Print version ISSN 1516-1439
Mat. Res. vol.15 no.2 São Carlos Mar./Apr. 2012 Epub Feb 29, 2012
Lorena Braga MouraI; Rodrigo Freitas GuimarãesI; Hamilton Ferreira Gomes de AbreuI,*; Hélio Cordeiro de MirandaI; Sérgio Souto Maior TavaresII
IDepartamento de Engenharia Metalúrgica e Materiais, Universidade Federal do Ceará - UFC, Campus do Pici, Bloco 702, CEP 60455-760, Fortaleza, CE, Brazil
IIDepartamento de Engenharia Mecânica, Universidade Federal Fluminense - UFF, Rua Passo da Pátria, 156, CEP 24210-240, Rio de Janeiro, RJ, Brazil
One method to face the effects of naphthenic acid corrosion in petroleun refining plants is to use alloys with good resistance to this kind of corrosion. For this purpose, molybdenum additions to chromium containing steels are specially recommended. In this work, experimental Fe-9Cr-xMo (x = 5, 7 and 9 wt. (%)) ingots were cast, forged and hot rolled before being tested in a naphthenic acid environment. Evolution of the mechanical properties with composition is presented. The precipitation and dissolution of phases were investigated with the help of Thermocalc and electron backscattered diffraction (EBSD).
Keywords: FeCrMo, molybdenum, naphthenic corrosion
Corrosion in crude heavy oil processing has become a major problem in the oil industry. Most oil plants have been designed and constructed to process oil with low levels of naphthenic acidity. Discovery of heavy oil reservoirs has transformed the processing of highly acidic oil into a major economic endeavor. Naphthenic corrosion is caused by carboxylic acid oxygenates. The two methods most frequently used to control naphthenic corrosion are blending with oils with low naphthenic acid content to dilute and reduce acidity to an acceptable level for a particular distillation unit and use of materials resistant to naphthenic corrosion. The most appropriate method depends on many factors such as economy, velocity, turbulence, temperature and concentration of acids1.
The presence of naphthenic acids and sulfur complexes considerably increases the corrosion of parts subjected to high temperatures during oil processing, typically higher than about 230 ºC. Two major areas are normally affected, distillation towers (atmospheric and vacuum) and piping circuits, including furnace tubes. According to reports from oil refineries, Mo-bearing stainless steels such as AISI 316 and AISI 317 are the preferred; low-carbon grades are usually specified to also resist wet corrosion during downtime periods. Increasing the Mo content increases naphthenic acid corrosion but a minimum of about 2.3 to 2.5% is indicated to be necessary. There are some works published in the literature that agree with this statement and also specify that the addition of Mo in the alloy considerably increases its resistance to naphthenic acid corrosion2-6.
The use of austenitic stainless steels in distillation units in oil refineries poses its own problems of stress corrosion cracking due to either chlorides during operation and downtime or polythionic acids during down time7. An alternative would be to use ferritic stainless steels that are not susceptible to these stress corrosion cracking mechanisms, with an addition of Mo, to make them resistant to naphthenic acid corrosion.
Another application would be in the tubes inside the furnaces used for oil refining. The material of these tubes is usually 5Cr0.5Mo or 9Cr-1Mo and they do not resist naphthenic acid corrosion. The choice of these steels is based on their resistance to sulfidation. Again, for this application, it would be attractive to refineries to use ferritic steel with a higher amount of Mo in its chemical composition to resist both sulfidation and naphthenic acid corrosion.
Mo stabilizes the chromium passive protective film, limiting naphthenic acid corrosion6. Molybdenum is a strong ferrite stabilizing element. Increasing the amount of molybdenum increases mechanical strength and hardness, and reduces toughness. When dealing with Fe-Cr-Mo steel, some deleterious phases could precipitate: sigma (σ), chi (χ) and mu (µ) during service at high temperature. These phases may occur at temperatures between 600 and 900 ºC, their formation is usually slow, requiring several hours, days, months, or years depending on the exposed temperature. These phases are diffusional and are in equilibrium in a phase diagram, and therefore high cooling rates are required after welding or heat treatment to prevent the formation of these phases8-12.
Rapid cooling keeps the allowing elements in solid solution but rapid cooling would normally harden some of the ferritic steels such as 5Cr0.5Mo and 9Cr1Mo steels. These steels are normally used in the normalized and tempered or in annealed condition.
σ phase is an intermetallic compound of nominal composition Fe-Cr and tetragonal unit cell, which is hard and nonmagnetic. The χ phase, in turn, has an ordered cubic structure, with composition Fe - 25% Cr - Ni 2% - 15-20% Mo - 1% Mn. Both phases, χ and σ, cause embrittlement below about certain temperature and, in particular room temperature, and are undesirable in a steel. They reduce ductility and toughness. Phase µ has a hexagonal crystal structure, a nominal formula Fe7Mo6 and is a hard and brittle intermetallic compound. Phase µ is detrimental to materials, because its precipitation causes depletion of alloying elements in the matrix. In addition, its needle shape facilitates crack propagation13-17.
