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Experimental and simulative failure analysis of reformer furnace outlet manifolds serviced in aromatic plant
Engineering Failure Analysis 57 (2015) 350–362
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Engineering Failure Analysis journal homepage: www.elsevier.com/locate/efa
Experimental and simulative failure analysis of reformer furnace outlet manifolds serviced in aromatic plant Xiaowei Wang, Jianqun Tang, Jianming Gong ⁎, Luyang Geng, Yong Jiang, Hao Liu School of Mechanical and Power Engineering, Nanjing Tech University, Nanjing 211816, China
a r t i c l e
i n f o
Article history: Received 30 March 2015 Received in revised form 6 August 2015 Accepted 10 August 2015 Available online 14 August 2015 Keywords: High temperature Failure Outlet manifold Creep Finite element method
a b s t r a c t Reformer furnace hot outlet manifold is a critical component in aromatic plants. Unfortunately, a hot outlet manifold used in Nanjing YPC Refining & Chemical Co., Ltd. failed after it has been in service for only 2 years. In the present paper, a failure analysis of the failed hot outlet manifold was performed by careful visual observation, mechanical properties tests, chemical content examination, microstructural analysis and microchemical analysis of fine precipitates. Finite element method (FEM) considering creep effect was also employed to define the evolution of critical region. Experimental results showed that the failure was a combination of overheating and high stress concentration. Exposing it to excessive temperatures at high stress significantly accelerated the creep rate, thus, leading to the premature failure. FEM analysis showed that hoop stress generated by the work inner pressure was the main reason leading to the crack initiated and propagated along axial direction of the manifold. Moreover, sequence of the failure region achieved by the simulation correlated well with the experimental analysis. To prevent the failure, better control of operating temperature was strongly recommended. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction It is well known that reformer furnace hot outlet manifold system is the critical component in aromatic plants. To keep the plant online and running effectively, it needs to be operated reliably without unplanned shutdowns. Generally, high temperature stainless steels, such as 20Cr32NiNb, Alloy 800 and its derivatives (Alloy 800H, Alloy 800HT), are used extensively in hot manifold components due to their good weldability, high ductility and excellent resistance to creep damage. However, long serviced time and high operating temperatures (in the range of 750 °C to 950 °C) make material of manifold easy to suffer from degradations, even though sufficient considerations have been taken in material selection. Up to now, a variety of factors leading to the failure have been detected [1–3] such as overheating, metallurgical aging, creep, oxidation, thermal fatigue, carburization and so on. Dini et al. [4] studied a failure of continuous-annealing furnace radiant tubes and results showed that the failure was a combination of creep damage and oxidation attack. Jaske [5] has addressed the issues associated with failures of reformer outlet manifold systems and welded joints between various components were the major concerns. Homan and Magnan [6] conducted a failure analysis of 20Cr32NiNb alloy at a hydrogen reforming plant and determined that G-phase was responsible for the failure which caused the ductility dip and liquation cracking. Besides, the embrittlement of cast 20Cr32Ni1Nb used in steam reformer applications was also investigated by Knowles et al. [7]. Consequently, it is valuable to understand the failure mechanism of hot outlet manifold in engineering practice.
⁎ Corresponding author. E-mail address: [email protected] (J. Gong).
http://dx.doi.org/10.1016/j.engfailanal.2015.08.014 1350-6307/© 2015 Elsevier Ltd. All rights reserved.
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2. Background Reformer furnace outlet manifolds used in Nanjing YPC Refining & Chemical Co., Ltd. were built in 2012 and their design lifetime is about 15 years. However, unexpectedly, there was an outlet manifold suffering from serious premature damage after in service of only about 2 years, which has led to significant economic losses. The failed manifold was made of centrifugally cast 20Cr33NiNb pipe and designed to operate at high temperature of 940 °C with the inner pressure of 3.5 MPa. Its thickness and outer diameter are 31 mm and 212 mm, respectively. Detail geometry of the failure manifold is shown in Fig. 1. As can be seen, there are twenty pigtails connected to the manifold. The connection between manifold and pigtails were conducted via manual arc welding method. In order to solve the issues associated with the failure and prevent them from happening again, detailed analysis was required to identify the cause of failure. In the present paper, mechanical properties test, optical microscopy (OM) and scanning electron microscopy (SEM) observation of failed region samples and energy-dispersive spectroscopy (EDS) analysis of fine precipitates were undertaken to understand the cause of failure. Finite element method considering the effect of creep was also employed to define the critical region and compute the variations of Mises stress, hoop stress and creep strain with time in these critical regions.
3. Experimental details and results Two pieces of manifolds were used for experimental analysis, which were Manifold A and Manifold B, as shown in Fig. 2(a) and (e), respectively. Both Manifolds A and B were extracted from a manifold (Fig. 1), thus Manifolds A and B have been in service for the same time. However, Manifold A suffered from much more serious damage, which was the failed region causing the unplanned shutdown. As for Manifold B, no visible cracks and damage could be found, as shown in Fig. 2(e). Chemical compositions were examined by spectroscopy chemical analysis method. Tensile properties and Charpy impact properties at different temperatures were respectively analyzed by Instron 5689 vertical tensile machine and BG-300 high temperature impact machine. Macrohardness was measured using Vickers hardness tester (HXD-1000TM). Microstructure of the material was observed by using Zeiss Axio Imager A1m OM and JSM-6360 SEM. EDS analysis system in conjunction with the SEM was employed to carry out the microchemical characterization of fine precipitates. 3.1. Visual observations Fig. 2(a) shows the general aspect of the failed manifold, evident expansion and deformation can be found. In combination with the condition that the manifold serviced at very high temperature, and thus the fracture should be ductile. Fractured surface reveals that all the cracks along the axial direction of the manifold seem to originate from the middle pigtail (Fig. 2(a)). Technically, according to the crack orientation, high hoop stress must have played a critical role in the failure. Fig. 2(b) shows the location of primary failure where inner surfaces of pigtail and manifold cross. It is worth to see that the initiation sites also propagate along the axial direction and locate at the geometric mutation region where stress concentration is very likely to happen. Cracks near the weldment connecting pigtails and manifold are shown in Fig. 2(c). Obviously, cracks propagated along the heat affected zone (HAZ) can be observed. Site 1 as marked by yellow arrow seems like the crack initiation site on the outer surface. Fig. 2(d) reveals the local fracture surface in Fig. 2(a), apparent tearing characteristic can be found on the fracture surface, also many material fragments are nearly tear down. Meanwhile, Fig. 2(e) shows the un-failed Manifold B which is used for the comparison with failed Manifold A.
3.2. Material and mechanical properties 3.2.1. Material chemical composition verification Traditionally, hot manifold components are usually manufactured from wrought Alloy 800, 800H, 800HT or similar cast 20Cr32NiNb alloys. In this study, the outlet manifolds were made of centrifugally cast heat-resistant austenitic stainless steel 20Cr33NiNb alloy which is made in Shanghai Zhuoran Engineering Technology Co. Ltd. Chemical compositions of Manifolds A and B were both examined by SPECTROMAXx Direct Reading Spectrometer to make sure the compositions are in agreement with the requirements. The chemical compositions are listed in Table 1 and are compared with the nominal chemical compositions of ASTM A351 Grade CT15C (20Cr32NiNb) castings. Results show that the compositions are within the specification for 20Cr32NiNb steel.
Fig. 1. Geometry of the manifold.
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Fig. 2. (a) General aspect showing the whole fracture surface and cracking location; (b) Macroscopic crack initiation sites; (c) Crack morphology near weldment; (d) Local crack morphology; (e) Un-failed Manifold B.
3.2.2. Tensile properties The tensile test specimens were prepared in accordance with the ASTM-E8M specification. Table 2 lists the tensile properties of Manifolds A and B at 20 °C (room temperature) and at high temperature of 950 °C to simulate service condition. Manifold A tensile specimens were taken from the failed region near pigtail with no visible cracks could be found. While as for Manifold B, specimens were taken from the same region as Manifold A. Results show that all the tensile properties in Manifold A are obviously lower than the Manifold B, especially for the elongation. In comparison with room temperature, high temperature leads to the smaller magnitude of ultimate tensile strength (UTS) and larger magnitude of elongation.
