sliding wear damage of 304L stainless steel at room temperature: metallurgical and micromechanical investigations

Wear 249 (2001) 37–49

Effect of test duration on impact/sliding wear damage of 304L stainless steel at room temperature: metallurgical and micromechanical investigations A. Van Herpen a , B. Reynier a,∗ , C. Phalippou b,1 a

b

ENSTA/LME/SPM, Chemin de la Hunière, 91761 Palaiseau Cedex, France CEA-DRN/DMT/SEMT/DYN C.E.N., Saclay Bât 607, 91191 Gif Sur Yvette Cedex, France

Received 13 May 1999; received in revised form 3 January 2001; accepted 14 January 2001

Abstract The control-rod cluster assemblies (RCCAs) can be damaged by impact-sliding wear due to flow-induced vibrations which generate contacts with their guidance devices (RCC guide tube). Impact/sliding methodological wear tests have been performed at room temperature on stainless steel claddings (304L). Only the duration variable has been selected to evaluate the wear effect on the material (other experimental conditions have been fixed for all tests). Some non-destructive examinations have been performed on the worn specimens, using weighing, scanning electron microscopy and 2D profilometry. The results show clearly a sensitive damage of the two contacting bodies. X-ray diffraction measurements have been made in order to follow the evolution of the initial microstructure and micromechanical state of the 304L stainless steel (strain-hardening, residual stresses and phase transformation induced by plasticity). The use of these techniques show that test duration has no effect on the behavior of material even if wear damage continues to progress. So these data brought us to the fore that the main wear mechanism at room temperature is an oxidation of the surface layers followed by an oxides detachment stage due to the impact-sliding motion. © 2001 Elsevier Science B.V. All rights reserved. Keywords: 304L stainless steel; Test duration; Damage; Impact/sliding; X-ray diffraction

1. Introduction In nuclear power plants (PWR), flow-induced vibrations generate wear which affects loosely support tubular structures such as control-rod cluster assemblies (RCCAs) and may lead to costly shutdowns with potentially serious consequences. The turbulence due to the flow outside the claddings gives rise to vibratory excitation which brings them into contact with guiding devices at intervals. The standard material selected for tubes and guidance is a 304L austenitic stainless steel. Many experiments [1–4] on these components have focused on the mechanical aspect of the damage of this material using different types of loading and motion (contact geometry, vibration amplitude or frequency and environmental parameters). The main objective of the ∗ Corresponding author. Tel.: +33-1-69-31-97-46; fax: +33-1-69-31-99-97. E-mail addresses: [email protected] (A. Van Herpen), [email protected] (B. Reynier), [email protected] (C. Phalippou). 1 Tel.: +33-1-69-08-63-94; fax: +33-1-69-08-76-19.

present study was not to complete the inventory of this works with new test series but to adopt a different approach to the problem emphasizing the role of the microstructure and mechanical properties of the material in its wear behavior. Moreover, it is interesting to relate the wear test results to the parameters used in tribology. Our experimental approach fixes physical parameters of wear such as contacting bodies, environment and dynamics, in varying test duration. A high value of the vibration generator excitation has been deliberately selected in order to create a “significant” damage within a reasonable time duration. Firstly, the initial metallurgical state of the material has to be well known before studying the influence of the wear tests on the stainless steel. Secondly, structural characterizations on wear scars will allow us to observe the changes of the microstructure generated by the impact/sliding motion during the test for different durations. Several techniques will be used: these include a descriptive analysis based on detailed examinations of worn surfaces (using scanning electron microscopy and 2D profilometry) and X-ray diffraction (XRD) analysis of the microstructure and the micromechanical state of the surface layers of the material.

0043-1648/01/$ – see front matter © 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 4 3 - 1 6 4 8 ( 0 1 ) 0 0 5 2 1 - X

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A. Van Herpen et al. / Wear 249 (2001) 37–49 Table 1 Normalized chemical composition of AISI 304L (average values) C

Mn

Si

P

S

Cr

Ni

Fe

≤0.03% ≤2.0% ≤1.0% ≤0.04% ≤0.03% 17–19% 9–11% Bal.

Fig. 1. Wear specimen (ring and tube), both in 304L stainless steel.

