EELS study of niobium carbo-nitride nano-precipitates in ferrite

Micron 37 (2006) 492–502 www.elsevier.com/locate/micron

EELS study of niobium carbo-nitride nano-precipitates in ferrite E. Courtois a, T. Epicier a,*, C. Scott b a

GEMPPM, umr CNRS 5510, Baˆt. B. Pascal, INSA de Lyon, F-69621 Villeurbanne cedex, France b Arcelor Research SA, Voie Romaine-BP30320, F-57283 Maizie`res-le`s-Metz cedex, France Received 26 July 2005; received in revised form 21 October 2005; accepted 21 October 2005

Abstract Micro-alloying steels allow higher strength to be achieved, with lower carbon contents, without a loss in toughness, weldability or formability through the generation of a fine ferrite grain size with additional strengthening being provided by the fine scale precipitation of complex carbonitride particles. Niobium is reported to be the most efficient micro-alloying element to achieve refinement of the final grain structure. A detailed microscopic investigation is one of the keys for understanding the first stages of the precipitation sequence, thus transmission electron microscopy (TEM) is required. Model Fe–(Nb,C) and Fe–(Nb,C,N) ferritic alloys have been studied after annealing under isothermal conditions. However the nanometre scale dimensions of the particles makes their detection, structural and chemical characterization delicate. Various imaging techniques have then been employed. Conventional TEM (CTEM) and high resolution TEM (HRTEM) were used to characterise the morphology, nature and repartition of precipitates. Volume fractions and a statistical approach to particle size distributions of precipitates have been investigated by energy filtered TEM (EFTEM) and high angle annular dark field (HAADF) imaging. Great attention was paid to the chemical analysis of precipitates; their composition has been quantified by electron energy loss spectroscopy (EELS), on the basis of calibrated ‘jump-ratios’ of C–K and N–K edges over the Nb–M edge, using standards of well-defined compositions. It is shown that a significant addition of nitrogen in the alloy leads to a complex precipitation sequence, with the co-existence of two populations of particles: pure nitrides and homogeneous carbo-nitrides respectively. q 2005 Elsevier Ltd. All rights reserved. PACS: 68.37.Lp; 79.20.Uv; 61.66.Dk; 64.60.My Keywords: TEM; EELS; Micro-alloyed steels; Carbide precipitation

1. Introduction Computer-assisted metallurgy is a recent trend that is more and more employed by manufacturers to help speed up the development of new kinds of alloys. It consists in modeling the final mechanical properties of a material designed to fulfill the requirements of an industrial application, as a function of the composition and applied thermo-mechanical treatments. To make this approach valid, preliminary studies are required. These studies basically consist in validating and/or fitting the predictions of thermodynamic models via an accurate and representative microstructural characterization. This is particularly needed in the case of processes involving precipitation in alloys. Here, models are generally based on the classical * Corresponding author. Tel.: C33 472 4384 94; fax: C33 472 4388 30. E-mail addresses: [email protected] (E. Courtois), thierry. [email protected] (T. Epicier).

0968-4328/$ - see front matter q 2005 Elsevier Ltd. All rights reserved. doi:10.1016/j.micron.2005.10.009

theory for diffusive phase transformation (Kampmann and Wagner, 1991), and treat simultaneously the nucleation, growth and ripening phenomena (Shercliff and Ashby, 1990). The particle size distribution, their number and volume fraction can be calculated (Deschamps and Brechet, 1999; Myhr and Grong, 2000) and from these values the effect of the precipitates on the mechanical properties is predicted (Perez and Deschamps, 2003; Maugis and Goune, 2005). Nevertheless, this theoretical approach needs experimental data to compare with and confirm the models. In the present work, an extensive experimental study of the precipitation of niobium carbo-nitrides in ferrite is presented. This work refers to micro-alloyed steels, which have received considerable interest over many years because they represent very good candidates for a wide range of industrial applications. In this respect, a small addition of niobium to steel is of special interest, since it is well known to yield significant improvements in mechanical properties (Palmiere et al., 1996; DeArdo, 1998; Charleux et al., 2001). Because of its strong affinity for carbon and nitrogen, niobium forms a fine

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Table 1 Composition of laboratory Fe–Nb–C–N steels [wt ppm] Alloy

Nb

C

N

S

Mn

Al

O

[C]/[N]a

[CCN]/[Nb]a

Steel A Steel B

790 843

120 59

10 64

10 23

20 13

60 20

13 44

14.0 1.1

1.28 1.0

a

Atomic fraction ratio.