3.2.3. Charpy impact properties Charpy impact properties of Manifolds A and B are shown in Fig. 3. At least three samples were measured for each test condition. It is obviously to see that impact properties of Manifold A are significantly lower than Manifold B. At different temperatures in the range of 20 °C to 950 °C, no marked difference of impact property related to test temperatures can be seen in Manifold A. However, for the Manifold B, significant difference can be observed. Evident increasing of Charpy impact magnitude is presented when temperature is no more than 600 °C, while as the temperature increases higher than 600 °C, the impact property shows a modest decrease with increasing temperature, and the toughness is low even at the temperature close to the service condition. Explanation for this phenomenon can be found in the microstructural analysis.
Table 1 Chemical compositions (wt.%) of 20Cr33NiNb and comparison with ASTM A351 alloy CT15C. Material
C
Si
Mn
P
S
Cr
Ni
Nb
Fe
20Cr33NiNb in Manifold A 20Cr33NiNb in Manifold B ASTM A351 alloy CT15C
0.080 0.074 0.05–0.15
0.79 0.77 0.5–1.5
0.72 0.72 0.15–1.5
0.023 0.019 0.03 max
0.0038 0.0031 0.03 max
19.91 19.91 19.0–21.0
32.04 31.91 31.0–41.0
1.00 0.94 0.5–1.5
Bal. Bal. Bal.
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Table 2 Tensile properties of 20Cr33NiNb alloy in Manifolds A and B. 20 °C (room temperature)
Manifold A Manifold B
950 °C
0.2% Proof (MPa)
UTS (MPa)
Elongation (%)
0.2% Proof (MPa)
UTS (MPa)
Elongation (%)
219.5 224.7
446 525
13 33
100 113
130 139
15 40
3.2.4. Hardness properties Vickers hardness values along the manifold thickness were measured at five different positions, as shown in Fig. 4. Measured positions were inner surface, 1/4 thickness from the inner surface, 1/2 thickness from the inner surface, 3/4 thickness from the inner surface and the outer surface as listed in Y-axis of Fig. 4. Vickers hardness values were measured with a 500 g load and four measurements for each position. Results show that hardness doesn't present obvious difference along thickness direction for each manifold, only slightly smaller values can be found in the inner surface. In addition, hardness values of the two manifolds along thickness present very similar distributions. Similar hardness values for both Manifolds A and B indicate that the two manifolds have experienced the same working conditions or same heat treatment. 3.3. Microstructural analysis Due to the difficulty in obtaining the as-received centrifugally cast 20Cr33NiNb alloy, microstructural analysis of the as-received material were compared with the study conducted by Knowles et al. [7], in which the microstructural photographs of the asreceived centrifugally cast 20Cr32NiNb alloy were presented, as shown in Fig. 5. As can be seen, typical microstructure of the alloy consists of complex interdendritic eutectic carbide networks in an austenitic matrix and the carbides present a lamellar or skeletal shape. In addition, some fine intra-dendritic carbides can also be seen in the austenitic matrix. It is these complex interdendritic eutectic carbides which are properly distributed throughout the austenite matrix that contribute to the excellent creep properties of this alloy. Whereas, many studies [8–13] have been extensively validated that long periods of aging of these similar alloys would led to the coarsening of M23C6 precipitate and the formation of many new phases, such as Si-rich G-phase, η′-phase and maybe Zphase. Among these precipitates, G-phase is the phase that always formed during long time aging at elevated temperatures and its molecular has been identified as Nb6(Ni,Fe)16(Si,Cr)7 [12]. Unfortunately, due to the low melting temperature of G-phase, its formation made the material encounter with severe in-service embrittlement and high susceptibility of the repair welding liquation cracking [11,14]. Technically, embrittlement of the service-exposed alloy caused by G-phase can be mitigated by a solution annealing heat treatment at temperatures in excess of 1100 °C [11]. Such a treatment will result in the transformation from G-phase to NbC particle and make the material recover from service-induced embrittlement [14]. Therefore, if the material has been serviced at elevated temperatures for a long time, but no G-phases could be found, we can thus infer that the material must have been conducted by solution annealing heat treatment or exposed to excessive temperature, at least higher than 1100 °C, which leads to the dissolution of new phases. Hence, evidence for overheating of the ex-service material can be provided with no G-phase detected. Fig. 6 shows the optical micrographs of Manifold A in different magnifications on the outer surface (Fig. 6(a), (c), (e)) and inner surface (Fig. 6(b), (d), (f)). Massive cracks exist at the manifold outer surface (Fig. 6(a)), and propagation of these cracks not only
Fig. 3. Impact properties at different temperatures.
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Fig. 4. Distribution of hardness along manifold thickness.
from the outer to inner surface but also from the inner to outer surface. However, for the manifold inner surface, only cracks propagating from the inner surface to outer surface can be observed in Fig. 6(b). Propagation direction of these cracks indicates that failure of the manifold first occurred on the inner surface, and outer surface following. Hence, cracks of the inner surface appear to cause the failure of outer surface. Fig. 6(c) and (d) respectively shows optical micrographs of Fig. 6(a) and (b) in higher magnification. According to the study performed by Dewar [15], centrifugal casting could produce a consistent columnar grain structure throughout. However, in the present study, columnar grains can only be seen near the outer surface, but change into equiaxed grains on the inner surface, which indicates that inner surface may suffer from higher temperatures. What's more, great number of creep voids can be observed along grain boundaries or near the triple grain boundaries. Creep voids on the outer surface present chain-like structure (Fig. 6(c)) and become larger or coalesce into micro intergranular cracks on the inner surface (Fig. 6(d)). In summary, all of these phenomena indicate that inner surface suffered from more damage. Hence, the final intergranular creep rupture occurs as the micro intergranular cracks link up. Unexpectedly, although the material has been experienced for a long time, microstructure of the alloy didn't change a lot, there were few eutectic carbides (lamellar or skeletal shape) still existing in the austenitic matrix and along the grain boundaries, as can be seen in Fig. 6(e) and (f). SEM images of Manifold A are presented in Fig. 7. Creep voids are obviously marked in Fig. 7(a), nearly all of the creep voids are originated from the triple grain boundaries. Thus it can be inferred that high stress concentration or dissolving of carbides in the triple grain boundaries during long term aging may facilitate the creep voids nucleation or growth. Fig. 7(b) shows the morphology of precipitates along the grain boundaries. It is interesting to find that the strings of interdendritic precipitates along gain boundaries tend to break up into discreet particles and the lamellar type structure seems to be destroyed. In addition, from the morphologies of interdendritic precipitates, less blocky morphology could be found. Thus, dissolution of the precipitates due to higher temperature can be deduced. To determine chemical compositions of these precipitates, EDS spectroscopy analysis method was used. It was of interest that only two phases could be detected within the precipitates. Fig. 7(c) and (d) respectively present the EDS analysis result of the two phases which were obviously marked by arrows with Point 1 and Point 2 remarks. Results show that size of precipitate Point 1 is much smaller than Point 2 size. What's more, Point 1 has much high concentration of Nb. Further investigation found
Fig. 5. Optical micrographs of as-received 20Cr32Ni1Nb alloy in reference [7].
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Fig. 6. Optical micrographs of Manifold A in different magnifications on the outer surface (a) ×25, (c) ×100, (e) ×500 and inner surface (b) ×25, (d) ×100, (f) ×500.
that the compositions of Point 1 were in agreement with the compositions of Nb(C, N) phase provided by Dewar [8] and Shi [11]. Thus, precipitate Point 1 should be Nb(C, N) phase. In addition, it is noteworthy that, for this material in as-received condition, large number density of fine Nb(C, N) can be produced from solution annealing [8]. While for Point 2 precipitate, a comparison between the analysis result with the result also given by Dewar [15] indicated that compositions of Point 2 precipitate were closest to ɤ-Fe phase. Unexpectedly, although detailed analysis of these precipitates has been performed, no G-phase could be detected. On the other hand, according to the discussion above, G-phase should have been formed during long period of aging. However, this is not the case. Hence, the conclusion that G-phase should be dissolved due to the excessive temperatures could be deduced. In order to find the microstructure difference in Manifolds A and B, Fig. 8 shows the optical micrographs of Manifold B. In comparison with the as-received microstructure (Fig. 5), the microstructure doesn't change in any major way. However, the density of the eutectic carbides seems to be smaller than Manifold A, especially for the region near the outer surface (Fig. 6(e)). Predictably, creep voids also can be observed in the microstructure and the inner surface also presents larger number and bigger size of creep voids. Similar to Manifold A, the inner surface also suffered from the more creep damage. To summarize, these observations indicate that Manifold B is in a relatively early stage of creep damage.
