不锈钢与镍基合金氢脆的对比(英)_图文

1、 International Journal of Hydrogen Energy27(2002793 800Comparison of hydrogen embrittlement of stainless steels andnickel-base alloysOmar A.El kebir,Andrzej SzummerFaculty of Material Science and Engineering,Warsaw University of Technology,Wo l oska141,02-507Warsaw,PolandAbstractSeveral stainless st

2、eels and Ni-base alloys were investigated to compare their hydrogen embrittlement susceptibility in annealed,cold-worked,and aged conditions.Cathodic hydrogen charging was used for introducing hydrogen,and reverse bending tests were applied in measuring hydrogen-induced ductility loss.It was found t

3、hat hydrogen embrittlement increased with increasing nickel content,cold-working and thermal aging at500C.Optical metallographic examination and X-ray di raction analysis showed substantial microstructural and structural changes in all the materials under investigation,but the standard and highly al

4、loyed stainless steels in cold-worked and cold-worked and aged cases were more resistant to hydrogen-induced ductility loss than highly corrosion resistant Ni-base alloys.?2002Published by Elsevier Science Ltd on behalf of the International Association for Hydrogen Energy.1.IntroductionHydrogen embr

5、ittlement of metals and alloys has been a concern to a number of industries for many years.It has become especially a problem to oil and gas production where high concentrations of H2S are frequently encoun-tered1.Austenitic stainless steels and Ni-base alloys are employed in large quantities in thi

6、s industry because they possess the required combination of strength and corrosion resistance.Austenitic stainless steels show good resistance to general corrosion in many aggressive aqueous envi-ronments,but can be susceptible to localized corrosion like pitting corrosion or stress corrosion cracki

7、ng.On the other hand Ni-base alloys provide excellent resistance to these forms of localized corrosion.It is well-known that both Ni-base alloys and austenitic stainless steels can be embrittled by hydrogen2in aqueous corrosive envi-ronments,but there is a lack of data about the hydrogen embrittleme

8、nt susceptibility of these materials in direct comparison.Corresponding author.Fax:00-48-22-628-1983.E-mail address:.pl(O.A.El kebir.It was the purpose of this work to determine and to compare the hydrogen embrittlement susceptibility of some Ni-base alloys and highly alloyed stainl

9、ess steels when investigated under the same conditions.2.Materials and experimental proceduresThe materials used in this study were Austenitic stainless steels,Ni-base alloys,and pure nickel as a reference ma-terial.The chemical composition of the alloys used in this study are given in Table1.The ma

10、terials were delivered in the formof sheets of di erent thickness,annealed(A,or cold-worked50%reduction of cross-section.The specimens were well prepared by using silicon car-bide abrasive papers of150,220,500,800mesh,respec-tively,and aluminum oxide forÿne polishing.To reveal the microstructur

11、e,the specimens were etched by Kings water(5ml of HNO3;15ml of HClwith or without glycerin.Cathodic hydrogen charging in a5%H2SO4solution with1mg=l AS2O3for20h at roomtem perature,and 1m g=1AS2O3added as a H-recombination poison un-der a current density of100mA=cm2was used for the0360-3199/02/$22.00

12、?2002Published by Elsevier Science Ltd on behalf of the International Association for Hydrogen Energy. PII:S0360-3199(0100151-3794O.A.El kebir,A.Szummer/International Journal of Hydrogen Energy27(2002793800Table1Chemical composition(in weight percentof the studied steelsElement Austenitic stainless

13、steel Hastelly Pure 316L20MOD20CB-3G C276NiC0.0250.010.060.030.020.08 Mn 1.780.9 2.0 1.37 1.00.2 P0.0180.0140.0350.016S0.0150.0060.0350.0050.005 Si0.590.4 1.000.450.080.2 Cr17.4521.719.021.022.2316.0Ni13.9626.332.535.042.0455.0599.5 Fe63.576946.3733.8740.3719.59 5.00.2 Mo 2.25 4.3 2.03.0 6.2816.0Cu0

14、.15 3.04.0 1.820.1 Co0.13 2.25 2.5Pb0.001Su0.002Cb0.0021 1.92N0.0500.028W0.65 4.0Al0.44B0.001Mg0.03Te0.05Zr0.02 V 0.35Fig.1.Reverse bending numbers for annealed materials before and after hydrogen roduction of hydrogen.Reverse bending tests were applied for well prepared specimens with d

15、imensions (25×4×0:19mm3to measure hydrogen-induced duc-tility loss for annealed,for50%cold-worked and for50% cold-worked and aged at500C specimens,before and after hydrogen charging.The structure of the alloys were examined using optical microscopy and scanning electron microscopy at di er