dispersion of niobium carbo-nitride precipitates, generally with a fcc NaCl type structure, in a ferrite (cc structure with aZ 0.2866 nm). This fine precipitation inhibits austenite grain coarsening during heating (Fossaert et al., 1995), and, in certain cases, suppresses austenite recrystallisation prior to the g/a transformation through the strain-induced precipitation of NbC (DeArdo, 1998). One of the most difficult aspects of any experimental characterization of the precipitation state is due to the nanometric size of the precipitates that are formed. Hence, transmission electron microscopy work has generally to be undertaken, but further difficulties arise: (i) in thin foils, the magnetic nature of the iron matrix makes any systematic analysis very difficult; (ii) even on extraction replicas, the classical use of carbon as a supporting film makes it almost impossible to quantify the chemistry of the carbo-nitrides, especially their carbon content. Several TEM works have already been published on such so-called high-strength low alloys (HSLA) (Hofer et al., 1996; Craven et al., 2000; Rainforth et al., 2002; Wilson and Craven, 2003; Beres et al., 2004). However, and owing to the previously mentioned difficulties, there is still a lack of comprehensive study of the evolution of the precipitation state as a function of the annealing time and temperature. Moreover, it is hard to find any statistical approach to the size distribution, volume fraction, crystallography and chemistry of the particles. Such an ambitious approach has been undertaken on two model (Fe–Nb–C–N) alloys, with significantly different nitrogen contents (Courtois, 2005a; Courtois, et al., 2005); a large part of this work has been devoted to quantitative measurements of size distributions, combining all available imaging techniques on both thin foils and replicas (e.g. conventional dark-field TEM, high-angle annular dark-field imaging in the STEM mode, high resolution TEM and energyfiltered TEM). In this paper, however, we will mainly focus on the quantitative analysis of the chemistry of the precipitates by electron energy-loss spectroscopy (EELS), which will be developed in Section 4, after a presentation of experimental methods (Section 2) and a brief survey of imaging in Section 3.

annealing times, e.g. 5 and 30 min, or in vacuum—quartz encapsulation—for longer annealing times, e.g. 300 min and 126 h), followed by water quenching. A final polygonal ferritic microstructure was obtained in both cases, with a regular grain size of a few mm3, as shown by Fig. 1. Thin foils were obtained by the conventional method of careful grinding to produce a thin disc, followed by final thinning to electron transparency by electropolishing using a solution of 60% methanol-35% butyoxyethanol-5% perchloric acid at K30 8C (Rainforth et al., 2002). Extraction replicas based on aluminum for carbon analysis were prepared by evaporating an amorphous alumina (AlOx, x!3/2) film onto a bulk sample that had been polished and lightly pre-etched. Depositions were carried out as described in (Scott et al., 2002). 2.2. Transmission electron microscopy and associated techniques Electron microscopy was performed using a JEOL 2010F field emission gun transmission electron microscope operating at 200 kV. The microscope was fitted with an Oxford EDX (energydispersive X-ray) analyser and a Gatan DigiPEELS spectrometer with a standard photodiode array detector. EELS spectra were acquired according to the following conditions (unless otherwise specified): probe convergent half-angle of 11 mrad, probe size of 1–2 nm, collection half-angle of 9.4 mrad, energy resolution (FWHM of the zero-loss peak) of 1.2C/K0.2 eV, acquisition times of individual core loss EELS spectra between 3 and 5 s. In order to avoid spurious carbon signals in the TEM, the aluminabased replicas were cooled down to K170 8C. High-resolution imaging was performed on both replicas and thin foils, which had previously been cleaned in a Gala

2. Experimental methods 2.1. Materials and sample preparation The chemical composition for the two steels investigated is reported in Table 1. Both A and B alloys were solution treated at 1250 8C to dissolve the precipitates, and then water quenched. Precipitation was induced through subsequent annealing at temperatures between 650 and 800 8C (performed in salt bath for short

Fig. 1. Grain structure of steel A in the initial state: (a) scanning electron micrograph showing the grain boundaries revealed by a short 3% nital attack; (b) TEM bright field image.

