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XRD investigation of ion irradiated Ti3Si0.90Al0.10C2
Nuclear Instruments and Methods in Physics Research B 268 (2010) 506–512
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XRD investigation of ion irradiated Ti3Si0.90Al0.10C2 Xingmin Liu a, Marion Le Flem a,*, Jean-Luc Béchade a, Fabien Onimus a, Théodore Cozzika a, Isabelle Monnet b a b
CEA Saclay, DEN/DMN/SRMA/LA2M, 91191 Gif sur Yvette Cedex, France Centre de recherche sur les Ions, les Matériaux et la Photonique, CEA/IRAMIS/CIMAP, 14076 Caen Cedex 5, France
a r t i c l e
i n f o
Article history: Received 1 October 2009 Received in revised form 17 November 2009 Available online 2 December 2009 Keywords: MAX phases Ion irradiation Nuclear reactors X-ray diffraction Microstructure
a b s t r a c t Ti3SiC2 is one of the most promising materials belonging to Mn+1AXn phases, which exhibit good damage tolerance, thermal stability and mechanical properties. Recently, in the frame of research on future gas cooled fast nuclear reactors, Ti3SiC2 has been considered as an innovative candidate material, which could be incorporated in some core components such as fuel cladding. At the present time, however, very few data are available concerning the behaviour of this material after irradiation. In this work, Ti3Si0.90Al0.10C2 samples were irradiated with high energy Kr and Xe ions and characterized by X-ray diffraction. Patterns were analysed in terms of change in peak intensity, peak position and width. Rietveld refinements were also performed. Increase in micro-strains and lattice parameter with irradiation dose was highlighted. The formation of b-Ti3SiC2, which has never been observed by experimental XRD on non irradiated material, was proposed for the highly irradiated samples. A partial recovery of the microstructure with temperature was found. Ó 2009 Elsevier B.V. All rights reserved.
1. Introduction Refractory materials are needed as structural components for future Generation IV nuclear power reactors [1], operating at much higher temperature compared to pressurized water reactors or previous fast breeders reactors. For example, structural material should be able to reach the ambitious goals of dose (60–90 dpa) and temperature (600–1000 °C up to 1650 °C in accidental scenarios) assigned to the gas fast reactor (GFR) core components. Refractory metallic alloys (W, Ta, Mo) are not compatible in terms of neutronic (large absorption cross section) which negatively impacts the reactor efficiency. Ceramics such as ZrC or SiC combine high melting temperature, good thermal conductivity, and neutronic compatibility with fast neutron spectrum. Therefore they have been identified to be used as core components in the GFRs [2,3]; consequently, in parallel to neutron irradiation experiments, many charged particles irradiations have been performed to understand the microstructure and mechanical properties change with dose and temperature [4–8]. The main drawback of ceramics is their very poor damage tolerance and catastrophic brittle failure resulting in a high failure risk, especially in thermal or mechanical shock. That is why, ceramic
* Corresponding author. Tel.: +33 1 69 08 40 98; fax: +33 1 69 08 71 30. E-mail addresses: [email protected] (X.M. Liu), marion.lefl[email protected] (M. Le Flem), [email protected] (J.-L. Béchade), [email protected] (F. Onimus), [email protected] (T. Cozzika), [email protected] (I. Monnet). 0168-583X/$ - see front matter Ó 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.nimb.2009.11.017
matrix composites have been considered for GFR components, especially for fuel cladding applications. Concurrently to the development of fiber reinforced (SiCf/SiC) composites, layered Mn+1AXn, ternary compounds (where n = 1, 2, or 3, M is an early transition metal, A is an A-group element, and X is either C or N [9]) recently attracted interest because, contrary to ceramics, they exhibit intrinsic damage tolerance properties linked to specific nanolayered structure of atomic planes allowing deformation via delamination and kink bands formation [10,11]. Among them, Ti3SiC2 [12] exhibits good thermal stability (between 1350 °C [13] and 1800 °C [14] depending on purity and atmosphere) that makes it potential candidate for GFR in-core applications as element of fuel cladding. The research upon Mn+1AXn phases aimed at reactor application is very limited. There is one report in the literature [15] dealing with corrosion behaviour of Ti2AlC and Ti3SiC2 in circulating molten lead at 650 °C and 800 °C for the potential application as cladding or structural materials in a lead-cooled fast reactor. There is almost no research on irradiation effects for Mn+1AXn phases in the literature; very recently, surface modification induced by charged particles irradiations was highlighted (change in roughness) [16]. The questionable point about Mn+1AXn phases is the following: how the specific layered microstructure, and consequently the related outstanding mechanical properties, is affected by irradiation? In general, materials become more brittle after irradiation depending on the irradiation temperature (amorphisation can also be possible at low temperature). Then, there is
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a risk that Mn+1AXn compounds lose their damage tolerance ability and turn to brittle monolithic ceramic; data concerning behaviour under irradiation must be gathered. That is why CEA launched in 2006 both experimental neutron irradiations of Ti3(Si, Al)C2 in French reactors and irradiation programs with charged particles [16,17]. In this study, Ti3Si1xAlxC2 (x usually equal to 0.10), a solid solution of Ti3SiC2 which was optimized material from phase purity and oxidation resistance point of view [18,19], was irradiated by 74 MeV Kr and 92 MeV Xe ions at room temperature, 300 °C and 500 °C. X-ray diffraction patterns were obtained and refined by Rietveld method [20]. Irradiation induced micro-structural changes were analyzed based on experiment and refined results. 2. Experimental procedures 2.1. Material manufacture The ceramic material was fabricated by Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China. It dealed with Ti3SiC2 with 10% Al substituted to Si, which allows to get rid of TiC secondary phase usually found in the conventional preparing methods, and then, to enhance oxidation resistance [19]. Bulk Ti3Si0.90Al0.10C2 ceramics were prepared by in situ hot pressing solid–liquid reaction of elemental powders. Ti, Si, Al, and graphite elemental powders (stoichiometry proportions) were mixed in a polyurethane mill for 20 h. The powder mixtures were then cold pressed in a graphite mould coated with a boron nitride layer on the inner surface. Solid–liquid synthesis reaction and simultaneous densification were performed in a furnace using graphite as the heating element in a flowing argon atmosphere. The compacted mixture was hot pressed at about 1520 °C during 1 h, applying a uniaxial pressure of 30 MPa. Detailed descriptions of the synthesis of Ti3Si1xAlxC2 ceramics can be found elsewhere [19]. 2.2. Irradiation of materials The samples were 10 5 4 mm3 bars machined from the same bulk Ti3Si0.90Al0.10C2 material. The 10 5 mm2 surface of the bars was mechanically grinded and polished down to
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0.25 lm (diamond spray). The polished samples were degreased and washed with ethanol and acetone. Irradiation of Ti3Si0.90Al0.10C2 ceramic was performed under high vacuum (106 Torr) in IRRSUD line of GANIL (National Large Accelerator of Heavy Ions, Caen, France). The ions were 74 MeV +20 Kr86 and 92 MeV +23Xe129. Penetration depth of the two ions was estimated to be about 7–8 lm based on calculation by SRIM software (location of the implantation peak), the six first microns being affected by electronic interactions only as shown in Fig. 1. The fluences ranged from 1012 to 2 1015 cm2 (irradiated sample will be referenced as sample 2E15 for concise and this notation applied to other samples). Irradiation experiments were performed at room temperature, 300 °C and 500 °C. When performing room temperature irradiation, flux was kept lower than 4.5 109 cm2 s1 to avoid temperature increase. For high fluences (1 1014 cm2 and higher) needing a higher flux, a sample stage with cooling water was used. In all cases, samples were directly glued on the stage by conductive carbon double faces and irradiation surface was perpendicular to ion beam. For high temperature irradiation, a dedicated stage was used. Samples were fixed on the stage by a piece of stainless steel. The heating resistance and thermocouple were placed below the copper plate to make sure a rapid heating rate and accurate measurement of sample temperature. Cooling air was purged through a stainless steel tube welded on the back of copper stage after irradiation to ensure a rapid decrease of samples temperature. XPS analyses showed that only 10–100 nm of the surface were affected by oxygen pollution. 2.3. XRD characterization and Rietveld refinement X-ray diffraction data were obtained using a Bruker D8 Advance diffractometer with Cu ka radiation in the 2h range between 29° and 63°, which covered most of intense diffraction peaks of Ti3SiC2. Both divergence slit and receiving slit were set at 0.2 mm. The step interval was 0.02° (2h) and counting time was 40 s for each steps to get strong signal from the rather small surface analyzed (surface area is 10 5 mm2). The instrument operated at 40 kV and 40 mA and the height difference between irradiated surface and reference surface of sample holder was kept less than 50 lm. The pattern of standard LaB6 powder was collected at the same conditions and was used to generate instrumental
Fig. 1. Damage profile induced by 92 MeV Xe in Ti3SiC2 for dose of 1 1015 cm2 (SRIM calculation).
