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Microchemical and microstructural studies in a PTA weld overlay of Ni–Cr–Si–B alloy on AISI 304L stainless steel
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Surface & Coatings Technology 202 (2008) 2103 – 2112 www.elsevier.com/locate/surfcoat
Microchemical and microstructural studies in a PTA weld overlay of Ni–Cr–Si–B alloy on AISI 304L stainless steel C. Sudha ⁎, P. Shankar, R.V. Subba Rao, R. Thirumurugesan, M. Vijayalakshmi, Baldev Raj Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam - 603 102, India Received 8 May 2007; accepted in revised form 28 August 2007 Available online 7 September 2007
Abstract Ni–Cr–Si–B alloy coatings deposited on AISI 304L stainless steel substrate using plasma transferred arc welding process were studied with respect to the microchemical and microstructural modifications that take place during the process. The coating was characterized for (1) interface integrity and dilution (2) type, morphology and distribution of secondary phases in the coating (3) elemental redistribution between substrate and coating and (4) different phases that form during the deposition process. Formation of a transition zone of only 760 μm width was understood in terms of interdiffusion of Fe and Ni across the interface. Preferential redistribution of carbon to the surface of the coating was observed. It was explained on the basis of difference in thermodynamic activity. Apart from γ-Ni solid solution, the other primary phases identified in Ni–Cr–Si–B alloy were Cr2B, Cr7C3 and Cr3C2. The observed changes in microstructure, microchemistry and hardness have been understood based on the phase transitions of the Ni rich alloy during solidification and cooling on the substrate. © 2007 Elsevier B.V. All rights reserved. Keywords: Chemical repartitioning; Chromium boride; Interdiffusion zone; Ni–Cr–Si–B alloy coating; PTA
1. Introduction Extensive demand from a wide spectrum of industrial applications has been the major driving force for the vibrant research on hardfacing materials, leading to the development of a number of iron, nickel and cobalt based hardfacing alloys. Cobalt based alloys (e.g. Stellite and Triballoy) have been the most popular hardfacing alloys for applications involving wear, corrosion and high temperature service conditions [1]. These alloys also exhibit excellent galling resistance and are therefore widely used in applications where metal to metal contact occurs. However, in nuclear industry the main difficulty in using cobalt based alloy is the induced activity due to the formation of long lived radioactive isotope, Co60 resulting in difficulties during
⁎ Corresponding author. Tel.: +91 44 27480306; fax: +91 44 27480202. E-mail address: [email protected] (C. Sudha). 0257-8972/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2007.08.063
maintenance and decommissioning operations [2]. Hence, for hardfacing nuclear core components, cobalt based alloys are increasingly being replaced with nickel based alloys, like Ni– Cr–Si–B alloy [3]. Investigations on the microstructural changes in Co–Cr–W–C and Ni–Cr–Si–B alloy coatings produced on AISI 304 stainless steel (SS) using high power CO2 laser showed that Ni–Cr–Si–B alloy can be considered as an alternative to Co based coatings [4]. Ni–Cr–Si–B alloys are the preferred hardfacing alloy coatings on components susceptible to galling, high temperature oxidation and corrosion. Weld cladding of these alloys (thickness ≥ 3 mm) can be performed by arc welding processes (GTAW or SAW), solid state welding processes (explosion cladding or roll cladding) or vacuum brazing [5]. In recent years, non-conventional deposition processes like Plasma Transferred Arc Welding (PTAW) and laser hardfacing are being evaluated since they offer the following advantages: (1) lower microchemical redistribution like ‘dilution effect’ in the deposit (2) excellent bonding with the substrate and (3) ability to give dense, defect-free deposits.
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The good abrasion resistance of Ni–Cr–Si–B alloy is achieved by the formation of hard secondary precipitates like borides and carbides. C.R. Das et al. [6] have shown that the high hardness of GTAW deposits of Ni–Cr–Si–B alloy on AISI 316LN is due to the formation of chromium borides and chromium carbides. Microstructure of laser cladded iron, nickel and cobalt base coatings on a steel substrate was studied by Jiang et al. [7]. It was shown that the Ni–Cr–Si–B cladding layer had many subgrains, M23C6 carbides and γ′-nickel silicides in the austenite matrix. Laser surface alloying of NiCrSiB on an austenitic stainless steel was found [8] to produce γ-NiFeCr as the major phase with a few secondary phases such as carbides or borides. Corrosion attack was found to be initiated along the interfaces of the secondary phases which are the potential sites for formation of pits in the presence of chloride ions. The effect of powder feed rate and translation speed of a laser beam on the microstructure of a laser clad Ni–Cr–Si–B alloy deposited over a low carbon steel has also been reported [9,10]. Wear resistance of the clad was found to be poor since the microstructure mainly consisted of primary γ nickel dendrites, interdendritic eutectics of γ nickel, nickel borides and nickel silicides. Conde et al. [11] used high power diode laser for obtaining boride containing nickel based hardfacing alloy coatings on carbon steel. Addition of Cr, C and B was found to increase the surface hardness of the coating due to the formation of Cr23C6 type chromium rich carbides and Ni3B type Ni rich borides. Even though the advantages of Plasma Transferred Arc (PTA) welding process are well established, it is found that very limited literature is available on deposits of Ni–Cr–Si–B alloy on stainless steel by PTA process. Sufficiently thick borided coatings (1 to 1.5 mm) have been obtained on low carbon steel using PTA process [12]. These borided coatings were found to have a wear rate four orders of magnitude lower than the carbon steel substrate. Guojian Xu et al. [13] have characterized a Ni based layer (containing ∼ 15% Cr–4.28% Si–3.8% Fe–3.37% B) formed by plasma cladding process on SUS316L stainless steel (∼ 17% Cr–14% Ni–2.5% Mo). The eutectic structure consisting of Ni + CrB or Ni + Cr7C3 was identified in the clad. In the present work PTA welding process is chosen to get defect-free, thick (N1 mm) adherent weld overlay of Ni–Cr–Si– B alloy on AISI 304L stainless steel which is the structural material used for the storage vessels in fast reactor nuclear plant. The microstructural characterization of the coating has been carried out with the main objective of investigating the (1) integrity and compatibility of the coating with the substrate material (2) type of secondary phases and (3) effect of dilution on the microstructure and microchemistry. The heterogeneous microstructures in different layers of the coating are understood in terms of possible phase transitions undergone by the virgin Ni–Cr–Si–B alloy powder during the coating process. The effect of deposition parameters on elemental repartitioning and Table 1a Chemical composition of AISI 304L steel plate (substrate) Element
C
Cr
Ni
Mn
Mo
P
S
Si
Fe
Wt.%
0.044
18.4
9.2
1.6
0.07
0.021
0.006
0.48
69.36
Table 1b Chemical composition of Ni–Cr–Si–B alloy powders (coating) Element
Ni
Cr
B
Si
Fe
C
Wt.%
72.15
14.00
3.20
4.50
4.50
0.55
formation of precipitates is discussed, in the light of the data available in literature. 2. Experimental Ni–Cr–Si–B hardfacing alloy was deposited on a 15 cm× 15 cm AISI 304L stainless steel plate using PTA welding process. Nominal composition of stainless steel substrate and Ni–Cr–Si–B alloy powders are given in Tables 1a and 1b respectively. The welding parameters used for the deposition of Ni–Cr–Si–B alloy powders are given in Table 2. PTA welding conditions as listed in Table 2 were chosen after several experiments for optimizing the process parameters in terms of achieving desired deposition thickness, minimum dilution and maximum surface hardness. Neutron radiography was carried out to ensure that the coating is defect-free. Specimens of 3 cm× 3 cm dimension were cut by using electrical discharge based wire cutting machine. Subsequently these were cut into smaller pieces (1 cm× 1 cm) by Buehler Isomet 2000 precision wire saw with SiC wheel using water as the coolant. The samples were prepared by conventional metallographic procedures for metallurgical examination. Microhardness measurements on the precipitates were performed using a Shimadzu HMV 2000 microhardness tester with an applied load of 25 g. Cross section hardness profile across coating/matrix interface was obtained using a Vicker's hardness tester using a load of 10 kg. The type of phases present in the Ni–Cr–Si–B coating were analyzed using Philips X-ray diffractometer (θ–2θ geometry) with Cu Kα radiation. Philips XL 30 ESEM attached with an Energy Dispersive Spectrometer (EDS) having Berylium window was used for microstructural and microchemical examination. Quantitative microchemical characterization of the coating was carried out using a Cameca SX50 Electron Probe Micro Analyzer (EPMA) to determine (1) nature of precipitates in coating (2) microchemical redistribution and (3) dilution profile at the interface. X-ray elemental line scans were obtained for Fe, Cr, Ni, Si, B and C across the cross section. The microchemistry
Table 2 Welding parameters used in PTA welding process Welding process Powder feed rate Powder gas flow rate Centre gas flow rate Tungsten electrode diameter Tube to work distance Shielding gas Shielding gas flow rate Polarity Voltage Current Medium of cooling
PTA 30 g/min 4 L/min 3 L/min 4 mm 12 mm Argon l2 L/min DC 24 V 150 A Vermiculate powder
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Fig. 1. SEM micrograph of original Ni–Cr–Si–B alloy powder showing particle size ranging from 50–200 μm.
of the precipitates was used for identifying the type of precipitates. The extent of diffusion of Fe from the substrate into the coating was studied using diffusion profiles. Accelerating voltage of 20 kV, beam current of 20 nA and a beam size of 1 μm were used for the analyses. The analyzing conditions were optimized in order to eliminate overestimation of carbon content due to contamination. Fe Kα, Cr Kα and Ni Kα were analysed using LiF crystal, Si Kα using TAP while PC2 pseudo crystal was used for the analysis of B Kα and C Kα. Quantitative chemical analysis was performed using ‘Quanta’ an automated program which takes into consideration the corrections for atomic number, absorption and fluorescence. 3. Results 3.1. Initial microstructure of the starting material, Ni–Cr–Si–B alloy powder The main objective of the study is to understand the precipitate and morphology evolution undergone by Ni–Cr–
Fig. 3. SEM micrographs showing (a) the entire cross section of Ni–Cr–Si–B alloy coating on AISI 304L substrate which reveals a sharp interface and (b) the absence of microcracks and delamination at the interface.