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Fig. 3. Dark field TEM micrographs (operating reflection gZ(111)fcc in a [001]Fe zone axis) showing: (a) precipitates along a grain-boundary; (b) alignment of precipitates along dislocations in alloy A (after a treatment of 30 min at 800 8C).

niobium elemental map using the main window centred on the Nb–N2,3 edge (onset at 34 eV) and two pre-edge windows (3 eV wide at 27 and 32 eV respectively in order to extrapolate the background under the ionization Nb–N2,3 edge). 3. Imaging techniques Fig. 2. EFTEM treatment for imaging carbo-nitrides within ferrite: the energy positions of the three windows for the background extrapolation (labels W2 and W1 respectively) and the Nb–N2,3 imaging (MAX) are indicated (microscope LEO 912).

model PlasmaPrep5 plasma cleaner (the gas used was 6% 02, 94% Ar). Further high angle dark field (HAADF) imaging was performed using a JEOL annular detector fitted on the microscope. Finally, energy filtered transmission electron microscopy (EFTEM) was performed on two instruments: (i) a LEO 912 fitted with an Omega filter and operating at 120 kV, (ii) a JEOL3010 operating with a LaB6 filament at 300 kV and equipped with a Gatan imaging filter (GIF). EFTEM was essentially used in order to image the precipitates either on replicas or within thin foils, with the intention to evaluate the possibilities of this technique regarding the measurement of the volume fraction of the precipitates. Spectroscopic images were thus obtained using the Fe–M3,2 and Nb–N2,3 edges. The threewindow method, illustrated by Fig. 2, was used to generate a

3.1. Conventional TEM Conventional TEM can be extensively used to study the general microstructure of the alloys, especially the precipitation state. In the particular case of metallic carbides within ferrite, the occurrence of an orientation relationship (see Section 3.2) between the precipitates and the matrix makes it easy to image the particles in dark field, as shown in Fig. 3. As previously observed in the system Fe–Nb–C(–N) (Rainforth et al., 2002; Perrard, 2004), these micrographs show that the precipitation is essentially heterogeneous, since most of precipitates are aligned along grain-boundaries (Fig. 3a) and dislocations (Fig. 3b). Although conventional TEM work can be undertaken for statistical studies (e.g. precipitation size and distribution (Davis and Strangwood, 2002), some limitations appear: (i) the spatial resolution as well as the detectability of the smallest particles is limited, (ii) relaxation of Laue conditions, due to the unavoidable bending of steel thin foils, may cause some of the precipitates to be

Fig. 4. h001iFe HRTEM images of niobium carbo-nitrides precipitated within alloy A (after a treatment of 30 min at 800 8C): (a) [100]Fe//[110]fcc azimuth (edge-on view of a lens-shaped particle); (b) [001]Fe//[001]fcc azimuth (plane view). The indexing of diffractograms confirms the Baker–Nutting orientation relationship (for simplicity, the precipitates are labeled as ‘NbC’). In both images, Moire´s occur due to overlapping of the matrix over the precipitates.

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Fig. 5. HRTEM images of small precipitates (all observed along a h110ifcc zone axis) successfully extracted from alloy A (after a treatment of 60 min at 600 8C).

‘missed’ in dark field images (Courtois, 2005). Hence, HRTEM and HAADF on extraction replicas have been performed for a more accurate study. 3.2. HRTEM work In the case of thin foils, HRTEM has clearly confirmed that most precipitates adopt a fcc NaCl type structure, with afcc varying from 0.437 to 0.447 nm (Ohmori, 1975) as observed

for bulk niobium nitride, carbide and carbo-nitride close to the composition Nb(C,N)z1 (see Fig. 4). Moreover, these precipitates adopt the expected Baker–Nutting orientation relationship (Baker and Nutting, 1959) with the cc ferrite matrix: ½001
Fe ==½001
fcc and ½100
Fe ==½110
fcc

(1)

From these images, it is shown that precipitates are lensshaped, with a symmetry axis along the [001]fcc direction.

Fig. 6. HAADF imaging of niobium carbo-nitrides (alloy A, after a treatment of 30 min at 800 8C): (a–b) same area imaged in conventional bright field and HAADF respectively; note that particles 3 and 4 are easily identified in the HAADF mode, whereas they appear in poor diffraction contrast in (a); (c) EDX analysis of particle 4 confirming its chemical nature (almost stoichiometric carbide NbCz1); (d) HAADF image of another area showing occasional particles lying at the edge of a cracked replica (arrow); (e) comparison of the EELS spectra obtained respectively on an overhanging particle (arrowed in (d), top spectrum) and on the AlOx replica (bottom spectrum).

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Fig. 7. EFTEM imaging of a carbo-nitride within ferrite (alloy A after a heat treatment of 30 min at 800 8C): (a) elastic image; (b) Fe–M2,3 negative image; (c) Nb– N2,3 image (microscope LEO 912).