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resolution function (IRF, via the Caglioti polynomial) to account for the experimental broadening. Rietveld refinements were performed using FullProf program [21]. Pseudo-Voigt (pV) function was selected to refine peak profile. The displacement of the sample (height out of the focussing circle) and a constant background (only one refined parameter) were refined. Details of refinement will be given in Section 3.2.1. It must be noticed that the calculated X-ray investigation depth for Ti3SiC2 is deeper than the stopping range of Xe ions. Contribution from non irradiated bulk was taken into account in the refinements.
3. Results 3.1. XRD patterns 3.1.1. Room temperature irradiation Fig. 2 gives XRD diffraction patterns of the virgin sample and samples irradiated with Xe ions; it should be pointed out that peak intensities of sample 5E14, 1E15 and 2E15 have been multiplied by 3 to make lecture easier. For sample irradiated at the lowest dose, sample 1E13, the change in peak positions and peak broadening could not be detected compared to the virgin sample. For 5E13 sample and higher fluence levels, a continuous drop of peak intensity and increase of peak width could be found, and for 2E15 sample, very broad peaks appeared, which could mean that the degree of crystallinity decreases. A systematic shift of (0 0 8) peak (and also of the weak (0 0 6) and (1 0 9) peaks) to low angle could also be noticed. The shifting increases with irradiation dose and at last are the same for 5E14 and 1E15 samples showing the saturation of the phenomenon. The shift of (0 0 8) related peaks, and to a less extend, shift of (1 0 9) peaks, is consistent with previous observations via grazing XRD [16] (no obvious peak shift could be observed for the (1 1 0) diffraction). However, no amorphisation is obvious. This shift can be linked to continuous expansion of the unit cell along c axis. The lattice parameters after irradiation will be given in the refinement part. Another important characteristic of these patterns is the increase of relative intensity of (1 0 2) and (1 0 3) to (1 0 1) peaks for samples irradiated up to 5 1014 ionscm2 and above (the
same conclusion was drawn with considering integral peak areas ratios obtained by TOPAS software). Especially for 1E15 sample, intensity of (1 0 2) diffraction is even larger than that of (1 0 1). This pattern feature is consistent with simulated b-Ti3SiC2 diffraction patterns by Farber et al. [22], suggesting presence of this polymorph after irradiation. Besides a strong broadening of peaks centred around (1 0 3), (1 0 5) and (1 1 0) diffraction positions of virgin sample could also be noticed in the 1E15 sample. These broad peaks could be the sum of diffraction response from different phases. These results suggests strong modifications in the Ti3Si0.90 Al0.10C2 microstructure after Xe irradiation, dealing with: (i) change in degree of crystallinity (pattern intensity), (ii) change in lattice (peak broadening and shifting) and (iii) phase transformation (peak relative intensity and unexpected important broadening). In addition, some narrow non-shifted peaks could be found for 5E14, 1E15 and 2E15 samples (located at the same angles as for virgin sample). For 2E15 sample, these peaks are more intense than the two main broad peaks. This means there is a contribution from small amount of material that has not been affected by irradiation. It should not come from surface of irradiated sample since other grazing XRD analysis did not show any non-shifted peaks [16]. These peaks should rather be generated from bulk material underneath the ion implantation zone. This contribution from bulk is effective for all the samples but detected for highly irradiated samples only because of large shift (no overlap) and low intensity of peaks from irradiated layer. For samples irradiated with Kr ions (lower energy), a similar evolution of diffraction peaks was observed. The peak intensity drops and (0 0 8) peaks shift started from 1E14 sample. The increase of peak breadth was not as large as that of samples irradiated at the same dose with Xe ions. This observation confirms that lower energy Kr ions induced less damage compared to high energy Xe ions. 3.1.2. High temperature irradiation A recovery of peak intensity and peak position were found for high temperature irradiated samples for both ions. An example is given in Fig. 3 for 1E15 sample irradiated with Xe ions at 20 °C, 300 °C and 500 °C. Thermal instability of irradiation produced defects should be the reason of this recovery (recombination of irradiation defects).
Fig. 2. XRD patterns of samples irradiated at various doses with Xe ions.
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Fig. 3. XRD patterns of 1E15 samples irradiated at various temperatures with Xe ions.
