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Corrosion of hot pressed Si3N4–TiN composite in sulphuric acid aqueous solution
Corrosion Science 45 (2003) 2525–2539 www.elsevier.com/locate/corsci
Corrosion of hot pressed Si3N4–TiN composite in sulphuric acid aqueous solution V. Medri, M. Bracisiewicz, A. Ruffini, A. Bellosi
*
CNR-ISTEC, Institute for Science and Technology for Ceramics, Via Granarolo 64, Faenza 48018, Italy Received 2 August 2002; accepted 4 March 2003
Abstract The corrosion behaviour of an electroconductive Si3 N4 –35 vol% TiN composite, hot pressed with the addition of Al2 O3 and Y2 O3 as sintering aids, was studied in 1.8 M sulphuric acid aqueous solution at RT, 40 and 70 °C up to 400 h. The corrosion follows linear kinetics at RT and at 40 °C, involving only the progressive chemical dissolution of grain boundary phases, in the system Al–Y–Si–Ti–O–N. Chemical attack of TiN occurs at 70 °C, while Si3 N4 is not affected by the selected corrosive environment. The effect of the corrosion on flexural strength, fracture toughness and electrical resistivity were investigated. Very high strength levels are maintained after corrosion for 400 h at room temperature, while the strength decreases of about 10% and 16% after the treatment at 40 and 70 °C, respectively. The electrical resistivity rises after corrosion at 40 and 70 °C, in line with the progressive chemical dissolution of the conductive TiN particles. Ó 2003 Elsevier Ltd. All rights reserved. Keywords: Ceramic matrix composite A; SEM B; Acid corrosion C
1. Introduction Si3 N4 based ceramics are among the most important materials for structural applications because of their superior properties such as strength, hardness, thermal and chemical stability. Since the 1990s the attention has been devoted to
*
Corresponding author. Tel.: +39-546-699-724; fax: +39-546-463-81. E-mail address: [email protected] (A. Bellosi).
0010-938X/$ - see front matter Ó 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0010-938X(03)00077-5
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particulate-reinforced ceramics: the addition of an electrically conductive secondary phase such as TiN to the brittle Si3 N4 ceramics has been developed in order to improve toughness and strength and to induce electrical conductivity for specific applications in the field of high temperature heaters, igniters and heat exchangers [1–10]. While the corrosion behaviour of Si3 N4 materials have been extensively studied at high temperature [11–17] and in some extent at intermediate and low temperature in aqueous corrosive environments [18–41], only few studies have been reported on high temperature oxidation and corrosion of Si3 N4 –TiN particulate composite [42– 47] and no information about the corrosion in aqueous environment of these composites was found in literature. It is well known that the chemical stability of Si3 N4 based ceramics in acid or basic environments is strongly affected by grain boundary phase, depending on its composition, amount, distribution and crystallinity [18,19,22,24,27,29–35,37,38,40,41], while Si3 N4 matrix is usually less affected by the corrosion. The addition of a secondary phase, such as TiN, influences the overall corrosion of the material due to two concurrent features: (i) TiN has individual and specific corrosion characteristic, (ii) it changes the composition of the grain boundary phase [2]. In this study, the corrosion behaviour of an electroconductive hot pressed Si3 N4 –35 vol% TiN composite containing Al2 O3 and Y2 O3 as sintering aids was investigated in 1.8 M sulphuric acid aqueous solution at room temperature, 40 and 70 °C up to 400 h. The corrosion kinetics were determined and the mechanism governing the corrosion was defined. The degradation of the flexural strength and toughness and the decrease of the electrical conductivity were also evaluated in relationship with the microstructural changes which are the consequence of the corrosion.
2. Experimental procedure A fully dense hot pressed Si3 N4 –35 vol% TiN composite containing 1.05 wt.% Al2 O3 and 2.63 wt.% Y2 O3 as sintering aids, was employed for corrosion tests. Processing procedures are described elsewhere [47]. Rectangular plates 10.0 10.0 1.0 mm3 , whose wider surfaces were polished up to 6 lm, were used for corrosion tests. The specimens were cleaned in ultrasonicated acetone bath, dried, weighed (accuracy 0.01 mg) and then sealed in a polyethylene tube containing 50 ml of 1.8 M sulphuric acid solution, previously thermostated at the desired test temperature. The corrosion tests were performed at room temperature (RT, 20 °C), 40 and 70 °C for holding time up to 400 h. After the planned exposure time, the specimens were removed from the tubes, rinsed in boiling deionizated water, dried and weighed. Then the weight loss DW =S (mg/cm2 ) was calculated. A Varian Liberty 200 inductively coupled plasma atomic emission spectrometer (ICP-AES) was used for the determination of the Ti, Si, Y and Al cations released into the corrosive solutions. The inert ‘‘V’’ groove nebulizer was chosen because of the corrosive matrix. The following analytical wavelengths were chosen (nm): Si 251.611, Al 396.152, Y 371.030, Ti 334.941. These wavelengths showed a good linear
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response between concentration and intensity, and were free from interferences in the range of concentrations used in the present study. The viewing height was optimized on SBR (signal background ratio). In order to obtain the best linearity in the calibration curve, matrix-matched multielemental standard solutions were prepared in the same concentration range of the solutions under analysis. Microstructural characteristics of the corroded specimens were analyzed through X-ray diffraction and scanning electron microscopy (SEM). In particular the exposed surfaces, the polished cross sections and the fracture surfaces were examined. Mechanical strength, fracture toughness and electrical resistivity were measured on specimens cut in the size and shape of bars 2 2.5 25 mm3 with the faces ground to a longitudinal surfaces roughness of 0.17 lm and the edges chamfered. The tests were performed on the starting material and after a permanence of 400 h in the corrosive solution at RT, 40 and 70 °C. Flexural strength ðrÞ was measured with an Instron mod 1195, in 4-pt bending with 20 and 10 mm as outer span and inner span, respectively, using a crosshead speed of 0.5 mm/min. Fracture toughness ðKIc Þ was evaluated using the chevron-notched beam (CNB) in flexure. The test bars were notched with a 0.08 mm diamond saw; a0 and a1 were about 0.12 and 0.80, respectively. The flexural tests were performed on a semiarticulated alumina four-pt jig with a lower span of 20 mm and an upper span of 10 mm on an universal screw-type testing machine Instron mod.1195. The specimens were deformed with a crosshead speed of 0.05 mm/min. The ‘‘slice model’’ equation of Munz et al. [48] was used for the calculation of KIc . The electrical resistivity ðqÞ was measured at room temperature on the starting material and RT corroded bars by means of a four probe DC method, and on the bars corroded at 40 °C using a two probe DC method, inducing in both cases a longitudinal current along the bars. The current and the voltage readings were made at the same time in two different digital high-resolution multimeters. The resistivity values were determined from the electrical resistance measurements taking into account of the test leads distance and cross section area of the samples. On the bars corroded at 70 °C, the above described technique did not allow to detect the high resistivity, due to the fact that the modifications of the specimen surface are not homogeneous, because of the presence of a thick superficial corroded layer. Therefore, volume resistivity measurements on samples treated at 70 °C were performed with a guarded resistivity cell in conjunction with a high resistance meter. In order to assure a good and uniform electrical contact with the two larger sample surfaces where the voltage was applied, the specimens were placed in the test chamber with two conductive rubber electrodes between the sample and the metal electrodes of the cell. The measurements were carried out by applying a voltage of 10 V and waiting 2 min before reading the electrical resistance value. The resistivity values were finally determined from the resistance measurements taking into account the surface area of the main electrodes and thickness of the samples.
