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Microstructure and electrochemical properties of the bonding zone of AISI 316L steel joined with a Fe-based amorphous foil
Journal of Materials Processing Technology 210 (2010) 1051–1060
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Journal of Materials Processing Technology journal homepage: www.elsevier.com/locate/jmatprotec
Microstructure and electrochemical properties of the bonding zone of AISI 316L steel joined with a Fe-based amorphous foil J.A. Verduzco a , J. González-Sánchez b,∗ , V.H. Verduzco a , J. Solís c , J. Lemus-Ruiz a a b c
Instituto de Investigaciones Metalúrgicas, (UMSNH), P.O. Box 888, Morelia, Michoacán 58000, Mexico Centre for Corrosion Research, Universidad Autónoma de Campeche, Av. Agustín Melgar s/n, Col. Buenvista, Campeche, Cam., C.P. 24039, Mexico SEP-DGEST-ITTLA, Av. Mario Colín, S/N, Tlalnepantla Edo. de Méx., 54070, Mexico
a r t i c l e
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Article history: Received 1 December 2009 Received in revised form 30 January 2010 Accepted 17 February 2010
Keywords: Fe-based amorphous alloys Diffusion bonding Stainless steel Localised corrosion
a b s t r a c t The effect of temperature and holding time on the microstructure and corrosion resistance of the junction zone of AISI 316L stainless steel (SS) bonded to itself with Fe75 Cr8 P10 B7 filler alloy was investigated. The brazing alloy was prepared in the laboratory in the form of amorphous ribbons and its melting temperature was determined by differential thermal analysis (DTA) to be 1571 K. The joining process was carried out in a chamber with controlled Ar atmosphere at two temperatures: 1173 and 1273 K for different holding times not exceeding 40 min. Joining took place by the mechanism of diffusion bonding. The joints produced at 1273 K for 40 min exhibited no porosity in the reaction zones and presented the best quality. Scanning electron microscopy (SEM) characterization of the bonding zone revealed an improvement in the quality of the joints brazed at 1173 K for 20 min and longer. These samples had continuous base metal–filler alloy interfaces with minimum porosity. At 1273 K the bonding interfaces diffused and for the samples held for 40 min completely vanished and porosity disappeared. Even the presence of particle precipitates the bonding zone showed acceptable resistance to localised corrosion in non-aggressive electrolytes. SEM study revealed that irregularly precipitated particles and other phases of about 10 m in size formed in the interlayer during the joining process. The presence of -phase in samples bonded at 1273 K promoted preferential dissolution in the bonding zone in NaCl solution. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Metallic glasses or metallic amorphous alloys have been subjects of abundant studies over the past 30 years regarding their manufacture, mechanical and physical properties, and have been used in applications such as sensors, magnetic heads for recording items as reported by Davies (1983). More recently, Rabinkin et al. (1998) used this kind of alloys as joining element because of their ductility, homogeneity of microstructure, low melting temperatures and their superior corrosion resistance. One of the advantages to use metallic amorphous alloys obtained by rapid solidification, (106 –1 K/s) besides the possibility of having self-fluxing is that they can be obtained in tiny thickness and they are ductile and chemically and structurally homogeneous as demonstrated by Kimura et al. (1996). Their melting temperatures are relatively low which make them suitable to be used as bonding alloys for diverse mate-
∗ Corresponding author. Tel.: +52 981 811 98 00x62804; fax: +52 981 811 98 00x62899. E-mail addresses: [email protected] (J.A. Verduzco), [email protected] (J. González-Sánchez), [email protected] (V.H. Verduzco), josesolis@infinitum.com.mx (J. Solís), [email protected] (J. Lemus-Ruiz). 0924-0136/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2010.02.015
rials and applications. Amorphous alloys have been used for joining sections of heat exchangers as reported by Rabinkin et al. (1998). Wu et al. (2000) used metallic glasses for bonding Inconel 718 to Inconel X-750. Currently amorphous alloys have been successfully used to join Ti-based biomedical alloys by Miura et al. (2007) and ceramic matrix composite to titanium metal by Singh et al. (2008). These studies have employed Ni, Mg, Zr, Cu or Pd based metallic glass alloys as the filler element but no studies regarding to the use of Fe-based metallic glass as the filler material to join materials such as stainless steel have been reported. Hashimoto et al. (1976) reported that the corrosion rate of Fe72 Cr8 P13 C7 amorphous alloy is negligible in 0.5 M NaCl. Amorphous alloys present higher corrosion resistance than those with crystalline microstructure of similar composition as demonstrated by Jayaraj et al. (2005). These alloys are used extensively in areas where resistance to both wear and corrosion is needed. However, during welding processes amorphous alloys undergo partial or complete recrystallisation, (devitrification process) and their corrosion resistance decreases as reported by Scully et al. (2007). They reported that crystalline structure involves the presence of surface defects and discontinuity of the passive film. An interesting technological option is the use of diffusion bonding process to join austenitic stainless steel with amorphous alloys at temperatures lower than that for recrystallisation. The fact that
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amorphous alloys possess high corrosion resistance is associated to the negligible quantity of defects present in a crystalline alloy, such as grain boundaries and dislocations. However, the actual electrochemical mechanisms of the high corrosion resistance are still not fully understood. González-Sánchez et al. (2007) have previously reported that the corrosion resistance of amorphous alloys disappear due to recrystallisation as a consequence of bonding process at temperatures higher than 50% the melting temperature of the alloy. In the case of diffusion bonding of AISI 304 stainless steels using amorphous Ni-based alloys, bonding process conducted in the temperature range from 923 to 1123 K (650–850 ◦ C) induced sensitisation of the steel and the joining zone became highly susceptible to intergranular corrosion in 3.5% NaCl solution as reported by González-Sánchez et al. (2007). Due to the variety of engineering applications of austenitic stainless steels, research has been done about alternative manufacturing processes including diffusion bonding. Austenitic steels such as AISI 316L have shown better resistance to localised corrosion and to sensitisation during soldering and heat treatment processes. The resistance of this austenitic steel to sensitisation is due, as reported and explained by Lacombe et al. (1993) and Marshall (1984), to their low carbon content which is <0.03 wt%. The present research focussed on the corrosion resistance and microstructural characterization of the bonding zone of AISI 316L stainless steel samples joined with an amorphous Fe75 Cr8 P10 B7 ribbon alloy when they are in contact with distilled water and with 3.5wt% NaCl solution. It is a fundamental aspect to produce flawless joints and to keep acceptable corrosion resistance of stainless steels when are in contact with chloride containing electrolytes during service life. The relationship between the microstructure regarding the diffusion process at the joint interface and the corrosion resistance of bonding zone in a chloride containing electrolyte was studied.
2. Experimental procedure
Fig. 1. Differential thermal analysis of the Fe75 Cr8 P10 B7 glassy ribbon.
