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Sintering and characterization of YAG dispersed ferritic stainless steels
Materials Science and Engineering A 492 (2008) 60–67
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Sintering and characterization of YAG dispersed ferritic stainless steels S.M. Tiwari, S. Balaji, A. Upadhyaya ∗ Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur, UP 208016, India
a r t i c l e
i n f o
Article history: Received 11 July 2007 Received in revised form 28 February 2008 Accepted 29 February 2008 Keywords: Stainless steels Metal matrix composites Polarization Sintering
a b s t r a c t The present study investigates the effect of yttrium aluminium garnet (YAG) addition on the densification, mechanical, tribological and corrosion behaviour of ferritic (434L) stainless steels. The composites were sintered at both solid-state (1200 ◦ C) and supersolidus (1400 ◦ C) sintering conditions. Supersolidus sintering results in superior densification, hardness and corrosion resistance of both straight 434L stainless steel as well as YAG reinforced 434L stainless steels. The addition of YAG to 434L stainless steels at supersolidus sintered conditions improves the strength and wear resistance of 434L stainless steels without significantly degrading the corrosion performance. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Ferritic stainless steels possess a good combination of radiation as well as corrosion resistance. However, the creep resistance of these alloys are inferior to that of the conventional austenitic stainless steels. The ferritic stainless steels with fine oxide dispersoids are reported to have comparable high temperature properties as that of austenitic stainless steels [1]. Powder metallurgy (P/M) processing has emerged as an economical way of fabricating particulate reinforced composites with better homogeneity of second phase distribution [2]. Stainless steels while processed through conventional solid-state sintering, exhibit inferior mechanical and corrosion behaviour as compared to their cast and wrought counterparts due to the presence of the inherent porosities. Besides being stress concentrators, pores also are electrochemically anodic relative to the bulk of the sintered compact. Hence, they provide active sites for corrosion attack [3]. This causes the decreased corrosion resistance for sintered alloys as compared to their wrought counterparts. Baran and Shaw [4] thoroughly studied the corrosion properties of P/M ferritic stainless steels and compared the results to those of wrought steels. They reported lower open circuit potential (OCP), higher current density and lower degree of passivity in P/M alloys relative to their wrought counterparts. The differences were attributed both to a crevice corrosion mechanism within the pore of P/M parts as well as to the formation of secondary compounds, such as chromium carbide during sintering.
∗ Corresponding author. Tel.: +91 512 2597672; fax: +91 512 2597505. E-mail address: [email protected] (A. Upadhyaya). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.02.046
The properties of sintered products are superior in liquid phase sintering than conventional solid-state sintering, as the diffusivity through melt allows for rapid densification. Recently, the use of prealloyed powder has led to the novel process in liquid phase sintering, called supersolidus liquid phase sintering (SLPS). SLPS involves sintering of a prealloyed powder at a temperature between the solidus and the liquidus, which results from the partial melting of the solid phase [5,6]. Depending on the powder internal microstructure, the liquid phase can nucleate either along the grain boundaries, or at the inter-particle contacts. Once a critical fraction of the grains are covered by the liquid phase, the particles lose their rigidity and densification occurs by secondary rearrangement, followed by solution reprecipitation, pore removal and contact flattening or grain shape accommodation. Superior creep rupture strength of oxide dispersion strengthened ferritic steel was reported by Ukai et al. [1] due to finely distributed Y2 O3 particle in the ferrite matrix. The effect of Al2 O3 dispersion on the sintering behaviour and mechanical properties of solid-state sintered 434L ferritic stainless steel was investigated in detail by Mukherjee and Upadhyaya [7]. They reported that the composites possessing 4–6 vol.% of Al2 O3 result in improved mechanical properties. The authors [7] conducted potentiodynamic polarization studies on the samples in 1N sulphuric acid and showed that the sample containing 8 vol.% of Al2 O3 had better corrosion resistance. This was attributed to the interaction between Cr2 O3 and Al2 O3 in the formation of the protective oxide. Shankar et al. [8,9] investigated the effect of sintering temperature and yttria addition on electrochemical behaviour of P/M ferritic and austenitic stainless steel in acidic and chloride environment through polarization methods. They showed that the corrosion behaviour of yttria-
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Table 1 Composition of the as-received 434L and YAG powders Powder
Cr
Mo
Si
Mn
C
S
P
Fe
Y
Al
Ti
O
434L YAG
17 –
1.0 –
0.71 0.014
0.2 –
0.023 0.08
0.02 –
0.02 –
Balance 0.014
– 45.05
– 22.5
– <0.01
– Balance
dispersed stainless steel is comparable that of the corresponding unreinforced stainless steels. Recently, Balaji et al. [10] reported the beneficial effects of Ni3 Al and Fe3 Al reinforcements on the densification and corrosion behaviour of ferritic 434L stainless steels. The beneficial effect of YAG addition on the densification and hardness of the ferritic and austenitic stainless steel is previously reported by Jain et al. [11]. Recently, Balaji et al. [10] have examined in detail the electrochemical response of the YAG added 316L stainless steels. As compared to austenitic steels, ferritic stainless steels are much cheaper. This study, therefore, investigates the effect of YAG addition and sintering temperature on the mechanical, tribological and corrosion behaviour of P/M ferritic 434L stainless steels. 2. Experimental procedure The chemical composition and powder characteristics of the prealloyed ferritic 434L stainless steel (Ametek speciality metal products, USA) and the second phase yttrium aluminium garnet (Treibacher Auermet Produktionsge, Austria) powders are summarized in Tables 1 and 2. The 434L powder was mixed with required proportions of YAG powder (5 and 10 wt.%) in a mortar and pestle for 30 min. The powders were uniaxially compacted at a pressure of 600 MPa. Zinc stearate was used as die wall lubricant. The sintering was carried out in a MoSi2 heated horizontal tubular furnace (model: OKAY 70T-7, supplier: Bysakh, Kolkata, India) at a heating rate of 5 ◦ C/min in pure hydrogen atmosphere (dew point: −35 ◦ C). The as-pressed compacts were held isothermally for 1 h at 1200 ◦ C in the case of solid-state sintering and at 1400 ◦ C for supersolidus sintering. The sintered densities were determined from dimensional and weight measurements. To take into account the influence of the initial as-pressed density, the compact sinterability was also expressed in terms of densification parameter which is calculated as follows: Densification parameter =
sintered density − green density
theoretical density − green density (1)
The theoretical densities of the composites were calculated using inverse rule of mixture. The mechanical properties were measured on the sintered flat tensile bars produced as per MPIF standard 10. Vickers bulk hardness measurements were performed on all the microstructural surfaces at a 2 kg load using automated
and pre-programmed hardness tester supplied by Akashi Corporation, Japan. The observed hardness values are averages of five readings taken at random locations throughout the sample. The dry sliding wear behaviour of the sintered composites was assessed using a pin-on-disc type wear-testing machine (TR-20, DUCOM, Bangalore, India) against a hardened steel disc (EN31) at a sliding velocity of 2 m s−1 under 20 N normal load. Prior to the electrochemical tests, the samples were polished to mirror finish and ultrasonically degreased in acetone. The potentiodynamic polarization scans were conducted using a flat cell with an aqueous 0.1N H2 SO4 solution at room temperature. A standard three-electrode (reference, counter and working electrode) technique was used for the measurement. The reference electrode used for the present experiment was an Ag/AgCl (saturated with KCl) electrode and the alloy under test was the working electrode. A platinum mesh acted as a counter electrode. The electrochemical experiments were performed using a DC corrosion measurement system (PC4 Potentiostat, Gamry Instruments, Inc., USA). A delay of 3600 s was given prior to the polarization test in order to allow for stabilization of open circuit potential. The potentiodynamic scan was carried out with an initial potential of −250 mV vs OCP till 2 V vs OCP for all samples at a scan rate of 0.1667 mV s−1 . The Tafel slopes were established from the active region of the corresponding anodic and cathodic curves. The critical parameters like corrosion potential (Ecorr ), corrosion current (Icorr ), anodic (ˇa ) and cathodic (ˇc ) slopes and corrosion rate, primary passivation potential (Epp ), active peak current density (Icrit ), passive state current density (Ip ), break down potential (Ebp ), zero current potential (Ezcp ) were evaluated from the polarization curve. The susceptibility of the stainless steel matrix composites to localized corrosion was electrochemically measured through cyclic polarization measurements in aqueous 0.1N HCl, using a flat cell. Duplicate specimens were used in each test to verify the reproducibility of the result. The OCP was monitored as a function of time with respect to Ag/AgCl electrode until the potential of the sample reached a stable value. After the sample attained a constant potential, cyclic polarisation was started with an initial potential below 200 mV of the corrosion potential at a scan rate of 1 mV s−1 . The sweep direction was reserved at potential above the pitting potential after reaching an anodic current density of 100 mA/cm2 until the reverse scan reached the passive region to obtain the protection potential. The protection potential (Eprot ) is the potential at which the current on the reverse scan returns to a low value and has been taken as the potential at which the reverse scan intersects the forward scan.