In the case of extraction replicas, HRTEM has clearly shown that precipitates as small as 3 nm can be successfully extracted from the matrix (Fig. 5). 3.3. HAADF on replicas Owing to the large difference in atomic number between niobium and aluminium chemical species (Z equal respectively to 41 and 13), the HAADF contrast of precipitates lying over an AlOx film is suitable for an easy detection of particles on extraction replicas, as illustrated by Fig. 6a–c. Then, precipitates quickly detected in HAADF were analysed in EELS as described in Section 4. It should be noted that some particles were lying by chance at the edge of cracked replica (see Fig. 6d), making it possible to perform EELS analysis without any spurious signal coming from the supporting film (top spectrum in Fig. 6e). However, the use of AlOx extraction replicas in the present study has represented a great advantage compared to studies where carbon extraction replicas have been used (e.g. Wilson and Craven, (2003)), since it was not necessary to restrict the analysis to such overhanging precipitates. It can further be shown that the relatively low thickness of the replica (between 5 and 15 nm) did not affect the quality of the signal for the edges of interest (see Fig. 6e). 3.4. EFTEM imaging According to the experimental procedure given in Section 2.2, Fe–M2,3 and Nb–N2,3 EFTEM imaging of precipitates within thin foils, as illustrated in Fig. 7, has been undertaken.

In order to estimate the volume fraction of precipitates, hundreds of particles have further been imaged over a total area of several tens mm2 within thin foils (see Fig. 8). From these results, the local volume fraction ri in each elementary image has been deduced from the following simple relation: ri Z

Viprecipitate Vimatrix

(2)

The precipitated volume Viprecipitate is obtained as: Viprecipitate Z

X 4prF3

(3)

3

Where rF is the mean radius (or Ferret radius, that is the radius of the circle, the area of which is equal to the projected area of the precipitate) measured for each particle. Thus the precipitate volume is approximated by that of an equivalent sphere of radius rF; although this can be considered as a rough estimate, this approximation gave reasonable results in the quantitative interpretation of small-angle neutron scattering experiments (Perrard, 2004). The matrix volume is determined according to the measurement of the thickness ti of each area by EELS. For this purpose, the inelastic mean free-path lp for ferrite at 300 kV has been calculated to be 120 nm according to Egerton’s formula (Egerton, 1996). Then for each image i, the matrix volume Vimatrix is simply the product of the viewing area multiplied by ti. Finally the total volume fraction rtotal is the average of all local volume fractions ri. The resulting value is reported in Table 2, and compared to other measurements by independent techniques. It can be seen from these results that the EFTEM method provides an excellent estimate of the precipitate volume fraction. Table 2 Precipitate volume fraction rtotal as measured by different techniques (alloy A treated 30 min at 800 8C)

Fig. 8. Typical montage of Nb–N2,3 EFTEM images obtained on alloy A (after a treatment of 30 min at 800 8C), and showing alignment of precipitates along a grain-boundary—upper part of the figure—and along dislocations—left-hand side (microscope JEOL 3010).

Measurement method

EFTEM

Electrolytic dissolutiona

Small-angle neutron scattering (Perrard, 2004)

Multi-preci simulations (Perrard, 2004)

Volume fraction (%)

0.084

0.082

0.08

0.08

a Bulk method: several mm3 of matrix are dissolved, and the carbide residue is measured (Maugis et al., 2002).

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Fig. 9. HAADF image of a replica of steel B treated 5 0 at 650 8C and an EELS spectrum from the arrowed particle.

4. EELS quantification of the carbon and nitrogen contents of precipitates 4.1. Methodology Fig. 9 shows a typical EELS spectrum of a 10 nm niobium carbo-nitride precipitate observed on an extraction replica. For each analyzed precipitate, an additional spectrum was acquired on the replica itself about 10 nm away in order to confirm the absence of any significant carbon content within or on the replica (this point will be re-discussed in Section 4.2, devoted to the estimation of the final accuracy in chemical composition measurements). In order to perform an elemental analysis of the relative C/Nb and N/Nb ratios, two main difficulties have to be overcome: (i) as can be seen from Fig. 6e or Fig. 9, the carbon and nitrogen K-edges (near 287 and 401 eV respectively) overlap the ‘delayed-type’ niobium M-edge (starting near 207 eV), which makes the classical background subtraction by a classical power-law very hazardous (especially for the C–K edge). (ii) the use of the well-known relationship (Egerton, 1996): NC sM ðb; DNb Þ !ICK ðb; DC Þ Z Nb M NNb sK C ðb; DC Þ !INb ðb; DNb Þ

(4)

