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Improvement of mechanical properties of stainless maraging steel laser weldments by post-weld ageing treatments
Materials and Design 40 (2012) 276–284
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Improvement of mechanical properties of stainless maraging steel laser weldments by post-weld ageing treatments Jian An ⇑, Fanyan Meng, Xiaoxia Lv, Haiyuan Liu, Xiaoxi Gao, Yuanbo Wang, You Lu Key Laboratory of Automobile Materials, Ministry of Education, Department of Materials Science and Engineering, Nanling Campus of Jilin University, Changchun 130025, People’s Republic of China
a r t i c l e
i n f o
Article history: Received 4 January 2012 Accepted 13 March 2012 Available online 28 March 2012 Keywords: A. Ferrous metals and alloys D. Welding E. Mechanical
a b s t r a c t The response of stainless maraging steel weldments to post-weld ageing treatment has been investigated. Post-weld ageing was performed at five different temperatures, viz., 420 °C, 460 °C, 500 °C, 540 °C, and 580 °C. Metallographic characterization of weldment revealed three zones, namely fusion zone, heataffected zone (HAZ) and unaffected parent metal zone. Hardness and tensile properties were evaluated after ageing at different temperatures. Hardness in HAZ and fusion zone varied with ageing temperature differently from that of the parent metal; it became higher in HAZ and fusion zone than in parent metal zone above 420 °C. Among the applied ageing treatments, ageing at 460 °C achieved the highest tensile strength. A graph was constructed for determination of fracture location and post-weld heat treatment efficiency based on experimental results, using hardness ratio of HAZ to the treated parent material and hardness ratio of HAZ to the as-received parent material. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction For several decades, two major types of maraging steels, Fe–Ni system and Fe–Cr–Ni system alloys, have been developed and used in defense, power, and other commercial industries due to their superior properties such as ultra-high strength, excellent fracture toughness, good machining properties and weldability [1,2]. The most commonly used Fe–Ni maraging steel composition is based on 18 wt.% Ni with significant levels of Co (8–13 wt.%), while the subsequently developed Co-free maraging steels contain Fe, Ni, Mo and Ti [3–7]. The composition of Fe–Cr–Ni maraging steels is based on 10.2–16 wt.% Cr with considerable levels of Ni (3.9–10.8 wt.%) and other elements such as Mo, Ti, Al and Cu. Fe–Cr–Ni maraging steel usually possesses good corrosion resistance [8–12]. It is well known that the ultra-high strength of maraging steels is mainly attributed to the precipitation of densely distributed fine intermetallic precipitates in a martensitic matrix, which strengthen the alloys through Orowan-type mechanism. Typically, this kind of microstructure is adjusted by a simple two-step heat-treatment process which consists of solution annealing to produce a complete martensitic matrix, and subsequently ageing at intermediate temperatures (400 °C to 600 °C) to cause precipitation hardening. The ageing behaviour of maraging steels has been extensively investigated over recent decades. By means of high-resolution methods, such as atom probe ⇑ Corresponding author. Tel.: +86 431 85095874; fax: +86 431 85095876. E-mail address: [email protected] (J. An). 0261-3069/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2012.03.020
tomography (APT), transmission electron microscopy and smallangle neutron scattering, the size and chemical composition of nanometer-sized cluster of precipitates are identified in numerous maraging steels. For example, 18% Ni maraging steels exhibit ordered and coherent phases such as l, S and X (<450 °C), Ni3 (Mo, Ti), and Fe2Mo phases (450–500 °C), and austenite precipitation (500 °C-As) [13–16]. For Fe–Cr–Ni–Al–Ti stainless maraging steel, which contains in addition 0.5–1.0 at.% Si, a spherical G-phase with the chemical composition of Ti6Si7Ni16, and a rodshaped Ni3(Ti,Al) phase were observed [17,18]. Weldability is a unique property which makes maraging steels especially suitable for processing in aviation and space industry, since welding of maraging steels can be performed, without preheating, by processes such as electron beam, laser welding, and submerged arc [19–22]. Ramana et al. [23] studied the microstructure and residual stress distribution of dissimilar electron beam weld of 18 Ni maraging steel and medium alloy medium carbon steel, it was observed that the residual stress distribution would be more compressive if the maraging steel was in soft condition. Among these processes, gas tungsten-arc welding is widely utilized in view of the consistency of weld equality and overall efficiency [24]. There are various reports published recently on gas tungsten-arc welding of maraging steels. As the 18 Ni maraging steels represent majority of maraging steel products, the associated low tensile property and toughness rising from premature of reverted austenite, cracking in heat-affected zone (HAZ) immediately adjacent to fusion line associated with gas tungsten-arc welding were the subjects of several reports in the literature. Shamantha et al.
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[25] investigated microstructure changes during welding and subsequent post-weld ageing of 18 Ni maraging steel at various temperatures and concluded that pronounced segregation of Ti and Mo elements along interdendrite and intercellular boundaries in the weld metal led to austenite reversion at temperatures much lower than in unweld material. Similarly, Reddy et al. [26] reported that solutionized treatment of 18 Ni maraging steel weldments at 815 °C could not remove the segregation of Ti and Mo elements, only homogenization treatment at 1150 °C succeeded, and consequently resulted in optimum tensile properties. In most cases, welding 18 Ni maraging steels in the solutionized condition followed by post-weld ageing treatment at a relatively low temperature can attain weld efficiency over 90%. Laser welding is a well established technique that possesses remarkable characteristics such as high welding speed, low heat input, low distortion, deep penetration, narrow HAZ and free from filler wire. Therefore laser welding of maraging steels is a useful process for critical application such as aeronautics and space technology equipments. However, up to date, investigations on postweld heat treatment of maraging steel laser welds still lack systemically experimental data and need further understanding the responses from Fe–Cr–Ni system maraging steels other than 18 Ni system alloys. Reddy et al. [20] investigated the role of reverted austenite through various post-weld heat treatments on the microstructure and mechanical properties of 18% Ni (250 grade) maraging steel laser welds, and found that the homogenized + solutionized + aged weldments exhibited higher tensile properties and toughness compared to directly aged and solutionized + aged weldments due to the reduction in micro-segregation by diffusion of alloying elements such as Mo and Ti from the cell boundaries and the resulting absence of reverted austenite. To the best of our knowledge, post-weld heat treatments are mostly performed on 18 Ni maraging steels which were usually in solutionized condition, hardly on Fe–Cr–Ni system maraging steels. The considerable difference in chemical composition between 18 Ni maraging steels and Fe–Cr–Ni system maraging steels may result in quite different response to post-weld heat treatment in some aspects of reverted austenite formation, alloying element distribution and mechanical properties. Therefore, the present paper focus on microstructural evolution and tensile properties of Fe–Cr–Ni maraging stainless steel laser weldments through post-weld ageing, the studied material was in aged condition prior to welding. This process is more practical and economical for fabrication of large structures compared with post-weld heat treatment applied to maraging steel in solutionized condition. For this state of parent material, only localized area surrounding the weld, instead of the whole structure, needs post-weld heat treatment, which can be accomplished easily by using small-sized heat-treatment equipment. The experiments were designed to study the effect of post-weld ageing on tensile properties in a wide temperature range from 420 °C to 580 °C. A useful graph was established for assessment of fracture location and improvement in tensile yield strength after post-weld treatment using two parameters indicating hardness ratios between different regions in laser weldments. 2. Experimental details The test material is a type of Fe–Cr–Ni–Co–Mo–Ti–Cu stainless maraging steel in the form of 3.0-mm-thick sheet with composi-