Table 2 Characteristics of the as-received powders Property
3. Results and discussion
Powder 434L
YAG
Processing technique Powder shape
Gas atomization Spherical
Chemical reduction Rounded
Cumulative powder size, m D10 D50 D90
8.5 35.3 75.1
0.5 1.5 2.1
Apparent density (g/cm3 ) Flow rate, (s/50 g) Theoretical density (g/cm3 )
2.6 28 7.86
0.70 98 4.50
3.1. Densification response Fig. 1a compares the sintered density variation of 434L stainless steels and 434L–YAG composites with sintering temperature. Supersolidus sintering at 1400 ◦ C results in significant improvement in sintered density as compared to the compacts sintered at 1200 ◦ C. Elsewhere, Shankar et al. [8,9] and Balaji et al. [10] too observed a similar kind of improvement for the particulate reinforced stainless steel composites. This can be attributed to the liquid formation that aid in faster diffusion kinetics at the super-
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Fig. 1. Effect of sintering temperature and YAG addition on the (a) sintered density and (b) densification parameter of ferritic 434L stainless steel.
solidus sintering conditions. Besides enhanced diffusion kinetics, melt formation along the grain boundaries result in densification enhancement through capillary induced pore-filling and grain rearrangement [6]. The addition of YAG to 434L stainless steels did not significantly degraded the sintered density. In fact 5 wt.% YAG added 434L composites showed slightly higher sintered density than straight 434L stainless steels in both solid-state as well as supersolidus sintered conditions. This can be attributed to the grain inhibition due to the addition of the YAG dispersoids, which in turn, will increase the propensity for grain boundary added diffusion and thereby result in slightly greater densification in 5 wt.% YAG added stainless steel compacts. Elsewhere, Jain et al. [11] have attributed the improvement in sintered density with optimal amount of YAG additions to the chemical interaction between the YAG and stainless steel matrix. This has been proved in a previous report from our group [11] whereby EPMA analysis was performed at the grain boundary and in intragranular regions of sintered 434L–YAG compacts. At lower YAG content there is chemical interaction with the matrix phase, this is confirmed by EPMA analysis of samples. In case of higher YAG composites (10% YAG), because of fine particle size of YAG there is probability of YAG–YAG interaction and formation of its agglomerates at grain boundaries and thereby reducing metal–YAG chemical interaction this causes a decrease in sintered density. The variation of densification parameter with sintering temperature and YAG content is shown in Fig. 1b. Densification parameter also follows a similar trend as sintered density with respect to sintering temperature and YAG content. Fig. 2 shows the variation in axial and radial shrinkage after sintering of 434L stainless steels and 434L–YAG composites. In Fig. 2, the line drawn along the diagonal represents the isotropic shrinkage line in which the axial and radial shrinkage values are equal (slope = 1). As expected the supersolidus sintered compacts showed higher shrinkage than that of the solid-state sintered compacts due to their higher densification parameter. Despite their higher shrinkage values, the supersolidus sintered straight 434L stainless steels and 434L–5YAG composites lie close to the isotropic shrinkage line. The optical and scanning electron micrographs of the of the straight 434L stainless steel and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C are shown in Figs. 3 and 4, respectively. There is a significant difference in pore morphology and porosity content between the solid-state sintered and supersolidus sintered compacts. The solid-state sintered compacts revealed larger irreg-
ular pores predominantly along the grain boundaries, whereas, in the case of supersolidus sintered compacts the pores were well rounded and predominantly intragranular in nature. However, decrease in the amount of porosity with increase in sintering temperature is evident from the microstructures. The average grain size of the supersolidus sintered 434L stainless steel is larger than that of solid-state sintered one. It is evident from the microstructure that the addition of YAG to 434L stainless steels restricts the grain coarsening and retains the intergranular pores at higher sintering temperature. The scanning electron micrographs shown in Fig. 4 clearly indicate the YAG particles along the grain boundaries with bright contrast. At lower amount of YAG addition, the YAG particles are better linked with the matrix stainless steels. Whereas, a significant agglomeration and segregation of the YAG particles along the grain boundaries were observed for the 434L–10YAG composites. 3.2. Mechanical properties Fig. 5 shows the bulk hardness variation of 434L stainless steels and 434L–YAG composites with sintering temperature. The hardness of both 434L stainless steel as well as 434L–YAG composites
Fig. 2. Effect of sintering temperature and YAG addition on the axial and radial shrinkages in the ferritic stainless steel compacts.
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Fig. 3. Optical micrographs of straight 434L stainless steels and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C.
Fig. 4. SEM micrographs of straight 434L stainless steels and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C.
sintered at 1400 ◦ C were higher than that of solid-state sintered compacts. This can be attributed to the higher sintered density of the supersolidus sintered compacts. The influence of YAG addition on hardness is quite evident at the supersolidus sintered conditions. The reason for the increase in hardness is due to blocking of moving dislocations by the YAG dispersoids as well as inherent hardness of YAG dispersoids. Table 3 shows the tensile testing results of straight 434L stainless steels and 434L–YAG composites sintered at 1400 ◦ C. Addition of Table 3 Tensile properties of the supersolidus sintered straight 434L and 434L–YAG composites
Fig. 5. Effect of sintering temperature and YAG addition on the bulk hardness of the ferritic stainless steel compacts.