2. Experimental technique of wear tests 2.1. Wear specimen A wear test can be characterized by the type of loading and motion and by the nature of the wear pieces or first bodies (chemical composition and geometry). The nature of the wear problem itself has led us to consider these test variables. The wear couple consists of a tube specimen which is inserted into a thick annular ring (see Fig. 1). The contact geometry is thus conformal. In our study, the claddings (tube specimen) come from real RCCAs of 900 MW PWR (diameter of 9.7 mm, thickness of 0.5 mm) and are manufactured from austenitic stainless steel type AISI 304L (NF EN 10088-1 X02CrNi18-10). On the other hand, the counteracting parts (annular rings) are specially machined by the Commissariat à l’Énergie Atomique (CEA) from the same steel. All tests are performed using a radial clearance of 0.5 mm. A detailed composition of the 304L stainless steel is shown in Table 1. Some metallographic examinations have been realized on both tubes and rings in order to estimate the average grain size. Fig. 2 shows the microstructure of longitudinal section on selected wear specimens. The mean grain size of the tube’s

material is about 20–25 ␮m. Those of the rings is a little bit less (15–20 ␮m). The initial roughness has also been measured using an HOMMEL-WERKE T8D stylus profilometer in the axial direction of the tubular and annular wear specimens. The results are shown in Fig. 3. The ring exhibits a typical turning profile with an average surface roughness, Ra , equal to 1.6 ␮m. Peaks and valleys can be related to the depth of cut and to the speed rate selected for machining the pieces. On the other hand, the cladding exhibits a smaller roughness value (R a ∼ = 0.6 ␮m) with a lot of small valleys which are micro-cracks induced by the manufacturing process. 2.2. Type of movement Flow-induced movements between tubes and their system of guidance are assumed to be combinations of impact and sliding motions. We only consider typical dynamics in simplified environment conditions (in air and at room temperature). The wear test machine and selected conditions lead then to the path seen in Fig. 4. One observes two zones where the normal impact constitutes the major process according to the direction of excitation (big arrow). These two zones are connected by two areas where shock tangential component dominates. 2.3. Experimental device Fig. 5 shows the wear test machine, named CANDUSE, designed by Atomic Energy of Canada Limited (AECL).

Fig. 2. Microstructure of longitudinal section on tube (a) and ring (b) wear specimens.

A. Van Herpen et al. / Wear 249 (2001) 37–49

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Fig. 3. Initial roughness in the axial direction of the tubular (a) and annular (b) wear specimens.

It has been firstly developed for impact/sliding-wear tests at ambient conditions [5,6] and is used at CEA for studying the relationships between wear work-rates and vibratory motions. The test rig consists of a tube being excited with a vibration generator at the top, and a series of three clearance platforms, screwed to a four leg support structure. The legs are strongly attached to upper and lower plates and a containment tank may be installed for tests with water. Here, though, tests were in air. The tube wear specimen is attached to the excitation tube. The loose support wear specimen is incorporated into a force transducer assembly mounted on the instrument platform. Impact forces are measured using four miniature piezo-electric force transducers, located at 90◦ intervals around the circumference of the annular support specimen.

Relative motions are measured using a pair of eddy-current displacement probes mounted 90◦ to each other on the upper platform. The vibration generator is made of two stepping motors with eccentric masses m1 and m2 . A control unit makes them rotate at a common constant frequency, but in opposite directions. Various motions of the excitation tube are obtained by changing excitation parameters at the vibration generator and/or the relative tube-to-support position. Only the duration variable has been selected to evaluate the wear effect on the material. All the tests have been carried out in the same wear conditions. Severe conditions of excitation have been selected according to previous CEA experiments [7], in order to generate a significant wear within a reasonable test duration. Wear specimens (annular rings and tubes) were carefully cleaned at the beginning and end of each test (5 min in alcohol in ultrasonic cleaner). Accurate pre- and post-test weighing of all specimens is required to assess impact/ fretting-wear damage by specimen weight loss. The weight losses are evaluated using several weighings, the final uncertainty is about 2.10−5 g. The information provided by data acquisition system allows to characterize the dynamics of the test. The acquired data permit to calculate characteristic variables: mean work-rate, mean value of normal force, maximum value of normal force and contact duration. These data are used to check the stability of the excitation parameters during the tests. 2.4. Wear test procedure

Fig. 4. Record of the tube-to-support motion during wear test (50 ms; 0.1 mm/div).