3.2. Rietveld refinement and micro-structural analysis
e¼
3.2.1. Principles Pseudo-Voigt (pV(x)) function was selected to refine peak profile. Pseudo-Voigt function is an approximation of the Voigt function defined as the linear combination of a Lorentzian and a Gaussian functions (L(x) and G(x)). It has the following expression:
pVðxÞ ¼ gLðxÞ þ ð1 gÞGðxÞ
ð1Þ
with g ¼ g0 þ X2h
ð2Þ
where g0 and X are refined parameters to get Lorentzian mixing index g. The X coefficient is only refined for the LaB6 in order to measure accurately the instrumental broadening. For the refinement of the Ti3SiC2, the X coefficient was set to zero. For Pseudo-Voigt function, the full width at half maximum (FWHM) (H) has the following relationship with strain and size parameters:
H2 ¼ ðU þ D2ST Þ tan2 ðhÞ þ V tanðhÞ þ W þ
IG cos2 ðhÞ
ð3Þ
where U, V, W and DST are micro-strains effects parameters and IG is a domain size effects parameter. It is considered that V and W in this function should be zero when instrumental resolution function (IRF) was used and DST was also set to zero. So the evolution of the FWHM (with H in degrees) can be simply written as
H2 ¼ U tan2 ðhÞ þ
IG cos2 ðhÞ
ð4Þ
Basically, U and IG can be related to micro-strains and domain size through integral breadth (b in radians). From Stokes and Wilson [23] and Scherrer [24] function we know that
bstrain ¼ 4e tanðhÞ k bsize ¼ D cosðhÞ
ð5Þ ð6Þ
For Pseudo-Voigt function, integral breadth (b in radians) can be related to the FWHM (H in degrees) by
b¼
p p H pffiffiffiffiffiffiffiffiffiffiffiffiffi 180 2 g þ ð1 gÞ p ln 2
ð7Þ
So micro-strain e can be obtained by the refined parameter U (in degrees2) through the following function:
p p
pffiffiffiffi U
180 8 g þ ð1 gÞ
pffiffiffiffiffiffiffiffiffiffiffiffiffi p ln 2
ð8Þ
And coherently the domain size D could be written as a function of IG (in degrees2) through
D¼
180 2k g þ ð1 gÞ pffiffiffiffi p p IG
pffiffiffiffiffiffiffiffiffiffiffiffiffi p ln 2
ð9Þ
Based on a Hall and Williamson analysis not described here it has been chosen to refine only the micro-strains (U parameter) for the a-phase and only the domain size (IG parameter) for the b-phase. The parameter g0 was refined and in some case for high irradiation dose, it is fixed to 1 because the peaks are broad (Lorentzian profile). It must be noticed that the microstructure features should follow the damage profile caused by the passage of ion in the specimen. However, specific knowledge of microstructure profile requires extended TEM studies that are not available at the present time. Then, it sounds premature to take this into account in the refinement, i.e. increase the number of refinement parameters in Rietveld analysis. 3.2.2. Refinement results We found the agreement factors Rp and Rwp (reliability factor and weighted reliability factor, respectively) were too high when performing Rietveld refinement for highly irradiated samples if common experimental a-Ti3SiC2 phase only was considered. That is why the hypothesis was made to consider the high temperature b-Ti3SiC2 phase; the resulting refinements were considerably improved as shown in the following. b-Ti3SiC2 was first proposed by Farber et al. [22] through TEM observation: they suggested it was transformed from a-Ti3SiC2 during ion milling processing. But Yu [25] argued that low energy ion milling could not induce this phase transformation and what Farber et al. observed was actually still a-Ti3SiC2. Later, Wang and Zhou [26] made detailed first-principles total energy calculations and pointed out that a reversible polymorphic phase transition would occur when shear strain energy was large enough to overcome an energy barrier. Actually, the two polymorphs show the same space group symmetry, P63/mmc; the only difference is the Si occupancy in the (0 0 0 1) plane. In a-Ti3SiC2 the Si atoms occupy 2b Wyckoff position with fractional coordinates (0, 0, 1/4), while in b-Ti3SiC2, Si
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occupy 2d Wyckoff position with fractional coordinates (2/3, 1/ 3, 1/4). The crystal structures of the polymorphs are illustrated in Fig. 4. In the two sketches, Ti2 refers to atom bounded to Si atom. Ti2 and C atomic positions can be adjusted without breaking the symmetry restrictions. In this work, for a-Ti3Si0.90Al0.10C2 phase, the two positions were taken from literature [27]. For b-Ti3Si0.90Al0.10C2, since no experimental data could be found elsewhere, they were refined using XRD data for 5E14 sample, for which bphase formation was previously suggested in Section 3.1 (see Table 1). The values were then used for 1E15 and 2E15 samples. An example of quality improvement of the refinement by introducing b-phase is given in Fig. 5. Some of the parameters deduced by Rietveld refinement from several highly irradiated materials are listed in Table 2. The volume fraction of b-phase (which starts to form for 5E14 sample) increased with dose for Xe irradiation and, for 2E15 sample, almost whole sample was transformed into b-phase (a-phase is only about 1%). For Kr irradiation, the b-phase appeared for dose higher than 1 1015 ionscm2. Sample irradiated at 300 °C with Xe showed only 30% of b-phase, and only a-phase could be detected for both Xe and Kr irradiation at 500 °C. An increase in a-phase lattice parameter was highlighted. The increase of lattice parameter a (less than 0.2%) with increasing irradiation dose does not seem significant compared to that of c parameter. The change in c parameter with dose is given in Fig. 6. The increase in c parameter is a bit higher than previously reported by [16]: grazing XRD on Xe irradiated Ti3SiC2 (1 1015 ionscm2) lead to an increase of 1.07% linked to effect of electronic interactions only. An increase of 1.3% was found in this work via basic XRD. Actually, the microstructure features should change with the damage profile (Fig. 1), i.e. increase in c
Table 1 Structural parameters of a-Ti3Si0.90Al0.10C2 and b-Ti3SiC2. Formula
a-Ti3Si0.90Al0.10C2 [27]
b-Ti3SiC2
Space group Lattice parameters (nm)
P63/mmc (194) a = 3.0690 c = 17.6785
P63/mmc (194) a = 3.0344 c = 18.0178 [26]
(0, 0, 0) (2/3, 1/3, 0.1329) (0, 0, 0.25) (1/3, 2/3, 0.0724)
(0, 0, 0) (2/3, 1/3, 0.134) [28] (1/3, 2/3, 0.25) (1/3, 2/3, 0.071) [26]
Atomic positions Ti1 Ti2 Si C
Fig. 5. Rietveld refinement of XRD patterns of Xe irradiated 1E15 sample: (a) considering a-Ti3SiC2 only and (b) considering both a-and b-Ti3SiC2. In both cases, contribution from underneath non irradiated bulk material (a-Ti3SiC2) was introduced to refine the non-shifted peaks.
Fig. 4. Crystal structures of a- and b-Ti3SiC2 polymorphs. Species of atoms and corresponding Wyckoff positions are illustrated.
parameter with depth. Nevertheless, at this stage of the investigation, this average parameter is useful in the scope of comparative study of irradiation impact and it did not sound reasonable to increase the number of refinement parameters in Rietveld analysis. Furthermore, it seemed that no additional broadening of the (0 0 8), (0 0 6), (1 0 9) peak occurred suggesting that the profile in c parameter should have a limited impact on the XRD diagram. It seems that c parameter reaches a saturation value for 5E14 sample (the value dropped a little when irradiated to 1 1015 cm2). When irradiated at high temperature (500 °C), the
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X.M. Liu et al. / Nuclear Instruments and Methods in Physics Research B 268 (2010) 506–512 Table 2 Rietveld refinement results of Ti3Si0.90Al0.10C2 samples irradiated at various fluence levels. Ion
T (°C)
Dose (ions/ cm2)
Phase
a (Å)
c (Å)
U (degrees2)
g0
Kr Kr Kr Kr
RT RT RT RT
0 1E12 1E13 1E14 1E15
a a a a a
0.60 0.76 0.78 0.70 0.64
500 RT RT RT RT
1E15 1E13 5E13 1E14 5E14
0.0472 0.0167 0.0620 0.1101 0.5120
Xe
RT
1E15
17.6787 17.6791 17.6901 17.7394 17.9728 19.2389 17.7452 17.6917 17.7418 17.7836 17.9195 19.3658 17.9097 19.4151 19.6002 17.9253 19.6087 17.8504
0.0039 0.0054 0.0158 0.0821 0.3462
Kr Xe Xe Xe Xe
3.0678 3.0679 3.0698 3.0715 3.0689 3.0399 3.0695 3.0691 3.0717 3.0720 3.0691 3.0178 3.0669 3.0282 3.0340 3.0718 3.0232 3.0698
0.74 0.62 0.65 0.62 0.79 1 1 1 1 0.51
b
a a a a a b
Xe Xe
RT 300
2E15 1E15
a b b
a b
Xe
500
1E15
a
IG (degrees2)
1.62
1.8660
0.4697
2.56 3.42 3.66 1.01
0.1179
0.74
Fraction of phase (%)
Micro-strains in aphase (103)
100 100 100 100 68.9 ± 1.5 28.2 ± 1.3 100 100 100 100 60.5 ± 1.2 35.0 ± 1.2 30.4 ± 0.9 65.6 ± 1.6 95.7 ± 0.8 71.5 ± 1.6 28.5 ± 1.4 100
0.358 0.454 0.786 1.780 3.449
Domain size in bphase (nm)
Rp
Rwp
16.2 15.9 21.3 22.4 23.8
19.1 18.7 25.0 24.8 26.9
23.6 18.3 22.0 21.9 21.3
26.4 20.4 24.1 24.2 25.6
18.7
22.8
21.9 22.3
25.6 26.5
21.3
25.1
4.40 1.323 0.749 1.464 1.923 4.445 3.81 9.360 3.03 3.08 3.810
5.62 2.090
Note: Values with mean these data were fixed or that the patterns could not be refined.