Composition (wt.%)
Y.S. (MPa)
U.T.S. (MPa)
% Elongation
434L 434L–5YAG 434L–10YAG
235 265 147
378 411 167
29 13 2
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Fig. 6. Fractographs of the tensile tested surfaces of the supersolidus sintered straight 434L and 434L–YAG composites.
Fig. 6 shows the fractographs of the samples sintered at 1400 ◦ C. In case of the 434L–YAG composites, the brighter phase is YAG particles. From Fig. 6, it is evident that the straight 434L stainless steel shows distinct dimpled morphology, which is a typical characteristic of ductile failure. Increase in YAG content, changes the fracture mode from dimpled to intergranular mode, thereby resulting in lowering of ductility. This is attributed to lowering of sinterability with YAG addition and the presence of brittle additive phase at the 434L inter-particle interface. 3.3. Tribological response
Fig. 7. Variation of wear (m) with sliding distance for supersolidus sintered straight 434L and 434L–YAG composites.
YAG up to 5 wt.% into straight 434L ferritic stainless steel enhances both the yield and tensile strength. It is due to the combined effect of increase in sintered density and restriction of grain growth due to YAG addition. Higher amount (10 wt.%) of YAG addition reduces both tensile and yield strength drastically due to lower sintered density of 434L–10YAG composites and lesser interactions of dispersoids with the matrix phase.
Fig. 7 shows the effect of YAG dispersoids on the wear response of 434L ferritic stainless steel. The wear (m) vs sliding distance curves clearly indicate the better wear resistance of the 434L–5YAG composites as compared to the straight 434L stainless steels. This can be attributed to the higher hardness, better densification and restricted grain coarsening associated with the supersolidus sintered 434L–5YAG composites. Higher amount (10 wt.%) of YAG addition leads to higher wear due to the inferior densification. 3.4. Electrochemical response 3.4.1. OCP vs time The variation of open circuit potential with time for the sintered ferritic 434L stainless steel and 434L–YAG composites in 0.1N H2 SO4 and 0.1N HCl solution are shown in Fig. 8a and b, respectively. In 0.1N H2 SO4 solution the stabilization was achieved quickly for all
Fig. 8. OCP stabilization curves for the sintered 434L and 434L–YAG composites in (a) 0.1N H2 SO4 and (b) 0.1N HCl.
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Fig. 9. Potentiodynamic polarization curves for the straight 434L and 434L–YAG composites sintered at (a) 1200 ◦ C and (b) 1400 ◦ C in 0.1N H2 SO4 solution. Table 4 Passivity parameters obtained from the potentiodynamic polarization curves of 434L and 434L–YAG composites Composition
Sintering temperature (◦ C)
Icrit (A/cm2 )
Epp (mV)
Ipass (A/cm2 )
Ebp (mV)
434L
1200 1400 1200 1400 1200 1400
359 906 3402 340 4394 349
−429 −419 −313 −450 −271 −440
80 49 32 53 599 69
901 927 869 859 975 901
434L–5YAG 434L–10YAG
the samples, with the exception of solid-state sintered 434L–YAG composites which took relatively longer time for stabilization. All the samples get stabilized at a constant OCP (−520 ± 5 mV) in aqueous 0.1N H2 SO4 solution. In the case of 0.1N HCl there is significant difference in the stabilization voltages between the different composites. This variation in OCP value can be attributed to the different degree of attack by the aggressive Cl− ions on the various com-
posites. However, in both 0.1N H2 SO4 and 0.1N HCl solutions a continuous steady shift of the open circuit potential towards more cathodic value with time is evident. This trend could be related to selective anodic dissolution that occurs at the active surface of the composites. A similar trend in the free corrosion potential for sintered austenitic and ferritic stainless steel has been reported in the literature [3,4,13,14].
Fig. 10. Cyclic polarization curves for the straight 434L and 434L–YAG composites sintered at (a) 1200 ◦ C and (b) 1400 ◦ C in 0.1N HCl solution.
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Table 5 Tafel extrapolation data for 434L and 434L–YAG composites Composition
Sintering temperature (◦ C)
ˇa (mV/decade)
434L
1200 1400 1200 1400 1200 1400
166.8 112.8 274.0 267.7 1050 237.7
434L–5YAG 434L–10YAG
ˇc (mV/decade) 258.2 295.7 306.4 546.6 1192 432
Icorr (A/cm2 )
Ecorr (mV)
Corrosion rate (mmpy)
325 327 801 513 4570 409
−485 −502 −522 −526 −523 −522
3.594 3.234 8.875 5.103 50.64 4.151
Table 6 Pitting parameters obtained from cyclic polarization of 434L and 434L–YAG composites Composition
Sintering temperature (◦ C)
Ezcp (mV)
Ep (mV)
434L
1200 1400 1200 1400 1200 1400
−493 −525 −448 −472 −339 −336
−414 −134 −205 −42 −82 −156
434L–5YAG 434L–10YAG
3.4.2. Potentiodynamic polarization The potentiodynamic polarization curves, presented in Fig. 9a and b reveals the variation in current density as a function of applied voltage for pure 434L and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C, respectively. Prominent active–passive behaviour is shown by all the sintered components. The behaviour of the supersolidus sintered ferritic steel–YAG composites in the H2 SO4 environment is comparable to that of the sintered ferritic steel. The passivation data obtained from the polarization curve are presented in Table 4. The active peak current density (Icrit ) and the primary passivation potential (Epp ) do not show a definite trend. However, a significant difference between the critical current density of the solid-state sintered 434L–YAG composites and rest of the samples is evident from Table 4. This can be attributed to the heterogeneity produced by the YAG particle and the poor interaction of YAG with the matrix stainless steels. The passivation current density is similar for the ferritic stainless steels and its composites with the exception of solid-state sintered 434L–10YAG composites which showed an order of magnitude higher Ipass value. The potential for breakdown of the passivity is similar for all the samples and is around 900 mV except for the 434L–10YAG sample, which is 950 mV. The corrosion rate measurements were done using Tafel extrapolation and the data are given in Table 5. A low corrosion potential (Ecorr ) corresponds to inferior corrosion resistance
Eprot (mV) −408 −271 −356 −267 – –
Ep − Eprot (mV) 85 137 151 225 – –
[15,16]. The corrosion potential for the pure ferritic stainless steel is slightly higher than that of the respective composites. This suggests a slightly better corrosion resistance for the straight 434L stainless steels as compared to their YAG added counterparts. Shankar et al. [9] also reported similar corrosion potential values for the 434L–yttria composites processed at similar conditions. There is no significant difference in the Icorr values between the samples except the solid-state sintered 434L–10YAG composites. The corrosion rate too follows a similar trend as Icorr . The supersolidus sintered YAG added 434L composites exhibited comparable corrosion rates as that of unreinforced 434L stainless steels. Similar observations with corrosion rate of ferritic stainless steel–yttria composites have been reported by Shankar et al. [9]. The behaviour can be attributed to the higher sintered density, reduced grain boundary area due to grain coarsening, and more homogeneous microstructures achieved with the supersolidus liquid phase sintering. The literature on P/M stainless steels in acid solutions also reveals that corrosion resistance improves with increasing sintered density [17–19]. Due to the interaction between the YAG and 434L matrix at higher sintering temperatures, the YAG particles are better linked with the matrix leading to better passivation characteristics as compared to that of the solid-state sintered composites. The poor densification and presence of increased amount of reinforcements had locally weakened the passive layer of the solid-state
Fig. 11. SEM micrographs of the cyclic polarization tested straight 434L stainless steels and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C.