Wear tests have been done on 10 couples with six different durations. Four of them have been doubled (26, 50, 100 and 196 h). Two other durations have been added for a better

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A. Van Herpen et al. / Wear 249 (2001) 37–49

Fig. 5. CANDUSE wear test rig.

comprehension of the wear mechanisms (1 and 500 h). The results are summarized in Table 2. The tangential (Ft ) and normal forces (Fn ) are reproducible and lead to a ratio µ (friction coefficient) in accordance with the typical expected value for 304L/304L in air. Wear work-rates (Wr ) are also stabilized around 80±10 mW for all tests. The good behavior of the wear simulator implies that, in this case, test duration is the only relevant variable of the test. Mass losses (1mtube and 1mring ) are easily measurable for any duration (the 1 h test excepted) and correspond to significant wear of the test pieces.

3. Morphological aspects of wear tests 3.1. Macroscopic examination After the tests, we observed some dark brown marks on the circumference of tubes and on the internal ring surface (see Fig. 1). Some oxides debris, generated by wear, fell under the specimen couple but they were not collected for further analysis. A variation in the height of the contact area is also observed: its size measured along the axial direction increases with the test duration.

Table 2 Tangential and normal forces, sliding coefficient, mass losses and wear coefficients in function of test durationa,b Duration (h)

Fn (N)

Ft (N)

µ = |Ft |/Fn

Wr (mW)

1mtube (mg)

Ktube (Pa−1 )

1mring (mg)

Kring (Pa−1 )

26 26 50 50 100 100 196 196 500

11.60 11.80 11.70 11.70 9.99 12.40 10.40 10.60 11.50

4.6 4.3 4.0 4.0 3.8 3.8 3.4 4.2 4.4

0.4 0.36 0.34 0.34 0.38 0.34 0.33 0.4 0.38

94.40 84.80 81.10 76.80 72.70 86.40 58.40 79.10 81.90

1.17 1.01 1.02 0.38 0.98 2.40 2.29 2.99 4.44

1.66E−14 1.59E−14 8.90E−15 3.40E−15 4.70E−15 9.70E−15 6.40E−15 6.50E−15 6.50E−15

1.27 0.78 1.1 0.49 1.19 2.50 2.41 3.51 6.41

1.80E−14 1.23E−14 9.50E−15 4.40E−15 5.70E−15 9.71E−15 6.80E−15 7.70E−15 5.60E−15

Arithmetic mean

11.30

4.1

0.36

79.51

8.73E−15

8.86E−15

Fn (N) is the normal applied force, Ft (N) the tangential component of the applied force, µ = |Ft |/Fn the friction coefficient, Wr (mW) the wear work-rate, 1mtube (mg) the mass loss of the tube specimen, 1mring (mg) the mass loss of the ring specimen, Ktube (Pa−1 ) the specific wear coefficient of the tube specimen and Kring (Pa−1 ) is the specific wear coefficient of the ring specimen. b Specific wear coefficients were calculated using the classical Archard equation: K = (1mρ)/(W /1t). r a

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Fig. 6. Wear track profiles plotted for different durations.

3.2. Analysis by 2D profilometry We have estimated the contact width of the tube using the same profilometry technique as in Section 2.1. Data have been numerically filtered by arithmetic mean in order to extract the shape of wear tracks. All unidimensional profiles have been plotted for the same position of the tube (Fig. 6): namely the areas where maximum number of impacts occurred have been scanned. These curves attest to a real wear process increasing with test duration.

3.3. Scanning electron microscopy examination of wear scars The worn areas of tubes have been examined in details (Figs. 7–11) by scanning electron microscopy (JEOL JSM 840A). The worn area appears in light gray in the secondary electron images particularly at low magnification. Its shape is the same that optically observed. The upper limit of the worn area (Fig. 7a) is clearly distinct from the unworn zone. Because of the random nature of the contact which notably

Fig. 7. Typical view of wear scar by SEM. Detail of upper and lower limit of worn area.

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A. Van Herpen et al. / Wear 249 (2001) 37–49

Fig. 8. Different appearances of the oxide layer in the impact (a) or sliding (b) area.