2.2. Joining procedure The surface of the 316L SS samples to be joined (0.3 cm × 0.5 cm) were ground with up to a 600 grit (P 1200) emery papers and polished to mirror surface finish using a 1 m diamond paste. Two samples were put face to face in a sandwich type arrangement placing between them a piece of the Fe-based alloy ribbon. This assembly was then placed in a graphite die embedded in a boron nitride (99.5% pure) powder bed and subjected to mechanical pressure to keep the assembly together. The experimental set up used has been explained with detail elsewhere by González-Sánchez et al. (2007) and by Flores et al. (2006). Once each sample was assembled in the graphite die, it was positioned into the furnace filled with argon as shown in Fig. 2. The joining process was carried out at the two mentioned temperatures using holding times of 5, 10, 15, 20 and 40 min.
2.1. Materials Flat samples of AISI 316L stainless steel (316L SS) 1.0 cm × 0.5 cm and 3 mm thick were joined to themselves with a home-made Fe75 Cr8 P10 B7 amorphous alloy ribbon 40 m thick and 3 mm width. The bonding material was obtained by melt spinning in a controlled He atmosphere. Table 1 shows the chemical composition of both materials. The amorphous structure of the ribbon was confirmed by X-ray diffraction (XRD) using a SIEMENS D5000 diffractometer operated with Cu K␣ at 30 kV. The bonding and melting temperatures of the amorphous ribbon alloy were determined by differential thermal analysis (DTA) carried out in a SDT Q600 calorimeter which result is shown in Fig. 1. From the results of the DTA analysis it was decided to use two bonding temperatures to carry out the joining of the 316L steel samples. The bonding temperatures were 1173 and 1273 K, considering that the melting temperature of the amorphous ribbon was found to be around 1571 K.
Table 1 Chemical composition of the starting materials (at.%). Material
Amorphous alloy ribbon AISI 316L stainless steel
Chemical composition at.% Fe
Ni
Cr
P
Mo
B
75 64.2
– 11.4
8 20.8
10 –
– 3.4
7 –
Fig. 2. Setup of the joining.
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ples using scanning electron microscopy. Energy dispersive X-ray spectroscopy (EDS) punctual microanalysis was carried out at the interface and in the precipitate particles formed in samples joined at 1273 K regardless the holding time. 2.4. Evaluation of the electrochemical behaviour of the bonded samples
Fig. 3. Liquid-like X-ray diffraction pattern of the Fe75 Cr8 P10 B7 alloy ribbon.
2.3. Structural characterization Microstructural examination in the joining zones of the samples was performed on polished cross-sections of the experimental cou-
The corrosion resistance of the 316L SS jointed samples was evaluated by potentiodynamic polarisation applied to samples in 3.5 wt% NaCl solution (pH 8), and in distilled water (pH 5.8) at room temperature. The electrochemical tests were conducted in a conventional 3-electrode electrochemical cell formed by a saturated calomel electrode (SCE) as reference electrode, a platinum wire as auxiliary electrode and the bonded stainless steel samples as working electrodes in a glass container of 1000 ml of capacity. The 3.5 wt% NaCl solution was prepared with de-ionised water and NaCl salt of reagent grade purity. Polarisation was conducted on the different SS samples using a computer controlled potentiostat/galvanostat EG&G model 273 A. An anodic overpotential of 600 mV and a cathodic overpotential of 300 mV were applied with a scan rate of 10 mV min−1 .
Fig. 4. Cross-section of the interfaces of samples joined at 1173 K for joining times of: (a) 5 min, (b) 10 min, (c) 15 min, (d) 20 min and (e) 40 min.
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3. Results and discussion 3.1. Characterisation The result of the X-ray diffraction analysis is shown in Fig. 3 in which a liquid-like pattern for the as-cast alloy ribbon was observed indicating its amorphous structure. 3.2. Joints made at 1173 K Samples joined at 1173 K showed a continuous interface formed between the steel and the ribbon alloy with a number of pores that decreased as a function of the holding time. The metallographic analysis revealed clearly the different constituents of the bonding zone: the stainless steel sides, the alloy ribbon between them and a tiny phase (approximately 3 m thick) in the interface as can be observed in the micrographs of Fig. 4. This thin bonding phase formed in all cases regardless the holding time, separating clearly the stainless steel from the alloy ribbon, which at this magnification
did not present crystalline structure as no grain boundaries were observed. At this process temperature, bonding took place through the formation of a diffusion interface (the bonding phase) for all the holding times as the melting temperature of the amorphous alloy was about 400◦ higher. Punctual chemical microanalysis was performed on both sides of the joints including the interface. The chemical analysis is presented for different holding times as a function of position. Seven points were selected starting in the AISI 316L stainless steel side and finishing in the alloy ribbon passing through the bonding phase (point 4) as shown in Fig. 5. For all holding times, the concentration of Fe at the interface was higher than in both the stainless steel and the alloy ribbon sides. A decrease in Cr concentration was found in the bonding phase for holding times longer than 15 min but always higher than 10 at% with consequent increment of Cr in the stainless steel and in the alloy ribbon. This result indicated that at this holding time intense Cr diffusion took place. A minimal diffusion of Ni from the stain-
Fig. 5. Punctual chemical microanalysis of samples joined at 1173 K: (a) photomicrograph showing the points of the analysis, (b) 5 min, (c) 10 min, (d) 15 min, (e) 20 min and (f) 40 min holding times: = Fe, X = Cr, = Ni, 䊉 = P and ♦ = Mo.
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Fig. 6. EDS punctual chemical microanalysis of Fe, Cr, Ni and P as a function of holding time of points 2 and 6 of Fig. 5a, corresponding to the AISI 316L steel and alloy ribbon, respectively.