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sintered 434L–10YAG composites that lead to the poor passivation and corrosion performance. 3.4.3. Cyclic polarization The cyclic polarization curves for the 434L stainless steels and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C are shown in Fig. 10a and b, respectively, and the pitting corrosion data are tabulated in Table 6. The zero corrosion potential shows a shift towards the noble potential region with YAG dispersion. Pitting potential for 434L–YAG composites is nobler compared to the unreinforced 434L stainless steels. For straight 434L and 434L–5YAG composites, supersolidus sintering results in higher potential for pitting as compared to the solid-state sintering. However, at higher YAG content the behaviour is different. The protection potential also exhibits a similar trend. The protection potential for 434L–10YAG composite could not be obtained because the reverse scan was stopped after reaching the open circuit potential. These alloys are prone to pitting even at the open circuit potential in aqueous 0.1N HCl solution. The steel surface after exposure to chloride solution is observed in the scanning electron microscope and shown in Fig. 11. The attack is very severe in the chloride environment. Extensive metal dissolution has been observed in the solid sate sintered 434L stainless steel. The area of hysteresis loop is a direct measure of the pit propagation kinetics and the extent of localized attack depends on the nature of hysteresis loop [15,16]. It can be observed that the YAG dispersion in the ferritic steel has increased the pit propagation kinetics as can be confirmed from the increase in the hysteresis loop area during the cyclic scan. The dominating corrosion mechanism in these samples is crevice corrosion under the chloride-containing environment. At higher dispersoids content the behaviour is different. The microstructure reveals the segregation of YAG along the grain boundary region. This could be one of the possible reasons for the poor behaviour of the composite during the cyclic test. The mechanism proposed by Otero et al. [3] to justify the high corrosion kinetics observed in the P/M products is based on attack through crevice corrosion, which appears macroscopically as generalized attack. This starts in the pores and becomes more widespread. SEM micrograph shown in Fig. 11 confirms the localized nature of the attack which takes place through crevice corrosion. In reinforced materials the Cl− ion attack was localized on some specific areas preferably at the interface between the matrix and the reinforcements and the bulk material beneath dissolved. Variation in the reinforcement volume fraction has a marginal effect on the susceptibility of the composites to localized corrosion. The supersolidus sintered alloys show slightly lower corrosion susceptibility with relatively smaller hysteresis loop area in the chloride solution. This could be related to the lower porosity level achieved by supersolidus sintered alloys. The pores act as pre-existing crevices on
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the P/M surface [4]. Crevice corrosion situation can result from the consumption of dissolved oxygen within a pore. The oxygen deficiency and higher chloride concentration (due to influx of chloride ions in to the pore) associated with acid hydrolysis reaction with in the pore lead to localized dissolution of metal inside the pore. 4. Conclusions Supersolidus liquid phase sintering (SLPS) results in higher density in both ferritic 434L stainless steel and 434L–YAG composites. Addition of small amount (5 wt.%) of YAG to both the 434L stainless steel increases the densification of the composites. SLPS results in microstructural coarsening and intragranular pores with rounded morphology. However, the addition of fine yttrium aluminium garnet (YAG) powders to 434L restricts microstructural coarsening and the pores remain intergranular. The addition of YAG to 434L stainless steels improves the mechanical and tribological behaviour at supersolidus sintering conditions. Corrosion resistance of supersolidus sintered 434L and 434L–YAG composites is better as compared to that of solid-state sintered compacts in both 0.1N HCl and 0.1N H2 SO4 solutions. Corrosion behaviour of supersolidus sintered 434L–5YAG composites is comparable to that of sintered straight 434L compacts. Pitting corrosion behaviour (in 0.1N HCl solution) of supersolidus sintered compacts is better as compared to those processed through solid-state sintering. References [1] S. Ukai, S. Mizuta, T. Yoshitake, T. Okuda, M. Fujiwara, S. Hagi, T. Kobayashi, J. Nucl. Mater. 283–287 (2000) 702–706. [2] R.M. German, Powder Metallurgy Science, MPIF, Princeton, NJ, USA, 1994. [3] E. Otero, A. Pardo, E. Saenz, M.V. Utrilla, F.J. Perez, Corros. Sci. 40 (1998) 1421–1434. [4] M.C. Baran, B.A. Shaw, Int. J. Powder Metall. 36 (2000) 57–68. [5] L. Cambal, J.A. Lund, Int. J. Powder Metall. 8 (1972) 131–140. [6] R.M. German, Metall. Mater. Trans. A 28A (1997) 1553–1567. [7] S.K. Mukherjee, G.S. Upadhyaya, Int. J. Powder Metall. 19 (1983) 289. [8] J. Shankar, A. Upadhyaya, R. Balasubramaniam, Proceedings of the Advances in Powder Metallurgy Particulate Materials, vol. 13, MPIF, Princeton, NJ, USA, 2002, pp. 313–322. [9] J. Shankar, A. Upadhyaya, R. Balasubramaniam, Corros. Sci. 46 (2004) 487–498. [10] S. Balaji, G. Joshi, A. Upadhyaya, Scripta Mater. 56 (2007) 149–151. [11] J. Jain, A.M. Kar, A. Upadhyaya, Mater. Lett. 58 (2004) 2037–2040. [13] E. Otero, A. Pardo, E. Saenz, M.V. Utrilla, F.J. Perez, Mater. Charact. 35 (1995) 145–151. [14] E. Otero, A. Pardo, E. Saenz, M.V. Utrilla, F.J. Perez, Can. Metall. Quart. 36 (1997) 65–72. [15] A.J. Sedriks, Corrosion of Stainless Steels, John Wiley & Sons Inc., New York, NY, USA, 1996. [16] E.E. Stansbury, R.A. Buchanan, Fundamentals of Electrochemical Corrosion, ASM International, Ohio, USA, 2000. [17] G.H. Lei, R.M. German, H.S. Nayar, Powder Metall. Int. 15 (1983) 70–76. [18] E. Ahlberg, P. Engdahl, R. Johansson, Proceedings of the World Conference on Powder Metallurgy, vol. 1, London, 1990, pp. 419–432. [19] D.R. Gabe, Powder Metall. 4 (1977) 227–231.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Sintering and characterization of YAG dispersed ferritic stainless steels S.M. Tiwari, S. Balaji, A. Upadhyaya ∗ Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur, UP 208016, India
a r t i c l e
i n f o
Article history: Received 11 July 2007 Received in revised form 28 February 2008 Accepted 29 February 2008 Keywords: Stainless steels Metal matrix composites Polarization Sintering
a b s t r a c t The present study investigates the effect of yttrium aluminium garnet (YAG) addition on the densification, mechanical, tribological and corrosion behaviour of ferritic (434L) stainless steels. The composites were sintered at both solid-state (1200 ◦ C) and supersolidus (1400 ◦ C) sintering conditions. Supersolidus sintering results in superior densification, hardness and corrosion resistance of both straight 434L stainless steel as well as YAG reinforced 434L stainless steels. The addition of YAG to 434L stainless steels at supersolidus sintered conditions improves the strength and wear resistance of 434L stainless steels without significantly degrading the corrosion performance. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Ferritic stainless steels possess a good combination of radiation as well as corrosion resistance. However, the creep resistance of these alloys are inferior to that of the conventional austenitic stainless steels. The ferritic stainless steels with fine oxide dispersoids are reported to have comparable high temperature properties as that of austenitic stainless steels [1]. Powder metallurgy (P/M) processing has emerged as an economical way of fabricating particulate reinforced composites with better homogeneity of second phase distribution [2]. Stainless steels while processed through conventional solid-state sintering, exhibit inferior mechanical and corrosion behaviour as compared to their cast and wrought counterparts due to the presence of the inherent porosities. Besides being stress concentrators, pores also are electrochemically anodic relative to the bulk of the sintered compact. Hence, they provide active sites for corrosion attack [3]. This causes the decreased corrosion resistance for sintered alloys as compared to their wrought counterparts. Baran and Shaw [4] thoroughly studied the corrosion properties of P/M ferritic stainless steels and compared the results to those of wrought steels. They reported lower open circuit potential (OCP), higher current density and lower degree of passivity in P/M alloys relative to their wrought counterparts. The differences were attributed both to a crevice corrosion mechanism within the pore of P/M parts as well as to the formation of secondary compounds, such as chromium carbide during sintering.