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3. Result and discussion 3.1. Corrosion kinetics and mechanisms The microstructure of the as-hot pressed material consists of a fine-grained matrix of b-Si3 N4 grains surrounded by a glassy grain boundary phase and of well dispersed secondary phase (TiN) particles (Fig. 1). The grain boundary phase is composed mainly of amorphous phases in the systems Al–Y–Si–O–N and Y–Ti–O. The main part of alumina introduced as sintering aids enters in solid solution in b-Si3 N4 grains during the re-precipitation phenomena that occur in the liquid phase sintering process [49]. Moreover, Y2 TiO5 was identified among the re-crystallized grain boundary phases after annealing tests [47], indicating that during sintering of the mixture Si3 N4 –TiN, TiO2 present on the surfaces of TiN particles interacts with Y2 O3 leading to the formation of yttrium titanate [2]; this phase is therefore present in the amorphous state in the as hot pressed material. XRD patterns detected on the surfaces of the samples treated at 70 °C show a progressive decrease of the ratio R (1) between the heights of the TiN (200) and the Si3 N4 (101) peaks: R ¼ hTiN ð200Þ =hSi3 N4 ð101Þ
ð1Þ
from R ¼ 3:5 of the starting material to R ¼ 2:9 and 2.3 respectively after 100 and 400 h at 70 °C. This means that TiN amount decreases in the corroded samples compared to the starting material. On the contrary, the phase b-Si3 N4 does not undergo detectable variation. The microstructure of the corroded samples shows that the surface is nearly smooth but pitted, particularly after the tests at 70 °C. No masking layer was noticed after the tests under the various conditions, but rather an etched surface showing that the preferential corrosion sites are the grain boundary phase and the triple junctions, the TiN grains boundaries and small areas at the interface between the
Fig. 1. Secondary electron image of a polished and plasma etched surface of the hot pressed Si3 N4 –35 vol% TiN composite, showing the dispersion of TiN particles in the Si3 N4 matrix material.
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Fig. 2. Back scattered image of the surface corroded at 40 °C for 20 h: (1) grain boundary phase-depleted areas; (2) depleted TiN grain boundary; (3) leached areas between TiN and Si3 N4 .
TiN particles and the Si3 N4 matrix, leading to the formation of pits and leached areas (Fig. 2). The surface morphology is almost the same at room temperature and 40 °C: the surface voids and defects correspond to selective attack of the grain boundary phases (Fig. 3a and b). The leached areas in the samples treated at 70 °C increase with time: besides the grain boundary phases, TiN particles start to be attacked after short exposure time (Fig. 3c). After 400 h at 70 °C the exposed surface is extensively corroded (Fig. 3d): the superficial TiN grains are almost completely dissolved, as also confirmed by X-ray analysis. The weight loss of the samples in function of the corrosion time (Fig. 4) evidence that the corrosion behaviour at each test temperature and in the whole time range can be modelled by a simple linear rate law (2) [50]: DW =S ¼ kt
ð2Þ
The linear kinetics indicate that the rate controlling step is a chemical reaction, i.e. a dissolution of the reacting species at the reaction interface. The linear rate of the interface advance from the exposed surface towards the bulk of the specimen is confirmed by the analysis of the cross sections, where the portion of the material that is affected by corrosion is distinguished by the presence of porosity. The depth of penetration of the chemical attack and measured from the SEM micrographs, progressively increases with temperature and show a linear dependence with time (Fig. 5), with the same tendency exhibited by the weight loss curves. The values of the reaction rates, k (calculated both from the weight loss: mg cm2 h1 and from the corrosion layer thickness: lm h1 ) reported as a function of temperature in Arrhenius plot (Fig. 6) result in an apparent activation energy Ea ¼56 kJ/mol. The microstructures of the corroded layers in Figs. 7 and 8 highlight that the formation of porosity is due to the dissolution of the grain boundary phase and/or of TiN particles at 70 °C.
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Fig. 3. Secondary electron images of surfaces after corrosion at: (a) room temperature for 400 h; (b) 40 °C for 400 h; (c) 70 °C for 100 h; (d) 70 °C for 400 h.
At 20 °C the external layer of the sample, affected by chemical attack, is very thin and not continuous all over the surface up to about 200 h of exposure. A gradient of dissolution is present from the surface towards the bulk after all the tests performed. Moreover, because of the presence of some porosity in the bulk material mainly located around TiN grains, the contact between the reactants and TiN is favoured. It leads to improved partial dissolution of the external part of the TiN particles, this being evident after the corrosion at 70 °C. It could be thought that the boundaries of the TiN particles act as a preferential path for corrosion, which therefore propagates following the arrangement of the TiN network inside the composite material. The graded chemical attack is well evident in Fig. 8, that represent the microstructure modification after 400 h at 70 °C. Although the reaction interface advance is almost regular (the peak-valley distance measured on the front line is 5 lm), near the surfaces an extensive corrosion attack of both grain boundary phase and TiN
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7
k = 0.0150
RT
6
40°C ∆W/S, mg/cm2
5
70°C
4 3
k = 0.0035
2 1
k = 0.0005 0 0
100
200 Time, h
300
400
Fig. 4. Weight loss ðDW =SÞ against time plot, showing a perfect linear trend at each temperature in the whole time range.