This work involved investigation of experimental ingots with Cr varying from 9 to 17% and Mo from 5 to 7%. Aspects concerning resistance to naphthenic corrosion, mechanical properties, microstructure and kinetics of deleterious phases were analyzed. A detailed study of these experimental alloys using EBSD (electron backscattered diffraction) provided an understanding of the formation and dissolution of σ, µ, and χ phases.
2. Materials and Methods
Ingots of 90 kg, with chemical compositions shown in Table 1, were melted in an induction furnace. After casting, they were hot forged and hot rolled. The onset temperature for hot rolling ranged between 900 and 1000 ºC. The dimensions of plates after hot rolling were 200 mm wide and 12 mm thick. After hot rolling the alloys were solution annealing to dissolve alloying elements in the matrix homogeneously and to obtain uniform microstructure and properties. Samples were solution annealed for one hour at 1000 ºC (alloys with 5% of Mo), 1060 ºC (alloys with 7% Mo) and 1100 ºC (alloys with 9% Mo) and then water quenched to room temperature.
Samples for naphthenic corrosion were fist machined followed by gridding. After that, all samples were sand blasted. Then, they were weighed on a precision balance (0.1 mg). A reference weight 316 stainless steel was used during the weighing to minimize errors during this procedure.
Naphthenic corrosion tests were carried out in an autoclave (pressure of 30 lb/in2 or psig) and temperature of 340 ºC. Inside the autoclave was installed a rotating carousel where the specimens were fixed and rotated at 1500 rpm. The corrosive media was crude petroleum with high level of acidity (2.5 mg KOH/g). The uniform corrosion rates were determined according to ASTM® G1, with immersion cycles of 30 seconds, using Clark's solution for different compositions of Fe-Cr-Mo alloys. Samples of three commercial alloys, AISI 1020, AISI 316 and 9Cr1Mo, were placed in the autoclave in the same run of the experiment to be used as references.
Thermodynamic calculations had been done to determine possible phases that precipitate in the material. Thermo-Calc software and the TC-FE6 were used. Thermo-Calc is a software package that performs thermodynamic and phase diagrams calculations for multi-component systems. TC-FE6 is a database developed for Steel and Fe-alloys.
Tension tests were done at room temperature in an Instron 8802 tester according to ASTM® E8-03 standards. All tests were performed in samples annealed after hot rolling. Square cross section specimens were machined aligning the length of the specimen with the rolling direction of the sheet. Charpy impact experiments to determine absorbed energy and the fracture surface appearance were conducted according to ASTM® E23-07. Specimens of reduced size (5 ×10 × 55 mm) were machined with a v-notch parallel to the rolling direction.
Specimens for light optical microscopy (LOM) were prepared by a standard method including grinding, polishing and etching.
EBSD data was acquired in an Oxford Crystal 300 system attached to a SEM Philips XL-30. Channel-5 software from HKL was used for processing. The metallographic specimen were polished using colloidal silica for 30 minutes, and the EBSD measurements were performed on a SEM working with a tungsten filament electron gun operating at an acceleration voltage of 25 kV. The EBSD orientation maps have been acquired with a step size of 900 nm and by assuming the presence of four possible crystallographic phases: Fe-BCC for ferrite, Cr6Fe18Mo (χ-phase), Cr-Fe (σ phase) and Fe7Mo6 (µ phase).
3.1. Naphthenic acid corrosion resistance
Six different compositions of FeCrMo alloys, AISI 316 austenitic stainless steel, 9Cr-1Mo steel and AISI 1020 low carbon steel were selected to perform preliminary tests of naphthenic corrosion.
Figure 1 summarizes the naphthenic acid corrosion mass loss per year for six different FeCrMo alloys along with 9Cr1Mo, AISI 1020 steel and AISI 316 austenitic stainless steel. This is referred to as corrosion rate.
Increasing the Mo content from 1 to 5% in alloys with 9% of chromium resulted in a reduction of corrosion rate. However, a further increase to 9 Mo to 9% Cr steel did not further improve the corrosion resistance to this testing medium. The chromium content influenced the naphthenic acid corrosion resistance, as illustrated. Further addition of Mo to 9Cr steel resulted in some additional reduction of the rate of corrosion. Increasing the molybdenum content from 5 to 9% in alloys with 13 and 9% of chromium did not appear to exert a significant influence on the corrosion rates by the action of naphthenic acids. This contrasts the general belief that chromium addition increases the corrosion resistance to sulfidation but does very little contribution to increase the resistance to naphthenic acid corrosion.
