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Ultra-strong nickel aluminum bronze alloys with ultrafine microstructures by continuous heavy hot rolling
Journal Pre-proof Ultra-strong nickel aluminum bronze alloys with ultrafine microstructures by continuous heavy hot rolling Shuo Ma, Liming Fu, Xuedong Ma, Aidang Shan PII:
S1044-5803(19)32788-3
DOI:
https://doi.org/10.1016/j.matchar.2019.109986
Reference:
MTL 109986
To appear in:
Materials Characterization
Received Date: 13 October 2019 Revised Date:
27 October 2019
Accepted Date: 28 October 2019
Please cite this article as: S. Ma, L. Fu, X. Ma, A. Shan, Ultra-strong nickel aluminum bronze alloys with ultrafine microstructures by continuous heavy hot rolling, Materials Characterization (2019), doi: https:// doi.org/10.1016/j.matchar.2019.109986. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier Inc.
Ultra-strong nickel aluminum bronze alloys with ultrafine microstructures by continuous heavy hot rolling
Shuo Ma1, 2, 3, Liming Fu1, 2, 3, Xuedong Ma4, Aidang Shan1, 2, 3
1. School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai, 200240, People’s Republic of China 2. Collaborative Innovation Center for Advanced Ship and Deep-Sea Exploration (CISSE), Shanghai, 200240, People’s Republic of China 3. Shanghai Key Laboratory of High Temperature Materials and Precision Forming, Shanghai Jiao Tong University, Shanghai, 200240, People’s Republic of China 4. School of Mechanical Engineering and Automation, University of Science and Technology Liaoning, Anshan, 114031, People’s Republic of China Keywords: Nickel-Aluminum bronze (NAB); Heavy hot rolling; Ultrafine microstructures; Annealing twins; Ultra-strong
Abstract In this work, nickel aluminum bronze (NAB) alloys with superior mechanical properties were prepared by continuous heavy hot rolling at two different temperatures. When rolled at 950°C with a reduction ratio of 90%, slabs with an ultrahigh strength over 1300MPa and an elongation nearly 5% can be obtained, and the strength can be further improved by subsequent ageing process. The contribution of each strengthening factor to the yield stress was analysed using a semiquantitative model. The results show that grain refinement and dislocation strengthening are the two most important factors. Besides, a subsequent 700°C/5min heat treatment was carried out to improve the ductility of rolled slabs to nearly 15%, while the yield strength can be maintained at around 800MPa. After heat treatment at 700°C, high density annealling twins were generated. It was revealed that the intracrystalline twin boundaries helped to refine the microstructures and such refinement benefited the strength and ductility in many ways.
Corresponding authors at: No. 800 Dongchuan Road, Materials Science and Engineering Building A402, Minhang District, Shanghai 200240, China. E-mail addresses: [email protected] (L. Fu), [email protected] (A. Shan).