300
k = 0.0187
RT 250
40°C dCL, µm
200
70°C
150 100
k = 0.1706
50
k = 0.5790 0
0
100
200
300
400
Time, h
Fig. 5. Thickness of the corrode layer ðdCL Þ against time plot: dCL progressively increases with temperature and show a linear dependence with time as weight loss curves at each temperature in the whole time range.
phase takes place, while, deeper into the corroded layer, mainly grain boundary phase is dissolved. The degree of dissolution vi of each ion was calculated from the amount of Ti, Si, Y and Al released into the corrosive solution, with the following relationship (3): vi ¼ A=B
ð3Þ
where A and B are respectively the weight of the element released into the solution (measured by ICP-AES) and of the same elements present in the untreated material. The degrees of dissolution of the three cations Y, Al, Si, are represented in function of the corrosion time, in separate plots relative to each test temperature (Fig. 9). The
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from weight gain -1
from corroded layer thickness
Ea= 56.3 kJ/mol
-2
-1
[k, µm h ]
-ln k
-3 -4 -5
Ea= 56.2 kJ/mol -6
-2
-1
[k, mg cm h ]
-7 -8 0.0029
0.0032
0.0035
1/T, K-1
Fig. 6. The values of the reaction rates, k (calculated both from the weight loss: mg cm2 h1 and from the corrosion layer thickness: lm h1 ) reported as a function of temperature in Arrhenius plot result in an apparent activation energy Ea ¼56 kJ/mol.
Fig. 7. Back scattered images of the cross sections after corrosion for 400 h at room temperature (a) and 40 °C (b).
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Fig. 8. SEM micrographs of the cross section of the sample corroded at 70 °C for 400 h, showing the gradient of dissolution from the surface towards the bulk.
trend for the release of the these elements is constant: the degree of dissolution of yttrium, vY is always the highest, vAl is more then two times lower then vY , while vSi is the lowest one: it resulted 20 and 40 times lower than vAl and vY , respectively. The degrees of dissolution of the same cations relatively to the corroded volume of the ceramic composite were calculated with the relationship (4): vVi ¼ A=C
ð4Þ
where A and C are, respectively, the weight of the elements released in the solution and present in the corroded volume (estimated from the thickness of the corroded layer, dCL ). The data show that at 70 °C vVY 1, i.e. yttrium is totally leached in the corroded volume, while vVAl 0:46 and vVSi 0:03, i.e. the residual amounts of Al and Si in the corroded volume are respectively about 54% and 97%. The degree of dissolution of Ti relative to the three temperatures is shown in Fig. 10. vTi at 70 °C is significantly higher than that at room temperature and 40 °C. The results reported above indicate that the linear rate of the interface advance is consequent to the dissolution of the grain boundary phases at all the three tested temperatures and to the additional attack also on titanium nitride particles at 70 °C. A possible contribution to the corrosion from the dissolution of silicon nitride grains, particularly at 70 °C, as previously observed [40], has to be excluded due to
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V. Medri et al. / Corrosion Science 45 (2003) 2525–2539 0.04
RT
Y 0.03
Al
χi
Si 0.02
0.01
0 0
100
200
300
400
300
400
300
400
Time, h 0.20
40°C
χi
0.15
0.10
0.05
0 0
100
200
Time, h 0.6
70°C
χi
0.4
0.2
0 0
100
200
Time, h
Fig. 9. Degrees of dissolution vI of Y, Al and Si plotted as the function of time at room temperature, 40 °C and 70 °C.
the fact that only 3% of Si is removed from the corroded material and also because the ratios vY =vSi and vAl =vSi are constant for all the temperature and time of the tests. Therefore, at room temperature and at 40 °C, only the grain boundary phases in the system Al–Y–Si–O–N and Y2 TiO5 are leached. The difference in the values of
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0.04 RT
χ Ti
0.03
40°C 70°C
0.02
0.01
0 00
100 100
200 200
300 300
400 400
Time, h Fig. 10. Degree of dissolution vTi reported as the function of time at the three temperatures (RT, 40 and 70 °C).
Y and Al cations released in the solution is due to the fact that part of Al is within the b–Si3 N4 network [49], therefore only about 46% of Al introduced as alumina, as sintering aid, contributes to the formation of the grain boundary phases. The values of the weight of the cations released during the chemical dissolution of the grain boundary phase and measured in the corrosive solution by ICP-AES were compared to those of the weight gain for attack. There is an amount of light elements (N and O) solved from the grain boundary phase that cannot be measured, but only indirectly evaluated. The calculations, relative to the data from the tests at RT and at 40 °C, were based on the hypothesis that the Ti cations were released from the Y2 TiO5 compounds and that another phase contains: the residual Y, the measured Al and Si and an amount of N and O. The results show that the approximation of the ICP data and of the weight loss data is within 2% if a grain boundary phase with composition AlSi3 Y2 O4:5 N4 is present, in addition to Y2 TiO5 . The same evaluations are not possible on the data from the tests at 70 °C, because also the corrosion of TiN grains is massively involved. As above indicated, the corrosion of this Si3 N4 –TiN composite is controlled by the surface chemical reaction in H2 SO4 aqueous solution. The leachability of the grain boundary phases from acid aqueous environment arises from their own basic nature. In particular Y cation behave as a network modifier of the silicon-oxide-nitride grain boundary phase [51] and confer it a basic character: these compounds are known to be easily attacked by such environments [21,25,50]. Therefore, the corrosion of the composite under test in a sulphuric acid aqueous solution is due to the leaching of Y and Al ions, i.e. to the surface exchange between Hþ and exchangeable ions on or near the fresh grain boundary. Proton-induced corrosion of ceramics by sulphuric acid was previously reported [52,53]. Ion/proton exchange was indicated as the mechanisms determining chemical attack for lithium and barium aluminosilicates in H2 SO4 ; in this case ion exchanges induces phase transformations, accompanied by