∗ Corresponding author. Tel.: +91 512 2597672; fax: +91 512 2597505. E-mail address: [email protected] (A. Upadhyaya). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.02.046
The properties of sintered products are superior in liquid phase sintering than conventional solid-state sintering, as the diffusivity through melt allows for rapid densification. Recently, the use of prealloyed powder has led to the novel process in liquid phase sintering, called supersolidus liquid phase sintering (SLPS). SLPS involves sintering of a prealloyed powder at a temperature between the solidus and the liquidus, which results from the partial melting of the solid phase [5,6]. Depending on the powder internal microstructure, the liquid phase can nucleate either along the grain boundaries, or at the inter-particle contacts. Once a critical fraction of the grains are covered by the liquid phase, the particles lose their rigidity and densification occurs by secondary rearrangement, followed by solution reprecipitation, pore removal and contact flattening or grain shape accommodation. Superior creep rupture strength of oxide dispersion strengthened ferritic steel was reported by Ukai et al. [1] due to finely distributed Y2 O3 particle in the ferrite matrix. The effect of Al2 O3 dispersion on the sintering behaviour and mechanical properties of solid-state sintered 434L ferritic stainless steel was investigated in detail by Mukherjee and Upadhyaya [7]. They reported that the composites possessing 4–6 vol.% of Al2 O3 result in improved mechanical properties. The authors [7] conducted potentiodynamic polarization studies on the samples in 1N sulphuric acid and showed that the sample containing 8 vol.% of Al2 O3 had better corrosion resistance. This was attributed to the interaction between Cr2 O3 and Al2 O3 in the formation of the protective oxide. Shankar et al. [8,9] investigated the effect of sintering temperature and yttria addition on electrochemical behaviour of P/M ferritic and austenitic stainless steel in acidic and chloride environment through polarization methods. They showed that the corrosion behaviour of yttria-
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Table 1 Composition of the as-received 434L and YAG powders Powder
Cr
Mo
Si
Mn
C
S
P
Fe
Y
Al
Ti
O
434L YAG
17 –
1.0 –
0.71 0.014
0.2 –
0.023 0.08
0.02 –
0.02 –
Balance 0.014
– 45.05
– 22.5
– <0.01
– Balance
dispersed stainless steel is comparable that of the corresponding unreinforced stainless steels. Recently, Balaji et al. [10] reported the beneficial effects of Ni3 Al and Fe3 Al reinforcements on the densification and corrosion behaviour of ferritic 434L stainless steels. The beneficial effect of YAG addition on the densification and hardness of the ferritic and austenitic stainless steel is previously reported by Jain et al. [11]. Recently, Balaji et al. [10] have examined in detail the electrochemical response of the YAG added 316L stainless steels. As compared to austenitic steels, ferritic stainless steels are much cheaper. This study, therefore, investigates the effect of YAG addition and sintering temperature on the mechanical, tribological and corrosion behaviour of P/M ferritic 434L stainless steels. 2. Experimental procedure The chemical composition and powder characteristics of the prealloyed ferritic 434L stainless steel (Ametek speciality metal products, USA) and the second phase yttrium aluminium garnet (Treibacher Auermet Produktionsge, Austria) powders are summarized in Tables 1 and 2. The 434L powder was mixed with required proportions of YAG powder (5 and 10 wt.%) in a mortar and pestle for 30 min. The powders were uniaxially compacted at a pressure of 600 MPa. Zinc stearate was used as die wall lubricant. The sintering was carried out in a MoSi2 heated horizontal tubular furnace (model: OKAY 70T-7, supplier: Bysakh, Kolkata, India) at a heating rate of 5 ◦ C/min in pure hydrogen atmosphere (dew point: −35 ◦ C). The as-pressed compacts were held isothermally for 1 h at 1200 ◦ C in the case of solid-state sintering and at 1400 ◦ C for supersolidus sintering. The sintered densities were determined from dimensional and weight measurements. To take into account the influence of the initial as-pressed density, the compact sinterability was also expressed in terms of densification parameter which is calculated as follows: Densification parameter =
sintered density − green density
theoretical density − green density (1)
The theoretical densities of the composites were calculated using inverse rule of mixture. The mechanical properties were measured on the sintered flat tensile bars produced as per MPIF standard 10. Vickers bulk hardness measurements were performed on all the microstructural surfaces at a 2 kg load using automated
and pre-programmed hardness tester supplied by Akashi Corporation, Japan. The observed hardness values are averages of five readings taken at random locations throughout the sample. The dry sliding wear behaviour of the sintered composites was assessed using a pin-on-disc type wear-testing machine (TR-20, DUCOM, Bangalore, India) against a hardened steel disc (EN31) at a sliding velocity of 2 m s−1 under 20 N normal load. Prior to the electrochemical tests, the samples were polished to mirror finish and ultrasonically degreased in acetone. The potentiodynamic polarization scans were conducted using a flat cell with an aqueous 0.1N H2 SO4 solution at room temperature. A standard three-electrode (reference, counter and working electrode) technique was used for the measurement. The reference electrode used for the present experiment was an Ag/AgCl (saturated with KCl) electrode and the alloy under test was the working electrode. A platinum mesh acted as a counter electrode. The electrochemical experiments were performed using a DC corrosion measurement system (PC4 Potentiostat, Gamry Instruments, Inc., USA). A delay of 3600 s was given prior to the polarization test in order to allow for stabilization of open circuit potential. The potentiodynamic scan was carried out with an initial potential of −250 mV vs OCP till 2 V vs OCP for all samples at a scan rate of 0.1667 mV s−1 . The Tafel slopes were established from the active region of the corresponding anodic and cathodic curves. The critical parameters like corrosion potential (Ecorr ), corrosion current (Icorr ), anodic (ˇa ) and cathodic (ˇc ) slopes and corrosion rate, primary passivation potential (Epp ), active peak current density (Icrit ), passive state current density (Ip ), break down potential (Ebp ), zero current potential (Ezcp ) were evaluated from the polarization curve. The susceptibility of the stainless steel matrix composites to localized corrosion was electrochemically measured through cyclic polarization measurements in aqueous 0.1N HCl, using a flat cell. Duplicate specimens were used in each test to verify the reproducibility of the result. The OCP was monitored as a function of time with respect to Ag/AgCl electrode until the potential of the sample reached a stable value. After the sample attained a constant potential, cyclic polarisation was started with an initial potential below 200 mV of the corrosion potential at a scan rate of 1 mV s−1 . The sweep direction was reserved at potential above the pitting potential after reaching an anodic current density of 100 mA/cm2 until the reverse scan reached the passive region to obtain the protection potential. The protection potential (Eprot ) is the potential at which the current on the reverse scan returns to a low value and has been taken as the potential at which the reverse scan intersects the forward scan.