Fig. 6. Relative increase in lattice parameter c of a-Ti3Si0.90Al0.10C2 for samples irradiated with Xe and Kr ions at various temperatures (error bars are smaller than plot size).
Fig. 7. Micro-strains in a-Ti3Si0.90Al0.10C2 for samples irradiated with Kr and Xe ions at various temperatures.
4. Discussion and conclusions increase of c parameter is much smaller than that of room temperature irradiated sample, especially for Kr irradiated sample: this is consistent with microstructure recovery due to defect annealing. Same trend was observed for the change in unit cell volume. The broadening of a-Ti3Si0.90Al0.10C2 peaks was explained in terms of micro-strains contribution (suggested by a Hall and Williamson analysis not detailed here). Micro-strain values obtained at various doses and temperatures are presented in Fig. 7. Whereas no significant change is observed up to 1 1013 cm2, microstrains increased for higher dose for both Kr and Xe irradiated samples. It should be noticed for sample irradiated to 1 1015 cm2 at room temperature that the micro-strains of Xe irradiated sample are about three times of those irradiated with Kr. For samples irradiated at high temperature, a decrease in micro-strains is clearly seen for both ions: the higher the temperature, the larger the drop. A more pronounced drop could be expected for a higher irradiation temperature until complete recovery of the microstructure. It should be reminded that for b-phase, only domain size effect was considered as origin of peak broadening and IG in peak shape function was refined but U was set to zero.
Ti3Si0.90Al0.10C2 ceramics were irradiated at various doses (up to 2 1015 ionscm2) and various temperatures (20 °C, 300 °C and 500 °C) with high energy Kr and Xe ions. XRD patterns showed peak intensity decrease, peak width increase shift in peak position with increasing irradiation dose for both ions: this means progressive change in the microstructure, which is consistent with increase in irradiation damage. Rietveld refinement suggested that irradiation induced the following: Formation of b-Ti3Si0.90Al0.10 nano-domains above 1 1015 cm2 for Kr and 5 1014 cm2 for Xe ions irradiation. Large increase of lattice parameter c of a-Ti3Si0.90Al0.10. Appearance of micro-strains in a-Ti3Si0.90Al0.10. Defect annealing during high (300 °C, 500 °C) temperature irradiation resulting in moderate modification of microstructure. Such a recovery was suggested elsewhere by nanoindentation analysis [17]. Phase transitions induced by ion irradiation have already been observed in metal, alloys [29,30], oxide compounds [31,32]. But
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in other materials, such as in spinel compounds MgAl2O4 [33] and ZnAl2O4 [34], under low energy or swift heavy ions irradiation, a simple exchange of cations onto the spinel sites took place without any change of the space group. The present work suggests that ion irradiation of a-Ti3Si0.90Al0.10C2 results in the formation of b-Ti3Si0.90Al0.10C2 which has the same space group. The only difference between the two phases are the atom positions of Si. Calculation showed that, in Ti3SiC2 compounds, the bonding energy of Si with adjunction of Ti is quite low compared with Ti–C bond [26]. So under high energy irradiation Si atoms should be more easily displaced from their original position to form interstitial atoms. Since 2d sites are the largest holes in a-phase [34], Si atoms could hop from 2b sites to 2d sites to generate Frenkel defects (vacancies at 2b and interstitials at 2d). The number of Frenkel pairs should remain small when irradiation dose is low and may be randomly distributed in the matrix as separate defects or small defect clusters. With increasing irradiation dose, the number of Frenkel pairs increases leading to microdistorsions in the cell. At a threshold irradiation dose, enough interstitial atoms could line up between two layers of Ti and form small domains of b-phase: the adjacent Ti and C atoms would relax and found new positions (no strains), but should not break restrict of symmetry. The amount of b-phase would increase with dose due to the increasing number of interstitial atoms. When irradiations were performed at high temperature (300 °C and 500 °C) and high dose, the interstitial Si atoms should be easier to hop back to low energy 2b original site. So fewer defects would be formed in this case resulting in a lower amount of b-phase, which is consistent with the present results. Further investigations are in progress to directly highlight the presence of b-phase in the irradiated samples. Among them, XPS analysis should be able to detect Si–C bonds related to b-phase, as suggested by Yu et al. [28]. Raman spectrum and HRTEM could also give evidences of the presence of this phase and related works are in progress. According to lattice changes (parameter, strains), the recovery of the microstructure is well advanced after irradiation at 500 °C. It is reasonable to suppose that further recovery occurs at higher temperature. Then, in the operating temperature range of GFR (600–1000 °C), irradiation should have less influence on the properties related to the layered structure, e.g. embrittlement and damage tolerance. That means Ti3SiC2 ceramics could be a promising high temperature element of GFR core components.
Acknowledgements Authors are very grateful to Professor Yanchun Zhou from Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China, for having provided the samples. He is also thanked for collaborative work and fruitful discussions.