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Fig. 7. SEM images of Manifold A, (a) ×1000, (b) ×2000, (c) EDS spectra derived from point 1 grain boundary precipitates in (b), (d) EDS spectra derived from point 2 grain boundary precipitates in (b).
3.4. Fractography Fig. 9 illustrates the SEM fractographs of Manifold A near the outer surface. Great number of secondary cracks existing on the fracture surface and obvious characteristic of tearing can be seen as illustrated in Fig. 9(a). In order to determine the fracture mode, Fig. 9(b) shows the fractograph at higher magnification. Unfortunately, severe oxidation covers on the fracture surface, which makes the determination of fracture mode very difficult. However, although the serious oxidation of the fracture surface, many creep voids (as marked by arrows) can still be found in Fig. 9(b) and crack propagates along the creep voids can also be detected. Hence, creep cracking should be the main reason causing the failure.
Fig. 8. Optical micrographs of Manifold B (a) on the outer surface and (b) on the inner surface.
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Fig. 9. SEM fractographs of failed Manifold A on the outer surface.
As known to all, heat-resistant austenitic stainless steels are capable of developing a Cr2O3 protective oxide scale adherent to the material surface, thus, the steels should have good oxidation resistant properties. However, the material in the failure region suffered from serious oxidation as shown in Fig. 9(b). Thus, the protective oxide scale may have been broken during service exposure. Typically, excessive temperatures (at least larger than 1000 °C) [4] can break the protective oxide. Besides, the temperature is also consistent with the solution annealing temperatures. Therefore, the serious oxidation is another factor indicating that the failure was caused by excessive temperatures. As discussed in Section 3.2.3, Charpy impact properties of the Manifolds A and B have been investigated. It is interesting to find that impact energy for Manifold B tends to decreasing when test temperature excess 600 °C. To give an explanation of the phenomenon, Fig. 10 shows the SEM fractographs of Charpy impact specimen tested at temperature of 20 °C (Fig. 10(a)) and 900 °C (Fig. 10(c)). Obviously, ductile fracture mode can be seen for the two temperatures. Dimples in Fig. 10(c) seem to be larger than Fig. 10(a), also
Fig. 10. SEM fractographs of Charpy impact specimen tested at different temperatures, (a) 20 °C, (b) EDS spectra derived from precipitates at the bottom of dimple in (a), (c) 900 °C, (d) EDS spectra derived from precipitates at the bottom of dimple in (c).
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the dimple size is more homogeneous. More importantly, in comparison with Fig. 10(a) there are plenty of particles seem like carbides precipitates existing at the dimple's bottom, as shown in Fig. 10(b). As temperature increase higher than 600 °C, the increasing number of these particles may be the reason causing the impact energy decreasing. In order to determine these particles, EDS spectroscopy analysis method was also conducted, as shown in Fig. 10(b) and (d). According to the analysis result, the particle is Nb rich precipitate. Although compositions of the particles vary from one particle to another, the particle can still be recognized as Nb(C, N). Thus, the conclusion that higher temperature can favor the formation of Nb(C, N) precipitates can be drawn, and the result also correlates well with result from Shi and Lippold [11]. 4. Finite element analysis As described above, reformer furnace outlet manifolds serviced in aromatic plant are subjected to high inner pressure and high temperature. These extreme service conditions lead to the creep failure that is the main failure mode. In this section, finite element stress analysis was performed to help understand the cause of the outlet manifold cracking. The position of critical region and the evolution of stress at critical region were analyzed with the creep effect considered in the simulation. 4.1. Finite element model 3D finite element model of the outlet manifold was developed using ABAQUS finite element software. As experimental results reveled previously, cracking of the manifold first occurred at the place where pigtail crosses with manifold. Therefore, the finite element focused on the cross region and was simplified as shown in Fig. 11(a). Element type used in the stress analysis is C3D8R, 212,487 nodes and 187,365 elements were meshed and mesh size was fine enough to eliminate its influence. Mechanical boundary conditions were applied to reflect the actual restraint conditions and prevent rigid body motion, as shown in Fig. 11(b). On the front and back toroid of the manifold, symmetric boundary conditions in z-direction were applied and marked by BC-1 and BC-2, respectively. To prevent rigid body motion, encastre boundary conditions were applied at BC-3 and BC-4. Positions of BC-3 and BC-4 were located far away from the failure region to avoid the influence of these encastre boundary conditions on the failure region's calculation. In addition, a distributed work pressure (3.5 MPa) was imposed on the inner surface of manifold and pigtails to simulate actual working condition. From Fig. 2(c), it is obvious that the manifold and pigtails were assembled through welding. As the weld metal is very close to base metal, the base metal and weld metal were assumed to be same material. In the simulation, all the materials were assumed to be isotropic, linear elastic and viscoplastic. Values of Young's modulus, Poisson's ratio and yield strength at work temperature are listed in Table 3. The creep constitutive used in the simulation was assumed to follow Norton equation, which can be expressed as
ε c ¼ Bσ
n
ð1Þ
where ε c is the creep strain rate (h−1), σ is the stress (MPa), B and n are temperature dependent material constants, and their values can be seen in Table 3 [16]. In addition, during the creep calculation, the elastic strain was not included.
Fig. 11. Finite element model (a) element meshing; (b) boundary conditions.
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Table 3 Mechanical properties and creep parameters of 20Cr33NiNb at 940 °C. Young's modulus (GPa) 890
Poisson's ratio 0.3
Yield strength (MPa) 65
B 1.31 × 10
n −13
7.12
4.2. Results of modeling To check the material's strength and determine the critical region of the full model, Fig. 12 shows the simulation results of Mises stress and hoop stress at the initial time when the creep effect was not included. Fig. 12(a) shows that peak Mises stress is located at region A in axial direction of manifold, where inner surface of pigtail and manifold crossed. Value of the peak Mises stress is 28.8 MPa, obviously less than the material's yield stress, thus strength of the manifold meet the design criterion. As mentioned above, hoop stress is the main stress causing the cracking. Consequently, distribution of hoop stress is shown in Fig. 12(b). Peak hoop stress also locates at region A and its value is 27.0 MPa, similar to the peak Mises stress. In comparison with the result achieved by Knowles et al. [7], which recommended the pressure based hoop stress should typically below 10 MPa, the hoop stress is significantly higher. From the stress analysis results, it is not difficult to observe that region A is the stress concentration region and the most dangerous field of the full model. As is well known, creep always happens when working temperature larger than 0.5Tm (Tm is the molten temperature of material). In this analysis, high operating temperature of 940 °C makes the creep failure is the main factor that must not be neglected. Therefore, evolution of creep strain during service exposure should be paid special attention. Generally, for the high temperature furnace tube used in petrochemical plants, allowable creep strain is 5% [17]. That is to say, when the component creep strain reaches 5% or greater than the value, the component should be regarded as failure. Thus, Fig. 13 shows the stress distributions of Mises stress and hoop stress when creep strain of region A reaches 5%. Fig. 13(a) shows that in comparison with Fig. 12 location of the peak Mises stress doesn't change, but its value decrease from 28.8 MPa to 11.4 MPa, significant decreasing can be found due to the stress relaxation. Meanwhile, distribution of hoop stress is described in Fig. 13(b). Unexpectedly, location of the peak hoop stress changes from region A to region B, and its magnitude increase to 12.9 MPa. Obviously, significant creep strain is the reason leading to the redistribution of stresses. Therefore, according to the simulated results, region B should be the second dangerous field. In addition, as shown in Fig. 2(c), region B is HAZ of the weldment which joins the pigtails and the manifold together. Complex microstructure and welding residual stresses of the HAZ may be another reason causing the premature failure. From the FEM analysis result, failure region in the manifold has the following sequence: region A first and region B following. Failure of the region A appeared to cause the failure of region B, which correlates well with the experimental observation as discussed above.