less steel to the alloy ribbon was observed after 40 min of holding time. The diffusion of P from the alloy ribbon to the AISI 316L steel was negligible. However at the bonding phase the P concentration increased after 40 min. The diffusion of Mo from the SS to the alloy ribbon was absent at the joining temperature of 1173 K regardless holding time. In order to observe more clearly the depth of diffusion of Fe, Cr, Ni, and P, their concentrations at points labelled 2 (steel) and 6 (alloy ribbon) are plotted as a function of the holding time and presented in Fig. 6. The concentration of Fe and Cr in both the S31603 steel and amorphous ribbon presents the most intense changes as a function of the holding time. The microanalysis indicated that after a holding time of 15 min the concentration of Cr was the same in both metals (about 14 at%), which is high enough to provide a passive condition of both Fe-based alloys. The concentration of Fe at these two sites reached equal values after a holding time of approximately 23 min at 1173 K. In the case of Ni, important concentration changes took place only after a holding time of 20 min with a decrease in the steel and consequent increment in the ribbon alloy. The concentration of P in the SS was negligible for holding periods up to 10 min. For longer periods the concentration increased and remained constant around 3 at% after 20 and 40 min. 3.3. Joints made at 1273 K Joints produced at 1273 K did not showed a defined interface or bonding phase between the stainless steel and the alloy ribbon regardless the holding time as can be seen in Fig. 7. At this processing temperature different phases and precipitate particles formed in the ribbon alloy. These phases and precipitates formed at all 5 holding times. The boundary formed between the AISI 316L steel and the ribbon alloy was irregular for holding times longer than 5 min. The remnant ribbon alloy did not present grain boundaries
which could be an indication that even under that process conditions this alloy keep its amorphous structure. In order to determine the composition of the particles and diffusion of Fe, Cr, Ni, P, and Mo, punctual chemical analyses was carried out at and near the precipitate particles; the results are plotted as a function of the holding time in Fig. 8. It is seen that in the precipitated particles, as the holding time increased, Fe content decreased. Ni concentration in these particles also decreased and after a holding time of 15 min, Ni was not observed. On the other hand, Ni was present in the site tested that was supposed to be the ribbon alloy, which originally did not contain Ni. The Ni concentration found in the site near the precipitated particles is very similar to that for the AISI 316L steel. The P content decreased drastically from ∼13 at.% after 10 min to ∼2 at.% after 15 min which practically disappeared after 20 min and appeared again when the holding time was 40 min having ∼12 at.%. At this joining temperature, Mo diffused from the 316L SS firstly to the alloy ribbon and then to the precipitate particle, registering more concentration after bonding for period longer than 40 min. The non-clearly delimited interface formed between the AISI 316L steel and the alloy ribbon and the formation of precipitate particles in the bonding zone can be the consequence of presence of liquid states that might occur 1273 K. While for the joining process carried out at 1173 K, the SEM analysis revealed a bonding process by diffusion in the solid state because of the well-defined SS/alloy ribbon interfaces. The chemical composition of precipitated particles given in at% in Fig. 8 was converted to wt% and is presented in Table 2. Based on the isotherms at 923 K (650 ◦ C) of the Fe–Ni–Cr (Raynor and Rivlin, 1988) and Fe–Ni–Mo (Wada, 1973) ternary equilibrium diagrams shown in Fig. 9, the phases found in the bonding zone as a function of the holding time are proposed in Table 3. The predominating phase present in all holding times was -phase, which has a tetragonal unit cell as demonstrated (Barcik, 1983;
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Fig. 7. SEM images of the cross-sections of the unions carried out at 1273 K for (a) 5 min, (b) 10 min, (c) 15 min, (d) 20 min and (e) 40 min.
Marshall, 1984). Studies have reported its formation associated with the dissolution of M23 C6 , but also independently (Barcik, 1983). 3.4. Electrochemical behaviour and corrosion resistance Samples bonded at 1173 K presented resistance to localised dissolution in distilled water with no crevice corrosion or pitting attack developed in the bonding region during potentiodynamic polarisation. However, tests in 3.5% NaCl solution showed preferential dissolution in the zone surrounding the bonding region (bonding phase). The dissolution was due to crevice corrosion provoking the complete separation of stainless steel plates. The pores observed in the bonded phase on samples treated at 1173 K at different holding times were not the initiation sites for metal dissolution. The bonding phase suffered preferential dissolution because of Cr depletion with Cr concentration less than 12 at.% as shown in Fig. 5. Immersion test at open circuit potential in 3.5% NaCl solution showed that the preferential dissolution along the bonding phase took place after 10 days.
Samples bonded at 1273 K presented more noble corrosion potential (Ecorr vs SCE) in distilled water than samples joined at 1173 K. The anodic current density was lower for samples joined at 1273 K than for samples at 1173 K in distilled water except for samples bonded for 15 and 20 min, which presented equal and higher current densities respectively. Fig. 10 presents polarisation curves of samples bonded at 1173 and 1273 K immersed in distilled water and in 3.5 wt% NaCl after bonding times of 5, 10, 15, 20 and 40 min. In distilled water, the sample held for 5 min showed a passive region extended from −50 to 200 mV with a current density of 8.4 × 10−2 A/cm2 , which represents very low corrosion rate. 316L SS samples bonded with this amorphous alloy presented excellent corrosion resistance for applications involving non-aggressive electrolytes such as distilled water. For bonding times of 15, 20 and 40 min the samples treated at 1173 and at 1273 K presented similar behaviour in distilled water with an Ecorr around −200 mV and pseudo-passive behaviour which was characterised by a high polarised anodic reaction. This means that the increase of the applied anodic overpotential induced minimum increment of the anodic current density.
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Fig. 8. EDS punctual chemical microanalyses of Fe, Cr, Ni, P, and Mo as a function of the holding times of samples joined at 1273 K. The EDS were performed at 2000× at the points corresponding to the precipitate particles () and “alloy ribbon” (♦) zones of the SEM images of Fig. 7.
Fig. 9. Isotherms at 923 K (650 ◦ C) of ternary equilibrium diagrams phases of (a) Fe–Cr–Ni (Raynor and Rivlin, 1988) and (b) Fe–Cr–Mo (Wada, 1973). Symbols represent the → (Cr, Mo) + . phases: → ␣ + , 䊉 → ␥ + , → , (Cr, Mo) + , and
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Table 2 Chemical composition (wt%) at the precipitate particles formed at 1273 K for the different holding times. Holding time (min) 5 10 15 20 40
Chemical composition wt.% Fe Ni Cr
P
Mo
59.4 58.1 45.7 44.6 27.8
8.97 8.86 7.99 0.38 6.95
– 1.18 1.48 1.15 13.0
9.24 2.69 – – –
22.3 29.1 44.7 53.8 52.1
On the other hand, test conducted in 3.5 wt% NaCl solution showed that samples bonded at 1273 K presented in general more noble values of Ecorr than samples bonded at 1173 K. In this electrolyte, the anodic reaction was under charge transfer control
Table 3 Phases determined at the precipitate particles according to the isotherms at 650 ◦ C of ternary diagram phases of Fig. 9 for the different holding times. Holding time (min)
Phases
Symbol
5 10 15 20 40
␣+ ␥+ (Cr, Mo) + (Cr, Mo) +
䊉 ♦
kinetics with high anodic current densities and a change of mechanism around −200 mV (associated to the increase of the anodic curve slope). This indicated a fast metal dissolution taking place in the bonding zone. Austenitic stainless steel AISI 316L showed a pas-
Fig. 10. Polarisation curves for samples bonded at 1173 and 1273 K during 5, 10, 15, 20 and 40 min in distilled water and 3.5 wt% NaCl solution.