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the effects related to volume changes [52]. In our case, on the contrary, the glassy grain boundary phase dissolves: no barrier diminishes the effective contact area between reactants and the desorption of Y and Al at the interface is rate limiting. The evident degradation observed after the corrosion at 70 °C follows the dissolution of TiN particles; the consequent enhanced penetration towards the bulk of the reactant species also improves the dissolution of the grain boundary phases. The apparent activation energy calculated in this case is within a range of values reported in literature for the corrosion of similar systems [41] although in these cases the dissolution of the grain boundary phase was followed by a passivation process. Compared to the corrosion behaviour shown in literature of other silicon nitrides under similar testing conditions, the selected Si3 N4 –TiN composite has a rather good corrosion resistance in the sulphuric acid aqueous solutions, at least for relatively low temperatures, where the second phase, i.e. TiN, is not involved in the reaction. It has to be underlined that in our case the weigh loss for corrosion is about 5 times lower than that reported in literature for a silicon nitride sintered with the same sintering aid system. 3.2. Properties degradation The mechanical and electrical properties of Si3 N4 –35 vol% TiN composite, before and after corrosion tests at room temperature, 40 and 70 °C for 400 h, are reported in Table 1. The corrosive attack at room temperature doesnÕt lead to a degradation of properties, as flexural strength, fracture toughness and electrical resistivity are almost the same of the as-sintered material. Increasing the corrosion temperature an evident decrease of strength and of the electrical conductivity is observed (Table 1). The electrical resistivity improves of about 6 and 12 order of magnitude as a consequence of the attack at 40 and 70 °C, respectively. This is due to the porosity and progressively leak of percolation in the TiN electroconductive particle network that arises from dissolution of the grain boundary phase and of TiN particles in the corroded layer. The corrosive attack at 40 and 70 °C lowered the flexural strength of about 10% and 30%, respectively. This behaviour is consequent to the increased superficial and subsurface population of defects due to corrosion. Fig. 11 shows, as an example, a typical critical defect acting as fracture origin, in a test bar corroded at 70 °C, after the strength measurement. Table 1 Room temperature flexural strength, fracture toughness and electrical resistivity of the hot pressed Si3 N4 – 35 vol% TiN composite before and after corrosion tests Corrosion test
rF (MPa)
kIc (MPa m1=2 )
q (X cm)
Starting material RT, 400 h 40 °C, 400 h 70 °C, 400 h
835 116 848 69 759 27 597 24
5.67 0.28 5.61 0.56 Not measurable Not measurable
5.88 104 5–7 104 2–6 102 0.6–5 108
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Fig. 11. Secondary electron images of the fracture surface of a test bar corroded at 70 °C for 400 h: the fracture origin is located in the corroded layer near the interface with the bulk (a, b); in (c, a) particular of the corroded layer near the interface showing the high porosity as a typical critical defect acting as fracture origin.
After corrosion at 40 and 70 °C, it is impossible to evaluate the fracture toughness, due to the fact that the notches are not measurable, because the brittle and porous corroded surface layer of the specimens crumbles during notching and breaking.
4. Conclusion The corrosion in sulphuric acid aqueous solution at RT, 40 and 70 °C of an electroconductive Si3 N4 –35 vol% TiN composite, hot pressed with the addition of Al2 O3 and Y2 O3 as sintering aids, follows linear kinetics: there is linear advance of the reaction interface from the external sample surface towards the bulk and the surface chemical reaction is always rate limiting, with an apparent activation energy of 56 kJ/mol. The degree of dissolution for Y, Al, Si and Ti were measured on the solutions after the tests. The chemical attack at RT and at 40 °C involves only the progressive chemical dissolution of grain boundary phases, in the system Al–Y–Si–Ti–O–N, due to ion/ proton exchange in the glassy structure that induces the leaching of Y and Al ions. The attack of TiN occurs at 70 °C, while Si3 N4 is not affected by the selected corrosive environment. The effect of the corrosion on flexural strength, fracture toughness and electrical resistivity, investigated after a permanence of 400 h at each temperature in the corrosive solution, showed the following: 1. High strength levels are maintained after corrosion at RT and at 40 °C, while still good values (up to about 600 MPa) are detected after corrosion at 70 °C, notwithstanding the severe corrosion.
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2. Fracture toughness does not decrease after corrosion at RT, while the brittleness of the surface corroded layer makes it impossible to evaluate the toughness after corrosion at 40 and 70 °C. 3. No variation is revealed in the electrical resistivity after corrosion at RT; on the contrary, after corrosion at 40 and at 70 °C the materials behave like an insulator. Acknowledgements The work is supported by the European Project Research Training Network HPRN-CT-2000-00044 ‘‘Composite Corrosion’’. The research contract of M. Bracisiewicz is funded by the same Project. The Authors wish to thank their colleagues: S. Guicciardi and C. Melandri for the measurement of mechanical properties and G. Fabbri for the evaluation of the electrical properties.