Fig. 12. Results of simulation when creep strain is 0, (a) Mises stress; (b) Hoop stress.
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Fig. 13. Results of simulation when creep strain of location A reaches 5%, (a) Mises stress; (b) Hoop stress.
As regions A and B are the most dangerous place to crack, Fig. 14 gives the detailed description of Mises stress, hoop stress and creep strain variations with creep time in region A and region B. Fig. 14(a) shows that both Mises and hoop stress decrease drastically at the initial creep time and the time is about 111 h. As the creep time increasing, the two stresses quickly tend to steady until this region cracked. Linear relationship of creep strain variation with time can also be found, and this is reasonable when Norton creep constitutive model was adopted. Fig. 14(a) also shows the time of region A to reach 5% creep strain, the time is about 11,628 h which is the lifetime of region A. Fig. 14(b) shows the evolution of Mises stress, hoop stress and creep strain in region B. on the contrary, both Mises stress and hoop stress increase drastically at the initial creep time. As creep time increasing, Mises stress tends to steady quickly, while hoop stress presents slight increasing. Besides, hoop stress seems to be larger than Mises stress. Hence hoop stress should be the main failure contributor in region B. As can be seen in Fig. 14(b), when the creep time reaches 11,628 h, creep strain of region B is about 1%. Actually, the creep strain in the HAZ is underestimated, since the complex microstructure and residual stresses of HAZ weren't considered in the FEM model. 5. Discussion In the present paper, experimental and FEM analysis of the failed manifold was carried out in detail to achieve a possible cause of the failure. Mechanical properties examination and metallurgical analysis were both included in the experiments. Metallurgical
Fig. 14. Simulation results of hoop stress, Mises stress and creep strain variations with time at critical region of (a) region A and (b) region B.
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analysis illustrated that the failed manifold suffered from serious creep damage. Massive creep voids along grain boundaries can be found in the microstructural photographs as shown in Figs. 6(c), (d), (e), (f) and 7(a), (b). In addition, chain-like creep voids coalesced into intergranular creep cracks (Fig. 6(e), (f)) providing a good site for cracks originating and accelerating the fracture. Massive creep cracks in the manifold A significantly influence its mechanical properties. Thus can explain why the tensile elongation and impact energy in Manifold A are obviously lower than the Manifold B, as Table 2 and Fig. 3 show. It is known to all that achieving the sequence of failure is a very concerned factor for a failure analysis. Hence, combining of the results shown in Fig. 2(b) and (c) with the propagation direction of creep cracks on the manifold's inner surface (Fig. 6(a)) and outer surface (Fig. 6(b)) indicates that sequence of the failure is that crossed region of inner surface of pigtail and manifold first (Fig. 2(b)) and then HAZ weldment connecting pigtail and manifold follows (Fig. 2(c)). To give another explanation of this phenomenon FEM analysis was also employed. FEM results showed that the most critical place was region A (Figs. 12(a) and 13(a)), where the primary failure occurred. Simulated results showed that predicted location of crack initiation site coincided with the actually first failed region very well as illustrated in Fig. 2(b). As creep time increases, peak hoop stress of the manifold moves to region B (Fig. 13(b)) and thus region B should be the second failure place. Furthermore, region B is the HAZ (Fig. 2(c)), combination of the complex microstructure and residual stresses makes the region B more likely to fail. Therefore, FEM analysis considering creep effect predicted the practical failure sequence very well. To summarize, both metallurgical analysis and FEM simulation validated that creep damage is the main reason causing the premature failure. As it has been known, the failed manifold has been serviced for only 2 years, this is especially uncommon. Since the centrifugally cast 20Cr33NiNb alloy is a good heat-resistant austenitic stainless steel and selection of this material has been fully considered. In addition, the material has been examined. Thus serious creep damage must be attributed to excessive temperature or high stress concentration. It was well known [8–13] that long time aging of this heat-resistant austenitic stainless steel at elevated temperature would lead to the formation of G-phase. However, this was not the case in the study. No G-phase could be found in the failed manifold as revealed in Fig. 7. It has been well validated that G-phase can be dissolved due to high temperatures [11], thus overheating of the manifold should be the reason that no G-phase could be detected. In addition, dissolving of eutectic carbides (Fig. 7(b)) and sever oxidation of fracture surface (Fig. 9(b)) were another indicators validating the manifold has been operated in excessive temperatures. On the other hand, FEM simulated results showed that high stress concentration at high temperatures also can significantly accelerated the creep damage. Furthermore, obvious deformation (Fig. 2(a)) and evident tearing characteristic (Fig. 2(d)) of the failure manifold in macroscopic also revealed that the manifold has been subjected to high hoop stress too. 5.1. Cause of failure Most of the evidence revealed that the cause of failure was a combination of exposure to an excessive high temperature and high hoop stress at the critical regions. Exposure to an excessive high temperature at high stress level could have three detrimental effects. Firstly, the creep rate is accelerated leading to the material sustained larger viscoplastic strain during a shorter time. Besides, ductility of the material tends to decrease significantly at higher temperatures because of the appearance of Nb(C, N) precipitates, as explained in Figs. 3 and 10, thus a premature failure happens. Secondly, the formation and growth of grain boundary creep voids are accelerated. In addition to the high stress, creep crack can easily formed due to the coalescing of creep voids as observed in Fig. 6(c) and (d). Thirdly, overheating break the protective Cr2O3 oxidation film, thereby causing an accelerated oxidation attack. Cracking of the hard and brittle oxidation layer provide a good site for small crack initiation, thus causing an accelerated fracture. 5.2. Prevention approaches As obvious overheating of the failure manifold, thus it is strongly recommended to strictly control the operating temperature. It is necessary to check the manifold operating temperature on-line and ensure that the temperature is lower than the design temperature. What's more, decreasing the stress magnitude in stress concentration region is also useful. 6. Conclusions In this work, a failure analysis of the failed hot outlet manifolds made of 20Cr33NiNb used in reformer furnace was carried out based on experimental method and FEM analysis. According to the results obtained above, it could be concluded that the failure can be attributed to the combination of overheating and high hoop stress at critical region. Service overheating can significantly decrease the material ductility and its yield stress. Besides, in combination with high stress concentration, creep rates of the manifold can thus be drastically promoted, and then creep failure happens. Consequently, better control of operating temperature is recommended to prevent the failure, and decreasing the stress concentration is also useful. Acknowledgment The authors acknowledge the support of China Nanjing YPC Refining & Chemical Co., Ltd. for providing samples. References [1] J.J. Hoffman, High temperature aging characteristics of 20Cr32Ni1Nb castings, CORROSION 2000, NACE International, 2000.