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Mater. 373 (1–3) (2008) 424–428. [16] J.C. Nappé, P. Grosseau, F. Audubert, B. Guilhot, M. Beauvy, M. Benabdesselam, I. Monnet, Damages induced by heavy ions in titanium silicon carbide: effects of nuclear and electronic interactions at room temperature, J. Nucl. Mater. 385 (2) (2009) 304–307. [17] X.M. Liu, M. Le Flem, J.L. Béchade, Y.C. Zhou, T. Cozzika, S. Doriot, P. Forget, L. Gosmain, I. Monnet, TEM observations and nanoindentation measurements of Ti3SiC2 irradiated by charged particles, in: Proceedings of the 2nd International Congress on Ceramics, Verona, Italy, June 29–July 4, 2008. [18] Y.C. Zhou, H.B. Zhang, M.Y. Liu, J.Y. Wang, Y.W. Bao, Preparation of TiC free Ti3SiC2 with improved oxidation resistance by substitution of Si with Al, Mater. Res. Innovat. 8 (2) (2004) 97–102. [19] H.B. Zhang, Y.C. Zhou, Y.W. Bao, M.S. Li, Mechanism for the enhanced oxidation resistance of Ti3SiC2 by forming a Ti3Si0.9Al0.1C2 solid solution, Acta Mater. 52 (12) (2004) 3631–3637. [20] H.M. Rietveld, A profile refinement method for nuclear and magnetic structures, J. Appl. Crystallogr. 2 (1969) 65–71. [21] J. Rodríguez-Carvajal, FullProf: a program for Rietveld refinement and pattern matching analysis, in: Abstracts of the Satellite Meeting on Powder Diffraction of the XVth Congress of the International Union of Crystallography, IUCr, Toulouse, 1990, p. 127. [22] L. Farber, I. Levin, M.W. Barsoum, T. El-Raghy, T. Tzenov, High-resolution transmission electron microscopy of some Tin+1AXn compounds (n = 1, 2; A = Al or Si; X = C or N), J. Appl. Phys. 86 (5) (1999) 2540–2543. [23] A.R. Stokes, A.J.C. Wilson, The diffraction of X rays by distorted crystal aggregates-I, Proc. Phys. Soc. 56 (3) (1944) 174–181. [24] P. Scherrer, Bestimmung der grosse und der inneren Struktur Von Kolloidteilchen mittels Röntgenstrahlen, Nachr. Ges. Wiss. Göttingen 2 (1918) 98–100. [25] R. Yu, Ph.D. Thesis. Inst. Met. Res. CAS, 2002. [26] J.Y. Wang, Y.C. Zhou, Polymorphism of Ti3SiC2 ceramic: first-principles investigations, Phys. Rev. B 69 (14) (2004) 4108–4121. [27] H.B. Zhang, Ph.D. Thesis. Inst. Met. Res. CAS, 2006. [28] R. Yu, X.F. Zhang, L.L. He, H.Q. Ye, Topology of Charge Density and Elastic Properties of Ti3SiC2 Polymorphs, Paper LBNL-55853, 2004.. [29] G. Martin, P. Bellon, Driven alloys, Solid State Phys. 50 (1996) 189–331. [30] P. Bellon, G. Martin, Cascade effects in a nonequilibrium phase transition with metallurgical relevance, Phys. Rev. B 39 (4) (1989) 2403–2410. [31] D. Simeone, G. Baldinozzi, D. Gosset, S. LeGaër, L. Mazerolles, Impact of radiation defects on the structural stability of pure zirconia, Phys. Rev. B 70 (13) (2004) 134116. [32] A. Chartier, C. Meis, J.P. Crocombette, W.J. Weber, L.R. Corrales, Molecular dynamic simulation of disorder induced amorphization in pyrochlore, Phys. Rev. Lett. 94 (2) (2005) 025505. [33] D. Simeone, C. Dodane-Thiriet, D. Gosset, P. Daniel, M. 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Nuclear Instruments and Methods in Physics Research B journal homepage: www.elsevier.com/locate/nimb
XRD investigation of ion irradiated Ti3Si0.90Al0.10C2 Xingmin Liu a, Marion Le Flem a,*, Jean-Luc Béchade a, Fabien Onimus a, Théodore Cozzika a, Isabelle Monnet b a b
CEA Saclay, DEN/DMN/SRMA/LA2M, 91191 Gif sur Yvette Cedex, France Centre de recherche sur les Ions, les Matériaux et la Photonique, CEA/IRAMIS/CIMAP, 14076 Caen Cedex 5, France
a r t i c l e
i n f o
Article history: Received 1 October 2009 Received in revised form 17 November 2009 Available online 2 December 2009 Keywords: MAX phases Ion irradiation Nuclear reactors X-ray diffraction Microstructure
a b s t r a c t Ti3SiC2 is one of the most promising materials belonging to Mn+1AXn phases, which exhibit good damage tolerance, thermal stability and mechanical properties. Recently, in the frame of research on future gas cooled fast nuclear reactors, Ti3SiC2 has been considered as an innovative candidate material, which could be incorporated in some core components such as fuel cladding. At the present time, however, very few data are available concerning the behaviour of this material after irradiation. In this work, Ti3Si0.90Al0.10C2 samples were irradiated with high energy Kr and Xe ions and characterized by X-ray diffraction. Patterns were analysed in terms of change in peak intensity, peak position and width. Rietveld refinements were also performed. Increase in micro-strains and lattice parameter with irradiation dose was highlighted. The formation of b-Ti3SiC2, which has never been observed by experimental XRD on non irradiated material, was proposed for the highly irradiated samples. A partial recovery of the microstructure with temperature was found. Ó 2009 Elsevier B.V. All rights reserved.
1. Introduction Refractory materials are needed as structural components for future Generation IV nuclear power reactors [1], operating at much higher temperature compared to pressurized water reactors or previous fast breeders reactors. For example, structural material should be able to reach the ambitious goals of dose (60–90 dpa) and temperature (600–1000 °C up to 1650 °C in accidental scenarios) assigned to the gas fast reactor (GFR) core components. Refractory metallic alloys (W, Ta, Mo) are not compatible in terms of neutronic (large absorption cross section) which negatively impacts the reactor efficiency. Ceramics such as ZrC or SiC combine high melting temperature, good thermal conductivity, and neutronic compatibility with fast neutron spectrum. Therefore they have been identified to be used as core components in the GFRs [2,3]; consequently, in parallel to neutron irradiation experiments, many charged particles irradiations have been performed to understand the microstructure and mechanical properties change with dose and temperature [4–8]. The main drawback of ceramics is their very poor damage tolerance and catastrophic brittle failure resulting in a high failure risk, especially in thermal or mechanical shock. That is why, ceramic
* Corresponding author. Tel.: +33 1 69 08 40 98; fax: +33 1 69 08 71 30. E-mail addresses: [email protected] (X.M. Liu), marion.lefl[email protected] (M. Le Flem), [email protected] (J.-L. Béchade), [email protected] (F. Onimus), [email protected] (T. Cozzika), [email protected] (I. Monnet). 0168-583X/$ - see front matter Ó 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.nimb.2009.11.017
matrix composites have been considered for GFR components, especially for fuel cladding applications. Concurrently to the development of fiber reinforced (SiCf/SiC) composites, layered Mn+1AXn, ternary compounds (where n = 1, 2, or 3, M is an early transition metal, A is an A-group element, and X is either C or N [9]) recently attracted interest because, contrary to ceramics, they exhibit intrinsic damage tolerance properties linked to specific nanolayered structure of atomic planes allowing deformation via delamination and kink bands formation [10,11]. Among them, Ti3SiC2 [12] exhibits good thermal stability (between 1350 °C [13] and 1800 °C [14] depending on purity and atmosphere) that makes it potential candidate for GFR in-core applications as element of fuel cladding. The research upon Mn+1AXn phases aimed at reactor application is very limited. There is one report in the literature [15] dealing with corrosion behaviour of Ti2AlC and Ti3SiC2 in circulating molten lead at 650 °C and 800 °C for the potential application as cladding or structural materials in a lead-cooled fast reactor. There is almost no research on irradiation effects for Mn+1AXn phases in the literature; very recently, surface modification induced by charged particles irradiations was highlighted (change in roughness) [16]. The questionable point about Mn+1AXn phases is the following: how the specific layered microstructure, and consequently the related outstanding mechanical properties, is affected by irradiation? In general, materials become more brittle after irradiation depending on the irradiation temperature (amorphisation can also be possible at low temperature). Then, there is