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[25] D.E. Clark, B.K. Zoitos (Eds.), Corrosion of Glass, Ceramics and Ceramic Superconductors, Noyes Publications, Park Ridge, NJ, 1992. [26] A. Okada, S. Ilo, T. Asano, M. Yoshimura, J. Ceram. Soc. Jpn. Int. Edition 100 (1992) 80. [27] K.A. Nielsen, J.E. Engell, N. Thomsen, in: G. Ziegler, H. Hausne (Eds.), Euro-Ceramics II, vol. 2, Deutsche Keramische, Gesellschaft, Koln, Germany, 1992, p. 1309. [28] F.A. Costa Olivera, R.J. Fordham, J.H.W. De Wit, in: G. Ziegler, H. Hausne (Eds.), Euro-Ceramics II, vol. 2, Deutsche Keramische, Gesellschaft, Koln, Germany, 1992, p. 1351. [29] K.A. Nielsen, H.G. Krarup, J.E. Engell, G. Pritzl, in: P. Duran, J.F. Fernandez (Eds.), Third EuroCeramics, vol. 3, Faenza Editrice Iberica, SL, Spain, 1993, p. 985. [30] K. Oda, T. Yoshio, Y. Miyamoto, M. Koizumi, J. Am. Ceram. Soc. 75 (5) (1993) 1365. [31] S. Ilo, J. Otsuka, T. Niwa, Y. Tajima, M. Watanabe, in: M. Mizutani (Ed.), Advanced MaterialsÕ93, I/A: Ceramics, Powders, Corrosion of Advanced Processing, vol. 14A, Trans. Mater. Res. Soc. J., Elsevier, Amsterdam, 1994, p. 215. [32] T. Yoshio, K. Oda, Corros. Engng. 44 (1995) 481. [33] M. Arai, A. Kohno, Corros. Engng. 44 (1995) 325. [34] A. Bellosi, T. Graziani, F. Monteverde, Key Engng. Mater. 113 (1996) 215. [35] A. Okada, M. Yoshimura, Key Engng. Mater. 113 (1996) 227. [36] S.W. Sharkawy, A.M. El Aslabi, Corros. Sci. 40 (7) (1998) 1119. [37] K. Komeya, T. Meguro, S. Atago, C.-H. Lin, Y. Abe, M. Komatsu, Key Engng. Mater. 161–163 (1999) 235. [38] K.R. Mikeska, S.J. Bennison, S.L. Grise, J. Am. Ceram. Soc. 83 (5) (2000) 1160. [39] S. Somiya, Mater. Chem. Phys. 67 (2001) 157. [40] F. Monteverde, C. Mingazzini, M. Giorgi, A. Bellosi, Corros. Sci. 43 (2001) 1851. [41] J. Schilm, M. Herrmann, G. Michael, Key Engng. Mater. 206–213 (2002) 1117. [42] T. Graziani, D.J. Baxter, A. Bellosi, Key Engng. Mater. 113 (1996) 123. [43] T. Graziani, D.J. Baxter, A. Bellosi, Key Engng. Mater. 108–110 (1995) 227. [44] A. Bellosi, A. Tampieri, L.Y. Zhen, Mater. Sci. Engng. A 127 (1990) 115. [45] F. Deschaux-Beaume, T. Cutard, N. Frety, C. Levaillant, in: P. Vincenzini (Ed.), Proceedings of 9th Cimtec World Ceramic Congress, Techna srl, 1999, p. 389. [46] Y.G. Gogotsi, in: G. Ziegler, M. Hausner (Eds.), Proceedings Europ. Ceram. Soc. 2nd Conf., vol. 2, Ausbourg, 1991, p. 1517. [47] M. Bracisiewicz, V. Medri, A. Bellosi, Appl. Surf. Sci. 202 (2002) 139–149. [48] D.G. Munz, J.L. Shannon Jr., R.T. Bubsey, Int. J. Fract. 16 (1980) R137–R141. [49] A. Bellosi, in: Y.G. Gogotsi, R.A. Andriewski (Eds.), Materials Science of Nitrides, Borides, Carbides, NATO-ARW Series, Kluwer Academic Publisher, Doordtecht, The Netherlands, 1999, p. 285. [50] R.A. McCauley, Corrosion of Ceramics, Marcel Dekker, New York, 1995. [51] M. Ohasi, K. Nakamura, K. Hirao, S. Kanzaki, S. Hampshire, J. Am. Ceram. Soc. 78 (1) (1995) 71. [52] T. Scanu, J. Guglielmi, Ph. Colomban, Solid State Ionics 70/71 (1994) 109. [53] Ph. Colomban, J. Mol. Struct. 270 (1992) 407.
Corrosion of hot pressed Si3N4–TiN composite in sulphuric acid aqueous solution V. Medri, M. Bracisiewicz, A. Ruffini, A. Bellosi
*
CNR-ISTEC, Institute for Science and Technology for Ceramics, Via Granarolo 64, Faenza 48018, Italy Received 2 August 2002; accepted 4 March 2003
Abstract The corrosion behaviour of an electroconductive Si3 N4 –35 vol% TiN composite, hot pressed with the addition of Al2 O3 and Y2 O3 as sintering aids, was studied in 1.8 M sulphuric acid aqueous solution at RT, 40 and 70 °C up to 400 h. The corrosion follows linear kinetics at RT and at 40 °C, involving only the progressive chemical dissolution of grain boundary phases, in the system Al–Y–Si–Ti–O–N. Chemical attack of TiN occurs at 70 °C, while Si3 N4 is not affected by the selected corrosive environment. The effect of the corrosion on flexural strength, fracture toughness and electrical resistivity were investigated. Very high strength levels are maintained after corrosion for 400 h at room temperature, while the strength decreases of about 10% and 16% after the treatment at 40 and 70 °C, respectively. The electrical resistivity rises after corrosion at 40 and 70 °C, in line with the progressive chemical dissolution of the conductive TiN particles. Ó 2003 Elsevier Ltd. All rights reserved. Keywords: Ceramic matrix composite A; SEM B; Acid corrosion C
1. Introduction Si3 N4 based ceramics are among the most important materials for structural applications because of their superior properties such as strength, hardness, thermal and chemical stability. Since the 1990s the attention has been devoted to
*
Corresponding author. Tel.: +39-546-699-724; fax: +39-546-463-81. E-mail address: [email protected] (A. Bellosi).
0010-938X/$ - see front matter Ó 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0010-938X(03)00077-5
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particulate-reinforced ceramics: the addition of an electrically conductive secondary phase such as TiN to the brittle Si3 N4 ceramics has been developed in order to improve toughness and strength and to induce electrical conductivity for specific applications in the field of high temperature heaters, igniters and heat exchangers [1–10]. While the corrosion behaviour of Si3 N4 materials have been extensively studied at high temperature [11–17] and in some extent at intermediate and low temperature in aqueous corrosive environments [18–41], only few studies have been reported on high temperature oxidation and corrosion of Si3 N4 –TiN particulate composite [42– 47] and no information about the corrosion in aqueous environment of these composites was found in literature. It is well known that the chemical stability of Si3 N4 based ceramics in acid or basic environments is strongly affected by grain boundary phase, depending on its composition, amount, distribution and crystallinity [18,19,22,24,27,29–35,37,38,40,41], while Si3 N4 matrix is usually less affected by the corrosion. The addition of a secondary phase, such as TiN, influences the overall corrosion of the material due to two concurrent features: (i) TiN has individual and specific corrosion characteristic, (ii) it changes the composition of the grain boundary phase [2]. In this study, the corrosion behaviour of an electroconductive hot pressed Si3 N4 –35 vol% TiN composite containing Al2 O3 and Y2 O3 as sintering aids was investigated in 1.8 M sulphuric acid aqueous solution at room temperature, 40 and 70 °C up to 400 h. The corrosion kinetics were determined and the mechanism governing the corrosion was defined. The degradation of the flexural strength and toughness and the decrease of the electrical conductivity were also evaluated in relationship with the microstructural changes which are the consequence of the corrosion.