Si–B alloy upon coating on AISI 304L austenitic stainless steel. Hence it is essential to uniquely characterize the virgin Ni–Cr– Si–B alloy, which is used for the PTA deposition process. These microstructural details would serve as the ‘reference’ structure and enable the identification of the possible phase transition during the coating process. Fig. 1 shows the SEM micrograph of Ni–Cr–Si–B alloy powder. It is seen from the SEM micrograph that the initial charge of Ni–Cr–Si–B alloy is granular with particle size in the range of 50–200 μm. Conventional X-ray diffraction analysis of these powders (Fig. 2) revealed that the major phase present is Ni rich γ phase, in addition to minor amounts of chromium carbides (Cr7C3, Cr3C2) and chromium boride (Cr2B). Based on the above evidence it may be inferred that the reference structure consists of the following: γ(Ni) + Cr7C3 + Cr3C2 + Cr2B. 3.2. Microstructural observation of the coating
Fig. 2. XRD pattern obtained from Ni–Cr–Si–B alloy powder.
Fig. 3(a) shows the SEM micrograph of the interface between Ni–Cr–Si–B alloy coating and AISI 304L substrate. Average thickness of the coating is ∼5.9 ± 0.9 mm. The interface between AISI 304L substrate and Ni–Cr–Si–B alloy coating is found to be sharp as depicted clearly from the micrograph. Fig. 3(b)
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Fig. 4. X-ray diffraction pattern obtained from the surface of Ni–Cr–Si–B alloy coating.
shows that the interface is free from microcracks or delamination features. The surface of the Ni–Cr–Si–B alloy coating was examined by neutron radiography and it was found to be free of cracks or thickness variations. X-ray diffraction pattern obtained from the surface of the weld overlay of Ni–Cr–Si–B alloy (Fig. 4) showed peaks corresponding to Ni rich γ phase along with other phases like Cr2B, Cr3C2 and Cr7C3. The type of phases present in the coating was same as that of the original powder except for a significant increase in the intensity for chromium carbide (Cr3C2) phase. The lattice parameter for Ni rich γ phase was found to be 0.3529 nm (both in the original powder and the coating) which is nearly equal to that of pure Ni (0.3524 nm) [JSPDS no. 04-0850]. The hardness of AISI 304L substrate was around 150 VHN while the macrohardness profile taken across the cross section (Fig. 5) showed a systematic increase in hardness value from the interface to the surface of the coating. Hardness value increased from a minimum of 250 VHN to a maximum of 700 VHN. Three regions with different range of hardness values could be distinguished in the hardness profile namely region 1 close to
Fig. 5. Hardness profile taken across the cross section of PTA weld overlay of Ni–Cr–Si–B alloy on AISI 304L substrate showing three different regions in the coating.
Fig. 6. SEM micrograph showing a typical microstructure of region 1 (up to ∼760 μm from the interface).
the interface with hardness value around 300 VHN, region 2 (450–500 VHN) and region 3 close to the surface with a very high hardness of around 650–700 VHN. Detailed microstructural investigation was carried out, in order to correlate the above changes in hardness value with the possible phase transition in the corresponding regions of the coating. 3.2.1. Microstructure of region 1 The microstructure of Ni–Cr–Si–B alloy coating varied considerably as a function of distance from the interface. Fig. 6 shows the SEM micrograph obtained from region 1 (∼ 300 VHN), very close to the interface. This region consists of high volume fraction of thin lamellae resembling irregular flake like eutectic structures. To identify the constituents of the lamellar product EPMA line scan was obtained across them. From the elemental line scan (Fig. 7) the lamellae were identified to be chromium borides. The regions in between the lamellae showed enrichment for Ni and Fe. Line scans were obtained from at least 6–7 precipitates of the same morphology to arrive at the above observations. Table 3 gives the result of quantitative chemical analysis performed on the secondary phases. Since the size of the
Fig. 7. EPMA line scan across the lamellar dendrites showing enrichment for Cr and B.
C. Sudha et al. / Surface & Coatings Technology 202 (2008) 2103 –2112 Table 3 Chemical composition of precipitate in region 1 Morphology
Irregular eutectic flakes (contribution from matrix possible)
Wt.% of elements Ni
Fe
Cr
B
C
Si
37.05
15.39
33.40
10.26
1.74
2.16
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substrate and the deposit showed a gradual increase in average Si content from the interface to the surface of the coating. The observations in this section can be summarized as follows: region 1 which is very close to the interface has the structure of lamellar dendrites with the lamellae showing enrichment for Cr and B. Region 2 shows the presence of secondary phases like chromium rich borides and carbides of four different morphologies and region 3 has predominantly needle shaped chromium carbides in addition to chromium borides and γ-Ni matrix.
lamellae is about one micron matrix contribution effects in EPMA measurements need to be considered.
3.3. Microchemical redistribution across the interdiffusion zone
3.2.2. Microstructure of region 2 The microstructure beyond region 1 of the coating had precipitates of four different morphology in the γ matrix as shown in Fig. 8: needle like ‘A’, bulky globular ‘B’, irregular flake like eutectic structures ‘C’ and cuboidal precipitates ‘D’ surrounded by ‘C’. All the four morphologically different precipitates had similar microchemistry as revealed by X-ray elemental mapping (Fig. 9): enrichment of Cr and depletion of Ni and Fe. Since light element analysis was not possible using EDAX with Be window used in the present study, elemental line scans across the precipitates were obtained using electron microprobe analyzer to unambiguously identify the type of precipitates. Fig. 10 shows the line scan across a bulky globular precipitate (B) which is found to be rich in chromium and carbon. Both the precipitates ‘A’ and ‘B’ were found to be of the same kind. It is possible that the bulky globular precipitate represents the cross section of a needle like precipitate. The line scan obtained across ‘C’ was similar to Fig. 7 showing enrichment for Cr and B on the lamellae. The cuboidal precipitate ‘D’ was also found to be a chromium rich boride as seen from Fig. 11. Table 4 gives a correlation between the observed morphology and microchemistry of the precipitates observed in region 2.
In order to examine if there is a significant dilution, X-ray spectra were obtained using the Scanning Electron Microscope (SEM) at various distances from the interface. From the X-ray spectra given in Fig. 15 it can be observed that on the Ni–Cr– Si–B alloy side up to a distance of around 62 μm from the interface, the Fe content was about 30.52 wt.% which is much higher than the actual Fe content in Ni–Cr–Si–B alloy (4.5 wt. %). As the distance from the interface further increases on the Ni–Cr–Si–B alloy side the average Fe content gradually decreases (Fig. 16) up to a distance of 760 μm beyond which the composition matches with that of reference Ni–Cr–Si–B alloy powder. On the other side of the interface i.e. AISI 304L substrate side a similar effect due to Ni diffusion could be noticed but only up to a distance of 20–30 μm. X-ray elemental line scan obtained using an electron microprobe thus further confirmed that significant microchemical redistribution has taken place up to a distance of 760 μm from the interface. In Tables 3 and 5 if the contribution from matrix and the precipitate phase to the total chemical composition of the irregular eutectic flakes was taken into consideration it was observed that borides in region 1 have higher iron content than those present in region 2. Thus it is possible to distinguish between the borides in region 1 and 2 based on the Fe content.
3.2.3. Microstructure of region 3 Fig. 12 reveals high volume fraction of needle like precipitates in addition to the irregular flake like eutectic structures at the surface of the coating where the Vicker's hardness was very high in the range of ∼ 700 VHN. Similar analyses as explained in the previous section were carried out in order to identify the type of precipitates. In Fig. 13, an X-ray elemental line scan showed that the needle like precipitates present in region 3 was chromium carbides and the eutectic structures were chromium borides. The hardness of chromium borides was around 2575 VHN and that of chromium carbides was 1670 VHN. The quantitative chemical analysis of all secondary phases was carried out and results are given in Table 5. From the results it is observed that boron solubility in carbides is negligible whereas carbon has limited solubility in borides. Ni content in borides is lower than that in carbides and both Ni and Si are found to selectively redistribute away from carbides and borides. Higher Si content in type ‘C’ precipitates was only due to matrix effect as evident from the higher Ni content in these precipitates. The X-ray elemental line scan (Fig. 14) obtained for Si across the
4. Discussion The results presented above on the formation of secondary phases of different morphology and microchemistry in a
Fig. 8. SEM micrograph showing the typical microstructure of region 2.