Table 2 Characteristics of the as-received powders Property
3. Results and discussion
Powder 434L
YAG
Processing technique Powder shape
Gas atomization Spherical
Chemical reduction Rounded
Cumulative powder size, m D10 D50 D90
8.5 35.3 75.1
0.5 1.5 2.1
Apparent density (g/cm3 ) Flow rate, (s/50 g) Theoretical density (g/cm3 )
2.6 28 7.86
0.70 98 4.50
3.1. Densification response Fig. 1a compares the sintered density variation of 434L stainless steels and 434L–YAG composites with sintering temperature. Supersolidus sintering at 1400 ◦ C results in significant improvement in sintered density as compared to the compacts sintered at 1200 ◦ C. Elsewhere, Shankar et al. [8,9] and Balaji et al. [10] too observed a similar kind of improvement for the particulate reinforced stainless steel composites. This can be attributed to the liquid formation that aid in faster diffusion kinetics at the super-
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Fig. 1. Effect of sintering temperature and YAG addition on the (a) sintered density and (b) densification parameter of ferritic 434L stainless steel.
solidus sintering conditions. Besides enhanced diffusion kinetics, melt formation along the grain boundaries result in densification enhancement through capillary induced pore-filling and grain rearrangement [6]. The addition of YAG to 434L stainless steels did not significantly degraded the sintered density. In fact 5 wt.% YAG added 434L composites showed slightly higher sintered density than straight 434L stainless steels in both solid-state as well as supersolidus sintered conditions. This can be attributed to the grain inhibition due to the addition of the YAG dispersoids, which in turn, will increase the propensity for grain boundary added diffusion and thereby result in slightly greater densification in 5 wt.% YAG added stainless steel compacts. Elsewhere, Jain et al. [11] have attributed the improvement in sintered density with optimal amount of YAG additions to the chemical interaction between the YAG and stainless steel matrix. This has been proved in a previous report from our group [11] whereby EPMA analysis was performed at the grain boundary and in intragranular regions of sintered 434L–YAG compacts. At lower YAG content there is chemical interaction with the matrix phase, this is confirmed by EPMA analysis of samples. In case of higher YAG composites (10% YAG), because of fine particle size of YAG there is probability of YAG–YAG interaction and formation of its agglomerates at grain boundaries and thereby reducing metal–YAG chemical interaction this causes a decrease in sintered density. The variation of densification parameter with sintering temperature and YAG content is shown in Fig. 1b. Densification parameter also follows a similar trend as sintered density with respect to sintering temperature and YAG content. Fig. 2 shows the variation in axial and radial shrinkage after sintering of 434L stainless steels and 434L–YAG composites. In Fig. 2, the line drawn along the diagonal represents the isotropic shrinkage line in which the axial and radial shrinkage values are equal (slope = 1). As expected the supersolidus sintered compacts showed higher shrinkage than that of the solid-state sintered compacts due to their higher densification parameter. Despite their higher shrinkage values, the supersolidus sintered straight 434L stainless steels and 434L–5YAG composites lie close to the isotropic shrinkage line. The optical and scanning electron micrographs of the of the straight 434L stainless steel and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C are shown in Figs. 3 and 4, respectively. There is a significant difference in pore morphology and porosity content between the solid-state sintered and supersolidus sintered compacts. The solid-state sintered compacts revealed larger irreg-
ular pores predominantly along the grain boundaries, whereas, in the case of supersolidus sintered compacts the pores were well rounded and predominantly intragranular in nature. However, decrease in the amount of porosity with increase in sintering temperature is evident from the microstructures. The average grain size of the supersolidus sintered 434L stainless steel is larger than that of solid-state sintered one. It is evident from the microstructure that the addition of YAG to 434L stainless steels restricts the grain coarsening and retains the intergranular pores at higher sintering temperature. The scanning electron micrographs shown in Fig. 4 clearly indicate the YAG particles along the grain boundaries with bright contrast. At lower amount of YAG addition, the YAG particles are better linked with the matrix stainless steels. Whereas, a significant agglomeration and segregation of the YAG particles along the grain boundaries were observed for the 434L–10YAG composites. 3.2. Mechanical properties Fig. 5 shows the bulk hardness variation of 434L stainless steels and 434L–YAG composites with sintering temperature. The hardness of both 434L stainless steel as well as 434L–YAG composites
Fig. 2. Effect of sintering temperature and YAG addition on the axial and radial shrinkages in the ferritic stainless steel compacts.
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Fig. 3. Optical micrographs of straight 434L stainless steels and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C.
Fig. 4. SEM micrographs of straight 434L stainless steels and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C.
sintered at 1400 ◦ C were higher than that of solid-state sintered compacts. This can be attributed to the higher sintered density of the supersolidus sintered compacts. The influence of YAG addition on hardness is quite evident at the supersolidus sintered conditions. The reason for the increase in hardness is due to blocking of moving dislocations by the YAG dispersoids as well as inherent hardness of YAG dispersoids. Table 3 shows the tensile testing results of straight 434L stainless steels and 434L–YAG composites sintered at 1400 ◦ C. Addition of Table 3 Tensile properties of the supersolidus sintered straight 434L and 434L–YAG composites
Fig. 5. Effect of sintering temperature and YAG addition on the bulk hardness of the ferritic stainless steel compacts.