tion given in Table 1, which was received in the short-term 420°C/1 h/air cooling aged condition. Rectangular sheets in the dimensions of 200 100 3 mm3 were machined from the as-received sheets and were autogenously laser beam welded using a CO2 laser. The parameters used for laser welding are listed in Table 2. The welds were subjected to ageing treatments at a temperature range varying from 420 °C to 580 °C for 2 h to examine the effect of ageing temperature on the microstructure and mechanical properties of welds. Metallographic microstructures of the weldments in the as-weld and post-weld aged conditions were characterized using a LEXT-OLS3000 confocal scanning laser microscope. Vilella’s reagent (picric acid 1 g + HCl 5 ml + ethanol 100 ml) was used for etching fusion zone, HAZ, and parent material to observe the microstructure. Mechanical properties including hardness and tensile properties of welds in different aged conditions were evaluated. Microhardness survey was conducted in the midsection of weld coupons with 200 g load on a Vickers hardness indentation tester. Tensile tests were carried out on a MTS 810 tensile testing system using specimens having a reduced cross-section of 5.0 mm 3.0 mm and a gauge length of 30.0 mm, which were machined from various heat-treated weld coupons with the welds located in the middle position. For each heat treatment temperature level, at least three specimens are tested to obtain an average value of tensile properties. Revered austenite contents in the welds and parent material in various heat-treated conditions were estimated using a Rigaku X-ray diffractometer (XRD) using Cu Ka radiation. The volume fraction of austenite was estimated from measurements of the integrated intensities of martensite and austenite peaks, assuming they are the only phases present. 3. Results and discussion 3.1. Microstructure and hardness dependence on ageing temperature for parent material The as-received material had a microstructure of complete martensite with fine grain size about 5–10 lm observed on optical microscope, as shown in Fig. 1. XRD analysis further revealed no austenite phase present, as shown in Fig. 2. Hardness dependence of this metal on ageing temperature and period was examined by heating it to a conventional solution temperature of 850 °C for 1 h followed by ageing in temperature range of 380–580 °C for various periods. From hardness curves in Fig. 3, it can be seen that the hardness depends greatly on the ageing temperature and period. At ageing temperatures lower than 500 °C, the hardness of the steel increased quickly first and then slowly as the ageing period increased, and it finally reached a constant value within about 165 min, whereas at the ageing temperature of 500 °C, it only took 45 min to reach a constant hardness value. The maximal hardness was only about 435HV and 450HV for the test material after ageing at 350 °C and 380 °C, respectively. Among those ageing treatments, the hardness of the material reached a peak value of 490HV after ageing at 420 °C for165 min. Therefore, in order to achieve high tensile strength and clarify the influence of reverted austenite formation on tensile properties of parent material and weld, the post-weld ageing temperature is selected to be in the range of 420–580 °C.
Table 1 Chemical composition of maraging steel (wt.%). Element
Cr
Ni
Co
Mo
Ti
Cu
Si
C
Fe
Composition
13.62
6.82
0.19
0.76
0.34
0.76
1.43
0.050
Balance
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Table 2 Laser welding condition. Laser power
Travel speed
Focal point
Beam diameter
Shielding gas flow rate
3.0 kW
2.0 m/min
Surface
0.6 mm
Ar: 1.05 m3/h + He: 0.45 m3/h
Fig. 1. Microstructure of as-received material. Fig. 3. Variation in hardness with ageing period at different temperatures for experimental maraging steel solutionized at 850 °C.
Fig. 2. XRD patterns of as-received material and parent material zone of weldments aged at different temperatures.
3.2. Microstructure and hardness distribution across weldment prior to ageing treatments The low magnitude optical cross-sectional microstructure of the weld of the stainless maraging steel as shown in Fig. 4a, clearly exhibits three regions, i.e. the middle fusion zone, the HAZ surrounding the weld, and the unaffected parent metal zone during the welding thermal cycle. The microstructure at centre of the weld represented the most coarse equiaxed martensite grains (Fig. 4b), which were transformed from austenite during cooling in the fusion zone. At the location near the interface between fusion zone and HAZ, the martensite grain size became smaller, and there are signs of dendritic structure growing perpendicular to the fusion zone/HAZ interface (Fig. 4c. The martensitic grains in HAZ, adjoining the weld, are slightly coarser than those of other part of HAZ, at about 60 lm thick (also Fig. 4c). This was attributed to the fact that during welding the parent material was heated at
high temperature in the austenite region where grain growth occurred; upon cooling, the austenite transformed to martensite, which inherited the coarse austenite grain. Away from the weld, the martensitic grain size decreased, and become even finer than the grains in unaffected material zone due to the combined effect of rapid heating to appropriate temperature in austenite region and quick cooling during laser welding. The microstructure in HAZ (Fig. 4d) next to the unaffected parent material has a dark etching region where the martensite phase experienced peak temperatures in the range of 590–730 °C [20]. The dark etching region exhibited two-phase microstructure, which was identified by XRD analysis, having a small amount of reverted austenite (about 3.6% volume fraction) formed within the martensite matrix. Fig. 4e shows the martensitic microstructure of the unaffected parent material during the welding thermal cycle. The micro-hardness profile across the weld revealed a W-shaped distribution, as shown in Fig. 5. The central fusion zone had a hardness range of 360–365HV corresponding to the hardness of martensite, about 100HV lower than the unaffected parent material. The two valleys in the curve correspond to the hardness values of HAZs, with the lowest hardness being about 320HV. The lower hardness value in HAZs than that in fusion zone is attributed to the formation of a small amount of reverted austenite during laser welding. 3.3. Microstructure and hardness distribution across the weldments after ageing treatments For as-weld microstructure of weldment, after post-weld ageing at a temperature range of 420–500 °C, the lathy feature of martensite and prior austenite grain boundaries could be observed much more clearly within the three regions. However, the sub-microstructural change occurred within martensite matrix at this stage could not be observed due to the optical microscopy resolution. A great difference in microstructures occurred after ageing at 540 °C and 580 °C since reverted austenite was formed at these temperatures. Fig. 6 presents microstructures in different regions
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279
Fig. 4. Microstructures of weldment of maraging steel: (a) low magnitude microstructure, (b) fusion zone, (c) fusion zone/HAZ interface, (d) HAZ and (e) unaffected parent material zone.