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a risk that Mn+1AXn compounds lose their damage tolerance ability and turn to brittle monolithic ceramic; data concerning behaviour under irradiation must be gathered. That is why CEA launched in 2006 both experimental neutron irradiations of Ti3(Si, Al)C2 in French reactors and irradiation programs with charged particles [16,17]. In this study, Ti3Si1xAlxC2 (x usually equal to 0.10), a solid solution of Ti3SiC2 which was optimized material from phase purity and oxidation resistance point of view [18,19], was irradiated by 74 MeV Kr and 92 MeV Xe ions at room temperature, 300 °C and 500 °C. X-ray diffraction patterns were obtained and refined by Rietveld method [20]. Irradiation induced micro-structural changes were analyzed based on experiment and refined results. 2. Experimental procedures 2.1. Material manufacture The ceramic material was fabricated by Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China. It dealed with Ti3SiC2 with 10% Al substituted to Si, which allows to get rid of TiC secondary phase usually found in the conventional preparing methods, and then, to enhance oxidation resistance [19]. Bulk Ti3Si0.90Al0.10C2 ceramics were prepared by in situ hot pressing solid–liquid reaction of elemental powders. Ti, Si, Al, and graphite elemental powders (stoichiometry proportions) were mixed in a polyurethane mill for 20 h. The powder mixtures were then cold pressed in a graphite mould coated with a boron nitride layer on the inner surface. Solid–liquid synthesis reaction and simultaneous densification were performed in a furnace using graphite as the heating element in a flowing argon atmosphere. The compacted mixture was hot pressed at about 1520 °C during 1 h, applying a uniaxial pressure of 30 MPa. Detailed descriptions of the synthesis of Ti3Si1xAlxC2 ceramics can be found elsewhere [19]. 2.2. Irradiation of materials The samples were 10 5 4 mm3 bars machined from the same bulk Ti3Si0.90Al0.10C2 material. The 10 5 mm2 surface of the bars was mechanically grinded and polished down to
507
0.25 lm (diamond spray). The polished samples were degreased and washed with ethanol and acetone. Irradiation of Ti3Si0.90Al0.10C2 ceramic was performed under high vacuum (106 Torr) in IRRSUD line of GANIL (National Large Accelerator of Heavy Ions, Caen, France). The ions were 74 MeV +20 Kr86 and 92 MeV +23Xe129. Penetration depth of the two ions was estimated to be about 7–8 lm based on calculation by SRIM software (location of the implantation peak), the six first microns being affected by electronic interactions only as shown in Fig. 1. The fluences ranged from 1012 to 2 1015 cm2 (irradiated sample will be referenced as sample 2E15 for concise and this notation applied to other samples). Irradiation experiments were performed at room temperature, 300 °C and 500 °C. When performing room temperature irradiation, flux was kept lower than 4.5 109 cm2 s1 to avoid temperature increase. For high fluences (1 1014 cm2 and higher) needing a higher flux, a sample stage with cooling water was used. In all cases, samples were directly glued on the stage by conductive carbon double faces and irradiation surface was perpendicular to ion beam. For high temperature irradiation, a dedicated stage was used. Samples were fixed on the stage by a piece of stainless steel. The heating resistance and thermocouple were placed below the copper plate to make sure a rapid heating rate and accurate measurement of sample temperature. Cooling air was purged through a stainless steel tube welded on the back of copper stage after irradiation to ensure a rapid decrease of samples temperature. XPS analyses showed that only 10–100 nm of the surface were affected by oxygen pollution. 2.3. XRD characterization and Rietveld refinement X-ray diffraction data were obtained using a Bruker D8 Advance diffractometer with Cu ka radiation in the 2h range between 29° and 63°, which covered most of intense diffraction peaks of Ti3SiC2. Both divergence slit and receiving slit were set at 0.2 mm. The step interval was 0.02° (2h) and counting time was 40 s for each steps to get strong signal from the rather small surface analyzed (surface area is 10 5 mm2). The instrument operated at 40 kV and 40 mA and the height difference between irradiated surface and reference surface of sample holder was kept less than 50 lm. The pattern of standard LaB6 powder was collected at the same conditions and was used to generate instrumental
Fig. 1. Damage profile induced by 92 MeV Xe in Ti3SiC2 for dose of 1 1015 cm2 (SRIM calculation).
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resolution function (IRF, via the Caglioti polynomial) to account for the experimental broadening. Rietveld refinements were performed using FullProf program [21]. Pseudo-Voigt (pV) function was selected to refine peak profile. The displacement of the sample (height out of the focussing circle) and a constant background (only one refined parameter) were refined. Details of refinement will be given in Section 3.2.1. It must be noticed that the calculated X-ray investigation depth for Ti3SiC2 is deeper than the stopping range of Xe ions. Contribution from non irradiated bulk was taken into account in the refinements.
3. Results 3.1. XRD patterns 3.1.1. Room temperature irradiation Fig. 2 gives XRD diffraction patterns of the virgin sample and samples irradiated with Xe ions; it should be pointed out that peak intensities of sample 5E14, 1E15 and 2E15 have been multiplied by 3 to make lecture easier. For sample irradiated at the lowest dose, sample 1E13, the change in peak positions and peak broadening could not be detected compared to the virgin sample. For 5E13 sample and higher fluence levels, a continuous drop of peak intensity and increase of peak width could be found, and for 2E15 sample, very broad peaks appeared, which could mean that the degree of crystallinity decreases. A systematic shift of (0 0 8) peak (and also of the weak (0 0 6) and (1 0 9) peaks) to low angle could also be noticed. The shifting increases with irradiation dose and at last are the same for 5E14 and 1E15 samples showing the saturation of the phenomenon. The shift of (0 0 8) related peaks, and to a less extend, shift of (1 0 9) peaks, is consistent with previous observations via grazing XRD [16] (no obvious peak shift could be observed for the (1 1 0) diffraction). However, no amorphisation is obvious. This shift can be linked to continuous expansion of the unit cell along c axis. The lattice parameters after irradiation will be given in the refinement part. Another important characteristic of these patterns is the increase of relative intensity of (1 0 2) and (1 0 3) to (1 0 1) peaks for samples irradiated up to 5 1014 ionscm2 and above (the
same conclusion was drawn with considering integral peak areas ratios obtained by TOPAS software). Especially for 1E15 sample, intensity of (1 0 2) diffraction is even larger than that of (1 0 1). This pattern feature is consistent with simulated b-Ti3SiC2 diffraction patterns by Farber et al. [22], suggesting presence of this polymorph after irradiation. Besides a strong broadening of peaks centred around (1 0 3), (1 0 5) and (1 1 0) diffraction positions of virgin sample could also be noticed in the 1E15 sample. These broad peaks could be the sum of diffraction response from different phases. These results suggests strong modifications in the Ti3Si0.90 Al0.10C2 microstructure after Xe irradiation, dealing with: (i) change in degree of crystallinity (pattern intensity), (ii) change in lattice (peak broadening and shifting) and (iii) phase transformation (peak relative intensity and unexpected important broadening). In addition, some narrow non-shifted peaks could be found for 5E14, 1E15 and 2E15 samples (located at the same angles as for virgin sample). For 2E15 sample, these peaks are more intense than the two main broad peaks. This means there is a contribution from small amount of material that has not been affected by irradiation. It should not come from surface of irradiated sample since other grazing XRD analysis did not show any non-shifted peaks [16]. These peaks should rather be generated from bulk material underneath the ion implantation zone. This contribution from bulk is effective for all the samples but detected for highly irradiated samples only because of large shift (no overlap) and low intensity of peaks from irradiated layer. For samples irradiated with Kr ions (lower energy), a similar evolution of diffraction peaks was observed. The peak intensity drops and (0 0 8) peaks shift started from 1E14 sample. The increase of peak breadth was not as large as that of samples irradiated at the same dose with Xe ions. This observation confirms that lower energy Kr ions induced less damage compared to high energy Xe ions. 3.1.2. High temperature irradiation A recovery of peak intensity and peak position were found for high temperature irradiated samples for both ions. An example is given in Fig. 3 for 1E15 sample irradiated with Xe ions at 20 °C, 300 °C and 500 °C. Thermal instability of irradiation produced defects should be the reason of this recovery (recombination of irradiation defects).
Fig. 2. XRD patterns of samples irradiated at various doses with Xe ions.