Composition (wt.%)
Y.S. (MPa)
U.T.S. (MPa)
% Elongation
434L 434L–5YAG 434L–10YAG
235 265 147
378 411 167
29 13 2
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Fig. 6. Fractographs of the tensile tested surfaces of the supersolidus sintered straight 434L and 434L–YAG composites.
Fig. 6 shows the fractographs of the samples sintered at 1400 ◦ C. In case of the 434L–YAG composites, the brighter phase is YAG particles. From Fig. 6, it is evident that the straight 434L stainless steel shows distinct dimpled morphology, which is a typical characteristic of ductile failure. Increase in YAG content, changes the fracture mode from dimpled to intergranular mode, thereby resulting in lowering of ductility. This is attributed to lowering of sinterability with YAG addition and the presence of brittle additive phase at the 434L inter-particle interface. 3.3. Tribological response
Fig. 7. Variation of wear (m) with sliding distance for supersolidus sintered straight 434L and 434L–YAG composites.
YAG up to 5 wt.% into straight 434L ferritic stainless steel enhances both the yield and tensile strength. It is due to the combined effect of increase in sintered density and restriction of grain growth due to YAG addition. Higher amount (10 wt.%) of YAG addition reduces both tensile and yield strength drastically due to lower sintered density of 434L–10YAG composites and lesser interactions of dispersoids with the matrix phase.
Fig. 7 shows the effect of YAG dispersoids on the wear response of 434L ferritic stainless steel. The wear (m) vs sliding distance curves clearly indicate the better wear resistance of the 434L–5YAG composites as compared to the straight 434L stainless steels. This can be attributed to the higher hardness, better densification and restricted grain coarsening associated with the supersolidus sintered 434L–5YAG composites. Higher amount (10 wt.%) of YAG addition leads to higher wear due to the inferior densification. 3.4. Electrochemical response 3.4.1. OCP vs time The variation of open circuit potential with time for the sintered ferritic 434L stainless steel and 434L–YAG composites in 0.1N H2 SO4 and 0.1N HCl solution are shown in Fig. 8a and b, respectively. In 0.1N H2 SO4 solution the stabilization was achieved quickly for all
Fig. 8. OCP stabilization curves for the sintered 434L and 434L–YAG composites in (a) 0.1N H2 SO4 and (b) 0.1N HCl.
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Fig. 9. Potentiodynamic polarization curves for the straight 434L and 434L–YAG composites sintered at (a) 1200 ◦ C and (b) 1400 ◦ C in 0.1N H2 SO4 solution. Table 4 Passivity parameters obtained from the potentiodynamic polarization curves of 434L and 434L–YAG composites Composition
Sintering temperature (◦ C)
Icrit (A/cm2 )
Epp (mV)
Ipass (A/cm2 )
Ebp (mV)
434L
1200 1400 1200 1400 1200 1400
359 906 3402 340 4394 349
−429 −419 −313 −450 −271 −440
80 49 32 53 599 69
901 927 869 859 975 901
434L–5YAG 434L–10YAG
the samples, with the exception of solid-state sintered 434L–YAG composites which took relatively longer time for stabilization. All the samples get stabilized at a constant OCP (−520 ± 5 mV) in aqueous 0.1N H2 SO4 solution. In the case of 0.1N HCl there is significant difference in the stabilization voltages between the different composites. This variation in OCP value can be attributed to the different degree of attack by the aggressive Cl− ions on the various com-
posites. However, in both 0.1N H2 SO4 and 0.1N HCl solutions a continuous steady shift of the open circuit potential towards more cathodic value with time is evident. This trend could be related to selective anodic dissolution that occurs at the active surface of the composites. A similar trend in the free corrosion potential for sintered austenitic and ferritic stainless steel has been reported in the literature [3,4,13,14].
Fig. 10. Cyclic polarization curves for the straight 434L and 434L–YAG composites sintered at (a) 1200 ◦ C and (b) 1400 ◦ C in 0.1N HCl solution.
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Table 5 Tafel extrapolation data for 434L and 434L–YAG composites Composition
Sintering temperature (◦ C)
ˇa (mV/decade)
434L
1200 1400 1200 1400 1200 1400
166.8 112.8 274.0 267.7 1050 237.7
434L–5YAG 434L–10YAG
ˇc (mV/decade) 258.2 295.7 306.4 546.6 1192 432
Icorr (A/cm2 )
Ecorr (mV)
Corrosion rate (mmpy)
325 327 801 513 4570 409
−485 −502 −522 −526 −523 −522
3.594 3.234 8.875 5.103 50.64 4.151
Table 6 Pitting parameters obtained from cyclic polarization of 434L and 434L–YAG composites Composition
Sintering temperature (◦ C)
Ezcp (mV)
Ep (mV)
434L
1200 1400 1200 1400 1200 1400
−493 −525 −448 −472 −339 −336
−414 −134 −205 −42 −82 −156
434L–5YAG 434L–10YAG
3.4.2. Potentiodynamic polarization The potentiodynamic polarization curves, presented in Fig. 9a and b reveals the variation in current density as a function of applied voltage for pure 434L and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C, respectively. Prominent active–passive behaviour is shown by all the sintered components. The behaviour of the supersolidus sintered ferritic steel–YAG composites in the H2 SO4 environment is comparable to that of the sintered ferritic steel. The passivation data obtained from the polarization curve are presented in Table 4. The active peak current density (Icrit ) and the primary passivation potential (Epp ) do not show a definite trend. However, a significant difference between the critical current density of the solid-state sintered 434L–YAG composites and rest of the samples is evident from Table 4. This can be attributed to the heterogeneity produced by the YAG particle and the poor interaction of YAG with the matrix stainless steels. The passivation current density is similar for the ferritic stainless steels and its composites with the exception of solid-state sintered 434L–10YAG composites which showed an order of magnitude higher Ipass value. The potential for breakdown of the passivity is similar for all the samples and is around 900 mV except for the 434L–10YAG sample, which is 950 mV. The corrosion rate measurements were done using Tafel extrapolation and the data are given in Table 5. A low corrosion potential (Ecorr ) corresponds to inferior corrosion resistance
Eprot (mV) −408 −271 −356 −267 – –
Ep − Eprot (mV) 85 137 151 225 – –
[15,16]. The corrosion potential for the pure ferritic stainless steel is slightly higher than that of the respective composites. This suggests a slightly better corrosion resistance for the straight 434L stainless steels as compared to their YAG added counterparts. Shankar et al. [9] also reported similar corrosion potential values for the 434L–yttria composites processed at similar conditions. There is no significant difference in the Icorr values between the samples except the solid-state sintered 434L–10YAG composites. The corrosion rate too follows a similar trend as Icorr . The supersolidus sintered YAG added 434L composites exhibited comparable corrosion rates as that of unreinforced 434L stainless steels. Similar observations with corrosion rate of ferritic stainless steel–yttria composites have been reported by Shankar et al. [9]. The behaviour can be attributed to the higher sintered density, reduced grain boundary area due to grain coarsening, and more homogeneous microstructures achieved with the supersolidus liquid phase sintering. The literature on P/M stainless steels in acid solutions also reveals that corrosion resistance improves with increasing sintered density [17–19]. Due to the interaction between the YAG and 434L matrix at higher sintering temperatures, the YAG particles are better linked with the matrix leading to better passivation characteristics as compared to that of the solid-state sintered composites. The poor densification and presence of increased amount of reinforcements had locally weakened the passive layer of the solid-state
Fig. 11. SEM micrographs of the cyclic polarization tested straight 434L stainless steels and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C.