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1. Introduction Among numerous copper-based alloys, nickel aluminum bronze (NAB) shows outstanding mechanical properties, including high mechanical strength, good corrosion resistance and high fracture toughness[1-3]. In particular, the unique combination of these three features makes it an even more competitive material. NAB alloys are widely used for marine equipment such as valves, pumps and propellers[2, 4]. Typical microstructure of as-cast NAB alloy is identified as the copper-rich α phase, the Cu3Al-rich martensite β′ phase and four kinds of intermetallic precipitates which are based on Fe3Al or NiAl[5-8]. The complex microstructure evolution reserves a large space for NAB alloys to improve their properties in many aspects. Because of the rapid development of industry and the increasingly application of NAB alloys, improving its mechanical properties has been paid much attention to in recent years[1, 4, 9-12]. Microstructure refinement is an effective way to enhance strength and ductility of the metallic material simultaneously. Among the methods to refine microstructures, plastic deformation technique is one of the most important [13, 14]. Many NAB alloys with excellent mechanical performance are prepared by various categories of plastic deformation techniques. Friction-stir processing (FSP) method can optimize the surface properties and prepare refined, homogenized microstructure, while the interior part is hard to be changed[12]. With ECAP (Equal channel angular pressing) methods, microstructures can be significantly refined and mechanical properties can be dramatically improved[10], while it is quite difficult to process large bulk samples[14]. Sarath. K et al[11] has prepared a series of NAB alloys with tensile strength of 450~650MPa, and strain to failure of 10~30% by hot rolling. Since the selection of rolling parameters led to grain growth and excessive decomposition of retained β, the strengthening effects were not significant. Lv et al[9] prepared series of sheet material with superior strength and ductility by isothermal cyclic hot rolling at 850°C. The authors found that the strength of rolled NAB alloy increases with degree of deformation degree. The grain sizes of α phases can be reduced to a few microns. The particle-stimulated nucleation (PSN) of grain recrystallization and martensite twins generated by rolling are also regarded as effective mechanisms. In recent years, a number of methods aimed at refining microstructures of metallic materials by rolling have been implemented and examined by our group [15-18]. All these attempts have achieved drastically improvement on mechanical properties of alloys. To enhance the effect of microstructure refinement and achieve a higher strength, a method of continuous heavy hot rolling without reheating is reported in present work. Heavy hot rolling (HHR) and warm rolling (HWR) are common methods to prepare ultrafine structure metallic materials with high-strength and satisfactory ductility. In medium carbon steel, refinement of microstructures during HHR can help to achieve a superb mechanical properties[16]. Micro-size lamella structures formed in rolling reportedly can generate strengthening effect[17]. In Al-Mg alloy, stronger cube texture developed by HHR was studied, which is attributable to the preferentially recovered structure of cube regions leading to the preferential nucleation of cube grains[19]. In NAB alloy, related studies are still rare. In this study, NAB alloy with ultrahigh strength and sufficient ductility was achieved by HHR and subsequent aging. Strength-plasticity balanced properties were realized by subsequent 700°C/5min heat treatment. Such mechanical properties have reached the level of the strongest copper alloy, beryllium bronze. The evolution of 2
microstructure and tensile properties was investigated by scanning electron microscope (SEM), transmission electron microscope (TEM) and electron backscatter diffraction (EBSD) and tensile tests, respectively. Contribution of various strengthening factors to strength were quantified and high density annealing twins were focused on as a beneficial factor to high ductility.
2. Experimental procedures
2.1 Materials and processing A hot forged ingot of a commercial NAB alloy was prepared in this study, which has a nominal chemical compositions listed in Table 1. Such high contents of iron and nickel can refrain the formation of the brittle phase γ′[9]. The 20mm thick slabs were cut from the ingot for rolling. Heavy hot rolling (HHR) were carried out respectively from 950°C and 850°C (850HR and 950HR). The samples were both kept at the starting rolling temperature for 1h before entering into the rolling mill. The samples were rolled with an accumulative reduction ratio of 90% by three successive passes. A quenching process was followed by the last rolling pass and no reheating treatment was implemented during the whole rolling process. The thickness reduction of three passes was respectively 8mm, 6mm and 4mm. The time spent for the whole rolling process was nearly 20s. For the parameters of the rolling mill, the rolls diameter is 180mm and rotation velocity is 33rpm. After heavily hot rolling, samples were subjected to two kinds of heat treatments, low temperature aging and mid-temperature short-time heating, and then quenched. 350°C for 2h and 700°C for 5min are reported as typical in this work. Table 1 The chemical composition of the experimental NAB aluminum bronze, wt.% Sn
Al
Zn
Mn
Fe
Pb
Si
Ni
P
Others
Cu
＜0.1
9.5-11
＜0.5
＜0.3
4.5~5.5
＜0.02
＜0.1
4.5~5.5
＜0.01
＜1.0
Bal.
2.2 Microstructure characterization Samples for microstructure characterization were cut directly from the rolled and subsequent heat treated slabs. For optical microscopy (OM) and scanning electron microscopy (SEM), the specimens were mechanically polished and etched by a solution of 5g FeCl3+10mL HCl+100mL H2O. OM and SEM observations was respectively conducted using a ZEISS optical microscope and a JEOL JSM-7800F microscope coupled with energy dispersive spectrometer (EDS). The misorientation angle and grain sizes of specimens were determined by electron backscatter diffraction (EBSD) method using the 7800F SEM. A JEM-2100F transmission electron microscopy (TEM) was used to explore the microstructure evolution. For TEM sample preparation, thin plates were first mechanically grinded and polished to around 50µm and then twin-jet electron-polished using a solution of 4% perchloric acid and 96% ethanol solution at -25°C. 3
Phase structures were identified at room temperature by a Rigaku D/max 2500 X-ray diffraction (XRD) instrument with Cu-Kα radiation. The scanning 2θ range is 20 to 100 degree and the scanning rate is 2 deg/min. Data of 2θ range of Cu(111) and Cu(222) were scrupulously scanned to calculate the dislocation density. The Warren-Averbach method was used[20]. Grain size and elastic strain were calculated by data fitting. Samples for tensile test were machined from the center region of the rolling slabs along the rolling direction. The dimension of the gauge is 15mm in length. Tensile tests were conducted on a Zwick/Roell Z020 testing machine under an initial strain rate of 5×10-4 s-1 at room temperature.
3. Results 3.1 Tensile Properties Fig. 1a represents the uniaxial stress-strain curve of as-received, HHRed and subsequent heat treated NAB samples. Specific mechanical properties data are shown in Table 2. When continuous heavy hot rolled, the yield stress (YS) and ultimate tensile stress (UTS) of the NAB alloys can be largely improved. When started rolling at 950°C, a UTS of over 1300MPa can be obtained. Such UTS can be further increased by aging at 350°C for 2h, while the aging will reduce the plasticity of the HHRed samples. Ultimate strength of over 1300MPa in NAB alloy has not been reported in other works. Furthermore, a 700°C/5min heat treatment to HHRed samples will lead to a synergetic strengthening effect, which has an elongation of nearly 14.5% and yield strength of over 780MPa for 950HR samples. It can be clearly seen that 950°C is a better starting rolling temperature than 850°C. For all kinds of processing technics, the improvements of yield strength are remarkable. A digest of mechanical properties of HHRed NAB alloys in this study is shown by red five-pointed stars in light red oval within Fig. 1b. Some properties of NAB alloys in other works mentioned above are shown in ovals with other colors. The large improvement relative to the hot forged state is obvious. A yield strength of nearly 1200MPa is the highest of all reported NAB alloys. Of particular note is that, the strength of HHRed NAB is even higher than a large portion of beryllium bronze[21, 22], a group of strongest copper alloys that contains poisonous, rare-beryllium.
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Fig. 1 Mechanical properties of rolled NAB aluminum bronze a) Engineering stress-strain curves of the samples in this study and b) summary of yield stress-elongation data of samples in other studies
Table 2 Tensile properties of the NAB alloy in this study As-re
950HR
850HR
950HR350°C2h
850HR350°C2h
950HR700°C5min
850HR700°C5min
TS(MPa)
806.5
1309.5
1114.8
1341.1
1190.6
865
827
YS(MPa)
395.2
1098.2
859.2
1187.3
960.3
784.5
684.8
EL (％ ％)
19.8
5.6
5.61
3.41
2.77
14.5
11.1
3.2 Microstructures Fig. 2 shows optical microstructures of the samples before and after rolling. Samples in all states were composed of α phase, β′ phase and many kinds of k phases. The light-etching areas of copper-rich FCC solid solution or α phase, dark-etching regions of β phase transformation products or β′ phase and the uniformly distributed blue-etching tiny intermetallic or k phases can be easily distinguished[23]. In actual rolling process, the β′ phase area shown in Fig. 2a and Fig. 2b should be β phase, which is a BCC solid solution and stable only at temperatures above the eutectoid temperature (approximately 800°C). The microstructure before 850HHR is shown in Fig. 2a, α phase is equiaxed with an average grain size of nearly 10µm. Most k phases distribute on phase boundaries. As seen from Fig. 2b, almost all α phases exhibit Widmanstatten morphology, which usually precipitated rapidly during the quenching process[4]. Therefore, it can be inferred that little α phase exists before rolling. After heavily rolling, microstructures of samples 850HR and 950HR were both refined. α phases were elongated along the rolling direction. The grain sizes of the α and k phases of 950HR were smaller than those of 850HR since the two phases in 950HR precipitated through eutectoid phase transformation (β→α+k) during the rolling process while they barely existed before rolling (Seen from Fig. 2c and Fig. 2d).
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Fig. 2 Optical microstructure of samples before HHR (quenched) a)850HR b)950HR and after HHR c)850HR d)950HR The typical SEM image of 850HR and its EDS mapping results is respectively shown in Fig. 3a and Fig. 3b. Lamella structures of α phase and retained β or β′ phase can be seen (Fig. 3a). After HHRed, some β′ phases penetrated into α phases along the rolling direction, leading to many thin crevices within the elongated primary α phases. Such phenomenon has been reported in FSP process on NAB alloy[24]. From the EDS mapping results (Fig. 3b), varieties of k phases can be distinguished. k2 phases distribute in the Fe-rich area with globular morphology and grain sizes of 1~2µm. Most of the k2 phase precipitates along α/β′ phase boundaries. From Al K and Ni K results of Fig. 3b, Al, Ni-rich laths can be observed, which corresponds to the β transformation products regions. There are many k3 phases precipitating from β phase during the cooling process through the eutectoid transformation. k4 phases are hard to be identified in Fig. 3 since it is too small. k1 phase does not exist in the experimental alloy because the Fe content is not high enough[4]. Fig. 3c shows the magnified microstructure of 850HR after 350°C/2h. After 350°C aging, the shape of α and β′ phase does not change much while the phase boundaries become smoother. Some tiny lamella k3 phases appeared, which means parts of β′ phase decomposed during the heat treatment. The magnified microstructure of 950HR after 700°C/5min heating is shown in Fig. 3d. After 700°C/5min treating, equiaxed α grains grew rapidly through recrystallization. Lots of defects such as dislocations were remained in the samples after rolling, so the nucleation process of recrystallization was apt to take place. There are also many globular nanocrystalline k phases can be seen, which are regarded as mixture of k phases based on NiAl and Fe3Al. 6
Fig. 3 EDS map scanning results of 850HR and the magnified microstructures of rolled samples after heat treatment a)magnified microstructure of 850HR b)EDS map scanning results of a) c)magnified microstructure of 850HR after 350°C/2h heating d) magnified microstructure of 950HR after 700°C/5min heating
3.3 X-ray diffraction pattern analysis Fig. 4 shows the XRD patterns of samples after normalization processing in this study. The main peak, Cu (111), of all samples are normalized to the same height. Bragg diffraction peaks of Cu (111)/ β′ martensite, Cu (200), Cu (220), Cu (311) and Fe3Al/NiAl (110) can be well detected. The analysis method refers to Zhao′s work[25]. Since the main peak of the Cu based solid solution or α phase and β′ martensite are coincident, it is hard to characterize phase transformation between α and β′ phase exactly. As shown by black arrows in Fig. 4, the integral area of the diffraction peak of k phase becomes larger after heavily rolling, which indicates the precipitation of k phases during the rolling process. Broadening of Cu (200), Cu (220), Cu (311) peaks may suggest that the significant reduction of grain (domain) size and the increase of the elastic strains occurred. The three peaks above, apart from the normalized main peak Cu (111), are all decreased, indicating the increasing of the texture in HHRed samples. After 350°C/2h heating, almost no difference from HHRed samples can be observed. While after 700°C/5min heating, some indication of recovery and recrystallization can be found.
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Fig. 4 XRD patterns of samples after normalization processing in this study
3.4 EBSD Characterization Fig. 5 shows EBSD results of 950HR samples. Inverse pole figure (IPF) grain maps and corresponding grain-to-grain misorientation distributions of α phases are presented. k phases are marked with black, while the martensite β′ phases and β decomposition products are hard to be detected. From Fig. 5a, numerous deformed α grains can be observed. After heavy plastic deformation, the fraction of low-angle boundaries (below 15°) is over 60%. Many cells divided by low-angle boundaries within ultrafine grains can be observed. Grain boundaries are distortional and consist of steps. The average grain size is reduced to nearly 300nm. Many sub-grains are generated in the primary equiaxed grains by HHR. After 350°C/2h heat treatment (Fig. 5d), some Σ3 annealing twins (60°) can be observed while the low-angle boundaries still account for the majority. The misoriention angle distribution became more uniform. Recovery and boundary migration happened during aging. The grain sizes are still maintained at around 300nm. When 700°C/5min heat treated (Fig. 5g), recrystallization took place in a large scale. The fraction of Σ3 twin boundaries increased to nearly 30%. Thanks to these twin boundaries, there were still many coherent boundaries within recrystallization grains. It can also be seen that the quantity and average grain size of k phases are both reduced after 700°C heating.
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Fig. 5 EBSD results of 950HR a) and b) 950HR c) and d) 950HR after 350°C/2h heating e) and f) 950HR after 700°C/5min heating
3.5 TEM characterization Fig. 6a shows the dislocation network and dislocation cells in HHRed samples. High density dislocations from many different slip systems tangled with each other to form dislocation cells during severely plastic deformation[26]. The initial grains are divided into many tiny sub-grains by dislocation cells. It can be predicted that these sub-grains will rotate to accommodate further deformation during subsequent tensile processes. Then grains with large-angle misorientation will be formed, which can further refine the microstructures[27]. Fig. 6b shows the large quantities of 9
high density dislocation area, indicating that work hardening make a prominent contribution to the strengthening effect of HHRed samples[9]. Interaction between k precipitations and dislocations can also be observed in Fig. 6b. Dispersed k phases with grain size of around 100nm can be observed in HHRed samples from Fig. 6c. After 350°C ageing, more nanocrystalline k phases precipitate (Fig. 6d).
Fig. 6 TEM images of high strength samples a) and b)950HR c)850HR d) 850HR after 350°C/2h heating Fig. 6 shows the microstructures of the HHRed samples after heating at 700°C for 5min. Recrystallization grains and high density annealing twins are identified. Many recrystallization grains with grain sizes of several hundred nanometers can be observed from Fig. 6a. Some thickness extinction coutours shows non-equilibrium grain boundaries[14]. After heating at 700°C, almost all dislocations within grains disappeared and many nanocrystalline k phases were generated. The selected area diffraction pattern (SADP) embedded in Fig. 6a presents an annular form, which is the typical feature of polycrystalline materials. As seen from Fig. 6b, annealing twin boundaries distribute in high density. SADP embedded in Fig. 6c proves the identification of twins. Nano k precipitations distribute both in α matrix (Fig. 6d) and twin (Fig. 6d). Fig. 6f shows 10
the FFT results of the high resolution diffraction pattern of the annealing twin boundary.
Fig. 7 TEM images of strength-ductility balanced samples (950HR after 700°C/5min heating). a)Recrystallization grains b)high density precipitations and twin boundaries c)bright filed images of annealing twin and its diffraction pattern d) and e)dark images of α matrix and annealing twins respectively f)FFT results of the high resolution diffraction pattern of the annealing twin boundary
4. Discussion 4.1 Strengthening mechanism analysis For the ultrahigh yield strength of over 1100MPa of the samples 950HR (with an ultimate strength over 1300 MPa), ultrahigh density dislocations, ultrafine microstructures and the dispersed ultrafine or nanoscale k phases are regarded as dominant factors. For the subsequent improvement by 350°C aging, k phase precipitation (Fig. 5c and Fig. 6d) and the annealing twin generated (Fig. 5d) through heat treatment are supposed to be main reasons. To demonstrate contribution of various kinds of strengthening factors, the total yield stress was split by the two formulas (1) and (2) [11]. Yield strength of sample 950HR will be practically computed to verify the strengthening mechanisms. = +
( ) = + + +
(1) (2)
From no less than 20 optical microscope images, the fraction of α and β′ phase, fα and fβ , was 11
measured to be 16.6% and 83.4%, respectively. The contribution of the dislocation density to the yield strength can be estimated by the relationship (3)[28], where b is Burgers’ vector (0.256nm[11]), µ is the shear modulus (28.5GPa for Cu-Al alloys[29]), ρ is the dislocation density and Mα is a constant (0.6 for NAB alloy[28]). Calculated by data of X-ray results[25, 30], ρ=3.95×1014m-2 was calculated by Warren-Averbach method. Therefore, σDislocation=274.9MPa. σ =
!"#
%$(3)
The well-known Hall-Petch formula[31] as shown below can be used to estimate the contribution of grain refinement to the yield strength, where d is the grain size, σ0 is the intrinsic strength of the material in single crystal form, including solid solution and dispersion strengthening contributions[32]. +
σ& = σ' + () *,
(4)
Approximating with Cu-15at. pct Al alloy, σ0=27.5MPa and k=750MPa µm1/2 are determined according to Varschavsky’ s work for α phase, and σ0=68.6MPa and k=705MPa µm1/2 of Cu-19at. pct Al alloy for β′ phase[31]. When dα=300nm and d??=2µm is taken, σy of α phase and β′ phase will be 1396.9MPa and 567.2MPa, respectively. When multiplied by a weighting efficient or volume fraction, the total contribution of two phases σy=705.75MPa. Here Ashby’ s theory[33] , formula (5), is used to predicted the dispersion hardening contribution to the yield strength. Where b is Burgers’ vector; µ is the shear modulus; r and l are the particle radius and spacing, respectively. The adjustable parameter r0 is taken equal to 4b[11]. σ =
...01 4 23
2
56
7
(5)
The size (2r) and spacing (l) of globular k4 in α phase and lamella k3 in β′ phase were measured from over 20 high magnification SEM images, which is 97nm and 312nm for α phase and 262nm and 502nm for β′ phase, respectively. After weighting arithmetic, σPrecipitation=15.12MPa. For contribution of solid solution, process of calculating will not be shown in detail, while a range of 20~30MPa can be speculated for σSolid Solution [11]. In conclusion, the σtotal is calculated as 274.9+705.75+15.12+25=1014.3MPa, nearly 70MPa lower than the practical yield stress. The reason of the deviation can be attributed to the error of the model, or some other ignored factors, such as synergistic strengthening effects among α, β′ and k phases. It is sure that the remarkable microstructure refinement is most important factor of the ultrahigh yield stress. 4.2 High yield strength with considerable ductility after 700℃/5min heat treatment When aged at 700°C for 5min, high yield strength with considerable ductility can be achieved, as shown in Table 3. From Fig. 5e and Fig. 5f, after 700°C heating, most distorted α grains generated by HHR recovered and recrystallized to be equiaxed grains, while high density low-angle grain boundaries still existed in some distorted grains. Remarkable quantities of Σ3 twin boundaries were generated in the recrystallization grain interior, which accounted for nearly 30% of whole boundaries. Such a composite structure with relative coarse grains embedded with 12
nanotwin bundles would exhibit a good combination of high strength and ductility[34]. It is believed that these twin boundaries benefits the improvement of both yield stress and ductility[35]. The outstanding comprehensive mechanical properties can also be attributed to the promotion of the dislocation storage capability of the recrystallization grains. Different from deformation twin observed in other studies[9, 25], twins were induced by heat treatment after deformation in this study. These annealing twins exhibit more coherency than deformation twins, facilitating the gliding of dislocations along twin boundaries[14]. In conventional studies, grain size is an important factor that determines annealing twin density[36, 37]. Pande[38] proposed a model that shows the twin density is negative related with grain sizes. The large number of annealing twins emerged in this study (Fig. 7b) may be due to the ultrafine grain sizes generated by HHR. There are multiple strengthening mechanisms of yield strength in regard to annealing twin. The increase of boundaries can divide a primary grain into many parts, which effectively reduces the average grain sizes[39]. In Cu-Al alloys, when using inter-twin spacing as the effective grain size, the calculated yield strength by Hall-Petch relation will be more in accordance with the true value[40]. Moreover, it can be learnt that dislocations need higher applied stress to glide on annealing twin boundaries than other incoherent boundaries during gliding by molecular dynamics simulations[41]. The outstanding properties of annealing twin that can boost the strength and ductility have been reported by Lu[34, 42]. Ductility can be improved with the improvement of containing capacity of dislocations due to twin boundaries[34]. From Fig. 7b, many k precipitations can be observed in the high density twin area, which can have a cooperative effect on blocking dislocations sliding and improving the yield strength. During the recrystallization process under 700°C heat treatment, the quantity of dislocations were rapidly reduced. Dislocations absorption by grain boundaries and boundary migration may be one kind of recovery mechanism. For the fracture elongation of 14.8%, recrystallization grains without dislocations is the main factor, which has a considerable capacity to accommodate dislocations. It is noteworthy that the grain sizes still stay at ultrafine scale. Both high density annealing twin and k precipitates (Fig. 7a) can block the grain boundary migration, which will suppress the grain growth. To sum up, the contribution of the four conventional metallic material strengthening factors to the ultrahigh yield strength of the 950HR samples were respectively calculated. The results show grain-refinement strengthening is the dominant strengthening mechanism, and dislocation hardening is an important factor. In this study, the emergence of the high density annealing twin is a remarkable phenomenon, which is beneficial to the improvement of the strength and can effectively stabilize the grain sizes of the samples. For further research, some other feasible strengthening mechanisms should also be paid attention to, such as the synergistic effects among phases or the strengthening effect of the lamella structure.
5. Conclusions 1. Ultra-strong NAB alloys with a yield strength over 1100MPa, an ultimate tensile strength over 1300MPa and a fracture elongation nearly 5% have been prepared by continuous heavy hot rolling. 13
2. According to strengthening model analysis, dislocations, grain-refinement, precipitations and solid solution all contribute to the strengthening effect. Among these components, grain-refinement strengthening is dominant strengthening mechanism in HHRed samples. 3. Synergetic mechanical properties can be achieved by a 700°C/5min heat treatment after rolled from 950°C, leading to a yield strength of 784MPa and a fracture elongation of 14.8%. High density annealing twins generated by heating are beneficial to the improvement of strength and ductility.
Acknowledgments This work was funded by the National Natural Science Foundation of China (No. 51775258) and financially supported by National Major Science and Technology Project of China (No.2014ZX07214-002). The author (L. Fu) is grateful to the financial support from Startup Fund for Youngman Research at SJTU. (SJTU.18X100040023).
Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.
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Highlight • • •
Nickel aluminum bronze alloys with an ultimate strength of over 1300MPa were prepared. Significant microstructure refinement was achieved by continuous heavy warm rolling. Strength model analysis was carried out to evaluate the contribution of four kinds of strengthening factors.
•
High density annealing twins were characterized and their strengthening mechanisms were discussed.
1
Conflict of Interest
There are no conflicts of interest.
Liming Fu and Aidang Shan
School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai, 200240, People’s Republic of China Tel: 86-21-54748974 Email: [email protected]
Fig. 1 Graphic showing: (a) the tensile engineering stress-strain curves of the NAB alloy for different processing conditions and (b) comparison of the tensile properties in this study with those of other existing results in other studies.
Fig. 2 Typical optical microstructures of the NAB alloy specimens. (a) 850°C annealing and quenched before HHR; (b) 950°C annealing and quenched before HHR; (c) 850HR; (d) 950HR.
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Fig. 3 SEM analysis of the HHRed NAB alloys. (a) 850HR; (b) chemical element distribution of the EDS map scanning of the image (a); (c) 850HR and 350°C/2h aging; (d) 950HR and 700°C/5min annealing.
2
Fig. 4
XRD patterns of the experimental specimens for different processing conditions.
3
Fig. 5 EBSD images including the inverse pole figure maps (IPF/RD) plus boundary distribution map (black line represents misorientation angle θ＞15° and red line represents 2°≤θ≤15°) and the statistical chart of misorientation angle: (a),(b) 950HR; 350°C/2h; (e),(f) 950HR and 700°C/5min annealing.
4
(c), (d) 950HR and
Fig. 6
Typical TEM images of the high strength NAB alloys. (a),(b) 950HR; (c) 850HR; (d) 850HR and 350°C/2h aging.
5
Fig. 7
TEM observation of the strength-ductility balanced specimens with 950HR and
700°C/5min annealing treatment. Typically graphic showing: (a) the recrystallization ultrafine grains , (b) high density precipitations and annealing twins, (c) bright field image of annealing twins and its corresponding diffraction pattern, (d) and (e) dark field images of α matrix and annealing twins, respectively , and (f) FFT results of the high resolution diffraction pattern of the annealing twin boundary.
6
S1044-5803(19)32788-3
DOI:
https://doi.org/10.1016/j.matchar.2019.109986
Reference:
MTL 109986
To appear in:
Materials Characterization
Received Date: 13 October 2019 Revised Date:
27 October 2019
Accepted Date: 28 October 2019
Please cite this article as: S. Ma, L. Fu, X. Ma, A. Shan, Ultra-strong nickel aluminum bronze alloys with ultrafine microstructures by continuous heavy hot rolling, Materials Characterization (2019), doi: https:// doi.org/10.1016/j.matchar.2019.109986. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier Inc.
Ultra-strong nickel aluminum bronze alloys with ultrafine microstructures by continuous heavy hot rolling
Shuo Ma1, 2, 3, Liming Fu1, 2, 3, Xuedong Ma4, Aidang Shan1, 2, 3
1. School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai, 200240, People’s Republic of China 2. Collaborative Innovation Center for Advanced Ship and Deep-Sea Exploration (CISSE), Shanghai, 200240, People’s Republic of China 3. Shanghai Key Laboratory of High Temperature Materials and Precision Forming, Shanghai Jiao Tong University, Shanghai, 200240, People’s Republic of China 4. School of Mechanical Engineering and Automation, University of Science and Technology Liaoning, Anshan, 114031, People’s Republic of China Keywords: Nickel-Aluminum bronze (NAB); Heavy hot rolling; Ultrafine microstructures; Annealing twins; Ultra-strong
Abstract In this work, nickel aluminum bronze (NAB) alloys with superior mechanical properties were prepared by continuous heavy hot rolling at two different temperatures. When rolled at 950°C with a reduction ratio of 90%, slabs with an ultrahigh strength over 1300MPa and an elongation nearly 5% can be obtained, and the strength can be further improved by subsequent ageing process. The contribution of each strengthening factor to the yield stress was analysed using a semiquantitative model. The results show that grain refinement and dislocation strengthening are the two most important factors. Besides, a subsequent 700°C/5min heat treatment was carried out to improve the ductility of rolled slabs to nearly 15%, while the yield strength can be maintained at around 800MPa. After heat treatment at 700°C, high density annealling twins were generated. It was revealed that the intracrystalline twin boundaries helped to refine the microstructures and such refinement benefited the strength and ductility in many ways.
Corresponding authors at: No. 800 Dongchuan Road, Materials Science and Engineering Building A402, Minhang District, Shanghai 200240, China. E-mail addresses: [email protected] (L. Fu), [email protected] (A. Shan).
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1. Introduction Among numerous copper-based alloys, nickel aluminum bronze (NAB) shows outstanding mechanical properties, including high mechanical strength, good corrosion resistance and high fracture toughness[1-3]. In particular, the unique combination of these three features makes it an even more competitive material. NAB alloys are widely used for marine equipment such as valves, pumps and propellers[2, 4]. Typical microstructure of as-cast NAB alloy is identified as the copper-rich α phase, the Cu3Al-rich martensite β′ phase and four kinds of intermetallic precipitates which are based on Fe3Al or NiAl[5-8]. The complex microstructure evolution reserves a large space for NAB alloys to improve their properties in many aspects. