloy (Fig. 5(c)) consists of ¾,£ A and ¡, but ¾ on the outermost surface is slightly thickerthan pure iron and Fe­1Mo. Furthermore, in the Fe­1Cr alloy(Fig. 5(d)), ¾ is observed on the outermost surface, while £ A

and ¡ are found on the inside, similarly to that of pure iron.Moreover, the Fe­1Al (Fig. 5(e)) alloys differ from pure ironwhere only £ A and ¡ are observed. Although surface structurein the Fe­1V alloy (Fig. 5(f )) is not clear, there is probablythin ¾ on the outermost surface according to X-ray diffractionpattern as shown in Fig. 4(g), and a thick £ A is formed inside.The surface phase structures observed in these phase mapscorrespond well with the X-ray diffraction pattern shownin Fig. 4.

In order to investigate the solid solution/precipitation stateof the alloying elements, the element distribution in thecompound layer and in the diffusion layer just beneath thecompound layer in the Fe­1Al, Fe­1Cr and Fe­1V alloysnitrided at 843K for 115.2 ks was analyzed by 3DAP.Figure 6 shows the three-dimensional atomic map of eachelement and nitrogen. Furthermore, a very fine Al enrichedregion is observed in the compound layer of Fe­1Al alloy(Fig. 6(a)), which is considered to be fine Al nitridesdistributed in £ A. On the other hand, coarser Al nitrides thanthat in the compound layer are formed in the diffusion layer.In the Fe­1Cr alloy (Fig. 6(b)), relatively coarse Cr enrichedregions are observed in the compound layer. Although it isnot clearly visible in the three-dimensional nitrogen map(Fig. 6(b)), composition analysis revealed that the Crenriched region has a higher nitrogen concentration thanthe surrounding area, therefore, this region should correspondto CrN precipitated in £ A. In the diffusion layer, a largeamount of CrN (Cr-enriched regions), that are finer than thosein the compound layer, are observed in the matrix, which isin contrast to the precipitation behavior in the Al-added

Fig. 4 X-ray diffraction patterns taken from the surface of the specimensnitrided at 843K for 115.2 ks of (a) Fe, (b) Fe­1Mo, (c) Fe­1Si, (d) Fe­1Mn, (e) Fe­1Cr, (f ) Fe­1Al and (g) Fe­1V alloys.

Fig. 3 Laser micrographs in cross-section of the specimens nitrided at843K for 115.2 ks in (a) Fe­1Mo, (b) Fe­1Si, (c) Fe­1Mn, (d) Fe­1Cr,(e) Fe­1Al and (f ) Fe­1V alloys.

F. Meng, G. Miyamoto and T. Furuhara598

specimen. Similar to the Fe­1Cr alloy, the Fe­1V alloy(Fig. 6(c)) exhibits that relatively coarse VN is dispersed inthe compound layer, and fine VN is densely precipitated inthe diffusion layer.

Figure 7 summarizes the effect of alloying elements on thehardness of the diffusion layer and compound layer measuredby the nano-indenter in pure iron and binary alloys nitrided at843K for 115.2 ks. Regarding the hardness of the diffusion

Fig. 6 3DAP elemental map of alloys and nitrogen in compound layer and diffusion zone, of the (a) Fe­1Al, (b) Fe­1Cr and (c) Fe­1Valloys nitrided at 843K for 115.2 ks.

Fig. 5 Phase map and corresponding nano-hardness depth profiles in the specimens nitrided at 843K for 115.2 ks in (a) Fe­1Mo, (b) Fe­1Si, (c) Fe­1Mn, (d) Fe­1Cr, (e) Fe­1Al and (f ) Fe­1V alloys.