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Fig. 9. X-ray elemental mapping obtained from a region in the BSE image showing Cr enrichment, Ni and Fe depletion in the precipitates and eutectic phases.
PTA weld overlay of Ni–Cr–Si–B alloy on stainless steel are discussed in the light of available information on phase diagram and possible non-equilibrium transformation in this system. The X-ray diffraction pattern (Fig. 2) obtained from Ni–Cr– Si–B alloy powder showed that the major precipitate phases present are Cr3C2 and Cr7C3 type chromium carbides and Cr2B type chromium borides in a matrix of γ-Ni. The macrohardness profile (Fig. 5) of Ni–Cr–Si–B alloy on AISI 304L stainless steel showed three regions in the coating with different Vicker's hardness values. Detailed investigation showed that the microstructure of Ni–Cr–Si–B alloy coating varied as a function of distance from the interface:
interface increases is a consequence of the process parameters influencing the thermal cycle which each region is subjected to, during coating. Region 1 which is adjacent to the interface exhibited the morphology of lamellar dendrites (Fig. 6) showing enrichment for Cr and B (Fig. 7). In the X-ray diffraction pattern (Fig. 4) obtained from the Ni–Cr–Si–B alloy coating the only boride phase observed is Cr2B. In an earlier work on laser clad and furnace melted Ni–Cr–Si–B alloy on AISI 1045 steel [14] presence of nickel boride in addition to chromium boride has been reported which was not observed in the present work. Cr2B is a eutectic solidification product in Cr–B binary phase
Region 1 Up to 760 μm from the interface showing solidification structure having a hardness of around 300 VHN. Region 2 Beyond region 1 having bulky globular and needle shaped chromium carbides and irregular eutectic shaped chromium borides in addition to γ-Ni matrix (macrohardness is ∼ 475 VHN). Region 3 On the surface of the coating predominantly needle shaped chromium carbides were observed in addition to the γ-Ni matrix with a hardness value of ∼700 VHN. The above variation in microstructure and hardness of Ni–Cr– Si–B alloy coating can be understood in terms of the following possible phase transitions during the process of coating. 4.1. Possible phase transitions in different regions of the Ni– Cr–Si–B alloy coating The observed change in the morphology, type and volume fraction of the secondary phases as the distance from the
Fig. 10. Elemental line scan obtained using an electron microprobe across a bulky globular precipitate ‘B’ showing enrichment for Cr and C.
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Fig. 11. Elemental line scan obtained using an electron microprobe from a cuboidal precipitate ‘D’ showing both the precipitate at the centre as well as the irregular eutectic flakes to be enriched in Cr and B.
diagram [15]. Campbell et al. [16] extrapolated the Cr–B binary phase diagram to Ni–Cr–B ternary system and predicted that the following invariant transformation may occur at 1495 K. L→Cr2 B þ αðCrÞ þ γðNiÞ
ð1Þ
α-Cr could not be detected as a separate phase possibly because it is submicron sized and occurs along with Cr2B precipitates and the eutectic structure was found to consist of only γ-Ni + Cr2B. The eutectic structure observed doesn't exhibit the lamellar morphology expected from normal eutectic phases, but rather had an irregular flake morphology similar to anomalous eutectic structures observed in Al–Si system [17,18]. However formation of an anomalous eutectic structure can be confirmed only after investigating further the entropy of solution, atom movement across non-isothermal interfaces and volume fraction as well as growth velocity of the faceting phase. During the initial stages of coating, the substrate acts as a heat sink as a result of which the temperature gradient existing at the interface is large. Under such circumstances a planar growth front is favoured where the lamellae grow opposite to the direction of heat flow [19]. It has been already observed [20,21] that both the interlamellar spacing and undercooling decreases upon increasing the imposed temperature gradient. This explains the smaller interlamellar spacing (∼ 0.18– 0.45 μm at about 100 μm) close to the interface compared to that in the regions well away from the interface (∼ 0.68– 1.38 μm at a distance of 3 mm from the interface). The lower hardness observed in region 1 may be due to lower volume
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Fig. 12. SEM micrograph showing needle shaped chromium carbides in region 3.
fraction of precipitates or dilution and the resultant effect of Fe on the hardness of Ni based matrix and lower hardness of iron rich chromium borides. The microstructure in region 2, at the centre of the coating, resembled a hypereutectic structure. Needle shaped and bulky globular precipitates were also seen in addition to the eutectic structures (Fig. 8). From EPMA and X-ray diffraction analysis (Figs. 10 and 4) these precipitates were identified as chromium carbides of either Cr3C2 or Cr7C3 type. Cuboidal precipitates present at the centre of the eutectic structures (phase ‘D’ as in Fig. 8) have been identified as Cr2B type chromium borides from XRD and EPMA analysis (Figs. 4 and 11). From the quantitative elemental analysis selective redistribution of Ni and Si away from the precipitate phases was observed which agrees with a similar observation made during the formation of iron borides in boriding of steels [22]. Even though Si was present up to 4.5%, a separate nickel silicide phase (Ni3Si) was not detected in XRD spectrum probably due to multiple peak overlaps. However it has been reported that [7,14] nickel silicides (Ni3Si) are observed as submicron sized precipitates
Table 4 Type of precipitates in region 2 Region in Fig. 8
Morphology
Type of precipitate
A B C D
Needle like Bulky globular Irregular eutectic flakes Cuboidal precipitate surrounded by eutectic flakes
Chromium carbide Chromium carbide Chromium boride Chromium boride
Fig. 13. Elemental line scan obtained using an electron microprobe from region 3 shows Cr and C enrichment across needle like precipitate and Cr, B enrichment on irregular eutectic flakes.
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Table 5 Average chemical composition of various precipitates in region 2 Region Morphology in Fig. 8
Wt.% of elements Ni
Fe
Cr
B
C
Si
A and B Needle shaped/bulky 5.48 15.94 68.35 0.07 10.11 0.05 globular C Irregular eutectic flakes 19.19 8.9 54.15 15.52 1.12 1.04 (contribution from matrix possible) D Cuboidal precipitate 1.86 6.27 75.86 13.83 2.13 0.05 surrounded by ‘C’
which is beyond the resolution limit of SEM and EPMA. Hence presence of Ni3Si phase in the present investigation cannot be ruled out. A small but gradual increase observed in the average Si content from the interface to the surface of the coating (Fig. 14) may be due to the difference in Si solute distribution at the solid liquid interface as reported in literature [14]. An attempt has been made to reconstruct the possible sequence of precipitate and morphology evolution in the coating, leading to the observed variation in microstructure in region 2. From the Cr– C binary phase diagram [23] the solidification temperatures for Cr3C2 and Cr7C3 are obtained as 2086 ± 285 K and 2041 ± 283 K respectively. Eutectic transformation to Cr2B occurs at even lower temperature of 1903 K [15]. The solidification temperature for γNi is around 1728 K [24] which decreases with increase in Cr concentration. Hence during solidification Cr3C2 and Cr7C3 type chromium carbides precipitate out from the melt initially. Owing to the fast diffusion of carbon in the melt, the primary carbides (phase A and B) are coarse as seen from Fig. 8. Further reduction in temperature results in Cr2B precipitation (phase D) followed by the eutectic transformation of the remaining liquid to γ-Ni + Cr2B (phase C). The coarse primary Cr2B precipitates (phase D) already present in the solidified structure act as easy nucleation sites for the γ-Ni + Cr2B eutectic phases. The secondary Cr2B (phase C) precipitates formed by solid state diffusion process are finer than the primary phase D precipitates due to the matrix strain and
Fig. 14. Elemental line scan for Si obtained across the interface of the substrate and the coating showing a gradual increase in average Si content from the interface to the surface of the coating.
Fig. 15. X-ray spectra obtained at various distances from the interface showing enhancement in Fe and Ni contents on Ni–Cr–Si–B alloy and AISI 304L substrate sides respectively.
diffusion constraints. The possible sequence of events during solidification of the plasma melted Ni–Cr–Si–B alloy powder can be given as: L1
Y
Cr3 C2 þ Cr7 C3 þ L2
Y
Cr3 C2 þ Cr7 C3 þ Cr2 B þ L3 A Cr3 C2 þ Cr7 C3 þ Cr2 B þ ðg Ni þ Cr2 BÞ
At the centre of the coating the supercooled melt acts as the heat sink, hence the temperature gradient is smaller and the undercooling is less. Therefore the interlamellar spacing of the eutectic
Fig. 16. EPMA line scan for Ni and Fe confirming the presence of dilution layer up to 760 μm.