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[2] J. Ahmad, J. Purbolaksono, L.C. Beng, Failure analysis on high temperature superheater Inconel® 800 tube, Eng. Fail. Anal. 17 (1) (2010) 328–333. [3] T.L. da Silveira, I. Le May, Reformer furnaces: materials, damage mechanisms, and assessment, Arab. J. Sci. Eng. 31 (2) (2006) 99–119. [4] G. Dini, S.M.M. Vaghefi, M. Lotfiani, M. Jafari, M. Safaei-Rad, M. Navabi, S. Abbasi, Computational and experimental failure analysis of continuous-annealing furnace radiant tubes exposed to excessive temperature, Eng. Fail. Anal. 15 (5) (2008) 445–457. [5] C.E. Jaske, Creep strength of welds in hydrogen reformer outlet manifolds, ASME 2009 Pressure Vessels and Piping Conference, American Society of Mechanical Engineers 2009, pp. 1385–1390. [6] J.J. Hoffman, J. Magnan, Cast 20Cr32Ni1Nb alloy aged mechanical property improvements via chemistry modifications, CORROSION 2003, NACE International, 2003. [7] D.M. Knowles, C.W. Thomas, D.J. Keen, Q.Z. Chen, In service embrittlement of cast 20Cr32Ni1Nb components used in steam reformer applications, Int. J. Press. Vessel. Pip. 81 (6) (2004) 499–506. [8] M.P. Dewar, A.P. Gerlich, Correlation between experimental and calculated phase fractions in aged 20Cr32Ni1Nb austenitic stainless steels containing nitrogen, Metall. Mater. Trans. A 44 (2) (2013) 627–639. [9] D.J. Powell, R. Pilkington, D.A. Miller, The precipitation characteristics of 20% Cr/25% Ni–Nb stabilised stainless steel, Acta Metall. 36 (3) (1988) 713–724. [10] B. Piekarski, Effect of Nb and Ti additions on microstructure, and identification of precipitates in stabilized Ni–Cr cast austenitic steels, Mater. Charact. 47 (3) (2001) 181–186. [11] S. Shi, J.C. Lippold, Microstructure evolution during service exposure of two cast, heat-resisting stainless steels—HP–Nb modified and 20–32Nb, Mater. Charact. 59 (8) (2008) 1029–1040. [12] Q.Z. Chen, C.W. Thomas, D.M. Knowles, Characterisation of 20Cr32Ni1Nb alloys in as-cast and Ex-Service conditions by SEM, TEM and EDX, Mater. Sci. Eng. A 374 (1) (2004) 398–408. [13] K. Shinozaki, H. Kuroki, Y. Nakao, K. Nishimoto, M. Inui, M. Takahashi, Deterioration of weldability of long-term aged HP heat-resistant cast steel containing Nb, Mo, and W, Weld. Int. 13 (1) (1999) 39–48. [14] S. Shi, J.C. Lippold, J. Ramirez, Hot ductility behavior and repair weldability of service-aged, heat-resistant stainless steel castings, Weld. J. 89 (10) (2010) 210–217. [15] M. Dewar, Characterization and Evaluation of Aged 20Cr32Ni1Nb Stainless Steels(Doctoral thesis) University of Alberta, 2012. [16] L. Shen, Damage Analysis and Life Prediction of Ethylene Cracking Furnace Tube Based on the Coupled Multi-factor Effects(Doctoral thesis) Nanjing Tech University, 2012. [17] T. Shan, High Temperature Structural Integrity, 2003 (Beijing).
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Engineering Failure Analysis journal homepage: www.elsevier.com/locate/efa
Experimental and simulative failure analysis of reformer furnace outlet manifolds serviced in aromatic plant Xiaowei Wang, Jianqun Tang, Jianming Gong ⁎, Luyang Geng, Yong Jiang, Hao Liu School of Mechanical and Power Engineering, Nanjing Tech University, Nanjing 211816, China
a r t i c l e
i n f o
Article history: Received 30 March 2015 Received in revised form 6 August 2015 Accepted 10 August 2015 Available online 14 August 2015 Keywords: High temperature Failure Outlet manifold Creep Finite element method
a b s t r a c t Reformer furnace hot outlet manifold is a critical component in aromatic plants. Unfortunately, a hot outlet manifold used in Nanjing YPC Refining & Chemical Co., Ltd. failed after it has been in service for only 2 years. In the present paper, a failure analysis of the failed hot outlet manifold was performed by careful visual observation, mechanical properties tests, chemical content examination, microstructural analysis and microchemical analysis of fine precipitates. Finite element method (FEM) considering creep effect was also employed to define the evolution of critical region. Experimental results showed that the failure was a combination of overheating and high stress concentration. Exposing it to excessive temperatures at high stress significantly accelerated the creep rate, thus, leading to the premature failure. FEM analysis showed that hoop stress generated by the work inner pressure was the main reason leading to the crack initiated and propagated along axial direction of the manifold. Moreover, sequence of the failure region achieved by the simulation correlated well with the experimental analysis. To prevent the failure, better control of operating temperature was strongly recommended. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction It is well known that reformer furnace hot outlet manifold system is the critical component in aromatic plants. To keep the plant online and running effectively, it needs to be operated reliably without unplanned shutdowns. Generally, high temperature stainless steels, such as 20Cr32NiNb, Alloy 800 and its derivatives (Alloy 800H, Alloy 800HT), are used extensively in hot manifold components due to their good weldability, high ductility and excellent resistance to creep damage. However, long serviced time and high operating temperatures (in the range of 750 °C to 950 °C) make material of manifold easy to suffer from degradations, even though sufficient considerations have been taken in material selection. Up to now, a variety of factors leading to the failure have been detected [1–3] such as overheating, metallurgical aging, creep, oxidation, thermal fatigue, carburization and so on. Dini et al. [4] studied a failure of continuous-annealing furnace radiant tubes and results showed that the failure was a combination of creep damage and oxidation attack. Jaske [5] has addressed the issues associated with failures of reformer outlet manifold systems and welded joints between various components were the major concerns. Homan and Magnan [6] conducted a failure analysis of 20Cr32NiNb alloy at a hydrogen reforming plant and determined that G-phase was responsible for the failure which caused the ductility dip and liquation cracking. Besides, the embrittlement of cast 20Cr32Ni1Nb used in steam reformer applications was also investigated by Knowles et al. [7]. Consequently, it is valuable to understand the failure mechanism of hot outlet manifold in engineering practice.
⁎ Corresponding author. E-mail address: [email protected] (J. Gong).
http://dx.doi.org/10.1016/j.engfailanal.2015.08.014 1350-6307/© 2015 Elsevier Ltd. All rights reserved.
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2. Background Reformer furnace outlet manifolds used in Nanjing YPC Refining & Chemical Co., Ltd. were built in 2012 and their design lifetime is about 15 years. However, unexpectedly, there was an outlet manifold suffering from serious premature damage after in service of only about 2 years, which has led to significant economic losses. The failed manifold was made of centrifugally cast 20Cr33NiNb pipe and designed to operate at high temperature of 940 °C with the inner pressure of 3.5 MPa. Its thickness and outer diameter are 31 mm and 212 mm, respectively. Detail geometry of the failure manifold is shown in Fig. 1. As can be seen, there are twenty pigtails connected to the manifold. The connection between manifold and pigtails were conducted via manual arc welding method. In order to solve the issues associated with the failure and prevent them from happening again, detailed analysis was required to identify the cause of failure. In the present paper, mechanical properties test, optical microscopy (OM) and scanning electron microscopy (SEM) observation of failed region samples and energy-dispersive spectroscopy (EDS) analysis of fine precipitates were undertaken to understand the cause of failure. Finite element method considering the effect of creep was also employed to define the critical region and compute the variations of Mises stress, hoop stress and creep strain with time in these critical regions.
3. Experimental details and results Two pieces of manifolds were used for experimental analysis, which were Manifold A and Manifold B, as shown in Fig. 2(a) and (e), respectively. Both Manifolds A and B were extracted from a manifold (Fig. 1), thus Manifolds A and B have been in service for the same time. However, Manifold A suffered from much more serious damage, which was the failed region causing the unplanned shutdown. As for Manifold B, no visible cracks and damage could be found, as shown in Fig. 2(e). Chemical compositions were examined by spectroscopy chemical analysis method. Tensile properties and Charpy impact properties at different temperatures were respectively analyzed by Instron 5689 vertical tensile machine and BG-300 high temperature impact machine. Macrohardness was measured using Vickers hardness tester (HXD-1000TM). Microstructure of the material was observed by using Zeiss Axio Imager A1m OM and JSM-6360 SEM. EDS analysis system in conjunction with the SEM was employed to carry out the microchemical characterization of fine precipitates. 3.1. Visual observations Fig. 2(a) shows the general aspect of the failed manifold, evident expansion and deformation can be found. In combination with the condition that the manifold serviced at very high temperature, and thus the fracture should be ductile. Fractured surface reveals that all the cracks along the axial direction of the manifold seem to originate from the middle pigtail (Fig. 2(a)). Technically, according to the crack orientation, high hoop stress must have played a critical role in the failure. Fig. 2(b) shows the location of primary failure where inner surfaces of pigtail and manifold cross. It is worth to see that the initiation sites also propagate along the axial direction and locate at the geometric mutation region where stress concentration is very likely to happen. Cracks near the weldment connecting pigtails and manifold are shown in Fig. 2(c). Obviously, cracks propagated along the heat affected zone (HAZ) can be observed. Site 1 as marked by yellow arrow seems like the crack initiation site on the outer surface. Fig. 2(d) reveals the local fracture surface in Fig. 2(a), apparent tearing characteristic can be found on the fracture surface, also many material fragments are nearly tear down. Meanwhile, Fig. 2(e) shows the un-failed Manifold B which is used for the comparison with failed Manifold A.