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sintered 434L–10YAG composites that lead to the poor passivation and corrosion performance. 3.4.3. Cyclic polarization The cyclic polarization curves for the 434L stainless steels and 434L–YAG composites sintered at 1200 ◦ C and 1400 ◦ C are shown in Fig. 10a and b, respectively, and the pitting corrosion data are tabulated in Table 6. The zero corrosion potential shows a shift towards the noble potential region with YAG dispersion. Pitting potential for 434L–YAG composites is nobler compared to the unreinforced 434L stainless steels. For straight 434L and 434L–5YAG composites, supersolidus sintering results in higher potential for pitting as compared to the solid-state sintering. However, at higher YAG content the behaviour is different. The protection potential also exhibits a similar trend. The protection potential for 434L–10YAG composite could not be obtained because the reverse scan was stopped after reaching the open circuit potential. These alloys are prone to pitting even at the open circuit potential in aqueous 0.1N HCl solution. The steel surface after exposure to chloride solution is observed in the scanning electron microscope and shown in Fig. 11. The attack is very severe in the chloride environment. Extensive metal dissolution has been observed in the solid sate sintered 434L stainless steel. The area of hysteresis loop is a direct measure of the pit propagation kinetics and the extent of localized attack depends on the nature of hysteresis loop [15,16]. It can be observed that the YAG dispersion in the ferritic steel has increased the pit propagation kinetics as can be confirmed from the increase in the hysteresis loop area during the cyclic scan. The dominating corrosion mechanism in these samples is crevice corrosion under the chloride-containing environment. At higher dispersoids content the behaviour is different. The microstructure reveals the segregation of YAG along the grain boundary region. This could be one of the possible reasons for the poor behaviour of the composite during the cyclic test. The mechanism proposed by Otero et al. [3] to justify the high corrosion kinetics observed in the P/M products is based on attack through crevice corrosion, which appears macroscopically as generalized attack. This starts in the pores and becomes more widespread. SEM micrograph shown in Fig. 11 confirms the localized nature of the attack which takes place through crevice corrosion. In reinforced materials the Cl− ion attack was localized on some specific areas preferably at the interface between the matrix and the reinforcements and the bulk material beneath dissolved. Variation in the reinforcement volume fraction has a marginal effect on the susceptibility of the composites to localized corrosion. The supersolidus sintered alloys show slightly lower corrosion susceptibility with relatively smaller hysteresis loop area in the chloride solution. This could be related to the lower porosity level achieved by supersolidus sintered alloys. The pores act as pre-existing crevices on
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the P/M surface [4]. Crevice corrosion situation can result from the consumption of dissolved oxygen within a pore. The oxygen deficiency and higher chloride concentration (due to influx of chloride ions in to the pore) associated with acid hydrolysis reaction with in the pore lead to localized dissolution of metal inside the pore. 4. Conclusions Supersolidus liquid phase sintering (SLPS) results in higher density in both ferritic 434L stainless steel and 434L–YAG composites. Addition of small amount (5 wt.%) of YAG to both the 434L stainless steel increases the densification of the composites. SLPS results in microstructural coarsening and intragranular pores with rounded morphology. However, the addition of fine yttrium aluminium garnet (YAG) powders to 434L restricts microstructural coarsening and the pores remain intergranular. The addition of YAG to 434L stainless steels improves the mechanical and tribological behaviour at supersolidus sintering conditions. Corrosion resistance of supersolidus sintered 434L and 434L–YAG composites is better as compared to that of solid-state sintered compacts in both 0.1N HCl and 0.1N H2 SO4 solutions. Corrosion behaviour of supersolidus sintered 434L–5YAG composites is comparable to that of sintered straight 434L compacts. Pitting corrosion behaviour (in 0.1N HCl solution) of supersolidus sintered compacts is better as compared to those processed through solid-state sintering. References [1] S. Ukai, S. Mizuta, T. Yoshitake, T. Okuda, M. Fujiwara, S. Hagi, T. Kobayashi, J. Nucl. Mater. 283–287 (2000) 702–706. [2] R.M. German, Powder Metallurgy Science, MPIF, Princeton, NJ, USA, 1994. [3] E. Otero, A. Pardo, E. Saenz, M.V. Utrilla, F.J. Perez, Corros. Sci. 40 (1998) 1421–1434. [4] M.C. Baran, B.A. Shaw, Int. J. Powder Metall. 36 (2000) 57–68. [5] L. Cambal, J.A. Lund, Int. J. Powder Metall. 8 (1972) 131–140. [6] R.M. German, Metall. Mater. Trans. A 28A (1997) 1553–1567. [7] S.K. Mukherjee, G.S. Upadhyaya, Int. J. Powder Metall. 19 (1983) 289. [8] J. Shankar, A. Upadhyaya, R. Balasubramaniam, Proceedings of the Advances in Powder Metallurgy Particulate Materials, vol. 13, MPIF, Princeton, NJ, USA, 2002, pp. 313–322. [9] J. Shankar, A. Upadhyaya, R. Balasubramaniam, Corros. Sci. 46 (2004) 487–498. [10] S. Balaji, G. Joshi, A. Upadhyaya, Scripta Mater. 56 (2007) 149–151. [11] J. Jain, A.M. Kar, A. Upadhyaya, Mater. Lett. 58 (2004) 2037–2040. [13] E. Otero, A. Pardo, E. Saenz, M.V. Utrilla, F.J. Perez, Mater. Charact. 35 (1995) 145–151. [14] E. Otero, A. Pardo, E. Saenz, M.V. Utrilla, F.J. Perez, Can. Metall. Quart. 36 (1997) 65–72. [15] A.J. Sedriks, Corrosion of Stainless Steels, John Wiley & Sons Inc., New York, NY, USA, 1996. [16] E.E. Stansbury, R.A. Buchanan, Fundamentals of Electrochemical Corrosion, ASM International, Ohio, USA, 2000. [17] G.H. Lei, R.M. German, H.S. Nayar, Powder Metall. Int. 15 (1983) 70–76. [18] E. Ahlberg, P. Engdahl, R. Johansson, Proceedings of the World Conference on Powder Metallurgy, vol. 1, London, 1990, pp. 419–432. [19] D.R. Gabe, Powder Metall. 4 (1977) 227–231.