Fig. 9. Oxide layer fragmented by wear process, showing large cracks.

depends on impact force, the lower limit exhibits a more progressive transition between the worn and unworn areas (see Fig. 7b). The brown color of the wear track can be related to the existence of an oxides layer. Nevertheless, the general appearance of this layer is variable according to observed zones and according to the type of tube-to-support contact. Two diametrically opposite wear zones have been observed. They correspond to the part of the tube path where the impact motion is dominant. These zones are characterized by flat-hammered oxidized areas resulting from repeated impacts and a severe damage

Fig. 10. Slip bands close to the upper limit of the worn area.

of the contact surface (Fig. 8a). The intermediate parts of the worn area exhibit a smoother oxidized surface which can be related to predominance of the sliding due to the tangential force component (Fig. 8b). These oxidized debris are trapped in the contact zone and form a variable thickness layer that is made up of conglomerates of very fine particles.

Fig. 11. Secondary electron image showing worn and unworn areas alternately near the bottom of the wear scar.

A. Van Herpen et al. / Wear 249 (2001) 37–49

These particles have been observed in great amount before the cleaning procedure that precedes the weighing of the wear samples. This oxide layer is then subjected to further wear cycles and fragmented. This locally entails the detachment by delamination and the removal of some big wear clusters. The large cracks of the oxidized layer (Fig. 9) and the fall of the debris during all the test duration accredit this scenario. Some slip bands (strain-hardened grains) located near the upper limit of the worn zone denote that plastic deformation of the cladding has occurred during the wear process (Fig. 10). Near the bottom of the wear scar (left side of Fig. 7), on the less worn area, we can see horizontal lines. At higher magnification (Fig. 11), these marks are corresponding to worn areas alternating with unworn areas due to the initial roughness of the ring. As we have noticed in Fig. 3, there is about 100 ␮m between two peaks, there is also the same distance between two worn area tracks.

4. X-ray diffraction analysis In addition to the previous morphological characteristics of the worn area, we studied the influence of the wear test on the microstructure and the micromechanical state of the surface layer in our stainless steel specimens. The relative proportion of the two phases and the residual stresses were mainly evaluated by XRD [8,9]. Details of residual stresses analysis are given in Appendix A. The effective X-ray penetration depth was about 5 ␮m for 50% of the total integrated intensity [10]. To reach a given depth inside the material, the surface layer was removed by electro-polishing which suppress residual stresses induced by mechanical polishing. No

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correction of the acquired data in order to account for the redistribution of the stresses during the polishing has been done due to the cylindrical geometry of the samples and to the local plasticity introduced by this kind of contact [11]. The XRD, as well as allowing estimation of the residual stresses, may also be used to evaluate the plastic strains [9,10,12]. The width of diffraction peak (often measured at 50% of the integrated peak and named full width at half maximum (FWHM)) depends mainly on the microstructure of the material, which is altered by the strain-hardening. It varies in the surface layers affected by wear and may be used to estimate the plastically deformed depth. 4.1. Initial metallurgical state This initial state has to be well known in order to understand the influence of the wear tests. Firstly, optical microscopy shows elongated nonmetallic inclusion in longitudinal cross-section (Fig. 2). At higher magnification, we can see a strain-hardened microstructure containing mechanical twins on cladding but not on rings (Fig. 12). Vickers micro-hardness of the claddings (Hv200 g = 255) and largest width of diffraction peak confirm this hardened state. The micromechanical state of subsurface layers has been investigated at different depths (using successive electrolytic polishing). Residual stresses (axial (σ A ) and hoop (σ H ) directions) profiles have been plotted. Though compressive residual stresses have been measured in the first layer of tubular sample, there is a tensile adaptation in the deep layer. Heterogeneous strain-hardening profiles has been shown by following the width of diffraction peak (Fig. 13b). Finally, no ␣0 -martensite has been detected in a sample before a wear test.

Fig. 12. Microstructures of transverse sections on cladding (a) and ring (b) wear specimen.

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Fig. 13. Initial state of tubular sample: in-depth profiles of residual stresses (a) and FWHM (b).