of weldment aged at 580 °C. The coarse martensitic microstructure in fusion zone showed lathy feature (see Fig. 6a and b) much more clearly than it did in the as-weld condition owing to easy etching from microstructural change. Furthermore, it noted that a rather thick network boundary region was formed surrounding martensitic grains after ageing at 580 °C, which identified by XRD analysis as the reverted austenite. With increasing ageing temperature and the formation of reverted austenite, the microstructures in HAZ and parent material zone also became increasingly dark and the grain boundary much more distinguished (see Fig. 6c and d). Micro-hardness surveys across the welds presented in Fig. 7 revealed that at all different ageing temperatures the hardness in HAZ was increased substantially to a level above 420HV. A comparison of hardness between HAZ and treated parent material zone revealed that: (a) At the ageing temperature of 420 °C, the hardness in HAZ is still lower than that in the treated parent material after the treatment; the hardness in HAZ is increased to 450–460HV, whereas it is 490HV in the treated parent material
zone. (b) HAZ has a little higher or similar hardness value compared to the treated parent material after 460 °C ageing; for example, it has a hardness of 465–470HV, which was a little higher than that of the treated parent material. This is because hardness increases in HAZ and decreases in the treated parent material zone with increasing ageing temperature. (c) Hardness of HAZ decreased but was still a little higher than that of the treated parent material after ageing at 540 °C and 580 °C; for example, HAZ has a hardness of about 10–20HV higher than that in the treated parent material zone. The reason that the hardness was higher in HAZ may be due to the difference in grain size and original sub-microstructural state between HAZ and the parent material. In addition, an interesting phenomenon was noted that after ageing at all temperatures, hardness values of the treated parent material zone conformed well to the hardness curves subjected to corresponding ageing treatment shown in Fig. 2; unlike the hardness of the parent material, the hardness of the fusion zone increased with increasing ageing temperature until 500 °C, then dropped. In order to explain
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Fig. 5. Hardness distribution across weldment.
the contradicting result, the as-received material was heated to 1050 °C for 10 min then quenched in oil to establish a quick cooling condition from higher solution temperature, which was similar to fusion zone during laser welding. The measured variation in hardness with ageing period after ageing treatments at different temperatures is shown in Fig. 8. It was found that the hardness in fusion zone conformed well to the hardness curves from 1050 °C quenching. This could be explained by the two main factors rising from higher solution temperature experienced by the fusion zone during welding: (a) Effective dissolution of inclusion such as Ti(Cx,N1x) during solidification into melt and subsequently into austenite, and increasing contents of Ti, C, and N in austenite, leading to a higher over-ageing temperature. Shmuleritsh et al. [27]
investigated the stability of Ti(CxN1x) inclusion in austenite in maraging 250 steel, and found that the concentration of Ti(CxN1x) in austenite decreased monotonically from 11.25 104 mole to 8.75 104 mole as the temperature increased from 900 °C to 1050 °C. A comparison of the inclusions on optical microscope, as shown in Fig. 9, revealed that both the density and size of the inclusions were remarkably reduced in fusion zone in relation to the unaffected parent material zone. Previous study also demonstrated that Ti-containing Fe–Cr–Ni–Al–Mo stainless maraging steel (PH13-8 Mo) treated with higher temperature solution exhibited a faster increase in hardness during different ageing treatments, that is, the rate of the effective hardness increase was much higher than Ti-free Fe–Cr–Ni–Al–Mo alloy treated with lower temperature solution due to more antiphase boundary formed by the pairing of the clustered solute species or larger coherent precipitates, and the formation of reverted austenite being delayed simultaneously [28]. (b) Homogenizing of alloying elements such as Ti, Mo and Ni within austenite is followed by the homogenization of martensite. The distribution of Ni, Ti and Mo elements in the microstructures of fusion zone and parent material zone in as-weld condition in the form of EPMA X-ray maps are shown in Figs. 10 and 11. No micro-segregation of Ni, Ti, and Mo was found in fusion zone during solidification by laser welding, whereas there was a slight micro-segregation of Ni, Ti, and Mo on prior austenite grain boundary in the parent material. The Reduction in microsegregation by diffusion of alloying elements can avoid the premature of reverted austenite at low ageing temperature, and thus lead to a decrease in the volume fraction of reverted austenite. The XRD analysis of the fusion zone and the treated parent material zone revealed that the reversion of austenite was not measurable until ageing temperature was above 500 °C, and the volume fractions of reverted austenite were 12.0% and 15.3% for fusion zone, 5.9% and 7.1% for the treated parent material zone after ageing at
Fig. 6. Microstructures of weldment after ageing at 580 °C: (a) fusion zone, (b) fusion zone/HAZ interface, (c) HAZ and (d) parent material zone.
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Fig. 7. Hardness distribution across the weldments after ageing at different temperatures: (a) 420 °C, (b) 460 °C, (c) 540 °C and (d) 580 °C.
tural evolution mechanisms rising from difference in alloying element contents within martensite. The different ageing behaviours between parent material zone and fusion zone are not fully clarified in detail in the present study, further investigation is needed. 3.4. Tensile properties of weldments after ageing treatment
Fig. 8. Variation in hardness with ageing time at different ageing temperatures for experimental maraging steel solutionized at 1050 °C.
540 °C and 580 °C, respectively. Therefore, the premature and the amount of reverted austenite were not the predominant factors influencing the variation of the different hardness with ageing temperature between fusion zone and the treated parent material zone. The dominant factor could be the different sub-microstruc-
Tests on the transverse tensile properties of the welds were performed, and the results of the tensile properties subject to various ageing treatments are presented in Table 3. For comparison purpose, the tensile properties of as-received parent material are also included. Previous studies have reported that in most cases, a welding efficiency over 90% was achieved for those maraging steels welds in the solutionized conditions. However, for the present material in aged condition with high-strength, laser welding consequently led a lower welding efficiency of 76.3%; fracture occurred at a location of HAZ where the hardness was about 320HV, much lower than the hardness of the unaffected parent material zone, 460HV. It is noticeable that all the applied postweld heat-treatments can significantly improve tensile properties of weldments and only those welds subjected to ageing in temperature range of 420–500 °C outperform the as-received material. In most cases of the present study, the tensile yield strength ratio of aged weldments to as-received parent material, which reflects how yield strength is improved after post-weld heat treatment, was over 100% except the welds aged at 540 °C and 580 °C; fracture occurred in the treated parent material zone instead of the weld.
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Tabor [29] proposed a simple relation between hardness H and yield strength Y, H = cY for ductile metals, where c is 3. From the results of hardness distribution and tensile tests aforementioned, we notice that the tensile properties of maraging steel laser weldments both before and after post-treatment depend greatly on hardness distribution across the welds, especially hardness in HAZ or parent material zone which is prone to fracture as it is usually the softest and weakest location. Therefore, we propose the use of two parameters (ratios) for the prediction of fracture location and improvement in yield strength as follows:
Fig. 9. Typical optical micrographs of the polished weldment: (a) parent material zone and (b) fusion zone. Inclusions are marked by the circles.
Therefore, a concept of joint efficiency is defined by the tensile yield strength ratio of aged weldments to as-received parent material. The tensile properties, such as yield strength and ultimate strength, of weldments aged at 420 °C and 460 °C, were much higher compared with as weld weldment. Based on the fact that treatment at 420 °C resulted in fracture occurring in HAZ and the rather low resulting elongation (2.4%), post-weld ageing treatments of the material in this study were, therefore, carried out at 460 °C.
