Structure change in Fe4N powders by mechanical milling: a new aspect and correction of our previous reports

Journal of Alloys and Compounds 282 (1999) 158–163

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Structure change in Fe 4 N powders by mechanical milling: a new aspect and correction of our previous reports a, a a a b c c K. Sumiyama *, H. Onodera , K. Suzuki , S. Ono , K.J. Kim , K. Gemma , Y. Nishi b

a Institute for Materials Research, Tohoku University, Sendai 980 -8577, Japan Department of Materials Research, Tokai University, Hiratsuka 259 -1207, Japan c Dongbu Advanced Research Institute, Taejon 305 -380, South Korea

Received 5 August 1998

Abstract We have prepared impurity-free Fex N powders (x53.6) and re-examined their structure change by mechanical milling and subsequent ¨ annealing. X-ray diffraction and Mossbauer spectroscopic studies of the milled powders indicate that the initial fcc (g9-) Fe 4 N phase is mechanically deformed and changes to the hcp (e-Fe 3 N type) phase, even though the Fe content is much higher than the stoichiometry of the latter compound. This phase transformation is due to the local structure similarity of Fe atoms in g9-Fe 4 N and e-Fe 3 N, and introduction of hcp type stacking faults along the most dense crystalline plane of (111) in the fcc lattice by shear force. The e-Fe 3 N phase is maintained after annealing at 2308C and it almost returns to the g9-Fe 4 N phase after annealing at 4208C. These results indicate that our previous reports for Fe x N powders (x$4) are not correct: the g9-Fe 4 N phase does not martensitically transform to the bct (a9-) Fe–N phase by mechanical milling and the milled Fe x N powders do not become a0-Fe 16 N 2 by subsequent annealing. The previous results were confused by the coexistence of bcc (a)-Fe, the oxidation in the powder specimens, and the very broad structure-less X-ray diffraction patterns of the powder specimens milled for less than 300 h.  1999 Elsevier Science S.A. All rights reserved. ¨ Keywords: Iron–nitrogen; Mechanical milling; X-ray diffraction; Mossbauer effect

1. Introduction Mechanical milling is a powerful method to produce nonequilibrium solid phases by mechanical energizing and quenching via simple solid-state reactions [1]. By milling metal powders in N 2 gas atmosphere, homogeneous nitrides have been obtained for refractory metal elements, Ti, Nb, etc., via solid–gas reactions [2,3]. However, since no homogeneous Fe nitride has been obtained by milling in an N 2 gas atmosphere [4], Fe nitrides were produced by milling Fe powders in an NH 3 gas atmosphere [5,6]. Ferromagnetic iron nitrides have been extensively studied since the discovery of Fe 16 N 2 which reveals a high magnetic flux density as well as good corrosion resistance for magnetic applications [7]. In a previous study, we | 4) by nitriding Fe powders in prepared Fe x N powders (x 5 an NH 3 gas flow atmosphere at 5508C and examined the structure and magnetic properties of Fe x N and Fe–Fe x Nmixed powders after mechanically milling and subsequent annealing [8–10]. We described that the fcc g9-Fe 4 N phase *Corresponding author.

martensitically transformed into the bct (a9-) Fe–N phase | 4) in these milled powders and the a9-Fe x N powders (x 5 became a0-Fe 16 N 2 by annealing. There have been a few papers which also described the a9-Fe–N phase formation in mechanically milled Fe powders in an N 2 atmosphere [11] and ion-implanted surfaces of bulk Fe [12], the a0-Fe 16 N 2 formation in bcc (a-)Fe1e-Fe 3 N powder mixtures obtained by mechanical milling and annealing [13,14], and the phase transformation of e-Fe 3 N (x.4) to a0-Fe 16 N 2 in an Fe x N thin plate (x.4) by quenching and annealing [15]. However, the other papers avoided or denied these statements: no martensitic transformation in mechanically milled a-Fe1g9-Fe 4 N and a-Fe1e-Fe 3 N powder mixtures, and no a0-Fe 16 N 2 formation by annealing of these milled powders [16–18]. In particular, Koyana et al. gave us the following critical comments (T. Koyana, T. Fukunaga, U. Mizutani, private communication). to our previous papers [8–10]: (1) the low index diffraction peaks and the very large hyperfine field component of 430 kOe which were observed in our previous milled Fe x N speci| 4) after annealing could not be ascribed to mens (x 5 formation of a0-Fe 16 N 2 , but they might correspond to the

0925-8388 / 99 / $ – see front matter  1999 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 98 )00813-5

K. Sumiyama et al. / Journal of Alloys and Compounds 282 (1999) 158 – 163 Table 1 Impurity composition (wt%) of the initial Fe powders

159

were observed at room temperature with a constant acceleration mode using 57 Co doped in a Rh matrix.

C

S

P

Si

Mn

Cu

N

O

0.0053

0.0053

0.0011

0.0032

0.0043

0.0002

0.0020

0.432

3. Results formation and / or contamination of magnetite, Fe 3 O 4 and (2) the a9-Fe–N phase formation in these milled powders was doubtful because the diffraction peaks at around the | 408 could be attributed to formation diffraction angle, 2u 5 of the hcp Fe 3 N (e-Fe) phase [19]. Since we were not particular about the Fe powders in the previous studies [8–10], our previous experiments might have suffered from oxidation and contamination. In | 4) the present study, we have prepared Fe x N powders (x 5 from impurity-free Fe powders and re-examined their structure change with milling and annealing through X-ray ¨ diffraction and Mossbauer effect measurements.

2. Experimental procedures Pure Fe powders 25–75 mm in diameter were supplied by Showadenko Co. Ltd. The results of chemical analyses shown in Table 1 indicate that the initial Fe powders are reasonably pure for the present purpose. These Fe powders were nitrided by a conventional method: they were annealed at 5508C for 30 h in an NH 3 flow atmosphere of 760 Torr. The nitrized powders were sealed in a milling pot with a powder-to-ball weight ratio of 1:7. A vibrating ball-mill system was operated at a frequency of 9.6 Hz at room temperature. The milled powders were sealed in glass tubes in a vacuum state of about 60 mTorr, and annealed at TA 5230 and 4208C. The chemical compositions of nitrogen and oxygen in milled powders were determined by a He-gas carrier fusion and heat conduction method. The results are shown in Table 2. The nitrogen content does not change by milling and contamination of oxygen is not so serious in the present powder specimens. The nitrogen content of about 22 at% is off-stoichiometric to the nitrogen-rich side in the present powders, whereas it was almost stoichiometric or off-stoichiometric to the Ferich side in the previous powders. The structures of the as-milled and annealed powders were characterized by X-ray diffraction with Cu Ka radia¨ tion and a graphite monochromator. Mossbauer spectra

3.1. X-ray diffraction [20] Fig. 1 shows the X-ray diffraction patterns of Fe x N powders (x53.6) milled for several periods, t M , together with those of the initial pure Fe powders and the asnitrided ones. Clear diffraction peaks corresponding to bcc Fe and no peaks of Fe oxides are detected in the initial Fe powders. Sharp diffraction peaks corresponding to g9-Fe 4 N and no peaks of a-Fe are detected for the as-nitrided powders with t M 50 h. After milling for t M 5200 h, very broad diffraction peaks are detectable in the region of 2u 535–508 and no clear peak is found at the higher angle range. It is difficult to determine the crystal structure of the milled powders in this milling stage. After milling for t M 5400 h, the diffraction peaks split and a low angle diffraction peak becomes distinct in the region of 2u 535– 508, being attributed to e-Fe 3 N. After milling for t M 5600 h, these diffraction peaks and those at a high-angle region of 2u .508 become more marked. Fig. 2 shows the X-ray diffraction patterns of Fe x N powders (x53.6) milled for t M 5400 h and annealed for 1

Table 2 Chemical composition (wt%) of nitrogen and oxygen in the initial Fe, as-nitrided Fe x N powders and Fe x N powders milled at several periods, t M Specimen

N

O

Initial Fe As-nitrided t M 5200 h t M 5400 h t M 5600 h

0.015 6.36 6.44 6.40 6.43

0.411 0.178 0.424 0.565 0.737

Fig. 1. X-ray diffraction patterns of pure Fe powders and Fe x N powders (x53.6) mechanically milled for several milling times, t M .

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K. Sumiyama et al. / Journal of Alloys and Compounds 282 (1999) 158 – 163

Fig. 3b. There are three broad peaks at around 290, 220 and 110 kOe, whereas a tiny peak below 40 kOe is probably due to an artifact in the analysis. ¨ Fig. 4a shows the Mossbauer spectra of Fe x N powders (x53.6) milled for t M 5400 h and annealed for 1 h at several TA . The sextet spectrum becomes narrower after annealing at TA 52308C. However, the Hhf distribution is still marked as shown in Fig. 4b: the distribution peaks at around 340, 290, 220 and 110 kOe become remarkable. On the other hand, well-split sextets are detected after annealing at TA 54208C, being attributed to the three components of g9-Fe 4 N as listed in Table 3. The intensity ratios of FeI:FeIIA:FeIIB521:51:28 are slightly different from the results for the as-nitrided powders [21].

4. Discussion

4.1. As-milled state

Fig. 2. X-ray diffraction patterns of Fe x N powders (x53.6) mechanically milled for t M 5400 h and subsequently annealed at TA for 1 h.

h at several temperatures, TA . Several diffraction peaks observed in the as-milled states become clearer after annealing at TA 52008C. The peak widths become sharp, indicating that the e-Fe 3 N structure is maintained but it partially changes to the g9-Fe 4 N structure. The diffraction peaks corresponding to g9-Fe 4 N again become predominant and those of e-Fe 3 N are slightly detectable after annealing at TA 54008C. However, no peak corresponding to a0-Fe 16 N 2 can be detected under the present preparation conditions. ¨ 3.2. Mossbauer effect ¨ Fig. 3 shows the Mossbauer spectra of Fe x N powders (x53.6) milled for t M together with those of the as-nitrided ones. In the as- nitrided powders (t M 50 h), well split sextets are detected, being interpreted as three sextet components. In the g9-Fe 4 N structure, there are crystallographically two non-equivalent Fe atom sites, FeI and FeII. The FeII site consists of two different electric field gradient components, A and B, with the intensity ratios of FeI:FeIIA:FeIIB51:2:1 [21]. The relative fraction of these three components obtained by a least-square fitting are identical to the reported results as listed in Table 3. Since the sextet spectra become very broad after milling and cannot be resolved into simple sextet components, we fitted the calculated spectrum with a distribution of hyperfine fields, Hhf , to the observed spectra by a least-square fitting [22]. We assume a proportional change of the isomer shift, IS with Hhf . The obtained distributions are shown in

As shown in Fig. 1, the crystal structure is severely distorted and cannot be determined from the present Fe x N powders (x53.6) milled for t M 5200 h, while the e-Fe 3 N structure is formed after milling of g9-Fe x N powders (x5 3.6) for t M 5400 h. However, Fig. 3 does not clearly show the e-Fe 3 N phase formation because of the wide distribution of Hhf in these milled powders. In order to understand this structure change schematically, we describe the crystalline units of the g9-Fe 4 N and e-Fe 3 N structures in Fig. 5. In the g9-Fe 4 N structure, Fe atoms at the corner of the cube contains six nearest neighbor (n.n.) unoccupied interstitial sites and those at the face center of the cube two n.n. occupied N interstitial sites and four n.n. unoccupied sites. In the e-Fe 3 N structure, all Fe atoms have equivalent nearest neighbors: six n.n. Fe atom sites and six octahedral interstitial sites, where the latter consists of two n.n. N atom sites and four unoccupied sites. In Fe x N with x.3, some of the n.n. N atom sites become vacant for x.3, while some n.n. unoccupied sites are filled by N atoms for x,3. Here, the environments of Fe atoms in the face centers of the g9-Fe 4 N and those of the e-Fe 3 N phases are similar. They both have the same number of n.n. atoms, i.e., Fe atoms, unoccupied and N-occupied interstitial sites. Only the directional arrangement of these neighbor is different. The hyperfine fields of Fe atoms sensitively depend on the number of n.n. N atoms, Nn , in the g9-Fe 4 N and e-Fe 3 N phases [21,23,24]: Hhf is 345 kOe for Nn 50 and 215 kOe for Nn 52 in g9-Fe 4 N, while Hhf is 295 kOe for Nn 51, 238 kOe for Nn 52 and 100 kOe for Nn 53 in e-Fe x N with x between 2 and 3. It is proposed that shear force is very important for alloy and compound formation by mechanical milling [25] because it enhances mixing and kneading of different elemental powders. During milling the g9-Fe x N powders | 4), the shear force induces displacement of Fe and N (x 5 atom planes on the most dense plane of (111) in the fcc