(where NX represent the number of atoms ‘X’ (XZC or Nb), sY X is the inelastic cross-section for the ‘Y’ edge (YZK or M), IXY is the X-edge area integrated over the energy-window DX and b is the effective collection angle) is not possible due to the

inaccuracy of any available calculated or parameterized crosssection for the Nb–M edge. Hence, a procedure involving reference spectra has been developed; for that purpose, five bulk ‘standards’ were used: one stoichiometric Nb1N1 powder (supplied by ABCR GmbHw) and four NbC1Kx compounds (see Table 3). Regarding the latter materials, the chemical composition has been calibrated using previously established relationships between the departure from stoichiometry (subscript x) and the lattice parameter afcc of the NaCl-type structure of the mono-carbide (Storms, 1967; Xu et al., 2000). The Nb6C5 powder is of particular interest since at this composition longrange ordering of carbon vacancies occurs, which makes it easy to determine confidently the departure from stoichiometry by electron diffraction and/or HRTEM (see Fig. 10), according to well-established crystallographic models for the M6C5 ordered phases (Epicier et al., 1989; Epicier, 1990). For each of these ‘standards’, the jump ratio of the K-edge of interest (C–K or N–K) has been measured as the ratio R of edge areas (respectively AC and ANb under the C–K and Nb–M edges) integrated within optimised energy windows as schematically illustrated by Fig. 11a in the case of niobium carbides. Then linear relationships were deduced between the measured jump-ratio and the chemical composition (see Fig. 11b for niobium carbides): RC Z 0:3048

½C
C 0:9606 ½Nb


=5= for niobium carbides

RN Z 0:2061

½N
C 0:9492 ½Nb


=5bis= for niobium nitrides

Table 3 Detail of niobium carbides ‘standards’ ‘Standard’

Origin

Lattice parameter (nm) deduced from X-ray diffractiona

Chemical composition deduced from the variation of afcc vs. x in NbC1Kx

NbC powder NbCz0.9 single crystal Nb6C5 powder NbCz0.8 single crystal

Commercial product (Johnsonw) (Kumashiro and Sakuma, 1981) (Landesman et al., 1985) (Kumashiro and Sakuma, 1981)

0.44682G0.0001 0.44601G0.0001 0.44568G0.005b 0.44440G0.0001

0.951G0.016 0.858G0.009 0.833G0.009 0.756G0.010

a b

lCu Ka1Z0.154056 nm, Rigaku diffractometer. The large imprecision arises from a z2‰ distortion of the fcc lattice due to carbon vacancy ordering (Landesman et al., 1985).

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Fig. 10. Crystallographic study of Nb6C5: (a) electron diffraction pattern showing superlattice reflections due to the ordering of carbon vacancies according to the monoclinic (subscript ‘m’) model (from Billingham and Lewis,(1972)) depicted in (b); (c) HRTEM superstructure image showing carbon-vacancy columns (bright dots) along the [1–21]fccZ[10–1]m azimuth ((111)fcc planes are horizontal).

Fig. 11. Quantitative analysis of the C–K ‘jump-ratio’ in the case of niobium carbides: (a) definition of integration energy windows (see text); (b) linear relationship between the jump-ratio factor R and the chemical composition of the five ‘standards’.

It should be mentioned that a mean-least square fitting (MLSF) procedure has also been applied in order to treat the (Nb,C,N) EELS spectra after a classical Fourier-ratio deconvolution of multiple scattering. The reference spectrum for the Nb–M edge was obtained from a pure niobium powder, and subsequently subtracted from NbC and NbN ‘standards’ spectra in order to obtain C–K and N–K reference edges respectively. Although this MLSF approach gave reasonable results (see for example Fig. 12), the jump-ratio method was preferred for the following reasons:

4.2. Estimation of the composition error Since one of the goals of the present study is to obtain quantitative information on the chemical composition of the carbo-nitride precipitates, it appears important to estimate the accuracy of the analysis.

(i) calculation of the jump-ratio is faster that the MLSF method; (ii) the jump-ratio appeared to be more robust with respect to the quality of plasmon deconvolution, since it is computed on a very narrow energy-window which is not much perturbed by the plasmon contribution; (iii) to a lesser extent, it also presents advantages with respect to the presence of some residual carbon contamination (see Section 4.2), since the C–K edge from amorphous carbon is dominated by the s* peak, which culminates at 291–292 eV, that is at the end of the energy range used to calculate the carbon contribution in the jump-ratio method (see Fig. 11a)1.

1

Obviously a reference spectrum of amorphous carbon can also be included in the MLSF procedure, but it will certainly appear to be unwise to add another fitting parameter in the present case.