3.2. Mechanical properties
Mechanical properties of Cr-Mo steels vary greatly depending on the thermal history. Comparison is thus valid only if they have similar heat treatment conditions. The results of tensile tests on alloys 17Cr5Mo, 15Cr5Mo, 9Cr5Mo, 9Cr7Mo and 9Cr9Mo in the annealed condition are shown in Figures 2a,b. Comparing values for yield and ultimate tensile stress of alloys 17Cr5Mo and 15Cr5Mo (ferritic stainless steel) with the commercial alloy K44 (AISI 444 - 18% Cr and 1.8% Mo) from Arcelor Mittal in the annealed condition, can be concluded that the experimental alloys have higher yield strength, 483.8 and 461.7 MPa against 380 MPa for alloy K44.
The variation in chromium content did not result in a significant change within the limits of yielding stress and ultimate tensile strength among the alloys 17Cr5Mo, 15Cr5Mo and 9Cr5Mo. On the other hand, the increase of Mo produces a large increase in yielding and ultimate stress in alloy 9Cr5Mo, 9Cr7Mo and 9Cr9Mo (Figure 2b). Molybdenum content presents a more relevant effect on mechanical properties than chromium.
Another aspect to analyze is toughness. Figures 3a,b show the values of energy absorbed in a Charpy-V reduced impact testing at 22 ºC (room temperature), and at 60 ºC, in specimens obtained from the alloys 17Cr5Mo, 15Cr5Mo 9Cr5Mo, 9Cr7Mo and 9Cr9Mo. As can be seen, regardless of the levels of chromium and molybdenum, no significant changes were observed in the values of absorbed energy at room temperature. The low values of absorbed energy obtained and appearance of the observed fractures after the Charpy-V impact test indicated that 20 ºC was below the transition temperature for these alloys where they exhibit a brittle-like behavior. However, increasing the test temperature to 60 ºC a considerable increase in placed these alloys above their ductile-brittle transition temperature as shown by a considerable increase in the corresponding absorbed energy. Tests at 60 ºC showed that increasing the Cr and Mo content resulted in decreasing of toughness of these alloys. According to these experiments, Mo has a greater influence on toughness than Cr. Figure 4 presents the appearance of the fracture after Charpy-V tests. The bright aspect of fractured surfaces is characteristic of brittle fracture type.
The determination of the amount of sheared surface was made in accordance to ASTM E 23 standard by measuring length and width of flat fracture region. Table 2 presents the results for shear percent for each composition at 25 and 60 ºC.
3.3. Thermodynamic simulations
Thermo-Calc is a tool that makes possible to predict phases at equilibrium and metastable equilibrium conditions based on the composition of the system, temperature and pressure. Thermodynamic modeling of the alloys was performed using the program TCW5 and database TCFE6. The chemical composition presented in Table 1 was used in modeling. Increasing the molybdenum content causes a reduction in the austenite field (Iron-γ), also causes an increase in the lower limit of austenite and a lower and upper limit of solubility of chromium in the austenite (γ-iron) to a maximum of 2.4% molybdenum (Table 3), when the austenitic field no longer exists. All compositions presented in Table 1 when heated from room temperature to liquid phase temperature do not precipitate austenite. They precipitate ferrite (Iron-α), σ, χ and µ. Figure 6 shows the variation of the weight fraction of the ferrite phase and deleterious phases σ, χ and µ in the alloys 9Cr5Mo and 9Cr9Mo when these alloys were heated from room temperature to liquid phase temperature. The sigma phase in alloy 9Cr5Mo is transformed into ferrite up to 400 ºC; the amount of µ phase remains constant up to 500 ºC, at which point it starts the precipitation of the χ phase. The χ phase reaches the maximum volume fraction at 650 ºC, when it starts its dissolution into ferrite. At 900 ºC this alloy is ferrite-BCC. The alloy with 9% Mo follows a similar pattern, showing differences in the volume fractions of phases and the temperatures at which transformations occur. Complete dissolution of all these alloys occurs above 1050 ºC, when the alloy is 100% ferrite.
Phase identification and the mechanism of dissolution of these phases can be seen on images obtained by EBSD. Figure 7 shows a map of phases of a sample of alloy 9Cr9Mo hot rolled atan initial temperature above 1100 ºC and cooled in air. The figure shows the presence of phases σ (red), µ (blue) and χ (green) at the ferrite (yellow) grain boundary. The phase map also reveals that these phases are concentrated at triple boundary junctions. Several researchers have observed this characteristic in the precipitation of the sigma phase in superferritic stainless steels18-20.
The χ phase is present in small quantities inside the mu and sigma phases (see green areas in Figure 7), and influences the formation of these later phases. Phases χ and µ are favored by the presence of molybdenum. The kinetics of the phases is related to temperature and exposure time. According to Park et al.18, the σ phase dissolves first, favoring the growth of µ and χ phases that continue to exist at higher temperatures. This behavior is in perfect agreement with the calculated amounts of the phases with Thermocalc shown in Figure 6.