16、ent magniÿcations.The phases com-position was determined by the X-ray di raction method using,Philips PW1830(Cu Kradiation,and Philips PW Fig.2.Reverse bending numbers for cold-worked materials before and after hydrogen charging.1140(Co Kwith parameters of(U=40KV;I=30mA di ractometers.3.Results

17、 and discussionFig.1shows the variation of the reverse bending numbers (RBNfor the materials tested in annealed condition be-fore and after hydrogen charging.The RBNs decrease with Ni content and are strongly a ected by cathodic hydrogenO.A.El kebir,A.Szummer /International Journal of Hydrogen Energ

18、y 27(2002793800795 Fig.3.Ductility loss for annealed materials after hydrogencharging. Fig.5.Fractographs of the annealed alloys:(a316L (13.96wt%Nibefore hydrogen charging;(b316L (13.96wt%Niafter hydrogen charging;(cHastelloy C-276(55.05wt%Nibefore hydrogen charging;(dHastelloy C-276(55.05wt%Nibefor

19、e hydrogencharging.Fig.4.Ductility loss for cold-worked materials after hydrogen charging.796O.A.El kebir,A.Szummer/International Journal of Hydrogen Energy27(2002793800 F ig.6.X-ray di raction patterns of cold-worked316L:(abefore hydrogen charging;(bafter hydrogen charging for20h at roomtem peratur

20、e. Fig.7.X-ray di raction patterns of cold-worked20CB-3:(abefore hydrogen charging;(bafter hydrogen charging for20h at room temperature.O.A.El kebir,A.Szummer/International Journal of Hydrogen Energy27(2002793800 797Fig.8.X-ray di raction patterns of solution annealed20Cb-3:(abefore hydrogen chargin

21、g;(bafter hydrogen charging for20h at room temperature.charging.Similarly for the50%cold-worked materials,after hydrogenation,the RBN decrease with Ni content,except for pure Ni(Fig.2.But taking into account the Ductility Loss parameter, measured as a ratio of the RBN according to the formulaDuctili

22、ty loss=RBN(RBNhRBN;where RBN is the reverse bending number before hydrogen charging,(RBNh is the reverse bending number after hy-drogen charging,it is seen that for high Ni content alloys this parameter decreases with Ni content and is the lowest for pure Ni(Fig.3.This indicates that for these allo

23、ys in annealed condition hydrogen ingress is relatively less de-structive.On the other hand for the50%cold-worked materials after hydrogen charging the Ductility Loss parameter shows a reverse tendency and increases remarkably with Ni content (Fig.4.It is interesting that the ductility is most a ect

24、ed for the Ni-base alloys Hastelloy G and Hastelloy C-276. The least of all a ected is the316stainless steel containing only about14%Ni.This means that after the mechanical treatments necessary to develop the required strength for employing these high Ni content alloys under high loads, these alloys

25、 will lose resistance in environments containing hydrogen ions.The same tendency was observed after thermo-mechanically treated(cold-worked and aged at 500Cmaterials.The phenomenon of higher susceptibility to hydrogen embrittlement of cold-worked materials,manifested in a lower absolute RBN(Figs.1an

26、d2,can be explained in terms of the density defects that are present in the heavily cold-worked materials.Cold-working produces many vacan-cies as well as a very high dislocation density.It is proba-ble that hydrogen lattice di usion is enhanced by the excess vacancy concentration3and pipe di usion

27、of hydrogen will occur along dislocations4,which makes the materials more a ected by hydrogen.The hydrogen embrittlement of the materials with di er-ent Ni content can also be compared by fracture surface ex-amination.The fracture modes in the reverse bending test of the14wt%Ni alloy(316Land the55wt

28、%Ni alloy(Hastelloy C-276before and after hydrogen charg-ing are shown in Fig.5.In the annealed alloys,fractures occurred predominantly in a ductile mode with small dim-ples(Fig.5a,b,and due to the hydrogen embrittlement the so-called quasi-cleavage fracture was observed only at the edge of the14wt%

29、Ni specimen(Fig.5b,while in the55wt%Ni alloy,brittle intergranular fracture pre-dominated,as shown in Fig.5d.This indicates that the in-crease in Ni content increases the fraction of intergranular fracture.These observations are in agreement with those798 O.A. El kebir, A. Szummer / International Jo

30、urnal of Hydrogen Energy 27 (2002 793 800 previously reported by Hinotani et al. 5, for NiCrFe alloys with di erent Ni content. The results of X-ray di raction studies, conducted on both the uncharged and the charged specimens of cold-worked 316L steel and 20Cb-3 alloy, are shown in Figs. 6 and 7. I