4.2. Non-destructive analysis on cladding worn specimen After each wear test, a non-destructive XRD analysis has been made on different parts of all oxidized areas. XRD spectra on worn areas are different from those measured on unworn cladding: a new body centered cubic phase with a unit cell size of 2.89 Å has been clearly identified (Fig. 14). This phase has been assumed to be a ␣0 -martensite induced from austenite by plastic deformation. This phenomenon is well known in such a stainless steel [13] and has been noticed in other wear experiments by several authors [14–17]. Predominant impact zones (PIZ) and predominant sliding zones (PSZ) have been separately analyzed showing different proportion of ␣0 -martensite. The PIZ (Fig. 15) exhibits up to 50% of ␣0 -martensite when a maximum amount of 15% have been found in the PSZ. In each case, the sampled depth is about 5 ␮m. The two previous zones have been analyzed on all tubular wear specimen. The test duration does not affect the relative proportion of each phase. Because XRD technique allows to reach the mechanical local state of one particular phase, hoop (σ H ) and axial (σ A ) residual stresses were determined in each phase for the same areas (PIZ and PSZ) as previously (Fig. 16). In spite of a relatively low confidence level of the experimental data, the four curves accredit the trends presented above. The mean width of the diffraction peaks used for the residual stress evaluation in austenite has been simultaneously measured, in order to estimate the strain-hardening of the material in the axial and hoop directions. An increase of this parameter (FWHM) shows clearly that surface layers have been hardened by the wear process. This result has been confirmed by Vickers micro-hardness. Strain-hardening seems to be quasi-isotropic because the evolution of FWHM is the same in the two directions. If this phenomenon and the martensite phase transformation of the initial austenite jointly occur, the equilibrium of the internal stresses is upset at the same time. Thus, surface axial residual stresses, initially slightly compressive (σ¯ A ≈ −70 MPa), become tensile up to 220 MPa.

On the other hand, hoop stresses are now compressive for all test durations, with a mean value of about −130 MPa. All the residual stresses evaluated on the ␣0 -martensite phase are compressive with a much higher value in the hoop direction (σ¯ H ≈ −630 MPa and σ¯ A ≈ −200 MPa). Therefore, attention has to be paid that results are strongly affected by all the experimental conditions (geometry of the wear specimen and the X-ray beam, rolling texture, grain size, etc.). Consequently, there is no significant evolution of the residual stresses with increasing test duration even if some smoothed curves are not perfectly horizontal. Residual stress measurements have been made mainly in the PIZ because the relative proportion of ␣0 -martensite is large enough to allow us to evaluate these parameters in the two phase at the same time. Nevertheless, the same measurements have been performed in the PSZ of one specimen tested during 100 h (Table 3). Values of same magnitude have been found even if the relative proportions of the two phases are very different. 4.3. Microstructural evolution of the subsurface layers Subsurface layers have been analyzed using the same XRD technique. To reach a given depth, successive layers of material have been removed by electro-polishing (see Section 4). FWHM, Vickers micro-hardness and residual stresses have been simultaneously estimated. In fact, no reliable residual stress values are given in this study because they vary very quickly and not linearly with depth of removal. This may be due to the wear process and type of contact which lead to local plasticity. Therefore, acquired data are difficult to be corrected in order to account for the redistribution of the stresses during the polishing, particularly for this tubular geometry except by using FEM calculations. Nevertheless, interesting results such as depth of strain-hardening and thickness of the ␣0 –␥ microstructure have been obtained (see Table 4). Superficial layers seems to be affected by austenitic transformation; the proportion of ␣0 -martensite decreases to zero through an external layer of about 15 ␮m. Despite absence

A. Van Herpen et al. / Wear 249 (2001) 37–49

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A. Van Herpen et al. / Wear 249 (2001) 37–49 Table 4 In-depth proportion and FWHM of ␣0 -martensite for 26, 100 and 196 h test duration Depth (␮m)

0 2 5 10 15 a

Fig. 15. Relative fraction of ␣0 -martensite vs. test duration in PSZ and PIZ.

of an initial value, the full width of the diffraction peak of the ␣0 -martensite ({2 1 1} lattice plane) is large and implies certainly a great crystalline defect density. Its value decreases also in the same surface layers (Table 4). Similar results have been obtained on the austenite diffraction peak ({2 2 0} lattice plane). After some more removals we find that the total depth affected by strain-hardening is more than 50 ␮m, and represents about 10% of initial thickness of the tube. This result has been checked by micro-hardness in-depth profile (Fig. 17). The strain-hardening of ␥-phase, estimated by these two methods does not evolve markedly with the test duration.