2. Experimental procedure A fully dense hot pressed Si3 N4 –35 vol% TiN composite containing 1.05 wt.% Al2 O3 and 2.63 wt.% Y2 O3 as sintering aids, was employed for corrosion tests. Processing procedures are described elsewhere [47]. Rectangular plates 10.0 10.0 1.0 mm3 , whose wider surfaces were polished up to 6 lm, were used for corrosion tests. The specimens were cleaned in ultrasonicated acetone bath, dried, weighed (accuracy 0.01 mg) and then sealed in a polyethylene tube containing 50 ml of 1.8 M sulphuric acid solution, previously thermostated at the desired test temperature. The corrosion tests were performed at room temperature (RT, 20 °C), 40 and 70 °C for holding time up to 400 h. After the planned exposure time, the specimens were removed from the tubes, rinsed in boiling deionizated water, dried and weighed. Then the weight loss DW =S (mg/cm2 ) was calculated. A Varian Liberty 200 inductively coupled plasma atomic emission spectrometer (ICP-AES) was used for the determination of the Ti, Si, Y and Al cations released into the corrosive solutions. The inert ‘‘V’’ groove nebulizer was chosen because of the corrosive matrix. The following analytical wavelengths were chosen (nm): Si 251.611, Al 396.152, Y 371.030, Ti 334.941. These wavelengths showed a good linear
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response between concentration and intensity, and were free from interferences in the range of concentrations used in the present study. The viewing height was optimized on SBR (signal background ratio). In order to obtain the best linearity in the calibration curve, matrix-matched multielemental standard solutions were prepared in the same concentration range of the solutions under analysis. Microstructural characteristics of the corroded specimens were analyzed through X-ray diffraction and scanning electron microscopy (SEM). In particular the exposed surfaces, the polished cross sections and the fracture surfaces were examined. Mechanical strength, fracture toughness and electrical resistivity were measured on specimens cut in the size and shape of bars 2 2.5 25 mm3 with the faces ground to a longitudinal surfaces roughness of 0.17 lm and the edges chamfered. The tests were performed on the starting material and after a permanence of 400 h in the corrosive solution at RT, 40 and 70 °C. Flexural strength ðrÞ was measured with an Instron mod 1195, in 4-pt bending with 20 and 10 mm as outer span and inner span, respectively, using a crosshead speed of 0.5 mm/min. Fracture toughness ðKIc Þ was evaluated using the chevron-notched beam (CNB) in flexure. The test bars were notched with a 0.08 mm diamond saw; a0 and a1 were about 0.12 and 0.80, respectively. The flexural tests were performed on a semiarticulated alumina four-pt jig with a lower span of 20 mm and an upper span of 10 mm on an universal screw-type testing machine Instron mod.1195. The specimens were deformed with a crosshead speed of 0.05 mm/min. The ‘‘slice model’’ equation of Munz et al. [48] was used for the calculation of KIc . The electrical resistivity ðqÞ was measured at room temperature on the starting material and RT corroded bars by means of a four probe DC method, and on the bars corroded at 40 °C using a two probe DC method, inducing in both cases a longitudinal current along the bars. The current and the voltage readings were made at the same time in two different digital high-resolution multimeters. The resistivity values were determined from the electrical resistance measurements taking into account of the test leads distance and cross section area of the samples. On the bars corroded at 70 °C, the above described technique did not allow to detect the high resistivity, due to the fact that the modifications of the specimen surface are not homogeneous, because of the presence of a thick superficial corroded layer. Therefore, volume resistivity measurements on samples treated at 70 °C were performed with a guarded resistivity cell in conjunction with a high resistance meter. In order to assure a good and uniform electrical contact with the two larger sample surfaces where the voltage was applied, the specimens were placed in the test chamber with two conductive rubber electrodes between the sample and the metal electrodes of the cell. The measurements were carried out by applying a voltage of 10 V and waiting 2 min before reading the electrical resistance value. The resistivity values were finally determined from the resistance measurements taking into account the surface area of the main electrodes and thickness of the samples.
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3. Result and discussion 3.1. Corrosion kinetics and mechanisms The microstructure of the as-hot pressed material consists of a fine-grained matrix of b-Si3 N4 grains surrounded by a glassy grain boundary phase and of well dispersed secondary phase (TiN) particles (Fig. 1). The grain boundary phase is composed mainly of amorphous phases in the systems Al–Y–Si–O–N and Y–Ti–O. The main part of alumina introduced as sintering aids enters in solid solution in b-Si3 N4 grains during the re-precipitation phenomena that occur in the liquid phase sintering process [49]. Moreover, Y2 TiO5 was identified among the re-crystallized grain boundary phases after annealing tests [47], indicating that during sintering of the mixture Si3 N4 –TiN, TiO2 present on the surfaces of TiN particles interacts with Y2 O3 leading to the formation of yttrium titanate [2]; this phase is therefore present in the amorphous state in the as hot pressed material. XRD patterns detected on the surfaces of the samples treated at 70 °C show a progressive decrease of the ratio R (1) between the heights of the TiN (200) and the Si3 N4 (101) peaks: R ¼ hTiN ð200Þ =hSi3 N4 ð101Þ
ð1Þ
from R ¼ 3:5 of the starting material to R ¼ 2:9 and 2.3 respectively after 100 and 400 h at 70 °C. This means that TiN amount decreases in the corroded samples compared to the starting material. On the contrary, the phase b-Si3 N4 does not undergo detectable variation. The microstructure of the corroded samples shows that the surface is nearly smooth but pitted, particularly after the tests at 70 °C. No masking layer was noticed after the tests under the various conditions, but rather an etched surface showing that the preferential corrosion sites are the grain boundary phase and the triple junctions, the TiN grains boundaries and small areas at the interface between the
Fig. 1. Secondary electron image of a polished and plasma etched surface of the hot pressed Si3 N4 –35 vol% TiN composite, showing the dispersion of TiN particles in the Si3 N4 matrix material.
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Fig. 2. Back scattered image of the surface corroded at 40 °C for 20 h: (1) grain boundary phase-depleted areas; (2) depleted TiN grain boundary; (3) leached areas between TiN and Si3 N4 .
TiN particles and the Si3 N4 matrix, leading to the formation of pits and leached areas (Fig. 2). The surface morphology is almost the same at room temperature and 40 °C: the surface voids and defects correspond to selective attack of the grain boundary phases (Fig. 3a and b). The leached areas in the samples treated at 70 °C increase with time: besides the grain boundary phases, TiN particles start to be attacked after short exposure time (Fig. 3c). After 400 h at 70 °C the exposed surface is extensively corroded (Fig. 3d): the superficial TiN grains are almost completely dissolved, as also confirmed by X-ray analysis. The weight loss of the samples in function of the corrosion time (Fig. 4) evidence that the corrosion behaviour at each test temperature and in the whole time range can be modelled by a simple linear rate law (2) [50]: DW =S ¼ kt
ð2Þ
The linear kinetics indicate that the rate controlling step is a chemical reaction, i.e. a dissolution of the reacting species at the reaction interface. The linear rate of the interface advance from the exposed surface towards the bulk of the specimen is confirmed by the analysis of the cross sections, where the portion of the material that is affected by corrosion is distinguished by the presence of porosity. The depth of penetration of the chemical attack and measured from the SEM micrographs, progressively increases with temperature and show a linear dependence with time (Fig. 5), with the same tendency exhibited by the weight loss curves. The values of the reaction rates, k (calculated both from the weight loss: mg cm2 h1 and from the corrosion layer thickness: lm h1 ) reported as a function of temperature in Arrhenius plot (Fig. 6) result in an apparent activation energy Ea ¼56 kJ/mol. The microstructures of the corroded layers in Figs. 7 and 8 highlight that the formation of porosity is due to the dissolution of the grain boundary phase and/or of TiN particles at 70 °C.