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product is larger than that observed at the interface (∼0.68– 1.38 μm). At the surface of the coating (region 3), large volume fraction of needle shaped precipitates was observed which were identified as chromium rich carbides from EPMA analysis (Fig. 13). Also the X-ray diffraction pattern obtained from the surface of Ni–Cr–Si– B alloy coating (Fig. 4) showed an increase in intensity for Cr3C2 phase. Such a chemical repartitioning of carbon to the surface of the coating has been reported earlier in the interrupted boriding of stainless steel [25] and also in boriding of a carburized layer [26,27]. A possible explanation for the above observed phenomena is as follows: solubility of carbon in borides is negligibly small [28] and it has also been reported [29] that the solubility of carbon decreases with increase in Ni content and is at its minimum at about 70% Ni. Ni has a positive interaction coefficient with carbon with a reported value [30] of interaction parameter being ɛCNi = +4.1 at 1273 K. In region 2 where the volume fraction of chromium borides is very high, carbon will segregate away from the borides and will tend to dissolve into the solid solution. From Fig. 13 it is clear that carbon content decreases in regions rich in Cr and B and vice versa. The presence of Ni in the solid solution increases the activity of carbon. This activity gradient offers the driving force for diffusion of carbon to neighboring regions 1 and 3. In region 3 since the existing temperature range is high, carbon diffusion is faster than in region 1 resulting in build up of higher concentration of carbon and consequently high volume fraction of chromium carbides. However presence of finer carbides in region 1 also cannot be ignored. Such a microchemical redistribution of carbon has been found [25] to minimize the strain energy introduced in the system during boriding. The hardness value of individual chromium boride was 2575 VHN and that of chromium carbide was 1670 VHN. Hence it is expected that the surface of the coating should have a low hardness compared to other regions. But a high hardness value of around 700 VHN is observed at the surface of the coating where the major phase is chromium carbide. The high hardness value obtained is only due to the large volume fraction of carbides at the surface of the coating. It is thus possible to identify different types of precipitates from their morphology and microchemistry. 4.2. Transition zone at the interface The main problems encountered in the weld deposition of nickel based Ni–Cr–Si–B hardfacing alloy are the high fluidity of the molten alloy, generation of residual stresses in the deposit that can lead to distortion/cracking and significant dilution of the deposit by the substrate material due to large difference in their melting point (Ni–Cr–Si–B alloy: 1223–1338 K, stainless steel: 1665–1717 K). Since dilution of the coating due to the diffusion of Fe from the substrate affects the microstructure of the coating especially the first few layers, dilution has to be kept to a minimum to preserve the integrity of the coating material. The interdiffusion zone was found to be present up to a distance of almost 760 μm as revealed from microhardness, SEM and EPMA observations (Figs. 5, 15 and 16). On AISI 304L substrate up to a distance of 20–30 μm the Ni content was found
2111
to be higher (29 wt.%). Thus Fe from stainless steel substrate has diffused to a distance of 760–800 μm into the Ni–Cr–Si–B coating, whereas Ni from Ni–Cr–Si–B alloy was found to have diffused only up to a distance of 20 μm. It is probably due to the fact that Ni is diffusing into solid substrate of AISI 304L stainless steel at lower temperature range while Fe is diffusing to a region which is experiencing higher temperatures. Dilution was also found to have an effect on the microchemistry of the precipitate phase present in region 1. Taking into consideration the contribution from the matrix as well as the secondary phase, the Fe content in the irregular eutectic flakes (phase C) in region 1 was found to be more than that present in region 2. Reduction in the hardness of the coating near the interface because of dilution effects has been already reported in literature [6]. It has been reported [11,12] that it is advantageous to have a thin transition zone and less gradient in hardness since it ensures perfect chemical bonding between the substrate and the coating. Hence observation of a lower gradient in hardness and smooth change in the Fe counts in the X-ray intensity profile (Fig. 16) suggest good adhesion of Ni–Cr–Si–B alloy with the substrate material, which agrees well with SEM observation of the defect-free interface. The width of the transition zone obtained in this study was lesser than that obtained for Gas Tungsten Arc Welded (GTAW) Ni–Cr–Si–B alloy on AISI 316 L(N) austenitic stainless steel where 2.5 mm thick transition zone was observed [6]. This may be because of the inherent advantages in PTA process (higher arc temperature and more concentrated heat pattern) or because of the optimization of the welding process with increased powder feed rate and translation speed of the torch, which would have resulted in lower degree of dilution [4]. It has been suggested that adopting a laser cladding process would further reduce the dilution effects [13]. Laser cladding of boron containing alloy powders (2.5– 3.32% B) was found to result in defect-free high quality coatings with heterogeneous microstructure consisting of primary dendrites of γ-Ni with interdendritic constituents of γ-Ni with borides, carbides or silicides [9,31,32]. However thick homogeneous coatings are difficult to obtain by laser welding process [33,34]. Bourithis et al. found [12] that PTA process was advantageous for getting thicker homogeneous coatings in low carbon steel. The transition zone thickness was in the range of 1000–1600 μm and the clad was found to consist of a eutectic structure of ferrite and iron borides. Increase in boron content was found to have a profound influence on the type of eutectic phases formed. On account of the lower heat input (approximately six times lesser than plasma weld overlay) and higher cooling rate, the primary phases in laser cladding process were finer (∼5 μm) than that of plasma cladding process (∼10–20 μm) [13]. Laser as well as plasma clad microstructure of a nickel based powder (3.37% B) showed a planar, eutectic (γ-Ni + CrB or γ-Ni + Cr7C3) and hypereutectic structures. Hardness was higher and wear resistance of the coating was considerably improved in a laser cladding process. 5. Conclusion (1) Adherent, defect-free Ni–Cr–Si–B alloy coating of maximum 7 mm thickness could be obtained on AISI
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304L substrate by suitable selection of process parameters using Plasma Transferred Arc process. (2) Ni–Cr–Si–B alloy coating was found to exhibit a heterogeneous microstructure consisting of three different regions as given below: Region 1 with a hardness value of ∼300 VHN having a eutectic solidification structure with γNi + Cr2B (phase C precipitates) eutectic constituents. Region 2 having Cr7C3 or Cr3C2 type chromium carbides (phase A or B) and Cr2B type chromium boride (phase C or D) in γ-Ni matrix with hardness around 475 VHN. Region 3 having a high volume fraction of needle shaped chromium carbides exhibiting a high hardness value of ∼700 VHN. The evolution of microstructure has been understood in terms of the mechanism of eutectic transformation, effect of temperature gradient and elemental redistribution. (3) Thickness of the interdiffusion zone obtained was ∼ 760 μm which was much lesser than that obtained using gas tungsten arc welding process (∼2.5 mm) which has been explained based on the change in process parameters. Acknowledgement The authors thank Dr. P.R Vasudeva Rao, Metallurgy and Materials Group Director for his keen interest and encouragement throughout the period of this project. Thanks are further due to Mr. C. Balasubramaniam, M/s Omplas System, Chennai for enthusiastic support in PTA deposition of the samples. References [1] C.R. Das, S.K. Albert, A.K. Bhaduri, G. Kempulraj, J. Mater. Process. Technol. 141 (2003) 60. [2] H. Ocken, Surf. Coat. Technol. 76–77 (1995) 456. [3] A.K. Bhaduri, S.K. Albert, V. Shankar, C.R. Das, G. Srinivasan, V. Ramasubbu, A. Dasgupta, A.L.E. Terrance, C. Sudha, P. Kuppusami, V.S. Raghunathan, Hardfacing and Surface Engineering for PFBR, Materials R & D for PFBR, IGCAR, Kalpakkam, January 1–2, 2003 (Conference proceedings), 2003, p. 249.