Fig. 12. MLSF analysis of the NbC powder ‘standard’ given in Table 3. The C– K edge reference has been obtained from the Nb6C5 ‘standard’ (after normalization); the deduced [C]/[Nb] ratio is 0.945, very close to 0.951 as measured from the calibration by X-ray diffraction (see Table 3).

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Fig. 13. EELS analysis of the AlOx replica: (a) comparison of carbon-free and carbon-containing areas (respectively top and bottom spectra); (b) quantitative determination of the [Camorphous]/[O] ratio in the case of the worst spectrum in (a).

In the case of NbC3 carbides, the final relative error for the carbon content D3/3 is the sum of three terms: (i) the error due to the X-ray calibration (about 1.7% as can be seen from Table 3); (ii) the error due to the departure from the linear relationship /5/ RCZf([C]/[Nb]): about 1.4% for the worst ‘standard’; (iii) the error introduced by any undesirable amorphous carbon contamination. The last term appears to be the most important: although EELS spectra were recorded at the liquid nitrogen temperature, and a great care was brought to reject any spectrum where a suspicious carbon edge was visually detected during the acquisitions, it appeared that very minor quantities of carbon could be measured from the AlOx replica itself while treating the spectra post-mortem. Fig. 13 illustrates the worst case, where a [Camorphous]/[O] ratio of 2% was measured from the replica. Assuming a 10 nm particle lying over a 10 nm-thick replica yields to a maximum error of about 3% for the carbon content to be attributed to the precipitate. All things considered, it can be claimed that the accuracy for the measurement of the carbon content within the carbo(-nitride) precipitates is about 6%. The accuracy for the nitrogen content is of the same order of magnitude, although significantly better (e.g. about 3%) since the N–K edge is not affected by the presence of spurious Camorphous signal.

4.3. Irradiation effects It is well known that non-stoichiometry in metallic carbides facilitates the displacement of carbon atoms due to knock-on damage under the incident electron beam in TEM studies (Venables and Lye, 1969). Thus disordering of long-range ordered phases can take place (V6C5 (Venables and Lye, 1969), V8C7 (Epicier, 1990)) and carbon loss may also occur (e.g. TiCz1, Das et al.,(1981)). This phenomenon is enhanced in nano-probe analysis with a FEG instrument, typically delivering a 1 nA current: Fig. 14 shows the evolution of the carbon and nitrogen contents of a 15 nm precipitate during EELS analysis with a 1 nm probe and successive acquisitions of 3 s. It is very clear that loss of nitrogen occurs first and is significant after about 30 s; carbon is less sensitive but is ejected out of the particle after about 60 s. The analysed composition, initially measured as Nb1C0.86N0.13 after the first acquisition, becomes Nb1C0.86N0.12 and Nb1C0.73N0.08 after 30 and 60 s respectively. The increase of the carbon content at the end of the sequence can be interpreted as a carbon redeposition as well as the beginning of contamination. According to this observation, the EELS acquisition time was limited to 3–5 s for all nano-analyses. 4.4. Results

1.2

normalized evolution of atomic ratios [C]/[Nb] ( ), [N]/[Nb] ( ) and [C+N]/[Nb] ( )

According to the above methodology, different precipitation states have been analysed for both A and B steels. A further

1

Table 4 Chemical composition of alloys A and B for different heat treatments (atomic fraction ratios)

0.8 0.6

Alloy

Treatment

Number of particles

[C]/[Nb]a

[N]/[Nb]a

[CCN]/ [Nb]a

Steel A

60 min at 800 8C 30 min at 650 8C 126 h at 650 8C

30

0.84

0.15

0.99

30

0.58

0.27

0.85

30

0.52

0.35

0.87

0.4 0.2

Steel B

cumulative irradiation time (s)

0 0

15

30

45

60

75

90

105

120

Fig. 14. EELS analysis of a 15 nm niobium carbo-nitride on a replica during an irradiation experiment.

Steel B a

The mean standard deviation varies between 0.05 and 0.08.

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Fig. 15. Line-scan EELS analysis of a carbo-nitride (replica from alloy A after a treatment of 30 min at 800 8C) with a probe size of 0.7 nm: (a) HRTEM image of the particle (the frame enlargement shows the (220)fcc lattice planes at 0.223 nm). The probe positions are indicated; (b) montage of the 18 EELS spectra. The first spectrum from the replica has been subtracted from all spectra. Note that the oxygen K-edge is almost completely absent within all spectra, and that no iron can be detected; (c) elemental analysis of the line-scan (the last spectrum has been omitted because of its poor signal-to-noise ratio).