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Fig. 3. Secondary electron images of surfaces after corrosion at: (a) room temperature for 400 h; (b) 40 °C for 400 h; (c) 70 °C for 100 h; (d) 70 °C for 400 h.
At 20 °C the external layer of the sample, affected by chemical attack, is very thin and not continuous all over the surface up to about 200 h of exposure. A gradient of dissolution is present from the surface towards the bulk after all the tests performed. Moreover, because of the presence of some porosity in the bulk material mainly located around TiN grains, the contact between the reactants and TiN is favoured. It leads to improved partial dissolution of the external part of the TiN particles, this being evident after the corrosion at 70 °C. It could be thought that the boundaries of the TiN particles act as a preferential path for corrosion, which therefore propagates following the arrangement of the TiN network inside the composite material. The graded chemical attack is well evident in Fig. 8, that represent the microstructure modification after 400 h at 70 °C. Although the reaction interface advance is almost regular (the peak-valley distance measured on the front line is 5 lm), near the surfaces an extensive corrosion attack of both grain boundary phase and TiN
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7
k = 0.0150
RT
6
40°C ∆W/S, mg/cm2
5
70°C
4 3
k = 0.0035
2 1
k = 0.0005 0 0
100
200 Time, h
300
400
Fig. 4. Weight loss ðDW =SÞ against time plot, showing a perfect linear trend at each temperature in the whole time range.
300
k = 0.0187
RT 250
40°C dCL, µm
200
70°C
150 100
k = 0.1706
50
k = 0.5790 0
0
100
200
300
400
Time, h
Fig. 5. Thickness of the corrode layer ðdCL Þ against time plot: dCL progressively increases with temperature and show a linear dependence with time as weight loss curves at each temperature in the whole time range.
phase takes place, while, deeper into the corroded layer, mainly grain boundary phase is dissolved. The degree of dissolution vi of each ion was calculated from the amount of Ti, Si, Y and Al released into the corrosive solution, with the following relationship (3): vi ¼ A=B
ð3Þ
where A and B are respectively the weight of the element released into the solution (measured by ICP-AES) and of the same elements present in the untreated material. The degrees of dissolution of the three cations Y, Al, Si, are represented in function of the corrosion time, in separate plots relative to each test temperature (Fig. 9). The
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from weight gain -1
from corroded layer thickness
Ea= 56.3 kJ/mol
-2
-1
[k, µm h ]
-ln k
-3 -4 -5
Ea= 56.2 kJ/mol -6
-2
-1
[k, mg cm h ]
-7 -8 0.0029
0.0032
0.0035
1/T, K-1
Fig. 6. The values of the reaction rates, k (calculated both from the weight loss: mg cm2 h1 and from the corrosion layer thickness: lm h1 ) reported as a function of temperature in Arrhenius plot result in an apparent activation energy Ea ¼56 kJ/mol.
Fig. 7. Back scattered images of the cross sections after corrosion for 400 h at room temperature (a) and 40 °C (b).
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Fig. 8. SEM micrographs of the cross section of the sample corroded at 70 °C for 400 h, showing the gradient of dissolution from the surface towards the bulk.
trend for the release of the these elements is constant: the degree of dissolution of yttrium, vY is always the highest, vAl is more then two times lower then vY , while vSi is the lowest one: it resulted 20 and 40 times lower than vAl and vY , respectively. The degrees of dissolution of the same cations relatively to the corroded volume of the ceramic composite were calculated with the relationship (4): vVi ¼ A=C
ð4Þ
where A and C are, respectively, the weight of the elements released in the solution and present in the corroded volume (estimated from the thickness of the corroded layer, dCL ). The data show that at 70 °C vVY 1, i.e. yttrium is totally leached in the corroded volume, while vVAl 0:46 and vVSi 0:03, i.e. the residual amounts of Al and Si in the corroded volume are respectively about 54% and 97%. The degree of dissolution of Ti relative to the three temperatures is shown in Fig. 10. vTi at 70 °C is significantly higher than that at room temperature and 40 °C. The results reported above indicate that the linear rate of the interface advance is consequent to the dissolution of the grain boundary phases at all the three tested temperatures and to the additional attack also on titanium nitride particles at 70 °C. A possible contribution to the corrosion from the dissolution of silicon nitride grains, particularly at 70 °C, as previously observed [40], has to be excluded due to
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V. Medri et al. / Corrosion Science 45 (2003) 2525–2539 0.04
RT
Y 0.03
Al
χi
Si 0.02
0.01
0 0
100
200
300
400
300
400
300
400
Time, h 0.20
40°C
χi
0.15
0.10
0.05
0 0
100
200
Time, h 0.6
70°C
χi
0.4
0.2
0 0
100
200
Time, h
Fig. 9. Degrees of dissolution vI of Y, Al and Si plotted as the function of time at room temperature, 40 °C and 70 °C.
the fact that only 3% of Si is removed from the corroded material and also because the ratios vY =vSi and vAl =vSi are constant for all the temperature and time of the tests. Therefore, at room temperature and at 40 °C, only the grain boundary phases in the system Al–Y–Si–O–N and Y2 TiO5 are leached. The difference in the values of
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0.04 RT
χ Ti
0.03
40°C 70°C
0.02
0.01
0 00
100 100
200 200
300 300
400 400
Time, h Fig. 10. Degree of dissolution vTi reported as the function of time at the three temperatures (RT, 40 and 70 °C).