RF ¼
HHAZ HTP
RI ¼
HTP ; HP
RI ¼
HHAZ ; HP
ð1Þ
if RF > 1
if RF < 1
ð2Þ
ð3Þ
where HHAZ is the hardness in HAZ, HTP is the hardness of the aged parent metal zone, and Hp is the hardness of as-received parent material. When RF, the hardness ratio of HAZ to the aged parent material zone, is lower than 1.0, fracture usually occurs in HAZ, and when it is higher than 1.0, fracture usually occurs in parent material zone. This method for estimating fracture location can also be applied to similar and dissimilar laser welds of other materials such as TRIP700 steel, UNC-C17200 copper beryllium alloy, magnesium alloy AZ31B, TWIP and TRIP steels [30–33], the fractures of these weldments occurred in soft HAZs or base metals, which agree well with Eq. (1). RI is the hardness ratio of heat-treated parent metal zone or HAZ to as-received parent metal; it can be used to estimate the improvement of tensile properties through weld-post heat treatment. If fracture occurs in the aged parent metal zone, i.e. RF greater than 1.0, Eq. (2) can be used, and RI over 1.0 means higher tensile properties of weldment than as received parent metal through heat treatment (here we refer to it as improvement) and vice versa (here refer to it as without improvement). If fracture happens to HAZ, Eq. (3) can be used to estimate improvement efficiency after weld-post treatment, RI over 1.0 means weldment having higher
Fig. 10. SEM micrograph and EDS mapping of fusion zone in as-weld condition: (a) SEM micrograph, (b), (c) and (d) EDS mapping for Ni, Mo and Ti, respectively.
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Fig. 11. SEM micrograph and EDS mapping of parent material zone in as-weld condition: (a) SEM micrograph, (b), (c) and (d) EDS mapping for Ni, Mo and Ti, respectively.
Table 3 Tensile properties of parent metal and transverse weld specimens.
a
Ageing temperature (°C)
0.2%YS (MPa)
UTS (MPa)
Elongation (%)
Failure location
Joint efficiencya (%)
420 460 500 540 580 As weld As-received material
1363.7 1493.9 1337.1 1239.2 1196.3 994.8 1303.8
1477.1 1512.8 1411.6 1366.8 1259.5 1272.1 1426.3
2.4 6.3 8.1 7.9 10.5 1.6 5.6
HAZ PM PM PM PM HAZ –
104.6 114.5 102.5 95.0 91.7 76.3 –
Yield strength ratio of aged weldment to as-received material.
tensile properties than parent metal after heat treatment and vice versa. Based on the two hardness ratios, a graph is established for determination of fracture location and post-weld heat-treatment efficiency of maraging steel laser weldments using the results from the present study, as shown in Fig. 12. It consists of four regions: I and II are parent material fracture-improvement region and HAZ fracture-improvement region, respectively. III and IV are HAZ fracture-without improvement region and parent fracture-without improvement region, respectively. In order to identify the validity and applicability of this graph, results of similar maraging steel in Ref. [20] were transformed into hardness ratios using Eqs. (1)–(3), and included in Fig. 12 as well. In Ref. [20] Reedy et al. investigated microstructure and mechanical properties of laser weldments of a 18 Ni (250 grade) maraging steel as-received in solutionized condition, which were subjected to three kinds of post-weld heat treatments including direct ageing at 480 °C, solutionizing at 815 °C + ageing at 480 °C, and homogenization at 1150 °C + solutionizing at 815 °C + ageing at 480 °C. The hardness across the weldments, tensile properties and fracture locations were systemically presented after each treatment, and in most cases the results can be well represented in Fig. 12. However, the fracture occurred in the weld (fusion zone) region instead of HAZ for weldment treated by homogenization at 1150 °C + solutionizing at 815 °C + ageing at 480 °C, this indicates that there is one limitation for proposed methodology that fracture could occur in the weld region if the hardness value is equal in fusion zone, HAZ, and treated parent material, that is, RF = 1.
Fig. 12. Graph for determination of fracture location and post-weld heat-treatment efficiency. The tensile efficiency data are put in parentheses.
4. Conclusions In the present investigation, the hardness profile and tensile properties of laser stainless maraging steel weldments subjected to post-weld ageing treatments at different temperatures were tested. Ageing enhanced hardness in HAZ and fusion zone in
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temperature range of 420–460 °C, consequently improved tensile efficiency. Hardness in fusion zone and HAZ varied differently from that of the treated parent material zone with ageing temperature; above 420 °C it was higher in fusion zone and HAZ than in the treated parent material zone, therefore fracture was located in the parent material zone. Based on the correlation between hardness distribution and tensile properties, a graph was constructed for determination of fracture location and post-weld heat treatment efficiency using two parameters, i.e. hardness ratio of HAZ to the treated parent material zone and hardness ratio of HAZ to as-received parent material. Experimental and reference results verified the validity and applicability of this graph.
Acknowledgements The authors gratefully acknowledge the financial support under the Project 985-automotive engineering of Jilin University and Technology Development Program of Jilin Province (Nos. 20050509, 20080314).
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[12] Ping DH, Ohnuma M, Hirakawa Y, Kadoya Y, Hono K. Microstructural evolution in 13Cr–8Ni–2.5Mo–2Al martensitic precipitation-hardened stainless steel. Mater Sci Eng A 2005;394A:285–95. [13] Lecomte JB, Servant C, Cizeron G. A comparison of the structural evolution occurring during anisothermal or isothermal treatments in the case of nickel and manganese type maraging alloys. J Mater Sci 1985;20:3339–52. [14] Tewari R, Mazumder S, Batra IS, Dey GK, Banerjee S. Precipitation in 18 wt.% Ni maraging steel of grade 350. Acta Mater 2000;48:1187–200. [15] Tavares SSM, Abreu HFG, Neto JM, Silva Da MR, Popa I. A thermomagnetic study of the martensite–austenite phase transition in the maraging 350 steel. J Alloys Compd 2002;358:152–6. [16] Peters DT. A study of austenite reversion during aging of maraging steels. Trans ASM 1968;61:62–74. [17] Schober M, Schnitzer R, Leitner H. Precipitation evolution in a Ti-free and Ticontaining stainless maraging steel. Ultramicroscopy 2009;109:553–62. [18] Gemperle A, Gemperlová J, Sha W, Smith GDW. Aging behaviour of cobalt free chromium containing maraging steels. Mater Sci Technol 1992;8:546–54. [19] Boniszewski T, Kenyon DM. Examination of electron beam welds in 18%Ni/Co/ Mo maraging steel sheet. Br Weld J 1966;13:415–35. [20] Reddy GM, Rao VV, Raju AV. The effect of post-weld heat treatments on the microstructure and mechanical properties of maraging steel laser weldments. Proc. ImechE Part L: J Mater: Des Appl 2011;223:149–59. [21] Duffey FD, Sutar W. Submerged arc welding of 18% Nickel maraging steel. Weld J 1965;44:251–63. [22] Sundaresan S, Manirajan M, Rao BN. On the fracture toughness evaluation in weldments of a maraging steel rocket