FWHM (␣0 {2 1 1} lattice plane)

Proportion of ␣0 -martensite (%) 26 h

100 h

196 h

26 h

100 h

196 h

48 42 28 6 0

46 44 8 – 0

34 –a 3 – 0

2.98 2.60 2.44 1.83 –

3.10 2.74 1.64 – –

3.10 – 2.13 – –

No available data.

5.1. Contact surface damage The first observations made during wear tests point out a severe oxidation of the contact area. Wear debris are also continuously expelled from the contact and their aspect — particularly their brown color — indicate that this origin is certainly related to the surface oxidation phenomenon. Generation of wear debris arises at the very beginning of the wear tests (even after only 1 h duration) and are going on along all the test duration. According to previous studies [2,17–19], under dynamic conditions, the oxide layer is continuously rubbed away from the subsurface material. So the metallic surface is always exposed to the oxidation process, promoted by ambient conditions. 5.2. Wear test results

5. Discussion The effect of test duration on the impact/sliding wear of 304L stainless steel has been investigated in air at room temperature.

The test dynamics are very stable for all tests (mean normal force |F¯n |, sliding coefficient µ and specific wear coefficient K) but we have noticed an increment of mass loss (1m) with increasing duration. Therefore, the mass loss is quite the same for the two counteracting bodies. Never-

Fig. 16. Surface residual stresses measured along axial (σ A ) and hoop (σ H ) directions, in austenite (␥) and martensite (␣0 ) phases, on PIZ wear scars.

Table 3 Surface residual stresses, mean values of 10 PIZ measurements and PSZ data for one 100 h sample PIZ PSZ

σ¯ A␣0 ≈ −200 MPa σA␣0 ≈ −240 MPa

σ¯ H␣0 ≈ −630 MPa σH␣0 ≈ −540 MPa

σ¯ A␥ ≈ 220 MPa σA␥ ≈ 120 MPa

σ¯ H␥ ≈ −130 MPa σH␥ ≈ −130 MPa

A. Van Herpen et al. / Wear 249 (2001) 37–49

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Fig. 17. Effect of test duration on FWHM and Vickers micro-hardness profiles.

theless, macroscopic view and profiles examinations of worn area show clearly a non-linear evolution of shape (Fig. 6). The track width has been plotted in function of test duration (Fig. 18a). This curve shows two different slopes. Firstly, there is a quick evolution for the shorter tests (<100 h). Secondly, the track width increases more slowly for test duration in a range of 100–500 h. The same two-step behavior has also been found on Fig. 18b where rate of mass loss has been plotted in function of duration. After 100 h, curves tend to become asymptotic. The wear process is then quasi-stabilized. A linear evolution of the mass loss can be expected for longer tests. From a mechanical point of view, the first stage of the wear process (running-in) is different. Actually, though the nominal impact force remains constant during the test, the surface of the contact increases continuously so contact stresses decrease at the same time. 5.3. Structural evolution A real evolution of the cladding microstructure has been observed (martensite formation, strain-hardening of ␣0 - and

␥-phases and upset of the residual stresses) but all the metallurgical analysis show that the test duration has very little influence on these parameters. These results attest that the real level of stresses in the contact zone remains above the initial yield strength of the material even if mean contact stresses are decreasing. Because the selected dynamics is impact/sliding motion (characterized by a contact duration of 70%), the maximal values of normal force recorded for all tests (127–143 N) are much higher than the mean value (F¯n ∼ = 11 N) calculated for the total duration. The high level of these instantaneous normal forces can be likely related to the microstructure modifications. These forces have been measured in the zone where impacts prevail (PIZ) but the tangential sliding is not negligible. These impact/sliding led to the most severe wear. Their effects are two-fold on this stainless steel (unstable austenite). Firstly, the normal component of impact forces induces high Hertz pressures which harden the surface layer (up to 50 ␮m) as we saw on FWHM and micro-hardness profiles (Fig. 17). Secondly, under combination of tangential sliding and normal impacting, shear stresses becomes very large and generate transformation of primary ␥-phase. This strain-induced martensite exhibits a

Fig. 18. Track width (a), tube specimen mass loss and wear kinetic (b) vs. test duration.