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sive behaviour during anodic polarisation in synthetic seawater in a potential range of about 600 mV as reported by González-Sánchez (2002). According to Pistorius and Burstein (1992) the nucleation of corrosion pits in stainless involves the breakdown of the passive film on a microscopic scale. Burstein et al. (1993) determined that this process induces small and sudden increments of anodic current, which are characterised by current spikes, leading to oxidation and dissolution of less than 0.01 m3 of metal. The nucleation process is unstable and in most cases it will stop by the regeneration of the passive film. Sedriks (1990) indicates the importance of different types of inhomogeneities at the material surface during pit initiation. Lattice defects at the surface, precipitates, intermetallic phases, second phases and non-metallic inclusions are advantageous sites for the nucleation of pitting corrosion due to their different electrochemical behaviour with respect the metal matrix. The presence of a passive film is required for pitting corrosion to take place in stainless steels in chloride containing electrolytes. However in the present case, the bonded samples did not present that passive region during anodic polarisation in the 3.5 wt% NaCl solution. Active dissolution took place instead as can be seen in the polarisation curves in Fig. 10 for samples bonded at the five different holding times. Clearly the anodic branch moved to the right for the case of test in NaCl solution compared to those in distilled water. No pitting corrosion was induced during anodic polarisation in the NaCl solution because this form of attack requires high anodic polarisation to provoke the breakdown of the passive film. As no passive film was present active dissolution took place instead. The bonding phase formed on samples joined at 1173 K with Cr concentration lower than 10 at.% acted as preferential site for metal dissolution. In the case of samples bonded at 1273 K, the dissolution of the bonding zone could be facilitated by the presence of the -phase and other precipitated particles. This induced anodic dissolution with anodic current densities of about 0.6 mA/cm2 . The presence of different types of discontinuities and defects at the bonding zone such as precipitates, intermetallic phases, second phases and non-metallic inclusions are preferential sites for interruption of the passive film as reported by Böhni (1992). As mentioned above, in the bonding zone there were several phases which inhibited the formation of a continuous passive film which induced preferential dissolution in the bonding zone. Rebak et al. (2008) reported that even micropartial devitrification could induce localised corrosion in amorphous Fe base alloys, which may be involved in the bonding process along with the formation of secondary phases. Samples bonded at the temperature of 1273 K showed in general more positive potentials and lower anodic current densities than samples bonded at 1173 K. The presence of -phase could induce the formation of micro-galvanic cells due to the different chemical composition of the phases. In distilled water such galvanic cell does not develop a significant potential gradient but in NaCl solution the micro-galvanic cells became very active. When joined by self-diffusion, the change to the non-crystalline structure of the amorphous alloy and the formation of different phases made the dissimilar nature of the alloys more significant and promoted selective dissolution coupled with crevice corrosion. The lower anodic current density showed by samples bonded at 1273 K compared to that for samples bonded at 1173 K could be due to the higher Cr concentration in the bonding zone reached at 1273 K even the presence of -phase in samples bonded at both temperatures. 4. Conclusions It was possible to produce 316L stainless steel joined samples through difussion bonding using a home-made Fe-based amorphous alloy ribbon as the joining element.
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At 1173 K the bonding process formed SS/alloy ribbon interfaces with clear differences in chemical composition and structure from the base metal and the alloy ribbon which indeced poor corrosion resistance in a chloride containing electrolyte. Samples bonded at 1273 K did not formed a defined bonded phase or zone, instead the ribbon alloy appeared surrounding austenitic grains of the AISI 316L stainles steel making an irregular bonding boundary. This behaviour was more evident as the holding time increased. The joining processes at 1173 K took place by self-diffusion bonding in solid state, whereas at 1273 K liquid state diffusion could take place partially forming some precipitate particles composed mainly of phase. When joined by self-diffusion, the change to the non-crystalline structure of the amorphous alloy and the formation of different phases made the dissimilar nature of the alloys more significant and promoted selective dissolution coupled with crevice corrosion. The bonding process of austenitic stainles steel to itself by using a Fe base amorphous alloy produce aceptable joints with minimal porosity and defects and with acceptable corrosion resistance in non-aggressive electrolytes. However due to the formation of different phases in the bonding zone localised corrosion along the bonding line is prompted in chloride containing electrolytes. Acknowledgements We would like to thank CONACYT, SEP-DGEST México and the CIC of the UMSNH for the financial support. JAV wishes to akcnowledge The University of Sheffield the facilities given to fabricate the Fe-based amorphous alloy ribbon. Also we would like to thank the Centre for Corrosion Research at the Universidad Autónoma de Capeche for the facilities given to JGS to carry out the electrochemical tests. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.jmatprotec.2010.02.015. References Barcik, J., 1983. Crystallographic characteristics of the process of -phase precipitation in chromium-nickel austenitic steels. J. Appl. Crystallogr. 16, 590–598. Böhni, H., 1992. Localised corrosion-mechanisms and methods. Mater. Sci. Forum. 111–112, 401–414. Burstein, G.T., Pistorius, P.C., Mattin, S.P., 1993. The nucleation and growth of corrosion pits on stainless steel. Corros. Sci. 35, 57–62. Davies, H.A., 1983. In: Luborsky, F.E. (Ed.), Amorphous Metallic Alloys. Butterworths, London, UK, pp. 8–25. Flores, J.G., Cervantes, J., Lemus-Ruiz, J., 2006. Joining of silicon nitride to metal (Mo and Ti) using a Cu–foil interlayer. Mater. Sci. Forum. 509, 99–104. González-Sánchez, J., Verduzco, J.A., Lemus-Ruiz, J., Téllez, M.G., Torres, R., 2007. Corrosion resistance of stainless steel joints bonded with a Ni-based amorphous alloy. Anti-Corrosion Methods Mater. 54/2, 68–73. González-Sánchez, J., 2002. PhD Thesis, Materials and Engineering Research Institute, Sheffield Hallam University, U.K. Hashimoto, K., Osada, K., Masumoto, T., Shimodaira, S., 1976. Characteristics of passivity of extremely corrosion resistant amorphous iron alloys. Corros. Sci. 16, 71–76. Jayaraj, J., Kim, Y., Kim, K., Seok, H., Fleury, E., 2005. Corrosion studies on Fe-based amorphous alloys in simulated PEM fuel cell environment. Sci. Technol. Adv. Mater. 6, 282–289. Kimura, H., Sasamori, K., Inoue, A., 1996. High strength Al–Ti–Fe alloys consisting of amorphous and fcc-Al phases prepared by rapid solidification. Mater. Trans. JIM 37, 1722–1725. Lacombe, P., Baroux, B., Beranger, G., 1993. Stainless steels. In: Les editions de physique. Les Ulis, France, pp. 331–336. Marshall, P., 1984. Austenitic stainless steels. In: Microstructure and Mechanical Properties. Elsevier Applied Science Publishers Ltd, Essex, England, pp. 23–64. Miura, E., Kato, H., Ogata, T., Nishiyama, N., Specht, E.D., Shiraihi, T., Inoue, A., Hisatsune, K., 2007. Mechanical property and corrosion resistance evaluations of Ti–6Al–7Nb alloy brazed with bulk metallic glasses. Mater. Trans., JIM 48, 2235–2243. Pistorius, P.C., Burstein, G.T., 1992. Detailed investigation of current transients from metastable pitting events on stainless steel: the transition to stability. Mater. Sci. Forum. 111–112, 429–452.