[4] M. Corchia, P. Delogu, F. Nenci, A. Belmondo, S. Corcoruto, W. Stabielli, Wear 119 (1987) 137. [5] ASM Committee on Hardfacing, Welding, Brazing and Soldering, Metals Handbook, vol. 6, ASM International, 1993, p. 789. [6] C.R. Das, S.K. Albert, A.K. Bhaduri, C. Sudha, A.L.E. Terrance, Surf. Eng. 21 (4) (2005) 290. [7] M. Jiang, X.P. Jiang, J.G. Huang, X.F. Sun, J.S. Zhang, Y.L. Ge, Z.Q. Hu, Mater. Lett. 7 (12) (1989) 453. [8] C.T. Kwok, F.T. Cheng, H.C. Man, Mater. Sci. Eng., A Struct. Mater.: Prop. Microstruct. Process. 290 (2000) 55. [9] Qian Ming, L.C. Lim, Z.D. Chen, Surf. Coat. Technol. 106 (1998) 174. [10] L.C. Lim, Qian Ming, Z.D. Chen, Surf. Coat. Technol. 106 (1998) 183. [11] A. Conde, F. Zubiri, y J. de Damborenea, Mater. Sci. Eng., A Struct. Mater.: Prop. Microstruct. Process. 334 (2002) 233. [12] L. Bourithis, S. Papaefthymiou, G.D. Papadimitriou, Appl. Surf. Sci. 200 (2002) 203. [13] Guojian Xu, Muneharu Kutsuna, Zhongjie Liu, Hong Zhang, Mater. Sci. Eng., A Struct. Mater.: Prop. Microstruct. Process. 417 (2006) 63. [14] Qiang Li, Dawei Zhang, Tingquan Lei, Chuanzhong Chen, Wenzhe Chen, Surf. Coat. Technol. 137 (2001) 122. [15] Thaddeus Massalski, Binary Alloy Phase Diagrams, ASM Handbook, vol. 1, ASM International, 1986, p. 473. [16] C.E. Campbell, U.R. Kattner, Calphad 26 (3) (2002) 477. [17] R. Elliott, S.M.D. Glenister, Acta Metall. 28 (1980) 1489. [18] R. Elliot, Int. Metals Rev. 219 (1977) 161. [19] W. Kurz, D.J. Fisher (Eds.), Fundamentals of Solidification, Trans Tech publications, Switzerland, 1986, ch. 4. [20] D.J. Fisher, W. Kurz, Acta Metall. 28 (1979) 777. [21] W. Kurz, D.J. Fisher, Acta Metall. 29 (1980) 11. [22] I. Uslu, H. Comert, M. Ipek, O. Ozdemir, C. Bindal, Mater. Des. 28 (2007) 55. [23] Thaddeus Massalski, Binary Alloy Phase Diagrams, ASM Handbook, vol. 1, ASM International, 1986, p. 838. [24] Thaddeus Massalski, Binary Alloy Phase Diagrams, ASM Handbook, vol. 1, ASM International, 1986, p. 1301. [25] P. Gopalakrishnan, P. Shankar, Balamurugan, S.S. Ramakrishnan, A.K. Tyagi, Mater. Sci. Technol. 10 (2004) 45. [26] M. Kulka, A. Pertek, L. Klimek, Mater. Charact. 56 (3) (2006) 232. [27] M. Kulka, A. Pertek. Appl. Surf. Sci. 214 (2003) 161. [28] C. Badini, C. Gianoglio, G. Pradelli, Surf. Coat. Technol. 30 (1987) 157. [29] T. Wada, H. Wada, J.F. Elliott, J. Chipman. Metall. Trans. 2 (1971) 2199. [30] E. Hsin Foo, C.H.P. Lupis, Acta Metall. 21 (1973) 1409. [31] C. Navas, R. Colaco, J. de Damborenea, R. Vilar, Surf. Coat. Technol. 200 (24) (2006) 6854. [32] C.P. Paul, Amit Jain, P. Ganesh, J. Neogi, A.K. Nath, Opt. Laser Eng. 44 (10) (2006) 1096. [33] I. Manna, J. Dutta Majumdar, B. Ramesh Chandra, S. Nayak, Narendra B. Dahotre, Surf. Coat. Technol. 201 (2006) 434. [34] J. Dutta Majumdar, B. Ramesh Chandra, I. Manna, Tribol. Int. 40 (2007) 146.
Surface & Coatings Technology 202 (2008) 2103 – 2112 www.elsevier.com/locate/surfcoat
Microchemical and microstructural studies in a PTA weld overlay of Ni–Cr–Si–B alloy on AISI 304L stainless steel C. Sudha ⁎, P. Shankar, R.V. Subba Rao, R. Thirumurugesan, M. Vijayalakshmi, Baldev Raj Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam - 603 102, India Received 8 May 2007; accepted in revised form 28 August 2007 Available online 7 September 2007
Abstract Ni–Cr–Si–B alloy coatings deposited on AISI 304L stainless steel substrate using plasma transferred arc welding process were studied with respect to the microchemical and microstructural modifications that take place during the process. The coating was characterized for (1) interface integrity and dilution (2) type, morphology and distribution of secondary phases in the coating (3) elemental redistribution between substrate and coating and (4) different phases that form during the deposition process. Formation of a transition zone of only 760 μm width was understood in terms of interdiffusion of Fe and Ni across the interface. Preferential redistribution of carbon to the surface of the coating was observed. It was explained on the basis of difference in thermodynamic activity. Apart from γ-Ni solid solution, the other primary phases identified in Ni–Cr–Si–B alloy were Cr2B, Cr7C3 and Cr3C2. The observed changes in microstructure, microchemistry and hardness have been understood based on the phase transitions of the Ni rich alloy during solidification and cooling on the substrate. © 2007 Elsevier B.V. All rights reserved. Keywords: Chemical repartitioning; Chromium boride; Interdiffusion zone; Ni–Cr–Si–B alloy coating; PTA
1. Introduction Extensive demand from a wide spectrum of industrial applications has been the major driving force for the vibrant research on hardfacing materials, leading to the development of a number of iron, nickel and cobalt based hardfacing alloys. Cobalt based alloys (e.g. Stellite and Triballoy) have been the most popular hardfacing alloys for applications involving wear, corrosion and high temperature service conditions [1]. These alloys also exhibit excellent galling resistance and are therefore widely used in applications where metal to metal contact occurs. However, in nuclear industry the main difficulty in using cobalt based alloy is the induced activity due to the formation of long lived radioactive isotope, Co60 resulting in difficulties during
⁎ Corresponding author. Tel.: +91 44 27480306; fax: +91 44 27480202. E-mail address: [email protected] (C. Sudha). 0257-8972/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2007.08.063
maintenance and decommissioning operations [2]. Hence, for hardfacing nuclear core components, cobalt based alloys are increasingly being replaced with nickel based alloys, like Ni– Cr–Si–B alloy [3]. Investigations on the microstructural changes in Co–Cr–W–C and Ni–Cr–Si–B alloy coatings produced on AISI 304 stainless steel (SS) using high power CO2 laser showed that Ni–Cr–Si–B alloy can be considered as an alternative to Co based coatings [4]. Ni–Cr–Si–B alloys are the preferred hardfacing alloy coatings on components susceptible to galling, high temperature oxidation and corrosion. Weld cladding of these alloys (thickness ≥ 3 mm) can be performed by arc welding processes (GTAW or SAW), solid state welding processes (explosion cladding or roll cladding) or vacuum brazing [5]. In recent years, non-conventional deposition processes like Plasma Transferred Arc Welding (PTAW) and laser hardfacing are being evaluated since they offer the following advantages: (1) lower microchemical redistribution like ‘dilution effect’ in the deposit (2) excellent bonding with the substrate and (3) ability to give dense, defect-free deposits.
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The good abrasion resistance of Ni–Cr–Si–B alloy is achieved by the formation of hard secondary precipitates like borides and carbides. C.R. Das et al. [6] have shown that the high hardness of GTAW deposits of Ni–Cr–Si–B alloy on AISI 316LN is due to the formation of chromium borides and chromium carbides. Microstructure of laser cladded iron, nickel and cobalt base coatings on a steel substrate was studied by Jiang et al. [7]. It was shown that the Ni–Cr–Si–B cladding layer had many subgrains, M23C6 carbides and γ′-nickel silicides in the austenite matrix. Laser surface alloying of NiCrSiB on an austenitic stainless steel was found [8] to produce γ-NiFeCr as the major phase with a few secondary phases such as carbides or borides. Corrosion attack was found to be initiated along the interfaces of the secondary phases which are the potential sites for formation of pits in the presence of chloride ions. The effect of powder feed rate and translation speed of a laser beam on the microstructure of a laser clad Ni–Cr–Si–B alloy deposited over a low carbon steel has also been reported [9,10]. Wear resistance of the clad was found to be poor since the microstructure mainly consisted of primary γ nickel dendrites, interdendritic eutectics of γ nickel, nickel borides and nickel silicides. Conde et al. [11] used high power diode laser for obtaining boride containing nickel based hardfacing alloy coatings on carbon steel. Addition of Cr, C and B was found to increase the surface hardness of the coating due to the formation of Cr23C6 type chromium rich carbides and Ni3B type Ni rich borides. Even though the advantages of Plasma Transferred Arc (PTA) welding process are well established, it is found that very limited literature is available on deposits of Ni–Cr–Si–B alloy on stainless steel by PTA process. Sufficiently thick borided coatings (1 to 1.5 mm) have been obtained on low carbon steel using PTA process [12]. These borided coatings were found to have a wear rate four orders of magnitude lower than the carbon steel substrate. Guojian Xu et al. [13] have characterized a Ni based layer (containing ∼ 15% Cr–4.28% Si–3.8% Fe–3.37% B) formed by plasma cladding process on SUS316L stainless steel (∼ 17% Cr–14% Ni–2.5% Mo). The eutectic structure consisting of Ni + CrB or Ni + Cr7C3 was identified in the clad. In the present work PTA welding process is chosen to get defect-free, thick (N1 mm) adherent weld overlay of Ni–Cr–Si– B alloy on AISI 304L stainless steel which is the structural material used for the storage vessels in fast reactor nuclear plant. The microstructural characterization of the coating has been carried out with the main objective of investigating the (1) integrity and compatibility of the coating with the substrate material (2) type of secondary phases and (3) effect of dilution on the microstructure and microchemistry. The heterogeneous microstructures in different layers of the coating are understood in terms of possible phase transitions undergone by the virgin Ni–Cr–Si–B alloy powder during the coating process. The effect of deposition parameters on elemental repartitioning and Table 1a Chemical composition of AISI 304L steel plate (substrate) Element
C
Cr
Ni
Mn
Mo
P
S
Si
Fe
Wt.%
0.044
18.4
9.2
1.6
0.07
0.021
0.006
0.48
69.36
Table 1b Chemical composition of Ni–Cr–Si–B alloy powders (coating) Element
Ni
Cr
B
Si
Fe
C
Wt.%
72.15
14.00
3.20
4.50
4.50
0.55
formation of precipitates is discussed, in the light of the data available in literature. 2. Experimental Ni–Cr–Si–B hardfacing alloy was deposited on a 15 cm× 15 cm AISI 304L stainless steel plate using PTA welding process. Nominal composition of stainless steel substrate and Ni–Cr–Si–B alloy powders are given in Tables 1a and 1b respectively. The welding parameters used for the deposition of Ni–Cr–Si–B alloy powders are given in Table 2. PTA welding conditions as listed in Table 2 were chosen after several experiments for optimizing the process parameters in terms of achieving desired deposition thickness, minimum dilution and maximum surface hardness. Neutron radiography was carried out to ensure that the coating is defect-free. Specimens of 3 cm× 3 cm dimension were cut by using electrical discharge based wire cutting machine. Subsequently these were cut into smaller pieces (1 cm× 1 cm) by Buehler Isomet 2000 precision wire saw with SiC wheel using water as the coolant. The samples were prepared by conventional metallographic procedures for metallurgical examination. Microhardness measurements on the precipitates were performed using a Shimadzu HMV 2000 microhardness tester with an applied load of 25 g. Cross section hardness profile across coating/matrix interface was obtained using a Vicker's hardness tester using a load of 10 kg. The type of phases present in the Ni–Cr–Si–B coating were analyzed using Philips X-ray diffractometer (θ–2θ geometry) with Cu Kα radiation. Philips XL 30 ESEM attached with an Energy Dispersive Spectrometer (EDS) having Berylium window was used for microstructural and microchemical examination. Quantitative microchemical characterization of the coating was carried out using a Cameca SX50 Electron Probe Micro Analyzer (EPMA) to determine (1) nature of precipitates in coating (2) microchemical redistribution and (3) dilution profile at the interface. X-ray elemental line scans were obtained for Fe, Cr, Ni, Si, B and C across the cross section. The microchemistry
Table 2 Welding parameters used in PTA welding process Welding process Powder feed rate Powder gas flow rate Centre gas flow rate Tungsten electrode diameter Tube to work distance Shielding gas Shielding gas flow rate Polarity Voltage Current Medium of cooling
PTA 30 g/min 4 L/min 3 L/min 4 mm 12 mm Argon l2 L/min DC 24 V 150 A Vermiculate powder
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Fig. 1. SEM micrograph of original Ni–Cr–Si–B alloy powder showing particle size ranging from 50–200 μm.