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A. Van Herpen et al. / Wear 249 (2001) 37–49

large volumic fraction (about 50%). On the contrary, the zones where sliding motion is predominant (PSZ) are less deformed and only up to 15% of ␣0 has been found. The initial level of residual stresses in austenite (hoop and axial directions) has been increased while strong compressive stresses have been generated in martensite. At least, wear process has generated a new stressed two-phase material under the contact area. This new heterogeneous microstructure can probably influence oxidation kinetics by increasing the oxygen diffusion inside the first layers of the cladding and by favoring the stress corrosion phenomenon.

Table 5 X-ray elastic constants of the austenite and martensite phases Phase

Plane {h k l}

(1/2)S2 (10−6 MPa−1 )

S1 (10−6 MPa−1 )

␥-Austenite ␣0 -Martensite

{2 2 0} {2 1 1}

6.05 5.76

−1.56 −1.25

obviously expected to be quite different, but the methodology developed here will be adapted to future works.

Appendix A 5.4. Wear mechanism Thus, the most likely wear mechanism of the claddings seems to be oxidation of the material. This process can be divided into two stages. Firstly, the surface protective film of the stainless steel (a thin layer of Cr2 O3 ) is rubbed away and an oxide layer begins to form on freshly exposed metallic surface. Secondly, wear kinetics (characterized by a periodic separation of the first bodies) can help breaking up the wear-oxide film as it reaches critical thickness. Finally, the wear cycle might be simply described as a continuous sequence of formation, detachment and ejection of oxides at the surface of the specimen. Our assumption might have been confirmed by qualitative analysis of the debris ejected during wear process. However, this latter analysis was not performed because only small particles were produced by each test and were difficult to collect within this kind of apparatus.

6. Conclusions Impact/sliding wear tests have been performed on controlrod claddings (austenitic stainless steel 304L). Morphological and microstructural analyses have been realized in order to study the influence of test duration on material damage at room temperature. Severe damage to contact zones have been evaluated by profilometry and by weighing after wear tests. Significant changes of the material surface micromechanical state (martensite formation, strain-hardening and residual stresses) have been measured but they cannot directly be related to the test duration. All these results show clearly that mechanisms of impact/sliding wear are specific and do not seem to be a simple addition of sliding laws plus impact laws. So impact/sliding tests are a necessary matter to understand in a more realistic way, wear of power plant components when wear simulators, operating in experimental conditions closed to the PWR primary system, are used. Concerning the influence of material microstructure on the wear mechanism, a complementary study would be necessary using other relevant test parameters (fixed duration with variable conditions of excitation). The temperature and real environmental PWR effects are

The residual stresses were measured in the Structure and Material Properties Research Group at École Nationale de Techniques Avancées (ENSTA/LME/SPM). The sample is set onto a four-axis (φ, χ, Ω, 2Θ) goniometer (SEIFERT MZ VI type) with a 360 mm radius, in ψ-assembly (i.e. the psi, ψ or tilt axis was χ, not Ω) [8,9], and connected to a micro-computer. Tubular samples were mounted in a cradle type sample holder and were strongly hold in place by two little screws. Sample alignment was performed using a dial gage probe (with an accuracy of ±10 ␮m). Goniometer alignment was ensured examining a plate containing annealed Fe powder in epoxy using ψ tilting in a range of −45◦ ≤ ψ ≤ +45◦ . The main experimental parameters were: {2 2 0} lattice planes of the ␥-austenite phase (Bragg angle 2θ ≈ 128.5◦ ) and {2 1 1} lattice planes of the ␣0 -martensite phase (2θ ≈ 154.5◦ ) using filtered Cr K␣ radiation (wavelength λ = 2.897 Å) at an exciting potential of 40 kV and a current of 35 mA (1.4 kW). A double pinhole collimator with a 2 mm opening was used with 0.5 mm receiving slit. Small X-ray beam and low divergence combined with masking to leave only a part of the surface exposed to the X-rays minimized the experimental errors due to the uncertainty regarding the true θ angle on a curved specimen. Stress measurements for biaxial analysis were performed at two angles φ (φ = 0 and 90◦ ) and nine angles ψ. An acquisition time of 500 s per peak was used for accurate peak shapes and good counting statistics. The stresses were calculated using generalized linear and elliptic least squares fitting methods of the measured strain distributions, i.e. the plot of d-spacing [8,10], (dψφ − d0 )/d0 versus sin2 ψ along the axial or hoop directions (dψφ and d0 denote the lattice spacing in the [ψ, φ] direction and the stress-free lattice spacing, respectively). The X-ray elastic constants, S1 and (1/2)S2 , for the two phases ␥-austenite and ␣0 -martensite, were derived assuming constant stress and strain in all grains in accordance with the Kröner model [20] and are shown in Table 5. References [1] J. Guinot, Etude bibliographique des travaux expérimentaux menés sur l’usure par impacts-glissements, Influence des principaux paramètres, Electricité de France, Direction des Etudes et Recherches, Report no. HT.22/89-22A, 1989.