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Rabinkin, A., Wenski, E., Ribaudo, A., 1998. Brazing stainless steel using a new MBFseries of Ni–Cr–B–Si amorphous brazing foils. Weld. J. 77, 66–75. Raynor, G.V., Rivlin, V.G., 1988. Phase Equilibria in Iron Ternary Alloys (Phase Diagrams of Ternary Iron Alloys IV). The Institute of Metals, London, pp. 485. Rebak, R., Daniel, D.S., Tiangan, L., Hailey, Ph., Farmer, J., 2008. Environmental testing of iron-based amorphous alloys. Metall. Mater. Trans. A 39A, 225–234. Scully, J., Gebert, A., Payer, J., 2007. Corrosion and related mechanical properties of bulk metallic glasses. J. Mater. Res. 22, 302–313. Sedriks, A.J., 1990. Effects of alloy composition and microstructure on the localised corrosion of stainless steels. In: Isaacs, H., Bertocci, U., Kruger, J., Smialowska,
S. (Eds.), Advances in Localised Corrosion (NACE-9). Orlando, Florida, pp. 253–262. Singh, M., Asthana, R., Shpargel, T.P., 2008. Brazing of ceramic–matrix composites to Ti and Hastealloy using Ni-base metallic glass interlayers. Mater. Sci. Eng. A 498, 19–30. Wada, T., 1973. In: Taylor, L. (Ed.), Iron–Molybdenum–Nickel. Metals Handbook, vol. 8. ASM Metals Park, OH, pp. 431–1431. Wu, X., Chandel, R.S., Li, H., Seow, H.P., Wu, S., 2000. Induction brazing of Inconel 718 to Inconel X-750 using Ni–Cr–Si–B amorphous foil. J. Mater. Process. Technol. 104, 34–43.
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Journal of Materials Processing Technology journal homepage: www.elsevier.com/locate/jmatprotec
Microstructure and electrochemical properties of the bonding zone of AISI 316L steel joined with a Fe-based amorphous foil J.A. Verduzco a , J. González-Sánchez b,∗ , V.H. Verduzco a , J. Solís c , J. Lemus-Ruiz a a b c
Instituto de Investigaciones Metalúrgicas, (UMSNH), P.O. Box 888, Morelia, Michoacán 58000, Mexico Centre for Corrosion Research, Universidad Autónoma de Campeche, Av. Agustín Melgar s/n, Col. Buenvista, Campeche, Cam., C.P. 24039, Mexico SEP-DGEST-ITTLA, Av. Mario Colín, S/N, Tlalnepantla Edo. de Méx., 54070, Mexico
a r t i c l e
i n f o
Article history: Received 1 December 2009 Received in revised form 30 January 2010 Accepted 17 February 2010
Keywords: Fe-based amorphous alloys Diffusion bonding Stainless steel Localised corrosion
a b s t r a c t The effect of temperature and holding time on the microstructure and corrosion resistance of the junction zone of AISI 316L stainless steel (SS) bonded to itself with Fe75 Cr8 P10 B7 filler alloy was investigated. The brazing alloy was prepared in the laboratory in the form of amorphous ribbons and its melting temperature was determined by differential thermal analysis (DTA) to be 1571 K. The joining process was carried out in a chamber with controlled Ar atmosphere at two temperatures: 1173 and 1273 K for different holding times not exceeding 40 min. Joining took place by the mechanism of diffusion bonding. The joints produced at 1273 K for 40 min exhibited no porosity in the reaction zones and presented the best quality. Scanning electron microscopy (SEM) characterization of the bonding zone revealed an improvement in the quality of the joints brazed at 1173 K for 20 min and longer. These samples had continuous base metal–filler alloy interfaces with minimum porosity. At 1273 K the bonding interfaces diffused and for the samples held for 40 min completely vanished and porosity disappeared. Even the presence of particle precipitates the bonding zone showed acceptable resistance to localised corrosion in non-aggressive electrolytes. SEM study revealed that irregularly precipitated particles and other phases of about 10 m in size formed in the interlayer during the joining process. The presence of -phase in samples bonded at 1273 K promoted preferential dissolution in the bonding zone in NaCl solution. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Metallic glasses or metallic amorphous alloys have been subjects of abundant studies over the past 30 years regarding their manufacture, mechanical and physical properties, and have been used in applications such as sensors, magnetic heads for recording items as reported by Davies (1983). More recently, Rabinkin et al. (1998) used this kind of alloys as joining element because of their ductility, homogeneity of microstructure, low melting temperatures and their superior corrosion resistance. One of the advantages to use metallic amorphous alloys obtained by rapid solidification, (106 –1 K/s) besides the possibility of having self-fluxing is that they can be obtained in tiny thickness and they are ductile and chemically and structurally homogeneous as demonstrated by Kimura et al. (1996). Their melting temperatures are relatively low which make them suitable to be used as bonding alloys for diverse mate-
∗ Corresponding author. Tel.: +52 981 811 98 00x62804; fax: +52 981 811 98 00x62899. E-mail addresses: [email protected] (J.A. Verduzco), [email protected] (J. González-Sánchez), [email protected] (V.H. Verduzco), josesolis@infinitum.com.mx (J. Solís), [email protected] (J. Lemus-Ruiz). 0924-0136/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2010.02.015
rials and applications. Amorphous alloys have been used for joining sections of heat exchangers as reported by Rabinkin et al. (1998). Wu et al. (2000) used metallic glasses for bonding Inconel 718 to Inconel X-750. Currently amorphous alloys have been successfully used to join Ti-based biomedical alloys by Miura et al. (2007) and ceramic matrix composite to titanium metal by Singh et al. (2008). These studies have employed Ni, Mg, Zr, Cu or Pd based metallic glass alloys as the filler element but no studies regarding to the use of Fe-based metallic glass as the filler material to join materials such as stainless steel have been reported. Hashimoto et al. (1976) reported that the corrosion rate of Fe72 Cr8 P13 C7 amorphous alloy is negligible in 0.5 M NaCl. Amorphous alloys present higher corrosion resistance than those with crystalline microstructure of similar composition as demonstrated by Jayaraj et al. (2005). These alloys are used extensively in areas where resistance to both wear and corrosion is needed. However, during welding processes amorphous alloys undergo partial or complete recrystallisation, (devitrification process) and their corrosion resistance decreases as reported by Scully et al. (2007). They reported that crystalline structure involves the presence of surface defects and discontinuity of the passive film. An interesting technological option is the use of diffusion bonding process to join austenitic stainless steel with amorphous alloys at temperatures lower than that for recrystallisation. The fact that
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amorphous alloys possess high corrosion resistance is associated to the negligible quantity of defects present in a crystalline alloy, such as grain boundaries and dislocations. However, the actual electrochemical mechanisms of the high corrosion resistance are still not fully understood. González-Sánchez et al. (2007) have previously reported that the corrosion resistance of amorphous alloys disappear due to recrystallisation as a consequence of bonding process at temperatures higher than 50% the melting temperature of the alloy. In the case of diffusion bonding of AISI 304 stainless steels using amorphous Ni-based alloys, bonding process conducted in the temperature range from 923 to 1123 K (650–850 ◦ C) induced sensitisation of the steel and the joining zone became highly susceptible to intergranular corrosion in 3.5% NaCl solution as reported by González-Sánchez et al. (2007). Due to the variety of engineering applications of austenitic stainless steels, research has been done about alternative manufacturing processes including diffusion bonding. Austenitic steels such as AISI 316L have shown better resistance to localised corrosion and to sensitisation during soldering and heat treatment processes. The resistance of this austenitic steel to sensitisation is due, as reported and explained by Lacombe et al. (1993) and Marshall (1984), to their low carbon content which is <0.03 wt%. The present research focussed on the corrosion resistance and microstructural characterization of the bonding zone of AISI 316L stainless steel samples joined with an amorphous Fe75 Cr8 P10 B7 ribbon alloy when they are in contact with distilled water and with 3.5wt% NaCl solution. It is a fundamental aspect to produce flawless joints and to keep acceptable corrosion resistance of stainless steels when are in contact with chloride containing electrolytes during service life. The relationship between the microstructure regarding the diffusion process at the joint interface and the corrosion resistance of bonding zone in a chloride containing electrolyte was studied.