limitation in the quantitative analysis arose for the shorter annealing times, due to the size of the extracted particles: although precipitates as small as 3 nm could be observed (see Fig. 5), the poor EELS signal-to-noise ratio of particles smaller than about 6 nm makes it impossible to get reliable results. A similar size of about 7 nm was reported by Wilson and Craven

(Wilson and Craven, 2003) in a recent STEM study of vanadium carbo-nitrides. Some compositional results are summarized in Table 4 and illustrated by Figs. 15 and 16. Fig. 15 shows a line-scan experiment performed with a 0.7 nm probe, and acquisition times of about 4 s. Although

Fig. 16. Co-existence of pure nitrides and carbo-nitrides in alloy B: (a) HAADF image of the replica after a treatment of 30 min at 650 8C: carbo-nitrides (CN) and pure nitrides (labelled N) are as indicated; (b) same as (a) for a treatment of 126 h at 650 8C; (c) EELS spectrum of nitride 1 in (b) (the top spectrum from the Nb1N1 ‘standard’ is shown for comparison); (d) EELS spectrum of carbo-nitride 2 in (b).

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compositional variations reach about 10% (which exceeds the values quoted in Section 4.2, owing to the smaller signal-tonoise ratio due to the smaller probe size), it can be stated that the composition of the carbo-nitrides appears to be homogeneous at a nanometric scale (in particular, no evidence can be obtained for a core-shell type structure in any particle analysed in the line-scan mode). Fig. 16 reveals an important feature in alloy B containing a significant amount of nitrogen (see Table 1): two distinct populations of precipitates exist, respectively pure niobium nitrides and sub-stoichiometric carbo-nitrides Nb(C,N)z0.86. 5. Discussion and conclusions 5.1. Precipitation sequence In this work, it has been shown that a quantitative chemical analysis of nanometric niobium nitrides and carbo-nitrides could be performed with EELS under conditions, which can be summarized as follows: (i) Precipitates as small as 3 nm can be successfully extracted from the ferrite matrix and observed on AlOx replicas. However, reliable EELS data could only be obtained on particles larger than 6 nm typically. (ii) Although the absence of carbon within the supporting film allows the precipitates to be easily analyzed (i.e. without any significant spurious carbon signal, as in the case of a classical carbon replica), it has been necessary to develop a method to quantify precisely the carbon and nitrogen content of the niobium-based particles. A treatment based on the jump-ratio for the C–K and N– K edges has enabled an accuracy D3/3 of less 6% for measuring the Nb(C,N)3 composition of precipitates in the 6–200 nm range. Great care has to be taken in order to avoid a significant loss of the light C and N elements from the precipitates, owing to irradiation effects in the electron microscope.

501

Satisfactory conditions have been established on the basis of dedicated irradiation experiments. Typically, EELS acquisition times less than 5 s are needed with nano-probes of 1– 2 nm, with a typical current of 1 nA, with a JEOL 2010F operating at 200 kV. Within these constraints, the results obtained in this work allow interesting features to be drawn regarding the scenario explaining the precipitation sequence of niobium carbonitrides in ferrite: (i) in alloy A, containing much more carbon than nitrogen ([C]/[N]Z14 as can be seen from Table 1), all precipitates have an homogeneous size and composition, and appear stoichiometric (½CC N
=½Nb
Z 0:99 after 60 min at 800 8C—see Table 4). (ii) in alloy B, containing equal amounts of carbon and nitrogen ([C]/[N]Z1.1 as can be seen from Table 1), two populations of precipitates are observed: pure stoichiometric nitrides which are bigger than sub-stoichiometric carbo-nitrides ([CCN]/[Nb]z0.86 as seen from Table 4). It can thus be stated that the presence of a significant amount of nitrogen in the steels promotes the formation of nitrides in addition to carbo-nitrides. HRTEM work on thin foils of both alloys treated at very short annealing times confirms this difference: Fig. 17 shows observations made on both alloys. It is clear that two types of objects can be identified in alloy B (containing more nitrogen): on the one hand, coherent atomic platelets are observed for the shortest annealing time investigated here (Fig. 17b). The HRTEM contrast resembles that obtained within nitrided iron alloys (Rockerby et al., 1986; Bor et al., 2002). And, on the other hand, already formed fcc precipitates are already present in the shortest annealing time investigated (Fig. 17c). Although EELS analysis could not be performed on such small objects within ferrite thin foils, it has been possible to check with qualitative EDX analysis that the platelets contain much less carbon than the precipitates. These findings are also fully consistent with previous field ion atom probe measurements on the same material (Bemont, 2003). Further experiments are still in progress with this technique in order to quantify more accurately the chemical composition of both types of particles. 5.2. Conclusions

Fig. 17. [100]Fe HRTEM images of alloys A and B showing the first stages of precipitation: (a) alloy A after a treatment of 30 min at 600 8C: evidence for a very thin platelet, the structure of which is compatible with a fcc structure in the Baker–Nutting orientation with respect to the ferrite matrix; (b) alloy B after a treatment of 5 min at 650 8C: evidence of mono-atomic platelets (e.g. arrow) enriched in niobium and nitrogen (as deduced from EDX); (c) same as (b) but evidence for a fcc precipitate in the Baker–Nutting orientation relationship and significantly enriched in carbon compared to the platelet in (b).