of the precipitates was used for identifying the type of precipitates. The extent of diffusion of Fe from the substrate into the coating was studied using diffusion profiles. Accelerating voltage of 20 kV, beam current of 20 nA and a beam size of 1 μm were used for the analyses. The analyzing conditions were optimized in order to eliminate overestimation of carbon content due to contamination. Fe Kα, Cr Kα and Ni Kα were analysed using LiF crystal, Si Kα using TAP while PC2 pseudo crystal was used for the analysis of B Kα and C Kα. Quantitative chemical analysis was performed using ‘Quanta’ an automated program which takes into consideration the corrections for atomic number, absorption and fluorescence. 3. Results 3.1. Initial microstructure of the starting material, Ni–Cr–Si–B alloy powder The main objective of the study is to understand the precipitate and morphology evolution undergone by Ni–Cr–
Fig. 3. SEM micrographs showing (a) the entire cross section of Ni–Cr–Si–B alloy coating on AISI 304L substrate which reveals a sharp interface and (b) the absence of microcracks and delamination at the interface.
Si–B alloy upon coating on AISI 304L austenitic stainless steel. Hence it is essential to uniquely characterize the virgin Ni–Cr– Si–B alloy, which is used for the PTA deposition process. These microstructural details would serve as the ‘reference’ structure and enable the identification of the possible phase transition during the coating process. Fig. 1 shows the SEM micrograph of Ni–Cr–Si–B alloy powder. It is seen from the SEM micrograph that the initial charge of Ni–Cr–Si–B alloy is granular with particle size in the range of 50–200 μm. Conventional X-ray diffraction analysis of these powders (Fig. 2) revealed that the major phase present is Ni rich γ phase, in addition to minor amounts of chromium carbides (Cr7C3, Cr3C2) and chromium boride (Cr2B). Based on the above evidence it may be inferred that the reference structure consists of the following: γ(Ni) + Cr7C3 + Cr3C2 + Cr2B. 3.2. Microstructural observation of the coating
Fig. 2. XRD pattern obtained from Ni–Cr–Si–B alloy powder.
Fig. 3(a) shows the SEM micrograph of the interface between Ni–Cr–Si–B alloy coating and AISI 304L substrate. Average thickness of the coating is ∼5.9 ± 0.9 mm. The interface between AISI 304L substrate and Ni–Cr–Si–B alloy coating is found to be sharp as depicted clearly from the micrograph. Fig. 3(b)
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Fig. 4. X-ray diffraction pattern obtained from the surface of Ni–Cr–Si–B alloy coating.
shows that the interface is free from microcracks or delamination features. The surface of the Ni–Cr–Si–B alloy coating was examined by neutron radiography and it was found to be free of cracks or thickness variations. X-ray diffraction pattern obtained from the surface of the weld overlay of Ni–Cr–Si–B alloy (Fig. 4) showed peaks corresponding to Ni rich γ phase along with other phases like Cr2B, Cr3C2 and Cr7C3. The type of phases present in the coating was same as that of the original powder except for a significant increase in the intensity for chromium carbide (Cr3C2) phase. The lattice parameter for Ni rich γ phase was found to be 0.3529 nm (both in the original powder and the coating) which is nearly equal to that of pure Ni (0.3524 nm) [JSPDS no. 04-0850]. The hardness of AISI 304L substrate was around 150 VHN while the macrohardness profile taken across the cross section (Fig. 5) showed a systematic increase in hardness value from the interface to the surface of the coating. Hardness value increased from a minimum of 250 VHN to a maximum of 700 VHN. Three regions with different range of hardness values could be distinguished in the hardness profile namely region 1 close to
Fig. 5. Hardness profile taken across the cross section of PTA weld overlay of Ni–Cr–Si–B alloy on AISI 304L substrate showing three different regions in the coating.
Fig. 6. SEM micrograph showing a typical microstructure of region 1 (up to ∼760 μm from the interface).
the interface with hardness value around 300 VHN, region 2 (450–500 VHN) and region 3 close to the surface with a very high hardness of around 650–700 VHN. Detailed microstructural investigation was carried out, in order to correlate the above changes in hardness value with the possible phase transition in the corresponding regions of the coating. 3.2.1. Microstructure of region 1 The microstructure of Ni–Cr–Si–B alloy coating varied considerably as a function of distance from the interface. Fig. 6 shows the SEM micrograph obtained from region 1 (∼ 300 VHN), very close to the interface. This region consists of high volume fraction of thin lamellae resembling irregular flake like eutectic structures. To identify the constituents of the lamellar product EPMA line scan was obtained across them. From the elemental line scan (Fig. 7) the lamellae were identified to be chromium borides. The regions in between the lamellae showed enrichment for Ni and Fe. Line scans were obtained from at least 6–7 precipitates of the same morphology to arrive at the above observations. Table 3 gives the result of quantitative chemical analysis performed on the secondary phases. Since the size of the
Fig. 7. EPMA line scan across the lamellar dendrites showing enrichment for Cr and B.
C. Sudha et al. / Surface & Coatings Technology 202 (2008) 2103 –2112 Table 3 Chemical composition of precipitate in region 1 Morphology
Irregular eutectic flakes (contribution from matrix possible)
Wt.% of elements Ni
Fe
Cr
B
C
Si
37.05
15.39
33.40
10.26
1.74
2.16
2107
substrate and the deposit showed a gradual increase in average Si content from the interface to the surface of the coating. The observations in this section can be summarized as follows: region 1 which is very close to the interface has the structure of lamellar dendrites with the lamellae showing enrichment for Cr and B. Region 2 shows the presence of secondary phases like chromium rich borides and carbides of four different morphologies and region 3 has predominantly needle shaped chromium carbides in addition to chromium borides and γ-Ni matrix.
lamellae is about one micron matrix contribution effects in EPMA measurements need to be considered.
3.3. Microchemical redistribution across the interdiffusion zone
3.2.2. Microstructure of region 2 The microstructure beyond region 1 of the coating had precipitates of four different morphology in the γ matrix as shown in Fig. 8: needle like ‘A’, bulky globular ‘B’, irregular flake like eutectic structures ‘C’ and cuboidal precipitates ‘D’ surrounded by ‘C’. All the four morphologically different precipitates had similar microchemistry as revealed by X-ray elemental mapping (Fig. 9): enrichment of Cr and depletion of Ni and Fe. Since light element analysis was not possible using EDAX with Be window used in the present study, elemental line scans across the precipitates were obtained using electron microprobe analyzer to unambiguously identify the type of precipitates. Fig. 10 shows the line scan across a bulky globular precipitate (B) which is found to be rich in chromium and carbon. Both the precipitates ‘A’ and ‘B’ were found to be of the same kind. It is possible that the bulky globular precipitate represents the cross section of a needle like precipitate. The line scan obtained across ‘C’ was similar to Fig. 7 showing enrichment for Cr and B on the lamellae. The cuboidal precipitate ‘D’ was also found to be a chromium rich boride as seen from Fig. 11. Table 4 gives a correlation between the observed morphology and microchemistry of the precipitates observed in region 2.