K. Sumiyama et al. / Journal of Alloys and Compounds 282 (1999) 158 – 163

161

¨ Fig. 3. (a) Mossbauer spectra at room temperature of Fe x N powders (x53.6) mechanically milled for several milling times, t M and (b) the distribution of hyperfine field fitted to the spectra for t M ±0 in (a).

lattice. The fcc stacking (ABCABC???) is converted to the hcp stacking (ABAB???) by the displacement along the k112l direction, changing the n.n. unoccupied interstitial sites to the n.n. N-occupied interstitial sites as shown in Fig. 5a. This leads to the populations of the Fe atom

components, Nn 51 and Nn 53, as well as the reduction of the Fe atom component, Nn 50 [23]. In Fig. 3b, the distribution peaks at around 290, 220 and 110 kOe support the formation of e-Fe x N (x53.6) in the mechanically milled specimens. It is worthwhile to call the result on

Table 3 ¨ Mossbauer parameters, the hyperfine field, Hhf , the isomer shift, IS, the quadrupole splitting, QS, the width of resonance lines and the intensity ratio of the individual sextet obtained by fitting the three sextets to the observed spectra for the as-milled Fe x N powders with x53.6 (see Table 2) and the ones annealed at 4208C Specimen

Fe atom site

H hf (kOe)

IS (mm / s)

QS (mm / s)

Width (mm)

Intensity (%)

As-milled Fe–N powders

FeI FeIIA FeIIB FeI FeIIA FeIIB

342 216 219 342 218 223

0.23 0.31 0.30 0.23 0.31 0.30

0.0 20.20 0.42 0.0 20.15 0.30

0.22 0.24 0.25 0.24 0.23 0.23

25 50 25 21 51 28

Annealed Fe–N powders

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K. Sumiyama et al. / Journal of Alloys and Compounds 282 (1999) 158 – 163

Fig. 5. Local structures of Fe atoms. (a) g9-Fe 4 N unit, (b) e-Fe 3 N and (c) e-Fe x N (x.3) from g 9-Fe 4 N by shear displacement along a k112l direction on a h111j plane.

¨ Fig. 4. (a) Mossbauer spectra at room temperature of Fe x N powders (x53.6) mechanically milled for t M 5400 h and subsequently annealed at TA for 1 h and (b) the distribution of hyperfine field fitted to the spectrum for TA 52308C in (a).

milled e-Fe x N powders (x52.3) [26]: the e-Fe 3 N phase is maintained even after milling for t M 5300 h. These results suggest that the e-Fe 3 N phase is more stable than the g9-Fe 4 N phase against mechanical stressing [5,6,18]. In the previous study, we milled g9-Fe x N powders | 4) and a-Fe and g9-Fe 4 N mixed powders up to t M 5 (x 5 300 h, and discussed their martensitic transformation to the a9-Fe–N phase by milling. In the region of 2u 535–508 of