31、mmediately after hydrogen charging, re ections arising from the hydride phase- are present at the small-angle sides of the re ections of the matrix h 6 13. The width of the matrix re ections after hydrogen charging is much broader than that without hydrogen charging and these re ections are also shi

32、fted to the smaller-angle side because of solid solution dilatation by hydrogen atoms. It should be noticed that the broadening of the matrix re ections of the annealed 20Cb-3 alloy because of the 50% cold-working (cf. Fig. 7(a with Fig. 8(a is much less than the broadening of these re ections becau

33、se of hydrogen charging (Fig. 8(b. This indicates that the matrix lattice is much more a ected by hydrogen ingress than by the 50% cold-working. The same phenomena were observed for all the materials tested in this work. Optical and scanning electron examinations have shown that in all hydrogen char

34、ged materials surface microcracks are formed inside the grains and along the grain boundaries or twin boundaries (Figs. 9 11. These surface cracks were never observed immediately after cessation of charging, but they formed gradually during aging at room temperature when desorption of hydrogen takes

35、 place 15. The original microstructures of the annealed 14 wt% Ni, 55 wt% Ni and pure Ni samples before hydrogenation are similar (Figs. 9(a 11(a, but after hydrogen charging the morphology of surface cracks changes signicantly with nickel content. In pure Ni and in 55 wt% Ni alloy C-276 the main su

36、rface cracks were mostly intergranular (Figs. 10(b and 11(b, whereas in the 14 wt% Ni steel 316L cracks lying inside the grains parallel to 1 1 1 planes 5 were more frequently observed in addition to the intergranular ones (Fig. 9(b. The parallel cracks usually were formed by twin interface opening.

37、 This is also in agreement with the observation that hydrogen charging frequently produced microtwins in addition to the increase in dislocation density 5. Figs. 9(c 11(c show the morphology of the hydrogen induced cracks in more detail. It should also be noted that surface crack formation was close

38、ly related to hydride formation. The results of X-ray di raction analysis showed that -hydride formed in all alloy examined 5,14 and the lattice constant of the hydride increases with increasing nickel content. According to the interpretation by Boniszewski and Smith 16 a uniform layer of hydride ph

39、ase during hydrogen charging produces compressive stresses in the surface layer due to volume expansion, which is the case of plastic deformation most distinctly visible in pure Ni samples (Fig. 11(b, (c. During the decomposition of the unstable hydride phase at room temperature, tensile stresses ar

40、e induced on the surfaces. These stresses and the simultaneous desorption of gaseous hydrogen result in grain boundary cracking. The Fig. 9. 316L stainless steel in annealed condition: (a original microstructure (optical microscopy; (b surface cracks after hydrogen charging (optical microscopy; (c s

41、urface parallel cracks inside grains after hydrogen charging (scanning electron microscopy. results of this work are basically in good agreement with this model. It should also be pointed out that grain boundaries di usion of hydrogen is faster than lattice di usion. This suggests that hydride layer

42、s also form preferentially at grain boundaries. Decomposition of the hydrides which occur during aging at room temperature will produce a large local tensile stress at the grain boundaries in a direction at right angle to the grain boundary planes and will open the grain boundaries at the surfaces,

43、as is visible in Figs. 10(b and 11(b. O.A. El kebir, A. Szummer / International Journal of Hydrogen Energy 27 (2002 793 800 799 Fig. 10. Hastelloy C-276 in annealed condition: (a original microstructure (optical microscopy; (b surface cracking after hydrogen charging (optical microscopy; (c surface

44、cracks along grain boundaries and inside the grains after hydrogen charging (scanning electron microscopy. Fig. 11. Pure Ni sample in annealed condition: (a original microstructure (optical microscopy; (b surface cracks after hydrogen charging (optical microscopy; (c surface cracks and deformations

45、after hydrogen charging (scanning electron microscopy. 4. Conclusions 1. The ductility of highly alloyed stainless steels and Ni-base alloys, measured as absolute numbers in reverse bending tests (RBN, decreased with Ni content due to cathodic hydrogen charging, both in annealed and in 50% cold-work

46、ed condition. 2. On the other hand the relative ductility loss parameter, measured as a ratio of the di erence of RBN before and after H-charging to the original RBN, decreased with Ni content for the annealed alloys and was the lowest for pure Ni, and increased for the 50% cold-worked alloys, indicating that after mechanical treatment the higher Ni content alloys lose resistance in environments containing hydrogen ions. 3. The higher Ni content enhances the intergranular hydrogen embrittlement