In order to examine if there is a significant dilution, X-ray spectra were obtained using the Scanning Electron Microscope (SEM) at various distances from the interface. From the X-ray spectra given in Fig. 15 it can be observed that on the Ni–Cr– Si–B alloy side up to a distance of around 62 μm from the interface, the Fe content was about 30.52 wt.% which is much higher than the actual Fe content in Ni–Cr–Si–B alloy (4.5 wt. %). As the distance from the interface further increases on the Ni–Cr–Si–B alloy side the average Fe content gradually decreases (Fig. 16) up to a distance of 760 μm beyond which the composition matches with that of reference Ni–Cr–Si–B alloy powder. On the other side of the interface i.e. AISI 304L substrate side a similar effect due to Ni diffusion could be noticed but only up to a distance of 20–30 μm. X-ray elemental line scan obtained using an electron microprobe thus further confirmed that significant microchemical redistribution has taken place up to a distance of 760 μm from the interface. In Tables 3 and 5 if the contribution from matrix and the precipitate phase to the total chemical composition of the irregular eutectic flakes was taken into consideration it was observed that borides in region 1 have higher iron content than those present in region 2. Thus it is possible to distinguish between the borides in region 1 and 2 based on the Fe content.
3.2.3. Microstructure of region 3 Fig. 12 reveals high volume fraction of needle like precipitates in addition to the irregular flake like eutectic structures at the surface of the coating where the Vicker's hardness was very high in the range of ∼ 700 VHN. Similar analyses as explained in the previous section were carried out in order to identify the type of precipitates. In Fig. 13, an X-ray elemental line scan showed that the needle like precipitates present in region 3 was chromium carbides and the eutectic structures were chromium borides. The hardness of chromium borides was around 2575 VHN and that of chromium carbides was 1670 VHN. The quantitative chemical analysis of all secondary phases was carried out and results are given in Table 5. From the results it is observed that boron solubility in carbides is negligible whereas carbon has limited solubility in borides. Ni content in borides is lower than that in carbides and both Ni and Si are found to selectively redistribute away from carbides and borides. Higher Si content in type ‘C’ precipitates was only due to matrix effect as evident from the higher Ni content in these precipitates. The X-ray elemental line scan (Fig. 14) obtained for Si across the
4. Discussion The results presented above on the formation of secondary phases of different morphology and microchemistry in a
Fig. 8. SEM micrograph showing the typical microstructure of region 2.
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Fig. 9. X-ray elemental mapping obtained from a region in the BSE image showing Cr enrichment, Ni and Fe depletion in the precipitates and eutectic phases.
PTA weld overlay of Ni–Cr–Si–B alloy on stainless steel are discussed in the light of available information on phase diagram and possible non-equilibrium transformation in this system. The X-ray diffraction pattern (Fig. 2) obtained from Ni–Cr– Si–B alloy powder showed that the major precipitate phases present are Cr3C2 and Cr7C3 type chromium carbides and Cr2B type chromium borides in a matrix of γ-Ni. The macrohardness profile (Fig. 5) of Ni–Cr–Si–B alloy on AISI 304L stainless steel showed three regions in the coating with different Vicker's hardness values. Detailed investigation showed that the microstructure of Ni–Cr–Si–B alloy coating varied as a function of distance from the interface:
interface increases is a consequence of the process parameters influencing the thermal cycle which each region is subjected to, during coating. Region 1 which is adjacent to the interface exhibited the morphology of lamellar dendrites (Fig. 6) showing enrichment for Cr and B (Fig. 7). In the X-ray diffraction pattern (Fig. 4) obtained from the Ni–Cr–Si–B alloy coating the only boride phase observed is Cr2B. In an earlier work on laser clad and furnace melted Ni–Cr–Si–B alloy on AISI 1045 steel [14] presence of nickel boride in addition to chromium boride has been reported which was not observed in the present work. Cr2B is a eutectic solidification product in Cr–B binary phase
Region 1 Up to 760 μm from the interface showing solidification structure having a hardness of around 300 VHN. Region 2 Beyond region 1 having bulky globular and needle shaped chromium carbides and irregular eutectic shaped chromium borides in addition to γ-Ni matrix (macrohardness is ∼ 475 VHN). Region 3 On the surface of the coating predominantly needle shaped chromium carbides were observed in addition to the γ-Ni matrix with a hardness value of ∼700 VHN. The above variation in microstructure and hardness of Ni–Cr– Si–B alloy coating can be understood in terms of the following possible phase transitions during the process of coating. 4.1. Possible phase transitions in different regions of the Ni– Cr–Si–B alloy coating The observed change in the morphology, type and volume fraction of the secondary phases as the distance from the
Fig. 10. Elemental line scan obtained using an electron microprobe across a bulky globular precipitate ‘B’ showing enrichment for Cr and C.
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Fig. 11. Elemental line scan obtained using an electron microprobe from a cuboidal precipitate ‘D’ showing both the precipitate at the centre as well as the irregular eutectic flakes to be enriched in Cr and B.
diagram [15]. Campbell et al. [16] extrapolated the Cr–B binary phase diagram to Ni–Cr–B ternary system and predicted that the following invariant transformation may occur at 1495 K. L→Cr2 B þ αðCrÞ þ γðNiÞ
ð1Þ
α-Cr could not be detected as a separate phase possibly because it is submicron sized and occurs along with Cr2B precipitates and the eutectic structure was found to consist of only γ-Ni + Cr2B. The eutectic structure observed doesn't exhibit the lamellar morphology expected from normal eutectic phases, but rather had an irregular flake morphology similar to anomalous eutectic structures observed in Al–Si system [17,18]. However formation of an anomalous eutectic structure can be confirmed only after investigating further the entropy of solution, atom movement across non-isothermal interfaces and volume fraction as well as growth velocity of the faceting phase. During the initial stages of coating, the substrate acts as a heat sink as a result of which the temperature gradient existing at the interface is large. Under such circumstances a planar growth front is favoured where the lamellae grow opposite to the direction of heat flow [19]. It has been already observed [20,21] that both the interlamellar spacing and undercooling decreases upon increasing the imposed temperature gradient. This explains the smaller interlamellar spacing (∼ 0.18– 0.45 μm at about 100 μm) close to the interface compared to that in the regions well away from the interface (∼ 0.68– 1.38 μm at a distance of 3 mm from the interface). The lower hardness observed in region 1 may be due to lower volume
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Fig. 12. SEM micrograph showing needle shaped chromium carbides in region 3.
fraction of precipitates or dilution and the resultant effect of Fe on the hardness of Ni based matrix and lower hardness of iron rich chromium borides. The microstructure in region 2, at the centre of the coating, resembled a hypereutectic structure. Needle shaped and bulky globular precipitates were also seen in addition to the eutectic structures (Fig. 8). From EPMA and X-ray diffraction analysis (Figs. 10 and 4) these precipitates were identified as chromium carbides of either Cr3C2 or Cr7C3 type. Cuboidal precipitates present at the centre of the eutectic structures (phase ‘D’ as in Fig. 8) have been identified as Cr2B type chromium borides from XRD and EPMA analysis (Figs. 4 and 11). From the quantitative elemental analysis selective redistribution of Ni and Si away from the precipitate phases was observed which agrees with a similar observation made during the formation of iron borides in boriding of steels [22]. Even though Si was present up to 4.5%, a separate nickel silicide phase (Ni3Si) was not detected in XRD spectrum probably due to multiple peak overlaps. However it has been reported that [7,14] nickel silicides (Ni3Si) are observed as submicron sized precipitates
Table 4 Type of precipitates in region 2 Region in Fig. 8
Morphology
Type of precipitate
A B C D
Needle like Bulky globular Irregular eutectic flakes Cuboidal precipitate surrounded by eutectic flakes
Chromium carbide Chromium carbide Chromium boride Chromium boride
Fig. 13. Elemental line scan obtained using an electron microprobe from region 3 shows Cr and C enrichment across needle like precipitate and Cr, B enrichment on irregular eutectic flakes.
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Table 5 Average chemical composition of various precipitates in region 2 Region Morphology in Fig. 8
Wt.% of elements Ni
Fe
Cr
B
C
Si
A and B Needle shaped/bulky 5.48 15.94 68.35 0.07 10.11 0.05 globular C Irregular eutectic flakes 19.19 8.9 54.15 15.52 1.12 1.04 (contribution from matrix possible) D Cuboidal precipitate 1.86 6.27 75.86 13.83 2.13 0.05 surrounded by ‘C’
which is beyond the resolution limit of SEM and EPMA. Hence presence of Ni3Si phase in the present investigation cannot be ruled out. A small but gradual increase observed in the average Si content from the interface to the surface of the coating (Fig. 14) may be due to the difference in Si solute distribution at the solid liquid interface as reported in literature [14]. An attempt has been made to reconstruct the possible sequence of precipitate and morphology evolution in the coating, leading to the observed variation in microstructure in region 2. From the Cr– C binary phase diagram [23] the solidification temperatures for Cr3C2 and Cr7C3 are obtained as 2086 ± 285 K and 2041 ± 283 K respectively. Eutectic transformation to Cr2B occurs at even lower temperature of 1903 K [15]. The solidification temperature for γNi is around 1728 K [24] which decreases with increase in Cr concentration. Hence during solidification Cr3C2 and Cr7C3 type chromium carbides precipitate out from the melt initially. Owing to the fast diffusion of carbon in the melt, the primary carbides (phase A and B) are coarse as seen from Fig. 8. Further reduction in temperature results in Cr2B precipitation (phase D) followed by the eutectic transformation of the remaining liquid to γ-Ni + Cr2B (phase C). The coarse primary Cr2B precipitates (phase D) already present in the solidified structure act as easy nucleation sites for the γ-Ni + Cr2B eutectic phases. The secondary Cr2B (phase C) precipitates formed by solid state diffusion process are finer than the primary phase D precipitates due to the matrix strain and
Fig. 14. Elemental line scan for Si obtained across the interface of the substrate and the coating showing a gradual increase in average Si content from the interface to the surface of the coating.