Y and Al cations released in the solution is due to the fact that part of Al is within the b–Si3 N4 network [49], therefore only about 46% of Al introduced as alumina, as sintering aid, contributes to the formation of the grain boundary phases. The values of the weight of the cations released during the chemical dissolution of the grain boundary phase and measured in the corrosive solution by ICP-AES were compared to those of the weight gain for attack. There is an amount of light elements (N and O) solved from the grain boundary phase that cannot be measured, but only indirectly evaluated. The calculations, relative to the data from the tests at RT and at 40 °C, were based on the hypothesis that the Ti cations were released from the Y2 TiO5 compounds and that another phase contains: the residual Y, the measured Al and Si and an amount of N and O. The results show that the approximation of the ICP data and of the weight loss data is within 2% if a grain boundary phase with composition AlSi3 Y2 O4:5 N4 is present, in addition to Y2 TiO5 . The same evaluations are not possible on the data from the tests at 70 °C, because also the corrosion of TiN grains is massively involved. As above indicated, the corrosion of this Si3 N4 –TiN composite is controlled by the surface chemical reaction in H2 SO4 aqueous solution. The leachability of the grain boundary phases from acid aqueous environment arises from their own basic nature. In particular Y cation behave as a network modifier of the silicon-oxide-nitride grain boundary phase [51] and confer it a basic character: these compounds are known to be easily attacked by such environments [21,25,50]. Therefore, the corrosion of the composite under test in a sulphuric acid aqueous solution is due to the leaching of Y and Al ions, i.e. to the surface exchange between Hþ and exchangeable ions on or near the fresh grain boundary. Proton-induced corrosion of ceramics by sulphuric acid was previously reported [52,53]. Ion/proton exchange was indicated as the mechanisms determining chemical attack for lithium and barium aluminosilicates in H2 SO4 ; in this case ion exchanges induces phase transformations, accompanied by
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the effects related to volume changes [52]. In our case, on the contrary, the glassy grain boundary phase dissolves: no barrier diminishes the effective contact area between reactants and the desorption of Y and Al at the interface is rate limiting. The evident degradation observed after the corrosion at 70 °C follows the dissolution of TiN particles; the consequent enhanced penetration towards the bulk of the reactant species also improves the dissolution of the grain boundary phases. The apparent activation energy calculated in this case is within a range of values reported in literature for the corrosion of similar systems [41] although in these cases the dissolution of the grain boundary phase was followed by a passivation process. Compared to the corrosion behaviour shown in literature of other silicon nitrides under similar testing conditions, the selected Si3 N4 –TiN composite has a rather good corrosion resistance in the sulphuric acid aqueous solutions, at least for relatively low temperatures, where the second phase, i.e. TiN, is not involved in the reaction. It has to be underlined that in our case the weigh loss for corrosion is about 5 times lower than that reported in literature for a silicon nitride sintered with the same sintering aid system. 3.2. Properties degradation The mechanical and electrical properties of Si3 N4 –35 vol% TiN composite, before and after corrosion tests at room temperature, 40 and 70 °C for 400 h, are reported in Table 1. The corrosive attack at room temperature doesnÕt lead to a degradation of properties, as flexural strength, fracture toughness and electrical resistivity are almost the same of the as-sintered material. Increasing the corrosion temperature an evident decrease of strength and of the electrical conductivity is observed (Table 1). The electrical resistivity improves of about 6 and 12 order of magnitude as a consequence of the attack at 40 and 70 °C, respectively. This is due to the porosity and progressively leak of percolation in the TiN electroconductive particle network that arises from dissolution of the grain boundary phase and of TiN particles in the corroded layer. The corrosive attack at 40 and 70 °C lowered the flexural strength of about 10% and 30%, respectively. This behaviour is consequent to the increased superficial and subsurface population of defects due to corrosion. Fig. 11 shows, as an example, a typical critical defect acting as fracture origin, in a test bar corroded at 70 °C, after the strength measurement. Table 1 Room temperature flexural strength, fracture toughness and electrical resistivity of the hot pressed Si3 N4 – 35 vol% TiN composite before and after corrosion tests Corrosion test
rF (MPa)
kIc (MPa m1=2 )
q (X cm)
Starting material RT, 400 h 40 °C, 400 h 70 °C, 400 h
835 116 848 69 759 27 597 24
5.67 0.28 5.61 0.56 Not measurable Not measurable
5.88 104 5–7 104 2–6 102 0.6–5 108
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Fig. 11. Secondary electron images of the fracture surface of a test bar corroded at 70 °C for 400 h: the fracture origin is located in the corroded layer near the interface with the bulk (a, b); in (c, a) particular of the corroded layer near the interface showing the high porosity as a typical critical defect acting as fracture origin.
After corrosion at 40 and 70 °C, it is impossible to evaluate the fracture toughness, due to the fact that the notches are not measurable, because the brittle and porous corroded surface layer of the specimens crumbles during notching and breaking.
4. Conclusion The corrosion in sulphuric acid aqueous solution at RT, 40 and 70 °C of an electroconductive Si3 N4 –35 vol% TiN composite, hot pressed with the addition of Al2 O3 and Y2 O3 as sintering aids, follows linear kinetics: there is linear advance of the reaction interface from the external sample surface towards the bulk and the surface chemical reaction is always rate limiting, with an apparent activation energy of 56 kJ/mol. The degree of dissolution for Y, Al, Si and Ti were measured on the solutions after the tests. The chemical attack at RT and at 40 °C involves only the progressive chemical dissolution of grain boundary phases, in the system Al–Y–Si–Ti–O–N, due to ion/ proton exchange in the glassy structure that induces the leaching of Y and Al ions. The attack of TiN occurs at 70 °C, while Si3 N4 is not affected by the selected corrosive environment. The effect of the corrosion on flexural strength, fracture toughness and electrical resistivity, investigated after a permanence of 400 h at each temperature in the corrosive solution, showed the following: 1. High strength levels are maintained after corrosion at RT and at 40 °C, while still good values (up to about 600 MPa) are detected after corrosion at 70 °C, notwithstanding the severe corrosion.
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2. Fracture toughness does not decrease after corrosion at RT, while the brittleness of the surface corroded layer makes it impossible to evaluate the toughness after corrosion at 40 and 70 °C. 3. No variation is revealed in the electrical resistivity after corrosion at RT; on the contrary, after corrosion at 40 and at 70 °C the materials behave like an insulator. Acknowledgements The work is supported by the European Project Research Training Network HPRN-CT-2000-00044 ‘‘Composite Corrosion’’. The research contract of M. Bracisiewicz is funded by the same Project. The Authors wish to thank their colleagues: S. Guicciardi and C. Melandri for the measurement of mechanical properties and G. Fabbri for the evaluation of the electrical properties.
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