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Fig. 3. XRD patterns of 1E15 samples irradiated at various temperatures with Xe ions.
3.2. Rietveld refinement and micro-structural analysis
e¼
3.2.1. Principles Pseudo-Voigt (pV(x)) function was selected to refine peak profile. Pseudo-Voigt function is an approximation of the Voigt function defined as the linear combination of a Lorentzian and a Gaussian functions (L(x) and G(x)). It has the following expression:
pVðxÞ ¼ gLðxÞ þ ð1 gÞGðxÞ
ð1Þ
with g ¼ g0 þ X2h
ð2Þ
where g0 and X are refined parameters to get Lorentzian mixing index g. The X coefficient is only refined for the LaB6 in order to measure accurately the instrumental broadening. For the refinement of the Ti3SiC2, the X coefficient was set to zero. For Pseudo-Voigt function, the full width at half maximum (FWHM) (H) has the following relationship with strain and size parameters:
H2 ¼ ðU þ D2ST Þ tan2 ðhÞ þ V tanðhÞ þ W þ
IG cos2 ðhÞ
ð3Þ
where U, V, W and DST are micro-strains effects parameters and IG is a domain size effects parameter. It is considered that V and W in this function should be zero when instrumental resolution function (IRF) was used and DST was also set to zero. So the evolution of the FWHM (with H in degrees) can be simply written as
H2 ¼ U tan2 ðhÞ þ
IG cos2 ðhÞ
ð4Þ
Basically, U and IG can be related to micro-strains and domain size through integral breadth (b in radians). From Stokes and Wilson [23] and Scherrer [24] function we know that
bstrain ¼ 4e tanðhÞ k bsize ¼ D cosðhÞ
ð5Þ ð6Þ
For Pseudo-Voigt function, integral breadth (b in radians) can be related to the FWHM (H in degrees) by
b¼
p p H pffiffiffiffiffiffiffiffiffiffiffiffiffi 180 2 g þ ð1 gÞ p ln 2
ð7Þ
So micro-strain e can be obtained by the refined parameter U (in degrees2) through the following function:
p p
pffiffiffiffi U
180 8 g þ ð1 gÞ
pffiffiffiffiffiffiffiffiffiffiffiffiffi p ln 2
ð8Þ
And coherently the domain size D could be written as a function of IG (in degrees2) through
D¼
180 2k g þ ð1 gÞ pffiffiffiffi p p IG
pffiffiffiffiffiffiffiffiffiffiffiffiffi p ln 2
ð9Þ
Based on a Hall and Williamson analysis not described here it has been chosen to refine only the micro-strains (U parameter) for the a-phase and only the domain size (IG parameter) for the b-phase. The parameter g0 was refined and in some case for high irradiation dose, it is fixed to 1 because the peaks are broad (Lorentzian profile). It must be noticed that the microstructure features should follow the damage profile caused by the passage of ion in the specimen. However, specific knowledge of microstructure profile requires extended TEM studies that are not available at the present time. Then, it sounds premature to take this into account in the refinement, i.e. increase the number of refinement parameters in Rietveld analysis. 3.2.2. Refinement results We found the agreement factors Rp and Rwp (reliability factor and weighted reliability factor, respectively) were too high when performing Rietveld refinement for highly irradiated samples if common experimental a-Ti3SiC2 phase only was considered. That is why the hypothesis was made to consider the high temperature b-Ti3SiC2 phase; the resulting refinements were considerably improved as shown in the following. b-Ti3SiC2 was first proposed by Farber et al. [22] through TEM observation: they suggested it was transformed from a-Ti3SiC2 during ion milling processing. But Yu [25] argued that low energy ion milling could not induce this phase transformation and what Farber et al. observed was actually still a-Ti3SiC2. Later, Wang and Zhou [26] made detailed first-principles total energy calculations and pointed out that a reversible polymorphic phase transition would occur when shear strain energy was large enough to overcome an energy barrier. Actually, the two polymorphs show the same space group symmetry, P63/mmc; the only difference is the Si occupancy in the (0 0 0 1) plane. In a-Ti3SiC2 the Si atoms occupy 2b Wyckoff position with fractional coordinates (0, 0, 1/4), while in b-Ti3SiC2, Si
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occupy 2d Wyckoff position with fractional coordinates (2/3, 1/ 3, 1/4). The crystal structures of the polymorphs are illustrated in Fig. 4. In the two sketches, Ti2 refers to atom bounded to Si atom. Ti2 and C atomic positions can be adjusted without breaking the symmetry restrictions. In this work, for a-Ti3Si0.90Al0.10C2 phase, the two positions were taken from literature [27]. For b-Ti3Si0.90Al0.10C2, since no experimental data could be found elsewhere, they were refined using XRD data for 5E14 sample, for which bphase formation was previously suggested in Section 3.1 (see Table 1). The values were then used for 1E15 and 2E15 samples. An example of quality improvement of the refinement by introducing b-phase is given in Fig. 5. Some of the parameters deduced by Rietveld refinement from several highly irradiated materials are listed in Table 2. The volume fraction of b-phase (which starts to form for 5E14 sample) increased with dose for Xe irradiation and, for 2E15 sample, almost whole sample was transformed into b-phase (a-phase is only about 1%). For Kr irradiation, the b-phase appeared for dose higher than 1 1015 ionscm2. Sample irradiated at 300 °C with Xe showed only 30% of b-phase, and only a-phase could be detected for both Xe and Kr irradiation at 500 °C. An increase in a-phase lattice parameter was highlighted. The increase of lattice parameter a (less than 0.2%) with increasing irradiation dose does not seem significant compared to that of c parameter. The change in c parameter with dose is given in Fig. 6. The increase in c parameter is a bit higher than previously reported by [16]: grazing XRD on Xe irradiated Ti3SiC2 (1 1015 ionscm2) lead to an increase of 1.07% linked to effect of electronic interactions only. An increase of 1.3% was found in this work via basic XRD. Actually, the microstructure features should change with the damage profile (Fig. 1), i.e. increase in c
Table 1 Structural parameters of a-Ti3Si0.90Al0.10C2 and b-Ti3SiC2. Formula
a-Ti3Si0.90Al0.10C2 [27]
b-Ti3SiC2
Space group Lattice parameters (nm)
P63/mmc (194) a = 3.0690 c = 17.6785
P63/mmc (194) a = 3.0344 c = 18.0178 [26]
(0, 0, 0) (2/3, 1/3, 0.1329) (0, 0, 0.25) (1/3, 2/3, 0.0724)
(0, 0, 0) (2/3, 1/3, 0.134) [28] (1/3, 2/3, 0.25) (1/3, 2/3, 0.071) [26]
Atomic positions Ti1 Ti2 Si C
Fig. 5. Rietveld refinement of XRD patterns of Xe irradiated 1E15 sample: (a) considering a-Ti3SiC2 only and (b) considering both a-and b-Ti3SiC2. In both cases, contribution from underneath non irradiated bulk material (a-Ti3SiC2) was introduced to refine the non-shifted peaks.