TEM and associated techniques have been applied to the study of niobium carbo-nitride precipitation in model steel alloys. The benefit of nano-probe EELS analysis has been demonstrated regarding the quantification of the composition of precipitates down to a size of about 6 nm. The potential application of EFTEM to get a reasonable estimation of the volume fraction of precipitates has also been explored. Further systematic work is required in order to get statistically significant data that can serve as an input to thermodynamic simulations. In this respect, the present results have a great implication for any attempt to model the

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precipitation of niobium carbo-nitrides in steel: it has been shown that a significant nitrogen content promotes the appearance of pure nitrides in addition to expected carbonitrides. Both populations exhibit different sizes and certainly act differently on the refinement of the final grain structure of the steels. Further work is in progress to understand this on a thermodynamic basis. Acknowledgements The authors gratefully acknowledge the CLYME (Consortium Lyonnais de Microscopie Electronique) for the access to the JEOL 2010F and Leo 912 microscopes, and the French Contrat de Programmes de Recherches (CPR) ‘CNRS-ArcelorAlcan/Pechiney-CEA’ for financial support. Thanks are also due to Patrick Barges (IRSID, Arcelor) for the help in the extraction replicas procedure, Philippe Maugis ((IRSID, Arcelor, and CIRIMAT-ENSIACET, F-Toulouse) for fruitful discussions, Be´atrice Vacher (LTDS, ECL, F-Ecully) and Pascale Bayle-Guillemaud (DRFMC, CEA, F-Grenoble) for assistance in the EFTEM experiments. One of the referees is gratefully acknowledged for pertinents comments. References Baker, R.G., Nutting, J., 1959. The tempering of a Cr–Mo–V–W and a Mo–V steel. In: Precipitation Processes in Steels. Iron and Steel Institute special report, London, pp. 1–22. Bemont, E., 2003. Thesis, University of Rouen, France. Beres, M., Weirich, T.E., Hulka, K., Mayer, J., 2004. TEM investigations of Fine Niobium Precipitates in HSLA Steel. Steel Res. Int. 75, 753–758. Billingham, J., Lewis, M.H., 1972. Superlattice with monoclinic symmetry based on the compound V6C5. Philos. Mag. 25-3, 661–671. Bor, T.C., Kempen, A.T.W, Tichelaar, F.D., Mittemeijer, E.J., 2002. Diffraction contrast analysis of misfit strains around inclusions in a matrix: VN particles in a-Fe. Philos. Mag. A82-5, 971–1001. Charleux, M., Poole, W.J., Militzer, M., Deschamps, A., 2001. Precipitation behavior and its effect on strengthening of an HSLA-Nb/Ti steel. Metall. Mater. Trans. A32-7, 1635–1647. Courtois, E., 2005. Thesis, INSA-Lyon, France. Courtois, E., Epicier, T., Scott, C., 2005. Characterisation of NbC and Nb(C,N) evolution within ferrite: contribution of TEM and advanced associated techniques. Mater. Sci. Forum. 500-501, 669–676. Craven, A.J., He, K., Garvie, L.A., Baker, T.N., 2000. Complex heterogeneous precipitation in Ti–Nb microalloyed Al-killed HSLA steels. Acta Mater. 48, 3857–3868 see also pages 3869–3878. Das, G., et al., 1981. Electron irradiation damage in TiC. J. Mater. Sci. 16, 3283–3291. Davis, C.L., Strangwood, M., 2002. Preliminary study of the inhomogeneous precipitate distributions in Nb-microalloyed plate steels. J. Mater. Sci. 37, 1083–1090. DeArdo, A.J., 1998. Microalloyed strip steels for the 21st century. Mater. Sci. Forum 284-286, 15–26. Deschamps, A., Brechet, Y., 1999. Influence of predeformation and ageing of an Al-Zn-Mg alloy: II. Modelling of precipitation kinetics and yield stress. Acta Mater. 47, 293–305.

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