Fig. 15. X-ray spectra obtained at various distances from the interface showing enhancement in Fe and Ni contents on Ni–Cr–Si–B alloy and AISI 304L substrate sides respectively.
diffusion constraints. The possible sequence of events during solidification of the plasma melted Ni–Cr–Si–B alloy powder can be given as: L1
Y
Cr3 C2 þ Cr7 C3 þ L2
Y
Cr3 C2 þ Cr7 C3 þ Cr2 B þ L3 A Cr3 C2 þ Cr7 C3 þ Cr2 B þ ðg Ni þ Cr2 BÞ
At the centre of the coating the supercooled melt acts as the heat sink, hence the temperature gradient is smaller and the undercooling is less. Therefore the interlamellar spacing of the eutectic
Fig. 16. EPMA line scan for Ni and Fe confirming the presence of dilution layer up to 760 μm.
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product is larger than that observed at the interface (∼0.68– 1.38 μm). At the surface of the coating (region 3), large volume fraction of needle shaped precipitates was observed which were identified as chromium rich carbides from EPMA analysis (Fig. 13). Also the X-ray diffraction pattern obtained from the surface of Ni–Cr–Si– B alloy coating (Fig. 4) showed an increase in intensity for Cr3C2 phase. Such a chemical repartitioning of carbon to the surface of the coating has been reported earlier in the interrupted boriding of stainless steel [25] and also in boriding of a carburized layer [26,27]. A possible explanation for the above observed phenomena is as follows: solubility of carbon in borides is negligibly small [28] and it has also been reported [29] that the solubility of carbon decreases with increase in Ni content and is at its minimum at about 70% Ni. Ni has a positive interaction coefficient with carbon with a reported value [30] of interaction parameter being ɛCNi = +4.1 at 1273 K. In region 2 where the volume fraction of chromium borides is very high, carbon will segregate away from the borides and will tend to dissolve into the solid solution. From Fig. 13 it is clear that carbon content decreases in regions rich in Cr and B and vice versa. The presence of Ni in the solid solution increases the activity of carbon. This activity gradient offers the driving force for diffusion of carbon to neighboring regions 1 and 3. In region 3 since the existing temperature range is high, carbon diffusion is faster than in region 1 resulting in build up of higher concentration of carbon and consequently high volume fraction of chromium carbides. However presence of finer carbides in region 1 also cannot be ignored. Such a microchemical redistribution of carbon has been found [25] to minimize the strain energy introduced in the system during boriding. The hardness value of individual chromium boride was 2575 VHN and that of chromium carbide was 1670 VHN. Hence it is expected that the surface of the coating should have a low hardness compared to other regions. But a high hardness value of around 700 VHN is observed at the surface of the coating where the major phase is chromium carbide. The high hardness value obtained is only due to the large volume fraction of carbides at the surface of the coating. It is thus possible to identify different types of precipitates from their morphology and microchemistry. 4.2. Transition zone at the interface The main problems encountered in the weld deposition of nickel based Ni–Cr–Si–B hardfacing alloy are the high fluidity of the molten alloy, generation of residual stresses in the deposit that can lead to distortion/cracking and significant dilution of the deposit by the substrate material due to large difference in their melting point (Ni–Cr–Si–B alloy: 1223–1338 K, stainless steel: 1665–1717 K). Since dilution of the coating due to the diffusion of Fe from the substrate affects the microstructure of the coating especially the first few layers, dilution has to be kept to a minimum to preserve the integrity of the coating material. The interdiffusion zone was found to be present up to a distance of almost 760 μm as revealed from microhardness, SEM and EPMA observations (Figs. 5, 15 and 16). On AISI 304L substrate up to a distance of 20–30 μm the Ni content was found
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to be higher (29 wt.%). Thus Fe from stainless steel substrate has diffused to a distance of 760–800 μm into the Ni–Cr–Si–B coating, whereas Ni from Ni–Cr–Si–B alloy was found to have diffused only up to a distance of 20 μm. It is probably due to the fact that Ni is diffusing into solid substrate of AISI 304L stainless steel at lower temperature range while Fe is diffusing to a region which is experiencing higher temperatures. Dilution was also found to have an effect on the microchemistry of the precipitate phase present in region 1. Taking into consideration the contribution from the matrix as well as the secondary phase, the Fe content in the irregular eutectic flakes (phase C) in region 1 was found to be more than that present in region 2. Reduction in the hardness of the coating near the interface because of dilution effects has been already reported in literature [6]. It has been reported [11,12] that it is advantageous to have a thin transition zone and less gradient in hardness since it ensures perfect chemical bonding between the substrate and the coating. Hence observation of a lower gradient in hardness and smooth change in the Fe counts in the X-ray intensity profile (Fig. 16) suggest good adhesion of Ni–Cr–Si–B alloy with the substrate material, which agrees well with SEM observation of the defect-free interface. The width of the transition zone obtained in this study was lesser than that obtained for Gas Tungsten Arc Welded (GTAW) Ni–Cr–Si–B alloy on AISI 316 L(N) austenitic stainless steel where 2.5 mm thick transition zone was observed [6]. This may be because of the inherent advantages in PTA process (higher arc temperature and more concentrated heat pattern) or because of the optimization of the welding process with increased powder feed rate and translation speed of the torch, which would have resulted in lower degree of dilution [4]. It has been suggested that adopting a laser cladding process would further reduce the dilution effects [13]. Laser cladding of boron containing alloy powders (2.5– 3.32% B) was found to result in defect-free high quality coatings with heterogeneous microstructure consisting of primary dendrites of γ-Ni with interdendritic constituents of γ-Ni with borides, carbides or silicides [9,31,32]. However thick homogeneous coatings are difficult to obtain by laser welding process [33,34]. Bourithis et al. found [12] that PTA process was advantageous for getting thicker homogeneous coatings in low carbon steel. The transition zone thickness was in the range of 1000–1600 μm and the clad was found to consist of a eutectic structure of ferrite and iron borides. Increase in boron content was found to have a profound influence on the type of eutectic phases formed. On account of the lower heat input (approximately six times lesser than plasma weld overlay) and higher cooling rate, the primary phases in laser cladding process were finer (∼5 μm) than that of plasma cladding process (∼10–20 μm) [13]. Laser as well as plasma clad microstructure of a nickel based powder (3.37% B) showed a planar, eutectic (γ-Ni + CrB or γ-Ni + Cr7C3) and hypereutectic structures. Hardness was higher and wear resistance of the coating was considerably improved in a laser cladding process. 5. Conclusion (1) Adherent, defect-free Ni–Cr–Si–B alloy coating of maximum 7 mm thickness could be obtained on AISI
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304L substrate by suitable selection of process parameters using Plasma Transferred Arc process. (2) Ni–Cr–Si–B alloy coating was found to exhibit a heterogeneous microstructure consisting of three different regions as given below: Region 1 with a hardness value of ∼300 VHN having a eutectic solidification structure with γNi + Cr2B (phase C precipitates) eutectic constituents. Region 2 having Cr7C3 or Cr3C2 type chromium carbides (phase A or B) and Cr2B type chromium boride (phase C or D) in γ-Ni matrix with hardness around 475 VHN. Region 3 having a high volume fraction of needle shaped chromium carbides exhibiting a high hardness value of ∼700 VHN. The evolution of microstructure has been understood in terms of the mechanism of eutectic transformation, effect of temperature gradient and elemental redistribution. (3) Thickness of the interdiffusion zone obtained was ∼ 760 μm which was much lesser than that obtained using gas tungsten arc welding process (∼2.5 mm) which has been explained based on the change in process parameters. Acknowledgement The authors thank Dr. P.R Vasudeva Rao, Metallurgy and Materials Group Director for his keen interest and encouragement throughout the period of this project. Thanks are further due to Mr. C. Balasubramaniam, M/s Omplas System, Chennai for enthusiastic support in PTA deposition of the samples. References [1] C.R. Das, S.K. Albert, A.K. Bhaduri, G. Kempulraj, J. Mater. Process. Technol. 141 (2003) 60. [2] H. Ocken, Surf. Coat. Technol. 76–77 (1995) 456. [3] A.K. Bhaduri, S.K. Albert, V. Shankar, C.R. Das, G. Srinivasan, V. Ramasubbu, A. Dasgupta, A.L.E. Terrance, C. Sudha, P. Kuppusami, V.S. Raghunathan, Hardfacing and Surface Engineering for PFBR, Materials R & D for PFBR, IGCAR, Kalpakkam, January 1–2, 2003 (Conference proceedings), 2003, p. 249.
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