3.2. Material and mechanical properties 3.2.1. Material chemical composition verification Traditionally, hot manifold components are usually manufactured from wrought Alloy 800, 800H, 800HT or similar cast 20Cr32NiNb alloys. In this study, the outlet manifolds were made of centrifugally cast heat-resistant austenitic stainless steel 20Cr33NiNb alloy which is made in Shanghai Zhuoran Engineering Technology Co. Ltd. Chemical compositions of Manifolds A and B were both examined by SPECTROMAXx Direct Reading Spectrometer to make sure the compositions are in agreement with the requirements. The chemical compositions are listed in Table 1 and are compared with the nominal chemical compositions of ASTM A351 Grade CT15C (20Cr32NiNb) castings. Results show that the compositions are within the specification for 20Cr32NiNb steel.
Fig. 1. Geometry of the manifold.
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Fig. 2. (a) General aspect showing the whole fracture surface and cracking location; (b) Macroscopic crack initiation sites; (c) Crack morphology near weldment; (d) Local crack morphology; (e) Un-failed Manifold B.
3.2.2. Tensile properties The tensile test specimens were prepared in accordance with the ASTM-E8M specification. Table 2 lists the tensile properties of Manifolds A and B at 20 °C (room temperature) and at high temperature of 950 °C to simulate service condition. Manifold A tensile specimens were taken from the failed region near pigtail with no visible cracks could be found. While as for Manifold B, specimens were taken from the same region as Manifold A. Results show that all the tensile properties in Manifold A are obviously lower than the Manifold B, especially for the elongation. In comparison with room temperature, high temperature leads to the smaller magnitude of ultimate tensile strength (UTS) and larger magnitude of elongation.
3.2.3. Charpy impact properties Charpy impact properties of Manifolds A and B are shown in Fig. 3. At least three samples were measured for each test condition. It is obviously to see that impact properties of Manifold A are significantly lower than Manifold B. At different temperatures in the range of 20 °C to 950 °C, no marked difference of impact property related to test temperatures can be seen in Manifold A. However, for the Manifold B, significant difference can be observed. Evident increasing of Charpy impact magnitude is presented when temperature is no more than 600 °C, while as the temperature increases higher than 600 °C, the impact property shows a modest decrease with increasing temperature, and the toughness is low even at the temperature close to the service condition. Explanation for this phenomenon can be found in the microstructural analysis.
Table 1 Chemical compositions (wt.%) of 20Cr33NiNb and comparison with ASTM A351 alloy CT15C. Material
C
Si
Mn
P
S
Cr
Ni
Nb
Fe
20Cr33NiNb in Manifold A 20Cr33NiNb in Manifold B ASTM A351 alloy CT15C
0.080 0.074 0.05–0.15
0.79 0.77 0.5–1.5
0.72 0.72 0.15–1.5
0.023 0.019 0.03 max
0.0038 0.0031 0.03 max
19.91 19.91 19.0–21.0
32.04 31.91 31.0–41.0
1.00 0.94 0.5–1.5
Bal. Bal. Bal.
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Table 2 Tensile properties of 20Cr33NiNb alloy in Manifolds A and B. 20 °C (room temperature)
Manifold A Manifold B
950 °C
0.2% Proof (MPa)
UTS (MPa)
Elongation (%)
0.2% Proof (MPa)
UTS (MPa)
Elongation (%)
219.5 224.7
446 525
13 33
100 113
130 139
15 40
3.2.4. Hardness properties Vickers hardness values along the manifold thickness were measured at five different positions, as shown in Fig. 4. Measured positions were inner surface, 1/4 thickness from the inner surface, 1/2 thickness from the inner surface, 3/4 thickness from the inner surface and the outer surface as listed in Y-axis of Fig. 4. Vickers hardness values were measured with a 500 g load and four measurements for each position. Results show that hardness doesn't present obvious difference along thickness direction for each manifold, only slightly smaller values can be found in the inner surface. In addition, hardness values of the two manifolds along thickness present very similar distributions. Similar hardness values for both Manifolds A and B indicate that the two manifolds have experienced the same working conditions or same heat treatment. 3.3. Microstructural analysis Due to the difficulty in obtaining the as-received centrifugally cast 20Cr33NiNb alloy, microstructural analysis of the as-received material were compared with the study conducted by Knowles et al. [7], in which the microstructural photographs of the asreceived centrifugally cast 20Cr32NiNb alloy were presented, as shown in Fig. 5. As can be seen, typical microstructure of the alloy consists of complex interdendritic eutectic carbide networks in an austenitic matrix and the carbides present a lamellar or skeletal shape. In addition, some fine intra-dendritic carbides can also be seen in the austenitic matrix. It is these complex interdendritic eutectic carbides which are properly distributed throughout the austenite matrix that contribute to the excellent creep properties of this alloy. Whereas, many studies [8–13] have been extensively validated that long periods of aging of these similar alloys would led to the coarsening of M23C6 precipitate and the formation of many new phases, such as Si-rich G-phase, η′-phase and maybe Zphase. Among these precipitates, G-phase is the phase that always formed during long time aging at elevated temperatures and its molecular has been identified as Nb6(Ni,Fe)16(Si,Cr)7 [12]. Unfortunately, due to the low melting temperature of G-phase, its formation made the material encounter with severe in-service embrittlement and high susceptibility of the repair welding liquation cracking [11,14]. Technically, embrittlement of the service-exposed alloy caused by G-phase can be mitigated by a solution annealing heat treatment at temperatures in excess of 1100 °C [11]. Such a treatment will result in the transformation from G-phase to NbC particle and make the material recover from service-induced embrittlement [14]. Therefore, if the material has been serviced at elevated temperatures for a long time, but no G-phases could be found, we can thus infer that the material must have been conducted by solution annealing heat treatment or exposed to excessive temperature, at least higher than 1100 °C, which leads to the dissolution of new phases. Hence, evidence for overheating of the ex-service material can be provided with no G-phase detected. Fig. 6 shows the optical micrographs of Manifold A in different magnifications on the outer surface (Fig. 6(a), (c), (e)) and inner surface (Fig. 6(b), (d), (f)). Massive cracks exist at the manifold outer surface (Fig. 6(a)), and propagation of these cracks not only
Fig. 3. Impact properties at different temperatures.
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Fig. 4. Distribution of hardness along manifold thickness.
from the outer to inner surface but also from the inner to outer surface. However, for the manifold inner surface, only cracks propagating from the inner surface to outer surface can be observed in Fig. 6(b). Propagation direction of these cracks indicates that failure of the manifold first occurred on the inner surface, and outer surface following. Hence, cracks of the inner surface appear to cause the failure of outer surface. Fig. 6(c) and (d) respectively shows optical micrographs of Fig. 6(a) and (b) in higher magnification. According to the study performed by Dewar [15], centrifugal casting could produce a consistent columnar grain structure throughout. However, in the present study, columnar grains can only be seen near the outer surface, but change into equiaxed grains on the inner surface, which indicates that inner surface may suffer from higher temperatures. What's more, great number of creep voids can be observed along grain boundaries or near the triple grain boundaries. Creep voids on the outer surface present chain-like structure (Fig. 6(c)) and become larger or coalesce into micro intergranular cracks on the inner surface (Fig. 6(d)). In summary, all of these phenomena indicate that inner surface suffered from more damage. Hence, the final intergranular creep rupture occurs as the micro intergranular cracks link up. Unexpectedly, although the material has been experienced for a long time, microstructure of the alloy didn't change a lot, there were few eutectic carbides (lamellar or skeletal shape) still existing in the austenitic matrix and along the grain boundaries, as can be seen in Fig. 6(e) and (f). SEM images of Manifold A are presented in Fig. 7. Creep voids are obviously marked in Fig. 7(a), nearly all of the creep voids are originated from the triple grain boundaries. Thus it can be inferred that high stress concentration or dissolving of carbides in the triple grain boundaries during long term aging may facilitate the creep voids nucleation or growth. Fig. 7(b) shows the morphology of precipitates along the grain boundaries. It is interesting to find that the strings of interdendritic precipitates along gain boundaries tend to break up into discreet particles and the lamellar type structure seems to be destroyed. In addition, from the morphologies of interdendritic precipitates, less blocky morphology could be found. Thus, dissolution of the precipitates due to higher temperature can be deduced. To determine chemical compositions of these precipitates, EDS spectroscopy analysis method was used. It was of interest that only two phases could be detected within the precipitates. Fig. 7(c) and (d) respectively present the EDS analysis result of the two phases which were obviously marked by arrows with Point 1 and Point 2 remarks. Results show that size of precipitate Point 1 is much smaller than Point 2 size. What's more, Point 1 has much high concentration of Nb. Further investigation found
Fig. 5. Optical micrographs of as-received 20Cr32Ni1Nb alloy in reference [7].