2. Experimental procedure
Fig. 1. Differential thermal analysis of the Fe75 Cr8 P10 B7 glassy ribbon.
2.2. Joining procedure The surface of the 316L SS samples to be joined (0.3 cm × 0.5 cm) were ground with up to a 600 grit (P 1200) emery papers and polished to mirror surface finish using a 1 m diamond paste. Two samples were put face to face in a sandwich type arrangement placing between them a piece of the Fe-based alloy ribbon. This assembly was then placed in a graphite die embedded in a boron nitride (99.5% pure) powder bed and subjected to mechanical pressure to keep the assembly together. The experimental set up used has been explained with detail elsewhere by González-Sánchez et al. (2007) and by Flores et al. (2006). Once each sample was assembled in the graphite die, it was positioned into the furnace filled with argon as shown in Fig. 2. The joining process was carried out at the two mentioned temperatures using holding times of 5, 10, 15, 20 and 40 min.
2.1. Materials Flat samples of AISI 316L stainless steel (316L SS) 1.0 cm × 0.5 cm and 3 mm thick were joined to themselves with a home-made Fe75 Cr8 P10 B7 amorphous alloy ribbon 40 m thick and 3 mm width. The bonding material was obtained by melt spinning in a controlled He atmosphere. Table 1 shows the chemical composition of both materials. The amorphous structure of the ribbon was confirmed by X-ray diffraction (XRD) using a SIEMENS D5000 diffractometer operated with Cu K␣ at 30 kV. The bonding and melting temperatures of the amorphous ribbon alloy were determined by differential thermal analysis (DTA) carried out in a SDT Q600 calorimeter which result is shown in Fig. 1. From the results of the DTA analysis it was decided to use two bonding temperatures to carry out the joining of the 316L steel samples. The bonding temperatures were 1173 and 1273 K, considering that the melting temperature of the amorphous ribbon was found to be around 1571 K.
Table 1 Chemical composition of the starting materials (at.%). Material
Amorphous alloy ribbon AISI 316L stainless steel
Chemical composition at.% Fe
Ni
Cr
P
Mo
B
75 64.2
– 11.4
8 20.8
10 –
– 3.4
7 –
Fig. 2. Setup of the joining.
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ples using scanning electron microscopy. Energy dispersive X-ray spectroscopy (EDS) punctual microanalysis was carried out at the interface and in the precipitate particles formed in samples joined at 1273 K regardless the holding time. 2.4. Evaluation of the electrochemical behaviour of the bonded samples
Fig. 3. Liquid-like X-ray diffraction pattern of the Fe75 Cr8 P10 B7 alloy ribbon.
2.3. Structural characterization Microstructural examination in the joining zones of the samples was performed on polished cross-sections of the experimental cou-
The corrosion resistance of the 316L SS jointed samples was evaluated by potentiodynamic polarisation applied to samples in 3.5 wt% NaCl solution (pH 8), and in distilled water (pH 5.8) at room temperature. The electrochemical tests were conducted in a conventional 3-electrode electrochemical cell formed by a saturated calomel electrode (SCE) as reference electrode, a platinum wire as auxiliary electrode and the bonded stainless steel samples as working electrodes in a glass container of 1000 ml of capacity. The 3.5 wt% NaCl solution was prepared with de-ionised water and NaCl salt of reagent grade purity. Polarisation was conducted on the different SS samples using a computer controlled potentiostat/galvanostat EG&G model 273 A. An anodic overpotential of 600 mV and a cathodic overpotential of 300 mV were applied with a scan rate of 10 mV min−1 .
Fig. 4. Cross-section of the interfaces of samples joined at 1173 K for joining times of: (a) 5 min, (b) 10 min, (c) 15 min, (d) 20 min and (e) 40 min.
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3. Results and discussion 3.1. Characterisation The result of the X-ray diffraction analysis is shown in Fig. 3 in which a liquid-like pattern for the as-cast alloy ribbon was observed indicating its amorphous structure. 3.2. Joints made at 1173 K Samples joined at 1173 K showed a continuous interface formed between the steel and the ribbon alloy with a number of pores that decreased as a function of the holding time. The metallographic analysis revealed clearly the different constituents of the bonding zone: the stainless steel sides, the alloy ribbon between them and a tiny phase (approximately 3 m thick) in the interface as can be observed in the micrographs of Fig. 4. This thin bonding phase formed in all cases regardless the holding time, separating clearly the stainless steel from the alloy ribbon, which at this magnification
did not present crystalline structure as no grain boundaries were observed. At this process temperature, bonding took place through the formation of a diffusion interface (the bonding phase) for all the holding times as the melting temperature of the amorphous alloy was about 400◦ higher. Punctual chemical microanalysis was performed on both sides of the joints including the interface. The chemical analysis is presented for different holding times as a function of position. Seven points were selected starting in the AISI 316L stainless steel side and finishing in the alloy ribbon passing through the bonding phase (point 4) as shown in Fig. 5. For all holding times, the concentration of Fe at the interface was higher than in both the stainless steel and the alloy ribbon sides. A decrease in Cr concentration was found in the bonding phase for holding times longer than 15 min but always higher than 10 at% with consequent increment of Cr in the stainless steel and in the alloy ribbon. This result indicated that at this holding time intense Cr diffusion took place. A minimal diffusion of Ni from the stain-
Fig. 5. Punctual chemical microanalysis of samples joined at 1173 K: (a) photomicrograph showing the points of the analysis, (b) 5 min, (c) 10 min, (d) 15 min, (e) 20 min and (f) 40 min holding times: = Fe, X = Cr, = Ni, 䊉 = P and ♦ = Mo.
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Fig. 6. EDS punctual chemical microanalysis of Fe, Cr, Ni and P as a function of holding time of points 2 and 6 of Fig. 5a, corresponding to the AISI 316L steel and alloy ribbon, respectively.