A. Van Herpen et al. / Wear 249 (2001) 37–49 [2] P.L. Ko, Wear of power plant components due to the impact and sliding, Appl. Mech. Rev. 50 (7) (1997) 387–411. [3] P.L. Ko, Wear due to Flow-induced Vibration, Technology for the ’90s, ASME Special Pub. # 100347, 1993, Chapter 8, pp. 865–896. [4] F. Axisa, Experimental study of tube/support impact forces in multi-span PWR steam generators tubes, ASME Symp. Flow-induced Vibrations 3 (1984) 139–148. [5] M.K. Weckwerth, N.J. Fischer, Technical manual for fretting-wear test machine for ambient conditions, Atomic Energy of Canada Ltd., Report AECL ET-FRET-31, 1991. [6] C. Phalippou, F. Hareux, Machine AECL pour essais d’usure en air et en eau froide: procédures d’utilisation, Commissariat à l’Energie Atomique Report CEA-DMT no. 91-609, 1991. [7] C. Phalippou, X. Delaune, Predictive analysis of wear work rates in wear test rigs, Pressure Vessel and Piping, ASME Publ., PVP-Vol. 328, Montréal, 21–27 Juillet 1996, pp. 247–256. [8] J. Lu, M. James, G. Roy, SEM Handbook of Measurements of Residual Stresses, Society for Experimental Mechanics, Bethel, 1996, pp. 71–132. [9] I.C. Noyan, J.B. Cohen, Residual Stress-Measurement by Diffraction and Interpretation, Springer, New York, 1987. [10] J.M. Sprauel, Etude par diffraction X des facteurs mécaniques influençant la corrosion sous contrainte, PhD Thesis, University of Paris VI, Jussieu, 1988.

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[11] M.G. Moore, W.P. Evans, Mathematical correction for stress in removed layers in X-ray diffraction residual stress analysis, SAE Trans. 66 (1958) 340–345. [12] S. Taira, X-ray Studies on Mechanical Behaviour of Materials, The Society of Materials Science, Kyoto, 1974. [13] G. Blanc, Deformation mechanisms of austenitic stainless steel, Les aciers inoxydables, Les Editions de Physique, Les Ulis-France, 1990, Chapter 17, pp. 611–628. [14] K. Hsu, T. Ahn, D. Rigney, Friction, wear and microstructure of un-lubricated austenitic stainless steel, Wear 74 (1980) 13–37. [15] C. Allen, A. Ball, B.E. Protheroe, The abrasive-corrosive wear of stainless steels, Wear 74 (1981/1982) 229–305. [16] A.F. Smith, The friction and sliding wear of unlubricated 316 stainless steel at room temperature in air, Wear 96 (1984) 301–318. [17] Z.Y. Yang, G.S. Naylor, D. Rigney, Sliding wear of 304 and 310 stainless steels, Wear 105 (1985) 73–86. [18] T.F.J. Quinn, Review of oxidational wear, Trib. Int. 16 (1983) 306–315. [19] T.F.J. Quinn, Oxidational wear modelling: Part II. The general theory of oxidational wear, Wear 175 (1994) 199–208. [20] E. Kröner, Berechnung der Elastischen Konstanten des Vielkristalls aus den Konstanten des Einkristalls, Zeitschn’ff für Physik 151 (1958) 504–518.