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Fig. 6. Optical micrographs of Manifold A in different magnifications on the outer surface (a) ×25, (c) ×100, (e) ×500 and inner surface (b) ×25, (d) ×100, (f) ×500.
that the compositions of Point 1 were in agreement with the compositions of Nb(C, N) phase provided by Dewar [8] and Shi [11]. Thus, precipitate Point 1 should be Nb(C, N) phase. In addition, it is noteworthy that, for this material in as-received condition, large number density of fine Nb(C, N) can be produced from solution annealing [8]. While for Point 2 precipitate, a comparison between the analysis result with the result also given by Dewar [15] indicated that compositions of Point 2 precipitate were closest to ɤ-Fe phase. Unexpectedly, although detailed analysis of these precipitates has been performed, no G-phase could be detected. On the other hand, according to the discussion above, G-phase should have been formed during long period of aging. However, this is not the case. Hence, the conclusion that G-phase should be dissolved due to the excessive temperatures could be deduced. In order to find the microstructure difference in Manifolds A and B, Fig. 8 shows the optical micrographs of Manifold B. In comparison with the as-received microstructure (Fig. 5), the microstructure doesn't change in any major way. However, the density of the eutectic carbides seems to be smaller than Manifold A, especially for the region near the outer surface (Fig. 6(e)). Predictably, creep voids also can be observed in the microstructure and the inner surface also presents larger number and bigger size of creep voids. Similar to Manifold A, the inner surface also suffered from the more creep damage. To summarize, these observations indicate that Manifold B is in a relatively early stage of creep damage.
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Fig. 7. SEM images of Manifold A, (a) ×1000, (b) ×2000, (c) EDS spectra derived from point 1 grain boundary precipitates in (b), (d) EDS spectra derived from point 2 grain boundary precipitates in (b).
3.4. Fractography Fig. 9 illustrates the SEM fractographs of Manifold A near the outer surface. Great number of secondary cracks existing on the fracture surface and obvious characteristic of tearing can be seen as illustrated in Fig. 9(a). In order to determine the fracture mode, Fig. 9(b) shows the fractograph at higher magnification. Unfortunately, severe oxidation covers on the fracture surface, which makes the determination of fracture mode very difficult. However, although the serious oxidation of the fracture surface, many creep voids (as marked by arrows) can still be found in Fig. 9(b) and crack propagates along the creep voids can also be detected. Hence, creep cracking should be the main reason causing the failure.
Fig. 8. Optical micrographs of Manifold B (a) on the outer surface and (b) on the inner surface.
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Fig. 9. SEM fractographs of failed Manifold A on the outer surface.
As known to all, heat-resistant austenitic stainless steels are capable of developing a Cr2O3 protective oxide scale adherent to the material surface, thus, the steels should have good oxidation resistant properties. However, the material in the failure region suffered from serious oxidation as shown in Fig. 9(b). Thus, the protective oxide scale may have been broken during service exposure. Typically, excessive temperatures (at least larger than 1000 °C) [4] can break the protective oxide. Besides, the temperature is also consistent with the solution annealing temperatures. Therefore, the serious oxidation is another factor indicating that the failure was caused by excessive temperatures. As discussed in Section 3.2.3, Charpy impact properties of the Manifolds A and B have been investigated. It is interesting to find that impact energy for Manifold B tends to decreasing when test temperature excess 600 °C. To give an explanation of the phenomenon, Fig. 10 shows the SEM fractographs of Charpy impact specimen tested at temperature of 20 °C (Fig. 10(a)) and 900 °C (Fig. 10(c)). Obviously, ductile fracture mode can be seen for the two temperatures. Dimples in Fig. 10(c) seem to be larger than Fig. 10(a), also
Fig. 10. SEM fractographs of Charpy impact specimen tested at different temperatures, (a) 20 °C, (b) EDS spectra derived from precipitates at the bottom of dimple in (a), (c) 900 °C, (d) EDS spectra derived from precipitates at the bottom of dimple in (c).
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the dimple size is more homogeneous. More importantly, in comparison with Fig. 10(a) there are plenty of particles seem like carbides precipitates existing at the dimple's bottom, as shown in Fig. 10(b). As temperature increase higher than 600 °C, the increasing number of these particles may be the reason causing the impact energy decreasing. In order to determine these particles, EDS spectroscopy analysis method was also conducted, as shown in Fig. 10(b) and (d). According to the analysis result, the particle is Nb rich precipitate. Although compositions of the particles vary from one particle to another, the particle can still be recognized as Nb(C, N). Thus, the conclusion that higher temperature can favor the formation of Nb(C, N) precipitates can be drawn, and the result also correlates well with result from Shi and Lippold [11]. 4. Finite element analysis As described above, reformer furnace outlet manifolds serviced in aromatic plant are subjected to high inner pressure and high temperature. These extreme service conditions lead to the creep failure that is the main failure mode. In this section, finite element stress analysis was performed to help understand the cause of the outlet manifold cracking. The position of critical region and the evolution of stress at critical region were analyzed with the creep effect considered in the simulation. 4.1. Finite element model 3D finite element model of the outlet manifold was developed using ABAQUS finite element software. As experimental results reveled previously, cracking of the manifold first occurred at the place where pigtail crosses with manifold. Therefore, the finite element focused on the cross region and was simplified as shown in Fig. 11(a). Element type used in the stress analysis is C3D8R, 212,487 nodes and 187,365 elements were meshed and mesh size was fine enough to eliminate its influence. Mechanical boundary conditions were applied to reflect the actual restraint conditions and prevent rigid body motion, as shown in Fig. 11(b). On the front and back toroid of the manifold, symmetric boundary conditions in z-direction were applied and marked by BC-1 and BC-2, respectively. To prevent rigid body motion, encastre boundary conditions were applied at BC-3 and BC-4. Positions of BC-3 and BC-4 were located far away from the failure region to avoid the influence of these encastre boundary conditions on the failure region's calculation. In addition, a distributed work pressure (3.5 MPa) was imposed on the inner surface of manifold and pigtails to simulate actual working condition. From Fig. 2(c), it is obvious that the manifold and pigtails were assembled through welding. As the weld metal is very close to base metal, the base metal and weld metal were assumed to be same material. In the simulation, all the materials were assumed to be isotropic, linear elastic and viscoplastic. Values of Young's modulus, Poisson's ratio and yield strength at work temperature are listed in Table 3. The creep constitutive used in the simulation was assumed to follow Norton equation, which can be expressed as
ε c ¼ Bσ
n
ð1Þ
where ε c is the creep strain rate (h−1), σ is the stress (MPa), B and n are temperature dependent material constants, and their values can be seen in Table 3 [16]. In addition, during the creep calculation, the elastic strain was not included.
Fig. 11. Finite element model (a) element meshing; (b) boundary conditions.