less steel to the alloy ribbon was observed after 40 min of holding time. The diffusion of P from the alloy ribbon to the AISI 316L steel was negligible. However at the bonding phase the P concentration increased after 40 min. The diffusion of Mo from the SS to the alloy ribbon was absent at the joining temperature of 1173 K regardless holding time. In order to observe more clearly the depth of diffusion of Fe, Cr, Ni, and P, their concentrations at points labelled 2 (steel) and 6 (alloy ribbon) are plotted as a function of the holding time and presented in Fig. 6. The concentration of Fe and Cr in both the S31603 steel and amorphous ribbon presents the most intense changes as a function of the holding time. The microanalysis indicated that after a holding time of 15 min the concentration of Cr was the same in both metals (about 14 at%), which is high enough to provide a passive condition of both Fe-based alloys. The concentration of Fe at these two sites reached equal values after a holding time of approximately 23 min at 1173 K. In the case of Ni, important concentration changes took place only after a holding time of 20 min with a decrease in the steel and consequent increment in the ribbon alloy. The concentration of P in the SS was negligible for holding periods up to 10 min. For longer periods the concentration increased and remained constant around 3 at% after 20 and 40 min. 3.3. Joints made at 1273 K Joints produced at 1273 K did not showed a defined interface or bonding phase between the stainless steel and the alloy ribbon regardless the holding time as can be seen in Fig. 7. At this processing temperature different phases and precipitate particles formed in the ribbon alloy. These phases and precipitates formed at all 5 holding times. The boundary formed between the AISI 316L steel and the ribbon alloy was irregular for holding times longer than 5 min. The remnant ribbon alloy did not present grain boundaries
which could be an indication that even under that process conditions this alloy keep its amorphous structure. In order to determine the composition of the particles and diffusion of Fe, Cr, Ni, P, and Mo, punctual chemical analyses was carried out at and near the precipitate particles; the results are plotted as a function of the holding time in Fig. 8. It is seen that in the precipitated particles, as the holding time increased, Fe content decreased. Ni concentration in these particles also decreased and after a holding time of 15 min, Ni was not observed. On the other hand, Ni was present in the site tested that was supposed to be the ribbon alloy, which originally did not contain Ni. The Ni concentration found in the site near the precipitated particles is very similar to that for the AISI 316L steel. The P content decreased drastically from ∼13 at.% after 10 min to ∼2 at.% after 15 min which practically disappeared after 20 min and appeared again when the holding time was 40 min having ∼12 at.%. At this joining temperature, Mo diffused from the 316L SS firstly to the alloy ribbon and then to the precipitate particle, registering more concentration after bonding for period longer than 40 min. The non-clearly delimited interface formed between the AISI 316L steel and the alloy ribbon and the formation of precipitate particles in the bonding zone can be the consequence of presence of liquid states that might occur 1273 K. While for the joining process carried out at 1173 K, the SEM analysis revealed a bonding process by diffusion in the solid state because of the well-defined SS/alloy ribbon interfaces. The chemical composition of precipitated particles given in at% in Fig. 8 was converted to wt% and is presented in Table 2. Based on the isotherms at 923 K (650 ◦ C) of the Fe–Ni–Cr (Raynor and Rivlin, 1988) and Fe–Ni–Mo (Wada, 1973) ternary equilibrium diagrams shown in Fig. 9, the phases found in the bonding zone as a function of the holding time are proposed in Table 3. The predominating phase present in all holding times was -phase, which has a tetragonal unit cell as demonstrated (Barcik, 1983;
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Fig. 7. SEM images of the cross-sections of the unions carried out at 1273 K for (a) 5 min, (b) 10 min, (c) 15 min, (d) 20 min and (e) 40 min.
Marshall, 1984). Studies have reported its formation associated with the dissolution of M23 C6 , but also independently (Barcik, 1983). 3.4. Electrochemical behaviour and corrosion resistance Samples bonded at 1173 K presented resistance to localised dissolution in distilled water with no crevice corrosion or pitting attack developed in the bonding region during potentiodynamic polarisation. However, tests in 3.5% NaCl solution showed preferential dissolution in the zone surrounding the bonding region (bonding phase). The dissolution was due to crevice corrosion provoking the complete separation of stainless steel plates. The pores observed in the bonded phase on samples treated at 1173 K at different holding times were not the initiation sites for metal dissolution. The bonding phase suffered preferential dissolution because of Cr depletion with Cr concentration less than 12 at.% as shown in Fig. 5. Immersion test at open circuit potential in 3.5% NaCl solution showed that the preferential dissolution along the bonding phase took place after 10 days.
Samples bonded at 1273 K presented more noble corrosion potential (Ecorr vs SCE) in distilled water than samples joined at 1173 K. The anodic current density was lower for samples joined at 1273 K than for samples at 1173 K in distilled water except for samples bonded for 15 and 20 min, which presented equal and higher current densities respectively. Fig. 10 presents polarisation curves of samples bonded at 1173 and 1273 K immersed in distilled water and in 3.5 wt% NaCl after bonding times of 5, 10, 15, 20 and 40 min. In distilled water, the sample held for 5 min showed a passive region extended from −50 to 200 mV with a current density of 8.4 × 10−2 A/cm2 , which represents very low corrosion rate. 316L SS samples bonded with this amorphous alloy presented excellent corrosion resistance for applications involving non-aggressive electrolytes such as distilled water. For bonding times of 15, 20 and 40 min the samples treated at 1173 and at 1273 K presented similar behaviour in distilled water with an Ecorr around −200 mV and pseudo-passive behaviour which was characterised by a high polarised anodic reaction. This means that the increase of the applied anodic overpotential induced minimum increment of the anodic current density.
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Fig. 8. EDS punctual chemical microanalyses of Fe, Cr, Ni, P, and Mo as a function of the holding times of samples joined at 1273 K. The EDS were performed at 2000× at the points corresponding to the precipitate particles () and “alloy ribbon” (♦) zones of the SEM images of Fig. 7.
Fig. 9. Isotherms at 923 K (650 ◦ C) of ternary equilibrium diagrams phases of (a) Fe–Cr–Ni (Raynor and Rivlin, 1988) and (b) Fe–Cr–Mo (Wada, 1973). Symbols represent the → (Cr, Mo) + . phases: → ␣ + , 䊉 → ␥ + , → , (Cr, Mo) + , and
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Table 2 Chemical composition (wt%) at the precipitate particles formed at 1273 K for the different holding times. Holding time (min) 5 10 15 20 40
Chemical composition wt.% Fe Ni Cr
P
Mo
59.4 58.1 45.7 44.6 27.8
8.97 8.86 7.99 0.38 6.95
– 1.18 1.48 1.15 13.0
9.24 2.69 – – –
22.3 29.1 44.7 53.8 52.1
On the other hand, test conducted in 3.5 wt% NaCl solution showed that samples bonded at 1273 K presented in general more noble values of Ecorr than samples bonded at 1173 K. In this electrolyte, the anodic reaction was under charge transfer control
Table 3 Phases determined at the precipitate particles according to the isotherms at 650 ◦ C of ternary diagram phases of Fig. 9 for the different holding times. Holding time (min)
Phases
Symbol
5 10 15 20 40
␣+ ␥+ (Cr, Mo) + (Cr, Mo) +
䊉 ♦
kinetics with high anodic current densities and a change of mechanism around −200 mV (associated to the increase of the anodic curve slope). This indicated a fast metal dissolution taking place in the bonding zone. Austenitic stainless steel AISI 316L showed a pas-
Fig. 10. Polarisation curves for samples bonded at 1173 and 1273 K during 5, 10, 15, 20 and 40 min in distilled water and 3.5 wt% NaCl solution.