motor case. Mater Des 2010;31:4921–6. [23] Ramana PV, Reddy GM, Mohandas T, Gupta. Microstructure and residual stress distribution of similar and dissimilar electron beam welds-maraging steel to medium alloy medium carbon steel. Mater Des 2010;31:749–60. [24] Nagarajan KV. Welding aspects of maraging steels. In: Proceedings of the conference on steels for engineering industries-trends in weldability, Indian Institute of Metals, Tiruchirapalli, India; 1998. p. 43–53. [25] Shamantha CR, Narayanan R, Iyer KJ, Radhakrishnan VM, Seashadri SK, Sundararajan S, et al. Microstructural changes during welding and subsequent heat treatment of 18 Ni (250-grade) maraging steel. Mater Sci Eng A 2000;287A:43–51. [26] Rao VV, Reddy GM, Raju AVS. Influence of post-weld heat treatments on microstructure and mechanical properties of gas tungsten arc maraging steel weldments. Mater Sci Technol 2010;26:1459–68. [27] Shmuleritsh M, Pinkces M, Weizman A, Frage N. Evolution of Ti-based nonmetallic inclusions during solution treatment of maraging 250 steel: Thermodynamic calculations and experimental verification. Metall Trans B 2011;42B:550–6. [28] Leitner H, Schober M, Schnitzer R, Zinner S. Strengthening behavior of Fe–Cr– Ni–Al–(Ti) maraging steels. Mater Sci Eng A 2011;528A:5264–70. [29] Tabor D. A simple theory of static and dynamic hardness. Proc R Soc A 1948;192:247–74. [30] Reisgen U, Schleser M, Mokrov O, Ahmed E. Numerical and experimental investigation of tensile behavior of laser beam weld TRIP700 steel. ISIJ Int 2011;51:429–34. [31] Akbari Mousavi SAA, Nikneijad ST. Study on the microstructure and mechanical properties of Nd: YAG pulsed laser beam weld of UNS-C17200 copper beryllium alloy. J Mater Process Technol 2010;210:1472–81. [32] Coelho RS, Kostka A, Pinto H, Riekehr S, Kocak M, Pyzalla AR. Microstructure and mechanical properties of magnesium alloy AZ31B laser beam welds. Mater Sci Eng A 2008;485A:20–30. [33] Mujica L, Weber S, Pinto H, Thomy C, Vollertsen F. Microstructure and mechanical properties of laser-welded joints of TWIP and TRIP steels. Mater Sci Eng A 2010;527A:2071–8.
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Improvement of mechanical properties of stainless maraging steel laser weldments by post-weld ageing treatments Jian An ⇑, Fanyan Meng, Xiaoxia Lv, Haiyuan Liu, Xiaoxi Gao, Yuanbo Wang, You Lu Key Laboratory of Automobile Materials, Ministry of Education, Department of Materials Science and Engineering, Nanling Campus of Jilin University, Changchun 130025, People’s Republic of China
a r t i c l e
i n f o
Article history: Received 4 January 2012 Accepted 13 March 2012 Available online 28 March 2012 Keywords: A. Ferrous metals and alloys D. Welding E. Mechanical
a b s t r a c t The response of stainless maraging steel weldments to post-weld ageing treatment has been investigated. Post-weld ageing was performed at five different temperatures, viz., 420 °C, 460 °C, 500 °C, 540 °C, and 580 °C. Metallographic characterization of weldment revealed three zones, namely fusion zone, heataffected zone (HAZ) and unaffected parent metal zone. Hardness and tensile properties were evaluated after ageing at different temperatures. Hardness in HAZ and fusion zone varied with ageing temperature differently from that of the parent metal; it became higher in HAZ and fusion zone than in parent metal zone above 420 °C. Among the applied ageing treatments, ageing at 460 °C achieved the highest tensile strength. A graph was constructed for determination of fracture location and post-weld heat treatment efficiency based on experimental results, using hardness ratio of HAZ to the treated parent material and hardness ratio of HAZ to the as-received parent material. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction For several decades, two major types of maraging steels, Fe–Ni system and Fe–Cr–Ni system alloys, have been developed and used in defense, power, and other commercial industries due to their superior properties such as ultra-high strength, excellent fracture toughness, good machining properties and weldability [1,2]. The most commonly used Fe–Ni maraging steel composition is based on 18 wt.% Ni with significant levels of Co (8–13 wt.%), while the subsequently developed Co-free maraging steels contain Fe, Ni, Mo and Ti [3–7]. The composition of Fe–Cr–Ni maraging steels is based on 10.2–16 wt.% Cr with considerable levels of Ni (3.9–10.8 wt.%) and other elements such as Mo, Ti, Al and Cu. Fe–Cr–Ni maraging steel usually possesses good corrosion resistance [8–12]. It is well known that the ultra-high strength of maraging steels is mainly attributed to the precipitation of densely distributed fine intermetallic precipitates in a martensitic matrix, which strengthen the alloys through Orowan-type mechanism. Typically, this kind of microstructure is adjusted by a simple two-step heat-treatment process which consists of solution annealing to produce a complete martensitic matrix, and subsequently ageing at intermediate temperatures (400 °C to 600 °C) to cause precipitation hardening. The ageing behaviour of maraging steels has been extensively investigated over recent decades. By means of high-resolution methods, such as atom probe ⇑ Corresponding author. Tel.: +86 431 85095874; fax: +86 431 85095876. E-mail address: [email protected] (J. An). 0261-3069/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2012.03.020
tomography (APT), transmission electron microscopy and smallangle neutron scattering, the size and chemical composition of nanometer-sized cluster of precipitates are identified in numerous maraging steels. For example, 18% Ni maraging steels exhibit ordered and coherent phases such as l, S and X (<450 °C), Ni3 (Mo, Ti), and Fe2Mo phases (450–500 °C), and austenite precipitation (500 °C-As) [13–16]. For Fe–Cr–Ni–Al–Ti stainless maraging steel, which contains in addition 0.5–1.0 at.% Si, a spherical G-phase with the chemical composition of Ti6Si7Ni16, and a rodshaped Ni3(Ti,Al) phase were observed [17,18]. Weldability is a unique property which makes maraging steels especially suitable for processing in aviation and space industry, since welding of maraging steels can be performed, without preheating, by processes such as electron beam, laser welding, and submerged arc [19–22]. Ramana et al. [23] studied the microstructure and residual stress distribution of dissimilar electron beam weld of 18 Ni maraging steel and medium alloy medium carbon steel, it was observed that the residual stress distribution would be more compressive if the maraging steel was in soft condition. Among these processes, gas tungsten-arc welding is widely utilized in view of the consistency of weld equality and overall efficiency [24]. There are various reports published recently on gas tungsten-arc welding of maraging steels. As the 18 Ni maraging steels represent majority of maraging steel products, the associated low tensile property and toughness rising from premature of reverted austenite, cracking in heat-affected zone (HAZ) immediately adjacent to fusion line associated with gas tungsten-arc welding were the subjects of several reports in the literature. Shamantha et al.