the diffraction patterns of the previous specimens, the superposition of a broad bcc Fe peak made our analysis vague [8,9]. In mechanically milled a-Fe1g9-Fe 4 N mixed powder specimens whose N content were lower than the stoichiometry of g9-Fe 4 N, such severely distorted diffraction patterns have been attributed to an a-Fe–N phase whose N content is much higher than the equilibrium a-Fe–N ferrite [13,16,26] or a nanocrystallline a-Fe phase with an amorphous-like phase [6,16,18]. In the present Fe x N (x53.6) powder specimens, since the N content is higher than the stoichiometry of g9-Fe 4 N and the g9-Fe 4 N phase is predominant without a-Fe coexistence, it is plausible that the partition of the conceivable phases, a-Fe, g9-Fe 4 N and e-Fe 3 N is different from those of the previous Fe-rich off-stoichiometric specimens. In the X- ray diffraction pattern of the present specimen milled for t M 5400 h, however, the e-Fe 3 N structure is formed, while neither a-Fe–N nor a9-Fe-N structures are detected. The distorted X-ray diffraction pattern does not reveal the exact crystal structure of the present specimen for t M 5200 h, being similar to the previous results [8– 10,13,14,16,18]. Moreover, the distribution of Hhf is similar for t M 5200 and 400 h in Fig. 3b. These results

K. Sumiyama et al. / Journal of Alloys and Compounds 282 (1999) 158 – 163

indicate the formation of the e-Fe 3 N structure by mechanical milling of g9-Fe 4 N even for t M 5200 and deny the martensitic transformation from the g9-Fe 4 N phase to the a9-Fe–N phase in milled powders (T. Koyana, T. Fukunaga, U. Mizutani, private communication). The previous experiments were confused by the coexistence of a-Fe with g9-Fe 4 N.

4.2. Annealed state ¨ The X-ray diffraction patterns and the Mossbauer spectra clearly demonstrate the phase change in the milled Fe x N powders (x53.6) by annealing. In Fig. 2, the e-Fe 3 N structure obtained for t M .200 h is maintained after annealed at 2308C for 1 h. However, it partially transforms to g9-Fe 4 N. The Hhf distribution in Fig. 4 estimated from ¨ the Mossbauer spectra in Fig. 4 also supports this result: the Hhf distribution curve reveals peaks at around 300, 220 and 110 kOe (ascribable to e-Fe 3 N) and no clear peak at around 240 and 340 kOe (ascribable to g9-Fe 4 N), where the peak at around 220 kOe probably overlaps with that at around 240 kOe. In Fig. 2, the e-Fe 3 N structure almost returns to g9-Fe 4 N after annealing at 4208C for 1 h [8–10,16,18]. As shown in Table 3, however, the intensity ratios of FeI:FeIIA:FeIIB5 21:51:28 for this specimen are different from those for the as-nitrided specimen [21]. The reduction of the FeI component and the increase of the FeII component are ascribed to slight retainment of the e-Fe 3 N phase, presence of a partial disordering and the off-stoichiometry in the annealed specimen. These results deny the results of previous experiments [8–10]: the a0-Fe 16 N 2 phase is not formed by annealing of mechanically milled g9-Fe x N powders. The previous confusion originates from the very broad X-ray diffraction peaks in the low-angle region (2u ,408) and the very broad hyperfine field component at around 430 kOe, which are mainly ascribed to the oxidation of powder specimens [13,14,18] (T. Koyana, T. Fukunaga, U. Mizutani, private communication). 5. Conclusion We prepared impurity-free g9-Fe x N powders (x53.6) ¨ and studied the X-ray diffraction and Mossbauer effect of these powders after mechanical milling and subsequent annealing. The mechanical stress induced by milling gives rise the shear displacement of the Fe and N atoms in the most dense plane of (111) fcc , leading to the structural transformation from the fcc (g9-Fe 4 N) to hcp (e-Fe 3 N) structures. The e-Fe 3 N phase produced by milling mainly transforms back into the g9-Fe 4 N phase by annealing. Therefore, we have to correct our previous conclusions as follows: there is no martensitic transformation from g9Fe 4 N to a9-Fe–N after prolonged milling and no partial formation of the a0-Fe 16 N 2 phase after annealing the mechanically milled powders.

163

Acknowledgements The authors wish to thank Dr K. Takada for chemical analyses of milled specimens and Mr Y. Kinoshita for supplying pure Fe powders. They also appreciate Dr T. Koyano in Tsukuba University, and Dr T. Fukunaga and Professor U. Mizutani in Nagoya University for their critical comments, which motivated the present study. This work was partially supported by a Grant-in-Aid for Scientific Researches (Grant No. 06452323) given by the Ministry of Education, Science, Culture and Sports, Japan.

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