Effects of Alloying Elements on Microstructure, Hardness and Growth Rate of Compound Layer in Gaseous-Nitrided Ferritic Alloys 599

layer just below the compound layer, the hardness of pureiron and Mo, Si, Mn added specimens increased by about1GPa compared to hardness in un-nitrided specimen,suggesting that alloy nitrides are not formed by the additionof those elements. On the other hand, as shown in Fig. 6,alloy nitrides are precipitated in the diffusion layers of theFe­1Cr, Fe­1Al and Fe­1V alloys, so the hardness increasessignificantly by nitriding. The hardness of the compoundlayer is affected greatly by the void formation, leading tolarge scatters even in the same specimen. Comparing themaximum hardness of each alloy which can be consideredas the true hardness without void effect, the hardness of £ Aincreases a little with the addition of any element, but theamount of hardening is smaller than that of the diffusionlayer. In addition, hardness for the Fe­1Mn suggests that ¾is harder than £ A. It is quite possible that even maximumhardness might have been underestimated due to the presenceof voids.

Figure 8 shows the relationship between the nitriding timeand the compound layer thickness of pure iron and all binaryalloys, Fe­1Mo, Fe­1Si, Fe­1Mn, Fe­1Cr, Fe­1Al and Fe­1V, nitrided at 843K for various times. The thickness shownin these graphs is a average thickness of each compoundlayer measured at 60 points by 50 µm intervals over the entirelength of the nitrided surface (3mm) of the specimen. In pureiron, the square of the compound layer thickness isproportional to the nitriding time. Therefore, it is clear thatthe growth of the compound layer is controlled by nitrogendiffusion. Similar to the pure iron, the parabolic law issatisfied in each of Mo, Si, Mn, Cr, and V-added specimens,while growth rate in those specimens is faster than pureiron. Conversely, in the Fe­1Al alloy, the parabolic law isnot satisfied. In the 28.8 ks specimen, the compound layer isthicker than pure iron, meanwhile it becomes thinner thanpure iron in the 115.2 ks specimen.

3.3 Effect of alloying elements on growth rate ofcompound layer

The nitrogen concentration in the compound layer wasmeasured by FE-EPMA in order to evaluate the effect ofalloying elements on the growth rate of the compound layer.Figure 9 shows (a) a nitrogen concentration profile and (b)enlarged profile of concentration distribution near theboundary between compound and diffusion layers in pureiron nitrided at 843K for 115.2 ks.

Fig. 7 Nano-hardness in (a) diffusion zone just below the compound layerand (b) compound layer in the alloys nitrided at 843K for 115.2 ks.

Fig. 8 Variations in squared thickness of compound layer in the alloysnitrided at 843K.

Fig. 9 (a) Nitrogen concentration profile in pure Fe after nitriding at 843Kfor 115.2 ks, (b) enlarged profile of compound layer. All the data takenalong three different lines are plotted.

F. Meng, G. Miyamoto and T. Furuhara600

Figure 9(a) demonstrates that nitrogen concentration onthe surface is about 25 at%, which decreases to about 20 at%as the depth down, and the concentration drops sharply inthe inner part of the specimen. On a basis of the EBSDmeasurement and the Fe­N binary phase diagram, it isconsidered that N content profile corresponds to the ¾, £ A, and¡ in the order from the surface. From the enlarged view ofthe compound layer front shown in Fig. 9(b), the nitrogenconcentration corresponds well with the ideal nitrogenconcentration of £ A (Fe4N: 20 at%).

Figure 10 summarizes the results of similar measurementsperformed on Si, Al, and V-added specimens nitrided at843K for 28.8 ks. Here, the ideal nitrogen concentration isshown as “ideal” in the figure, assuming that all the alloyingelements form alloy nitrides (Si3N4,AlN,VN) and all theiron atoms form Fe4N. As demonstrated previously, X-raydiffraction measurements (Fig. 4) revealed that ¾ and £ A areformed on the surface of Fe­1Si (Fig. 10(a)), Fe­1Al(Fig. 10(b)) and Fe­1V alloys (Fig. 10(c)). In the Fe­1Siand Fe­1Al alloys, although the EPMA measurement mighthave captured only £ A because the nitrogen enriched regionwas not found on the outermost surface, the surface regionconcentration is still higher than the ideal nitrogen

concentration assumed by the Si3N4 and AlN formation.Therefore, it indicates the presence of excess nitrogen inthese alloys. On the other hand, the nitrogen-enriched regionon the outermost surface in the Fe­1V alloy is considered tobe ¾. Focusing on the region corresponding to be £ A beneath¾, whose concentration is also still higher than the idealconcentration. Even after nitriding for 115.2 ks, excessnitrogen was still observed in the compound layer of theFe­1Si, Fe­1Al and Fe­1V alloys. And yet, the nitrogenconcentration distributions in the other alloys were similar tothat of the pure iron, and excess nitrogen was not observed. Itshould be noted that, in Fe­1Al and Fe­1Si alloys, a part ofthe compound layer grows spiky, and thus, the compoundlayer thickness is not constant even in the same sample.Therefore, the average thickness shown in Fig. 8 does notexactly coincide with the compound layer thickness in thenitrogen concentration distribution of Fig. 10.

Figure 11 shows effects of excess nitrogen on compoundlayer thickness nitrided at 843K for 28.6 ks and 115.2 ks. Theexcess nitrogen concentration was calculated by the differ-ence between the maximum nitrogen concentration in the £ Aregion and the ideal nitrogen concentration of each alloyobtained by EPMA measurement. In the 28.8 ks specimensof (a), it can be observed that the compound layer thicknessincreases together with the amount of excess nitrogen in theFe­1V, Fe­1Si, and Fe­1Al alloys. Even in the case of the115.2 ks specimens of (b), the compound layer thicknessincreases with the increase of excess nitrogen except for theFe­1Al alloy. Therefore, it is considered that excess nitrogenpromotes the growth of the compound layer. The reason forthis effect is explained qualitatively in Fig. 12. The result isconditioned by the crystal structure of £ A(Fe4N) in whichnitrogen is ordered at the body central position of Fe in thefcc unit cell. One of the four octahedral sites in the unit cell is

Fig. 10 Nitrogen concentration profiles of compound layer of (a) Fe­1Si,(b) Fe­1Al, (c) Fe­1V alloys after nitriding at 843K for 28.8 ks.

Fig. 11 Effects of excess nitrogen on compound layer thickness afternitriding at 843K for (a) 28.8 ks and (b) 115.2 ks.

Effects of Alloying Elements on Microstructure, Hardness and Growth Rate of Compound Layer in Gaseous-Nitrided Ferritic Alloys 601

the order site, while the remaining three octahedral sites areunstable anti-sites. The presence of excess nitrogen meansthat N also occupies anti-sites in addition to the order sites.As shown in Fig. 12, the activation energy of nitrogen jumpfrom anti-site to the adjacent site is considered to be lowerthan that from the ordered site. Therefore, the presence ofexcess nitrogen increases the diffusivity of N atoms in £ A andpromotes the growth of the compound layer. The Fe­1Sialloy was not measured using 3DAP because Si peak hasoverlapped with that of N in 3DAP measurement. Similarlyas the alloy nitride precipitation in the compound layer ofthe Fe­1Al and Fe­1Valloys, the presence of excess nitrogenseems to be related to the precipitation of alloy nitrides,however, further investigation is required to ascertain thishypothesis. On the other hand, even though the Fe­1Al alloyhas a large amount of excess nitrogen, the compound layerthickness is thinner in comparison with that in pure iron afterlong-term nitriding. This is presumably due to a large amountof AlN precipitation into the diffusion layer during long-termnitriding which consumes nitrogen from the compound layerand overcomes promotion of compound layer growthachieved by excess nitrogen.