Figure 8 shows orientation and phase maps for samples of the alloy 9Cr7Mo annealed at 1060 ºC for 10 minutes and one hour and cooling in air. As shown in Figure 8a (map of phase), χ, µ and σ phases have dissolved and the alloy became 100% α-BCC. The intermetallic phases surrounding ferritic grains that existed prior to annealing dissolves in the ferritic matrix after 10 minutes trial. The grain boundary regions in the ferritic matrix are now occupied by small grains of α-ferrite. With increasing annealing time, the crystallographic orientations of these grains will change tending to the orientation of the ferritic grain that surrounds them. The small ferrite grains are acquiring new orientations aligning with larger grains as indicated by the colors on the phase map in Figure 8b. Notice in the region marked by a circle on the right of Figure 8b that the grains present different shades of green, indicating the trend to form a larger grain with the same orientation. This process of reorientation of low angles grain boundaries and its subsequent grain growth during the transformation of the ferritic matrix is similar to the nucleation process during recrystallization. As subgrains are formed, they are able to grow on their neighbors by migrating sub-boundaries with the help of temperature. The subgrain absorbs dislocations in migration, decreasing the difference in orientation, its energy and its mobility until it becomes part of a larger grain21.
Another model observed during the test is known as nucleation by coalescence of subgrains. This mechanism is the coalescence of two neighboring subgrains equivalent to one rotation, making their crystalline lattices coincident. The coarsening of subgrains promotes grain growth, elimination of subgrains (decreasing the energy stored) and changes of orientation differences between the group affected by the subgrain coalescence and neighbors, which leads to the appearance of a high-angle boundary able to move at high speed21.
When the dissolution process is complete, all intermetallic phases are dissolved, resulting in a ferrite matrix with well-defined grain boundaries and grain orientations without the presence of any precipitation between grain boundaries. This is clearly shown in Figures 8c,d, respectively phase map and orientation map after 1 hour of annealing.
Figure 9 shows the experimental solution annealing temperatures for alloys with 9% Cr and Mo 5, 7 and 9% (blue curve) and alloys with 5% Mo and Cr - 9, 15 and 17% (red curve). Above these temperatures only α-BCC is present in the microstructure. Figure 9 shows that the increase of the amount of Mo has a greater effect on solution annealing temperature than increasing the amount of Cr.
Higher solution annealing temperature results in an increase in grain size. Larger grains reduce ductility and are also one of the factors responsible for increasing of ductile-brittle transition temperature of ferritic alloys. Figure 10 shows the microstructure of the alloy with 9% Cr and 5, 7 and 9% of Mo, respectively. The alloy 9Cr5Mo was solution annealed at a temperature of 1000 ºC; 9Cr7Mo at 1060 ºC; and 9Cr9Mo at 1100 ºC, resulting in very large grains.
It should be noted that these alloys were produced as very low-carbon alloys. Many researchers have observed that carbon in a solution slows the kinetics of recrystallization in iron due to the segregation of carbon in the grain boundaries, reducing grain growth22. In steels, small amounts of elements such as Al, V, Ti and Nb are traditionally used, to control grain refinement at high temperature or during recrystallization22,23. These are examples of actions that can improve grain size in a future study of these alloys.
In this study, Fe-Cr-Mo alloys with levels of 9 to 17% Cr and 5, 7 and 9% Mo were analyzed. The results allow the following conclusions:
Molybdenum content exceeding 5% by weight, besides leading to a series of metallurgical difficulties, does not promote improvements in the naphthenic corrosion resistance in crude petroleum with high level of acidity (2.5 mg KOH/g) in the alloys studied.
Increasing the level of chromium improved the naphthenic corrosion resistance in crude petroleum with high level of acidity (2.5 mg KOH/g) for alloys with 9 and 5% Mo.
The studied alloys showed a high yield strength and tensile strength, but low values of elongation.
At room temperature (25 ºC) alloys showed low values of energy absorbed, and brittle-like behavior. They exhibit ductile-like behavior for temperatures above 60 ºC.
Despite high levels of molybdenum favoring the formation of phases that are harmful to mechanical properties, it was possible to obtain microstructures without the presence of these phases through an adequate annealing.
The percentage of σ phase does not change significantly with the increase of Mo content. µ and χ phases increase with the increase of the Mo content in the alloy. Increasing Mo content also increases the solution annealing temperature of the alloys causing large grain sizes in samples with 9% Mo.
After dissolution of σ, χ and µ phases, a recrystallization process occurs with mechanisms of nucleation by coalescence of subgrains and reorientation of low angle grain boundaries.
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Received: August 10, 2011
Revised: January 23, 2012