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sive behaviour during anodic polarisation in synthetic seawater in a potential range of about 600 mV as reported by González-Sánchez (2002). According to Pistorius and Burstein (1992) the nucleation of corrosion pits in stainless involves the breakdown of the passive film on a microscopic scale. Burstein et al. (1993) determined that this process induces small and sudden increments of anodic current, which are characterised by current spikes, leading to oxidation and dissolution of less than 0.01 m3 of metal. The nucleation process is unstable and in most cases it will stop by the regeneration of the passive film. Sedriks (1990) indicates the importance of different types of inhomogeneities at the material surface during pit initiation. Lattice defects at the surface, precipitates, intermetallic phases, second phases and non-metallic inclusions are advantageous sites for the nucleation of pitting corrosion due to their different electrochemical behaviour with respect the metal matrix. The presence of a passive film is required for pitting corrosion to take place in stainless steels in chloride containing electrolytes. However in the present case, the bonded samples did not present that passive region during anodic polarisation in the 3.5 wt% NaCl solution. Active dissolution took place instead as can be seen in the polarisation curves in Fig. 10 for samples bonded at the five different holding times. Clearly the anodic branch moved to the right for the case of test in NaCl solution compared to those in distilled water. No pitting corrosion was induced during anodic polarisation in the NaCl solution because this form of attack requires high anodic polarisation to provoke the breakdown of the passive film. As no passive film was present active dissolution took place instead. The bonding phase formed on samples joined at 1173 K with Cr concentration lower than 10 at.% acted as preferential site for metal dissolution. In the case of samples bonded at 1273 K, the dissolution of the bonding zone could be facilitated by the presence of the -phase and other precipitated particles. This induced anodic dissolution with anodic current densities of about 0.6 mA/cm2 . The presence of different types of discontinuities and defects at the bonding zone such as precipitates, intermetallic phases, second phases and non-metallic inclusions are preferential sites for interruption of the passive film as reported by Böhni (1992). As mentioned above, in the bonding zone there were several phases which inhibited the formation of a continuous passive film which induced preferential dissolution in the bonding zone. Rebak et al. (2008) reported that even micropartial devitrification could induce localised corrosion in amorphous Fe base alloys, which may be involved in the bonding process along with the formation of secondary phases. Samples bonded at the temperature of 1273 K showed in general more positive potentials and lower anodic current densities than samples bonded at 1173 K. The presence of -phase could induce the formation of micro-galvanic cells due to the different chemical composition of the phases. In distilled water such galvanic cell does not develop a significant potential gradient but in NaCl solution the micro-galvanic cells became very active. When joined by self-diffusion, the change to the non-crystalline structure of the amorphous alloy and the formation of different phases made the dissimilar nature of the alloys more significant and promoted selective dissolution coupled with crevice corrosion. The lower anodic current density showed by samples bonded at 1273 K compared to that for samples bonded at 1173 K could be due to the higher Cr concentration in the bonding zone reached at 1273 K even the presence of -phase in samples bonded at both temperatures. 4. Conclusions It was possible to produce 316L stainless steel joined samples through difussion bonding using a home-made Fe-based amorphous alloy ribbon as the joining element.
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At 1173 K the bonding process formed SS/alloy ribbon interfaces with clear differences in chemical composition and structure from the base metal and the alloy ribbon which indeced poor corrosion resistance in a chloride containing electrolyte. Samples bonded at 1273 K did not formed a defined bonded phase or zone, instead the ribbon alloy appeared surrounding austenitic grains of the AISI 316L stainles steel making an irregular bonding boundary. This behaviour was more evident as the holding time increased. The joining processes at 1173 K took place by self-diffusion bonding in solid state, whereas at 1273 K liquid state diffusion could take place partially forming some precipitate particles composed mainly of phase. When joined by self-diffusion, the change to the non-crystalline structure of the amorphous alloy and the formation of different phases made the dissimilar nature of the alloys more significant and promoted selective dissolution coupled with crevice corrosion. The bonding process of austenitic stainles steel to itself by using a Fe base amorphous alloy produce aceptable joints with minimal porosity and defects and with acceptable corrosion resistance in non-aggressive electrolytes. However due to the formation of different phases in the bonding zone localised corrosion along the bonding line is prompted in chloride containing electrolytes. Acknowledgements We would like to thank CONACYT, SEP-DGEST México and the CIC of the UMSNH for the financial support. JAV wishes to akcnowledge The University of Sheffield the facilities given to fabricate the Fe-based amorphous alloy ribbon. Also we would like to thank the Centre for Corrosion Research at the Universidad Autónoma de Capeche for the facilities given to JGS to carry out the electrochemical tests. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.jmatprotec.2010.02.015. References Barcik, J., 1983. Crystallographic characteristics of the process of -phase precipitation in chromium-nickel austenitic steels. J. Appl. Crystallogr. 16, 590–598. Böhni, H., 1992. Localised corrosion-mechanisms and methods. Mater. Sci. Forum. 111–112, 401–414. Burstein, G.T., Pistorius, P.C., Mattin, S.P., 1993. The nucleation and growth of corrosion pits on stainless steel. Corros. Sci. 35, 57–62. Davies, H.A., 1983. In: Luborsky, F.E. (Ed.), Amorphous Metallic Alloys. Butterworths, London, UK, pp. 8–25. Flores, J.G., Cervantes, J., Lemus-Ruiz, J., 2006. Joining of silicon nitride to metal (Mo and Ti) using a Cu–foil interlayer. Mater. Sci. Forum. 509, 99–104. González-Sánchez, J., Verduzco, J.A., Lemus-Ruiz, J., Téllez, M.G., Torres, R., 2007. Corrosion resistance of stainless steel joints bonded with a Ni-based amorphous alloy. Anti-Corrosion Methods Mater. 54/2, 68–73. González-Sánchez, J., 2002. PhD Thesis, Materials and Engineering Research Institute, Sheffield Hallam University, U.K. Hashimoto, K., Osada, K., Masumoto, T., Shimodaira, S., 1976. Characteristics of passivity of extremely corrosion resistant amorphous iron alloys. Corros. Sci. 16, 71–76. Jayaraj, J., Kim, Y., Kim, K., Seok, H., Fleury, E., 2005. Corrosion studies on Fe-based amorphous alloys in simulated PEM fuel cell environment. Sci. Technol. Adv. Mater. 6, 282–289. Kimura, H., Sasamori, K., Inoue, A., 1996. High strength Al–Ti–Fe alloys consisting of amorphous and fcc-Al phases prepared by rapid solidification. Mater. Trans. JIM 37, 1722–1725. Lacombe, P., Baroux, B., Beranger, G., 1993. Stainless steels. In: Les editions de physique. Les Ulis, France, pp. 331–336. Marshall, P., 1984. Austenitic stainless steels. In: Microstructure and Mechanical Properties. Elsevier Applied Science Publishers Ltd, Essex, England, pp. 23–64. Miura, E., Kato, H., Ogata, T., Nishiyama, N., Specht, E.D., Shiraihi, T., Inoue, A., Hisatsune, K., 2007. Mechanical property and corrosion resistance evaluations of Ti–6Al–7Nb alloy brazed with bulk metallic glasses. Mater. Trans., JIM 48, 2235–2243. Pistorius, P.C., Burstein, G.T., 1992. Detailed investigation of current transients from metastable pitting events on stainless steel: the transition to stability. Mater. Sci. Forum. 111–112, 429–452.
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