Lastly, the thermodynamic effect of the alloying elementson the stability of £ A and ¾ nitride was calculated byThermoCalcμ. In FE-EPMA measurement, no substitutionalelement partitioning was found at the interface between thecompound layer and the diffusion layer therefore, para-equilibrium was assumed. In addition, since the thermody-namic database (TCFE7) used in this study did not includeSi, Al, V, and Mo data for nitrides, thus only Mn and Cr werecalculated. As a result, the addition of Mn and Cr decreasesthe nitrogen activity at the ¡/£ A and £ A/¾ phase boundaries,which clarifies that the addition of Mn and Cr stabilizes both£ A and ¾. It has been reported that Mo destabilizes £ A basedon the solubility measurement,10) but many other researcheshave reported that Mo stabilizes £ A based on gas equilibriummethod, internal friction method, and first-principle calcu-lation,11­13) hence, it is thought that Mo also stabilizes £ A.Therefore, the nitrogen diffusion-controlled growth of thecompound layer is also accelerated thermodynamically bythese nitride stabilizing elements.

In addition, when nitrogen atoms in nitride desorb intomore stable N2 gas, the voids are formed in the nitride.

Consequently, denitrification occurring due to void formationis considered to retard the compound layer growth, and thesuppression of void formation by the addition of Si, Al, andV is one of the factors that to promote the growth of thecompound layer.

4. Conclusion

The phase composition, hardness distribution, and growthrate of the compound layers of gas-nitrided pure iron andFe­M binary alloys were investigated and the followingresults were obtained.(1) In pure iron, ¾ and £ A beneath ¾ are formed together

with the formation of voids near the surface. Nano-indentation measurement shows that £ A is harder than¡, but the void formation causes softening near thesurface.

(2) The diffusion layer is hardened by precipitationstrengthening in Cr, Al, and V added alloys. Thehardness of the compound layer scatters greatly due tothe softening caused by the formation of voids, but themaximum hardness increases slightly by the additionof elements. Moreover, the addition of Si, Al, and Vsuppresses the formation of voids and leads to weakersoftening.

(3) The growth of the compound layer is accelerated bythe presence of excess nitrogen caused by the additionof Si, V, and Al. However, a large amount of AlNprecipitation occurs in the diffusion layer of the Al-added alloy during long-term nitriding, which retardsgrowth of the compound layer.

Acknowledgments

Part of this research was conducted in JST-CREST“Creation of innovative functions of materials based onelemental strategy”.

REFERENCES

1) G. Miyamoto, Y. Tomio, H. Aota, K. Oh-ishi, K. Hono and T.Furuhara: Mater. Sci. Technol. 27 (2011) 742­746.

2) G. Miyamoto, Y. Tomio, T. Furuhara and T. Maki: Mater. Sci. Forum492­493 (2005) 539­544.

3) M.H. Biglari, C.M. Brakman, E.J. Mittemeijer and S. van der Zwaag:Metall. Mater. Trans. A 26 (1995) 765­776.

4) Y. Tomio: Doctoral Thesis, Tohoku University, (2009).5) T. Takase, Y. Nakamura, M. Sumitomo, K. Kita and H. Moshino:

J. Japan Inst. Met. Mater. 40 (1976) 663­669.6) Y. Hiraoka, Y. Watanabe and O. Umezawa: J. Jpn. Soc. Heat Treatment

54 (2014) 313­318.7) S.S. Hosmani, R.E. Schacherl and E.J. Mittemeijer: Int. J. Mater. Res.

97 (2006) 1545­1549.8) S. Meka, R.E. Schacherl, E. Bischoff and E.J. Mittemeijer: Adv. Mater.

Res. 89­91 (2010) 371­376.9) D. Liedtke: Beitrag zum technisch ­ wirtschaftlichen Optimieren des

Nitrocarburierens von Bauteilen Dissertation, TU Berlin (1986).10) M. Sakamoto, K. Masumoto and Y. Imai: J. Japan Inst. Met. Mater. 37

(1973) 343­349.11) K. Kadoma, H. Sudo and K. Kumagai: Report of Committee on Heat-

resisting Metals and Alloys (1965) 61­75.12) W. Köster and W. Horn: Arch. Eisen 37 (1966) 155­160.13) C.S. Zhang: J. Alloy. Compd. 615 (2014) 854­862.

Fig. 12 (a) Crystal structure and a cross-sectional view of £ A and (b)schematic illustration of energy of order-site ( ) and anti-site (©).

F. Meng, G. Miyamoto and T. Furuhara602