Fig. 4. Crystal structures of a- and b-Ti3SiC2 polymorphs. Species of atoms and corresponding Wyckoff positions are illustrated.
parameter with depth. Nevertheless, at this stage of the investigation, this average parameter is useful in the scope of comparative study of irradiation impact and it did not sound reasonable to increase the number of refinement parameters in Rietveld analysis. Furthermore, it seemed that no additional broadening of the (0 0 8), (0 0 6), (1 0 9) peak occurred suggesting that the profile in c parameter should have a limited impact on the XRD diagram. It seems that c parameter reaches a saturation value for 5E14 sample (the value dropped a little when irradiated to 1 1015 cm2). When irradiated at high temperature (500 °C), the
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X.M. Liu et al. / Nuclear Instruments and Methods in Physics Research B 268 (2010) 506–512 Table 2 Rietveld refinement results of Ti3Si0.90Al0.10C2 samples irradiated at various fluence levels. Ion
T (°C)
Dose (ions/ cm2)
Phase
a (Å)
c (Å)
U (degrees2)
g0
Kr Kr Kr Kr
RT RT RT RT
0 1E12 1E13 1E14 1E15
a a a a a
0.60 0.76 0.78 0.70 0.64
500 RT RT RT RT
1E15 1E13 5E13 1E14 5E14
0.0472 0.0167 0.0620 0.1101 0.5120
Xe
RT
1E15
17.6787 17.6791 17.6901 17.7394 17.9728 19.2389 17.7452 17.6917 17.7418 17.7836 17.9195 19.3658 17.9097 19.4151 19.6002 17.9253 19.6087 17.8504
0.0039 0.0054 0.0158 0.0821 0.3462
Kr Xe Xe Xe Xe
3.0678 3.0679 3.0698 3.0715 3.0689 3.0399 3.0695 3.0691 3.0717 3.0720 3.0691 3.0178 3.0669 3.0282 3.0340 3.0718 3.0232 3.0698
0.74 0.62 0.65 0.62 0.79 1 1 1 1 0.51
b
a a a a a b
Xe Xe
RT 300
2E15 1E15
a b b
a b
Xe
500
1E15
a
IG (degrees2)
1.62
1.8660
0.4697
2.56 3.42 3.66 1.01
0.1179
0.74
Fraction of phase (%)
Micro-strains in aphase (103)
100 100 100 100 68.9 ± 1.5 28.2 ± 1.3 100 100 100 100 60.5 ± 1.2 35.0 ± 1.2 30.4 ± 0.9 65.6 ± 1.6 95.7 ± 0.8 71.5 ± 1.6 28.5 ± 1.4 100
0.358 0.454 0.786 1.780 3.449
Domain size in bphase (nm)
Rp
Rwp
16.2 15.9 21.3 22.4 23.8
19.1 18.7 25.0 24.8 26.9
23.6 18.3 22.0 21.9 21.3
26.4 20.4 24.1 24.2 25.6
18.7
22.8
21.9 22.3
25.6 26.5
21.3
25.1
4.40 1.323 0.749 1.464 1.923 4.445 3.81 9.360 3.03 3.08 3.810
5.62 2.090
Note: Values with mean these data were fixed or that the patterns could not be refined.
Fig. 6. Relative increase in lattice parameter c of a-Ti3Si0.90Al0.10C2 for samples irradiated with Xe and Kr ions at various temperatures (error bars are smaller than plot size).
Fig. 7. Micro-strains in a-Ti3Si0.90Al0.10C2 for samples irradiated with Kr and Xe ions at various temperatures.
4. Discussion and conclusions increase of c parameter is much smaller than that of room temperature irradiated sample, especially for Kr irradiated sample: this is consistent with microstructure recovery due to defect annealing. Same trend was observed for the change in unit cell volume. The broadening of a-Ti3Si0.90Al0.10C2 peaks was explained in terms of micro-strains contribution (suggested by a Hall and Williamson analysis not detailed here). Micro-strain values obtained at various doses and temperatures are presented in Fig. 7. Whereas no significant change is observed up to 1 1013 cm2, microstrains increased for higher dose for both Kr and Xe irradiated samples. It should be noticed for sample irradiated to 1 1015 cm2 at room temperature that the micro-strains of Xe irradiated sample are about three times of those irradiated with Kr. For samples irradiated at high temperature, a decrease in micro-strains is clearly seen for both ions: the higher the temperature, the larger the drop. A more pronounced drop could be expected for a higher irradiation temperature until complete recovery of the microstructure. It should be reminded that for b-phase, only domain size effect was considered as origin of peak broadening and IG in peak shape function was refined but U was set to zero.
Ti3Si0.90Al0.10C2 ceramics were irradiated at various doses (up to 2 1015 ionscm2) and various temperatures (20 °C, 300 °C and 500 °C) with high energy Kr and Xe ions. XRD patterns showed peak intensity decrease, peak width increase shift in peak position with increasing irradiation dose for both ions: this means progressive change in the microstructure, which is consistent with increase in irradiation damage. Rietveld refinement suggested that irradiation induced the following: Formation of b-Ti3Si0.90Al0.10 nano-domains above 1 1015 cm2 for Kr and 5 1014 cm2 for Xe ions irradiation. Large increase of lattice parameter c of a-Ti3Si0.90Al0.10. Appearance of micro-strains in a-Ti3Si0.90Al0.10. Defect annealing during high (300 °C, 500 °C) temperature irradiation resulting in moderate modification of microstructure. Such a recovery was suggested elsewhere by nanoindentation analysis [17]. Phase transitions induced by ion irradiation have already been observed in metal, alloys [29,30], oxide compounds [31,32]. But
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in other materials, such as in spinel compounds MgAl2O4 [33] and ZnAl2O4 [34], under low energy or swift heavy ions irradiation, a simple exchange of cations onto the spinel sites took place without any change of the space group. The present work suggests that ion irradiation of a-Ti3Si0.90Al0.10C2 results in the formation of b-Ti3Si0.90Al0.10C2 which has the same space group. The only difference between the two phases are the atom positions of Si. Calculation showed that, in Ti3SiC2 compounds, the bonding energy of Si with adjunction of Ti is quite low compared with Ti–C bond [26]. So under high energy irradiation Si atoms should be more easily displaced from their original position to form interstitial atoms. Since 2d sites are the largest holes in a-phase [34], Si atoms could hop from 2b sites to 2d sites to generate Frenkel defects (vacancies at 2b and interstitials at 2d). The number of Frenkel pairs should remain small when irradiation dose is low and may be randomly distributed in the matrix as separate defects or small defect clusters. With increasing irradiation dose, the number of Frenkel pairs increases leading to microdistorsions in the cell. At a threshold irradiation dose, enough interstitial atoms could line up between two layers of Ti and form small domains of b-phase: the adjacent Ti and C atoms would relax and found new positions (no strains), but should not break restrict of symmetry. The amount of b-phase would increase with dose due to the increasing number of interstitial atoms. When irradiations were performed at high temperature (300 °C and 500 °C) and high dose, the interstitial Si atoms should be easier to hop back to low energy 2b original site. So fewer defects would be formed in this case resulting in a lower amount of b-phase, which is consistent with the present results. Further investigations are in progress to directly highlight the presence of b-phase in the irradiated samples. Among them, XPS analysis should be able to detect Si–C bonds related to b-phase, as suggested by Yu et al. [28]. Raman spectrum and HRTEM could also give evidences of the presence of this phase and related works are in progress. According to lattice changes (parameter, strains), the recovery of the microstructure is well advanced after irradiation at 500 °C. It is reasonable to suppose that further recovery occurs at higher temperature. Then, in the operating temperature range of GFR (600–1000 °C), irradiation should have less influence on the properties related to the layered structure, e.g. embrittlement and damage tolerance. That means Ti3SiC2 ceramics could be a promising high temperature element of GFR core components.
Acknowledgements Authors are very grateful to Professor Yanchun Zhou from Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China, for having provided the samples. He is also thanked for collaborative work and fruitful discussions.
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