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[25] investigated microstructure changes during welding and subsequent post-weld ageing of 18 Ni maraging steel at various temperatures and concluded that pronounced segregation of Ti and Mo elements along interdendrite and intercellular boundaries in the weld metal led to austenite reversion at temperatures much lower than in unweld material. Similarly, Reddy et al. [26] reported that solutionized treatment of 18 Ni maraging steel weldments at 815 °C could not remove the segregation of Ti and Mo elements, only homogenization treatment at 1150 °C succeeded, and consequently resulted in optimum tensile properties. In most cases, welding 18 Ni maraging steels in the solutionized condition followed by post-weld ageing treatment at a relatively low temperature can attain weld efficiency over 90%. Laser welding is a well established technique that possesses remarkable characteristics such as high welding speed, low heat input, low distortion, deep penetration, narrow HAZ and free from filler wire. Therefore laser welding of maraging steels is a useful process for critical application such as aeronautics and space technology equipments. However, up to date, investigations on postweld heat treatment of maraging steel laser welds still lack systemically experimental data and need further understanding the responses from Fe–Cr–Ni system maraging steels other than 18 Ni system alloys. Reddy et al. [20] investigated the role of reverted austenite through various post-weld heat treatments on the microstructure and mechanical properties of 18% Ni (250 grade) maraging steel laser welds, and found that the homogenized + solutionized + aged weldments exhibited higher tensile properties and toughness compared to directly aged and solutionized + aged weldments due to the reduction in micro-segregation by diffusion of alloying elements such as Mo and Ti from the cell boundaries and the resulting absence of reverted austenite. To the best of our knowledge, post-weld heat treatments are mostly performed on 18 Ni maraging steels which were usually in solutionized condition, hardly on Fe–Cr–Ni system maraging steels. The considerable difference in chemical composition between 18 Ni maraging steels and Fe–Cr–Ni system maraging steels may result in quite different response to post-weld heat treatment in some aspects of reverted austenite formation, alloying element distribution and mechanical properties. Therefore, the present paper focus on microstructural evolution and tensile properties of Fe–Cr–Ni maraging stainless steel laser weldments through post-weld ageing, the studied material was in aged condition prior to welding. This process is more practical and economical for fabrication of large structures compared with post-weld heat treatment applied to maraging steel in solutionized condition. For this state of parent material, only localized area surrounding the weld, instead of the whole structure, needs post-weld heat treatment, which can be accomplished easily by using small-sized heat-treatment equipment. The experiments were designed to study the effect of post-weld ageing on tensile properties in a wide temperature range from 420 °C to 580 °C. A useful graph was established for assessment of fracture location and improvement in tensile yield strength after post-weld treatment using two parameters indicating hardness ratios between different regions in laser weldments. 2. Experimental details The test material is a type of Fe–Cr–Ni–Co–Mo–Ti–Cu stainless maraging steel in the form of 3.0-mm-thick sheet with composi-
tion given in Table 1, which was received in the short-term 420°C/1 h/air cooling aged condition. Rectangular sheets in the dimensions of 200 100 3 mm3 were machined from the as-received sheets and were autogenously laser beam welded using a CO2 laser. The parameters used for laser welding are listed in Table 2. The welds were subjected to ageing treatments at a temperature range varying from 420 °C to 580 °C for 2 h to examine the effect of ageing temperature on the microstructure and mechanical properties of welds. Metallographic microstructures of the weldments in the as-weld and post-weld aged conditions were characterized using a LEXT-OLS3000 confocal scanning laser microscope. Vilella’s reagent (picric acid 1 g + HCl 5 ml + ethanol 100 ml) was used for etching fusion zone, HAZ, and parent material to observe the microstructure. Mechanical properties including hardness and tensile properties of welds in different aged conditions were evaluated. Microhardness survey was conducted in the midsection of weld coupons with 200 g load on a Vickers hardness indentation tester. Tensile tests were carried out on a MTS 810 tensile testing system using specimens having a reduced cross-section of 5.0 mm 3.0 mm and a gauge length of 30.0 mm, which were machined from various heat-treated weld coupons with the welds located in the middle position. For each heat treatment temperature level, at least three specimens are tested to obtain an average value of tensile properties. Revered austenite contents in the welds and parent material in various heat-treated conditions were estimated using a Rigaku X-ray diffractometer (XRD) using Cu Ka radiation. The volume fraction of austenite was estimated from measurements of the integrated intensities of martensite and austenite peaks, assuming they are the only phases present. 3. Results and discussion 3.1. Microstructure and hardness dependence on ageing temperature for parent material The as-received material had a microstructure of complete martensite with fine grain size about 5–10 lm observed on optical microscope, as shown in Fig. 1. XRD analysis further revealed no austenite phase present, as shown in Fig. 2. Hardness dependence of this metal on ageing temperature and period was examined by heating it to a conventional solution temperature of 850 °C for 1 h followed by ageing in temperature range of 380–580 °C for various periods. From hardness curves in Fig. 3, it can be seen that the hardness depends greatly on the ageing temperature and period. At ageing temperatures lower than 500 °C, the hardness of the steel increased quickly first and then slowly as the ageing period increased, and it finally reached a constant value within about 165 min, whereas at the ageing temperature of 500 °C, it only took 45 min to reach a constant hardness value. The maximal hardness was only about 435HV and 450HV for the test material after ageing at 350 °C and 380 °C, respectively. Among those ageing treatments, the hardness of the material reached a peak value of 490HV after ageing at 420 °C for165 min. Therefore, in order to achieve high tensile strength and clarify the influence of reverted austenite formation on tensile properties of parent material and weld, the post-weld ageing temperature is selected to be in the range of 420–580 °C.
Table 1 Chemical composition of maraging steel (wt.%). Element
Cr
Ni
Co
Mo
Ti
Cu
Si
C
Fe
Composition
13.62
6.82
0.19
0.76
0.34
0.76
1.43
0.050
Balance
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Table 2 Laser welding condition. Laser power
Travel speed
Focal point
Beam diameter
Shielding gas flow rate
3.0 kW
2.0 m/min
Surface
0.6 mm
Ar: 1.05 m3/h + He: 0.45 m3/h
Fig. 1. Microstructure of as-received material. Fig. 3. Variation in hardness with ageing period at different temperatures for experimental maraging steel solutionized at 850 °C.
Fig. 2. XRD patterns of as-received material and parent material zone of weldments aged at different temperatures.
3.2. Microstructure and hardness distribution across weldment prior to ageing treatments The low magnitude optical cross-sectional microstructure of the weld of the stainless maraging steel as shown in Fig. 4a, clearly exhibits three regions, i.e. the middle fusion zone, the HAZ surrounding the weld, and the unaffected parent metal zone during the welding thermal cycle. The microstructure at centre of the weld represented the most coarse equiaxed martensite grains (Fig. 4b), which were transformed from austenite during cooling in the fusion zone. At the location near the interface between fusion zone and HAZ, the martensite grain size became smaller, and there are signs of dendritic structure growing perpendicular to the fusion zone/HAZ interface (Fig. 4c. The martensitic grains in HAZ, adjoining the weld, are slightly coarser than those of other part of HAZ, at about 60 lm thick (also Fig. 4c). This was attributed to the fact that during welding the parent material was heated at
high temperature in the austenite region where grain growth occurred; upon cooling, the austenite transformed to martensite, which inherited the coarse austenite grain. Away from the weld, the martensitic grain size decreased, and become even finer than the grains in unaffected material zone due to the combined effect of rapid heating to appropriate temperature in austenite region and quick cooling during laser welding. The microstructure in HAZ (Fig. 4d) next to the unaffected parent material has a dark etching region where the martensite phase experienced peak temperatures in the range of 590–730 °C [20]. The dark etching region exhibited two-phase microstructure, which was identified by XRD analysis, having a small amount of reverted austenite (about 3.6% volume fraction) formed within the martensite matrix. Fig. 4e shows the martensitic microstructure of the unaffected parent material during the welding thermal cycle. The micro-hardness profile across the weld revealed a W-shaped distribution, as shown in Fig. 5. The central fusion zone had a hardness range of 360–365HV corresponding to the hardness of martensite, about 100HV lower than the unaffected parent material. The two valleys in the curve correspond to the hardness values of HAZs, with the lowest hardness being about 320HV. The lower hardness value in HAZs than that in fusion zone is attributed to the formation of a small amount of reverted austenite during laser welding. 3.3. Microstructure and hardness distribution across the weldments after ageing treatments For as-weld microstructure of weldment, after post-weld ageing at a temperature range of 420–500 °C, the lathy feature of martensite and prior austenite grain boundaries could be observed much more clearly within the three regions. However, the sub-microstructural change occurred within martensite matrix at this stage could not be observed due to the optical microscopy resolution. A great difference in microstructures occurred after ageing at 540 °C and 580 °C since reverted austenite was formed at these temperatures. Fig. 6 presents microstructures in different regions
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279
Fig. 4. Microstructures of weldment of maraging steel: (a) low magnitude microstructure, (b) fusion zone, (c) fusion zone/HAZ interface, (d) HAZ and (e) unaffected parent material zone.
of weldment aged at 580 °C. The coarse martensitic microstructure in fusion zone showed lathy feature (see Fig. 6a and b) much more clearly than it did in the as-weld condition owing to easy etching from microstructural change. Furthermore, it noted that a rather thick network boundary region was formed surrounding martensitic grains after ageing at 580 °C, which identified by XRD analysis as the reverted austenite. With increasing ageing temperature and the formation of reverted austenite, the microstructures in HAZ and parent material zone also became increasingly dark and the grain boundary much more distinguished (see Fig. 6c and d). Micro-hardness surveys across the welds presented in Fig. 7 revealed that at all different ageing temperatures the hardness in HAZ was increased substantially to a level above 420HV. A comparison of hardness between HAZ and treated parent material zone revealed that: (a) At the ageing temperature of 420 °C, the hardness in HAZ is still lower than that in the treated parent material after the treatment; the hardness in HAZ is increased to 450–460HV, whereas it is 490HV in the treated parent material
zone. (b) HAZ has a little higher or similar hardness value compared to the treated parent material after 460 °C ageing; for example, it has a hardness of 465–470HV, which was a little higher than that of the treated parent material. This is because hardness increases in HAZ and decreases in the treated parent material zone with increasing ageing temperature. (c) Hardness of HAZ decreased but was still a little higher than that of the treated parent material after ageing at 540 °C and 580 °C; for example, HAZ has a hardness of about 10–20HV higher than that in the treated parent material zone. The reason that the hardness was higher in HAZ may be due to the difference in grain size and original sub-microstructural state between HAZ and the parent material. In addition, an interesting phenomenon was noted that after ageing at all temperatures, hardness values of the treated parent material zone conformed well to the hardness curves subjected to corresponding ageing treatment shown in Fig. 2; unlike the hardness of the parent material, the hardness of the fusion zone increased with increasing ageing temperature until 500 °C, then dropped. In order to explain
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Fig. 5. Hardness distribution across weldment.
the contradicting result, the as-received material was heated to 1050 °C for 10 min then quenched in oil to establish a quick cooling condition from higher solution temperature, which was similar to fusion zone during laser welding. The measured variation in hardness with ageing period after ageing treatments at different temperatures is shown in Fig. 8. It was found that the hardness in fusion zone conformed well to the hardness curves from 1050 °C quenching. This could be explained by the two main factors rising from higher solution temperature experienced by the fusion zone during welding: (a) Effective dissolution of inclusion such as Ti(Cx,N1x) during solidification into melt and subsequently into austenite, and increasing contents of Ti, C, and N in austenite, leading to a higher over-ageing temperature. Shmuleritsh et al. [27]
investigated the stability of Ti(CxN1x) inclusion in austenite in maraging 250 steel, and found that the concentration of Ti(CxN1x) in austenite decreased monotonically from 11.25 104 mole to 8.75 104 mole as the temperature increased from 900 °C to 1050 °C. A comparison of the inclusions on optical microscope, as shown in Fig. 9, revealed that both the density and size of the inclusions were remarkably reduced in fusion zone in relation to the unaffected parent material zone. Previous study also demonstrated that Ti-containing Fe–Cr–Ni–Al–Mo stainless maraging steel (PH13-8 Mo) treated with higher temperature solution exhibited a faster increase in hardness during different ageing treatments, that is, the rate of the effective hardness increase was much higher than Ti-free Fe–Cr–Ni–Al–Mo alloy treated with lower temperature solution due to more antiphase boundary formed by the pairing of the clustered solute species or larger coherent precipitates, and the formation of reverted austenite being delayed simultaneously [28]. (b) Homogenizing of alloying elements such as Ti, Mo and Ni within austenite is followed by the homogenization of martensite. The distribution of Ni, Ti and Mo elements in the microstructures of fusion zone and parent material zone in as-weld condition in the form of EPMA X-ray maps are shown in Figs. 10 and 11. No micro-segregation of Ni, Ti, and Mo was found in fusion zone during solidification by laser welding, whereas there was a slight micro-segregation of Ni, Ti, and Mo on prior austenite grain boundary in the parent material. The Reduction in microsegregation by diffusion of alloying elements can avoid the premature of reverted austenite at low ageing temperature, and thus lead to a decrease in the volume fraction of reverted austenite. The XRD analysis of the fusion zone and the treated parent material zone revealed that the reversion of austenite was not measurable until ageing temperature was above 500 °C, and the volume fractions of reverted austenite were 12.0% and 15.3% for fusion zone, 5.9% and 7.1% for the treated parent material zone after ageing at
Fig. 6. Microstructures of weldment after ageing at 580 °C: (a) fusion zone, (b) fusion zone/HAZ interface, (c) HAZ and (d) parent material zone.
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Fig. 7. Hardness distribution across the weldments after ageing at different temperatures: (a) 420 °C, (b) 460 °C, (c) 540 °C and (d) 580 °C.
tural evolution mechanisms rising from difference in alloying element contents within martensite. The different ageing behaviours between parent material zone and fusion zone are not fully clarified in detail in the present study, further investigation is needed. 3.4. Tensile properties of weldments after ageing treatment
Fig. 8. Variation in hardness with ageing time at different ageing temperatures for experimental maraging steel solutionized at 1050 °C.
540 °C and 580 °C, respectively. Therefore, the premature and the amount of reverted austenite were not the predominant factors influencing the variation of the different hardness with ageing temperature between fusion zone and the treated parent material zone. The dominant factor could be the different sub-microstruc-
Tests on the transverse tensile properties of the welds were performed, and the results of the tensile properties subject to various ageing treatments are presented in Table 3. For comparison purpose, the tensile properties of as-received parent material are also included. Previous studies have reported that in most cases, a welding efficiency over 90% was achieved for those maraging steels welds in the solutionized conditions. However, for the present material in aged condition with high-strength, laser welding consequently led a lower welding efficiency of 76.3%; fracture occurred at a location of HAZ where the hardness was about 320HV, much lower than the hardness of the unaffected parent material zone, 460HV. It is noticeable that all the applied postweld heat-treatments can significantly improve tensile properties of weldments and only those welds subjected to ageing in temperature range of 420–500 °C outperform the as-received material. In most cases of the present study, the tensile yield strength ratio of aged weldments to as-received parent material, which reflects how yield strength is improved after post-weld heat treatment, was over 100% except the welds aged at 540 °C and 580 °C; fracture occurred in the treated parent material zone instead of the weld.
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Tabor [29] proposed a simple relation between hardness H and yield strength Y, H = cY for ductile metals, where c is 3. From the results of hardness distribution and tensile tests aforementioned, we notice that the tensile properties of maraging steel laser weldments both before and after post-treatment depend greatly on hardness distribution across the welds, especially hardness in HAZ or parent material zone which is prone to fracture as it is usually the softest and weakest location. Therefore, we propose the use of two parameters (ratios) for the prediction of fracture location and improvement in yield strength as follows:
Fig. 9. Typical optical micrographs of the polished weldment: (a) parent material zone and (b) fusion zone. Inclusions are marked by the circles.
Therefore, a concept of joint efficiency is defined by the tensile yield strength ratio of aged weldments to as-received parent material. The tensile properties, such as yield strength and ultimate strength, of weldments aged at 420 °C and 460 °C, were much higher compared with as weld weldment. Based on the fact that treatment at 420 °C resulted in fracture occurring in HAZ and the rather low resulting elongation (2.4%), post-weld ageing treatments of the material in this study were, therefore, carried out at 460 °C.
RF ¼
HHAZ HTP
RI ¼
HTP ; HP
RI ¼
HHAZ ; HP
ð1Þ
if RF > 1
if RF < 1
ð2Þ
ð3Þ
where HHAZ is the hardness in HAZ, HTP is the hardness of the aged parent metal zone, and Hp is the hardness of as-received parent material. When RF, the hardness ratio of HAZ to the aged parent material zone, is lower than 1.0, fracture usually occurs in HAZ, and when it is higher than 1.0, fracture usually occurs in parent material zone. This method for estimating fracture location can also be applied to similar and dissimilar laser welds of other materials such as TRIP700 steel, UNC-C17200 copper beryllium alloy, magnesium alloy AZ31B, TWIP and TRIP steels [30–33], the fractures of these weldments occurred in soft HAZs or base metals, which agree well with Eq. (1). RI is the hardness ratio of heat-treated parent metal zone or HAZ to as-received parent metal; it can be used to estimate the improvement of tensile properties through weld-post heat treatment. If fracture occurs in the aged parent metal zone, i.e. RF greater than 1.0, Eq. (2) can be used, and RI over 1.0 means higher tensile properties of weldment than as received parent metal through heat treatment (here we refer to it as improvement) and vice versa (here refer to it as without improvement). If fracture happens to HAZ, Eq. (3) can be used to estimate improvement efficiency after weld-post treatment, RI over 1.0 means weldment having higher
Fig. 10. SEM micrograph and EDS mapping of fusion zone in as-weld condition: (a) SEM micrograph, (b), (c) and (d) EDS mapping for Ni, Mo and Ti, respectively.
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Fig. 11. SEM micrograph and EDS mapping of parent material zone in as-weld condition: (a) SEM micrograph, (b), (c) and (d) EDS mapping for Ni, Mo and Ti, respectively.
Table 3 Tensile properties of parent metal and transverse weld specimens.
a
Ageing temperature (°C)
0.2%YS (MPa)
UTS (MPa)
Elongation (%)
Failure location
Joint efficiencya (%)
420 460 500 540 580 As weld As-received material
1363.7 1493.9 1337.1 1239.2 1196.3 994.8 1303.8
1477.1 1512.8 1411.6 1366.8 1259.5 1272.1 1426.3
2.4 6.3 8.1 7.9 10.5 1.6 5.6
HAZ PM PM PM PM HAZ –
104.6 114.5 102.5 95.0 91.7 76.3 –
Yield strength ratio of aged weldment to as-received material.
tensile properties than parent metal after heat treatment and vice versa. Based on the two hardness ratios, a graph is established for determination of fracture location and post-weld heat-treatment efficiency of maraging steel laser weldments using the results from the present study, as shown in Fig. 12. It consists of four regions: I and II are parent material fracture-improvement region and HAZ fracture-improvement region, respectively. III and IV are HAZ fracture-without improvement region and parent fracture-without improvement region, respectively. In order to identify the validity and applicability of this graph, results of similar maraging steel in Ref. [20] were transformed into hardness ratios using Eqs. (1)–(3), and included in Fig. 12 as well. In Ref. [20] Reedy et al. investigated microstructure and mechanical properties of laser weldments of a 18 Ni (250 grade) maraging steel as-received in solutionized condition, which were subjected to three kinds of post-weld heat treatments including direct ageing at 480 °C, solutionizing at 815 °C + ageing at 480 °C, and homogenization at 1150 °C + solutionizing at 815 °C + ageing at 480 °C. The hardness across the weldments, tensile properties and fracture locations were systemically presented after each treatment, and in most cases the results can be well represented in Fig. 12. However, the fracture occurred in the weld (fusion zone) region instead of HAZ for weldment treated by homogenization at 1150 °C + solutionizing at 815 °C + ageing at 480 °C, this indicates that there is one limitation for proposed methodology that fracture could occur in the weld region if the hardness value is equal in fusion zone, HAZ, and treated parent material, that is, RF = 1.
Fig. 12. Graph for determination of fracture location and post-weld heat-treatment efficiency. The tensile efficiency data are put in parentheses.
4. Conclusions In the present investigation, the hardness profile and tensile properties of laser stainless maraging steel weldments subjected to post-weld ageing treatments at different temperatures were tested. Ageing enhanced hardness in HAZ and fusion zone in
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temperature range of 420–460 °C, consequently improved tensile efficiency. Hardness in fusion zone and HAZ varied differently from that of the treated parent material zone with ageing temperature; above 420 °C it was higher in fusion zone and HAZ than in the treated parent material zone, therefore fracture was located in the parent material zone. Based on the correlation between hardness distribution and tensile properties, a graph was constructed for determination of fracture location and post-weld heat treatment efficiency using two parameters, i.e. hardness ratio of HAZ to the treated parent material zone and hardness ratio of HAZ to as-received parent material. Experimental and reference results verified the validity and applicability of this graph.
Acknowledgements The authors gratefully acknowledge the financial support under the Project 985-automotive engineering of Jilin University and Technology Development Program of Jilin Province (Nos. 20050509, 20080314).
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