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Table 3 Mechanical properties and creep parameters of 20Cr33NiNb at 940 °C. Young's modulus (GPa) 890
Poisson's ratio 0.3
Yield strength (MPa) 65
B 1.31 × 10
n −13
7.12
4.2. Results of modeling To check the material's strength and determine the critical region of the full model, Fig. 12 shows the simulation results of Mises stress and hoop stress at the initial time when the creep effect was not included. Fig. 12(a) shows that peak Mises stress is located at region A in axial direction of manifold, where inner surface of pigtail and manifold crossed. Value of the peak Mises stress is 28.8 MPa, obviously less than the material's yield stress, thus strength of the manifold meet the design criterion. As mentioned above, hoop stress is the main stress causing the cracking. Consequently, distribution of hoop stress is shown in Fig. 12(b). Peak hoop stress also locates at region A and its value is 27.0 MPa, similar to the peak Mises stress. In comparison with the result achieved by Knowles et al. [7], which recommended the pressure based hoop stress should typically below 10 MPa, the hoop stress is significantly higher. From the stress analysis results, it is not difficult to observe that region A is the stress concentration region and the most dangerous field of the full model. As is well known, creep always happens when working temperature larger than 0.5Tm (Tm is the molten temperature of material). In this analysis, high operating temperature of 940 °C makes the creep failure is the main factor that must not be neglected. Therefore, evolution of creep strain during service exposure should be paid special attention. Generally, for the high temperature furnace tube used in petrochemical plants, allowable creep strain is 5% [17]. That is to say, when the component creep strain reaches 5% or greater than the value, the component should be regarded as failure. Thus, Fig. 13 shows the stress distributions of Mises stress and hoop stress when creep strain of region A reaches 5%. Fig. 13(a) shows that in comparison with Fig. 12 location of the peak Mises stress doesn't change, but its value decrease from 28.8 MPa to 11.4 MPa, significant decreasing can be found due to the stress relaxation. Meanwhile, distribution of hoop stress is described in Fig. 13(b). Unexpectedly, location of the peak hoop stress changes from region A to region B, and its magnitude increase to 12.9 MPa. Obviously, significant creep strain is the reason leading to the redistribution of stresses. Therefore, according to the simulated results, region B should be the second dangerous field. In addition, as shown in Fig. 2(c), region B is HAZ of the weldment which joins the pigtails and the manifold together. Complex microstructure and welding residual stresses of the HAZ may be another reason causing the premature failure. From the FEM analysis result, failure region in the manifold has the following sequence: region A first and region B following. Failure of the region A appeared to cause the failure of region B, which correlates well with the experimental observation as discussed above.
Fig. 12. Results of simulation when creep strain is 0, (a) Mises stress; (b) Hoop stress.
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Fig. 13. Results of simulation when creep strain of location A reaches 5%, (a) Mises stress; (b) Hoop stress.
As regions A and B are the most dangerous place to crack, Fig. 14 gives the detailed description of Mises stress, hoop stress and creep strain variations with creep time in region A and region B. Fig. 14(a) shows that both Mises and hoop stress decrease drastically at the initial creep time and the time is about 111 h. As the creep time increasing, the two stresses quickly tend to steady until this region cracked. Linear relationship of creep strain variation with time can also be found, and this is reasonable when Norton creep constitutive model was adopted. Fig. 14(a) also shows the time of region A to reach 5% creep strain, the time is about 11,628 h which is the lifetime of region A. Fig. 14(b) shows the evolution of Mises stress, hoop stress and creep strain in region B. on the contrary, both Mises stress and hoop stress increase drastically at the initial creep time. As creep time increasing, Mises stress tends to steady quickly, while hoop stress presents slight increasing. Besides, hoop stress seems to be larger than Mises stress. Hence hoop stress should be the main failure contributor in region B. As can be seen in Fig. 14(b), when the creep time reaches 11,628 h, creep strain of region B is about 1%. Actually, the creep strain in the HAZ is underestimated, since the complex microstructure and residual stresses of HAZ weren't considered in the FEM model. 5. Discussion In the present paper, experimental and FEM analysis of the failed manifold was carried out in detail to achieve a possible cause of the failure. Mechanical properties examination and metallurgical analysis were both included in the experiments. Metallurgical
Fig. 14. Simulation results of hoop stress, Mises stress and creep strain variations with time at critical region of (a) region A and (b) region B.
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analysis illustrated that the failed manifold suffered from serious creep damage. Massive creep voids along grain boundaries can be found in the microstructural photographs as shown in Figs. 6(c), (d), (e), (f) and 7(a), (b). In addition, chain-like creep voids coalesced into intergranular creep cracks (Fig. 6(e), (f)) providing a good site for cracks originating and accelerating the fracture. Massive creep cracks in the manifold A significantly influence its mechanical properties. Thus can explain why the tensile elongation and impact energy in Manifold A are obviously lower than the Manifold B, as Table 2 and Fig. 3 show. It is known to all that achieving the sequence of failure is a very concerned factor for a failure analysis. Hence, combining of the results shown in Fig. 2(b) and (c) with the propagation direction of creep cracks on the manifold's inner surface (Fig. 6(a)) and outer surface (Fig. 6(b)) indicates that sequence of the failure is that crossed region of inner surface of pigtail and manifold first (Fig. 2(b)) and then HAZ weldment connecting pigtail and manifold follows (Fig. 2(c)). To give another explanation of this phenomenon FEM analysis was also employed. FEM results showed that the most critical place was region A (Figs. 12(a) and 13(a)), where the primary failure occurred. Simulated results showed that predicted location of crack initiation site coincided with the actually first failed region very well as illustrated in Fig. 2(b). As creep time increases, peak hoop stress of the manifold moves to region B (Fig. 13(b)) and thus region B should be the second failure place. Furthermore, region B is the HAZ (Fig. 2(c)), combination of the complex microstructure and residual stresses makes the region B more likely to fail. Therefore, FEM analysis considering creep effect predicted the practical failure sequence very well. To summarize, both metallurgical analysis and FEM simulation validated that creep damage is the main reason causing the premature failure. As it has been known, the failed manifold has been serviced for only 2 years, this is especially uncommon. Since the centrifugally cast 20Cr33NiNb alloy is a good heat-resistant austenitic stainless steel and selection of this material has been fully considered. In addition, the material has been examined. Thus serious creep damage must be attributed to excessive temperature or high stress concentration. It was well known [8–13] that long time aging of this heat-resistant austenitic stainless steel at elevated temperature would lead to the formation of G-phase. However, this was not the case in the study. No G-phase could be found in the failed manifold as revealed in Fig. 7. It has been well validated that G-phase can be dissolved due to high temperatures [11], thus overheating of the manifold should be the reason that no G-phase could be detected. In addition, dissolving of eutectic carbides (Fig. 7(b)) and sever oxidation of fracture surface (Fig. 9(b)) were another indicators validating the manifold has been operated in excessive temperatures. On the other hand, FEM simulated results showed that high stress concentration at high temperatures also can significantly accelerated the creep damage. Furthermore, obvious deformation (Fig. 2(a)) and evident tearing characteristic (Fig. 2(d)) of the failure manifold in macroscopic also revealed that the manifold has been subjected to high hoop stress too. 5.1. Cause of failure Most of the evidence revealed that the cause of failure was a combination of exposure to an excessive high temperature and high hoop stress at the critical regions. Exposure to an excessive high temperature at high stress level could have three detrimental effects. Firstly, the creep rate is accelerated leading to the material sustained larger viscoplastic strain during a shorter time. Besides, ductility of the material tends to decrease significantly at higher temperatures because of the appearance of Nb(C, N) precipitates, as explained in Figs. 3 and 10, thus a premature failure happens. Secondly, the formation and growth of grain boundary creep voids are accelerated. In addition to the high stress, creep crack can easily formed due to the coalescing of creep voids as observed in Fig. 6(c) and (d). Thirdly, overheating break the protective Cr2O3 oxidation film, thereby causing an accelerated oxidation attack. Cracking of the hard and brittle oxidation layer provide a good site for small crack initiation, thus causing an accelerated fracture. 5.2. Prevention approaches As obvious overheating of the failure manifold, thus it is strongly recommended to strictly control the operating temperature. It is necessary to check the manifold operating temperature on-line and ensure that the temperature is lower than the design temperature. What's more, decreasing the stress magnitude in stress concentration region is also useful. 6. Conclusions In this work, a failure analysis of the failed hot outlet manifolds made of 20Cr33NiNb used in reformer furnace was carried out based on experimental method and FEM analysis. According to the results obtained above, it could be concluded that the failure can be attributed to the combination of overheating and high hoop stress at critical region. Service overheating can significantly decrease the material ductility and its yield stress. Besides, in combination with high stress concentration, creep rates of the manifold can thus be drastically promoted, and then creep failure happens. Consequently, better control of operating temperature is recommended to prevent the failure, and decreasing the stress concentration is also useful. Acknowledgment The authors acknowledge the support of China Nanjing YPC Refining & Chemical Co., Ltd. for providing samples. References [1] J.J. Hoffman, High temperature aging characteristics of 20Cr32Ni1Nb castings, CORROSION 2000, NACE International, 2000.
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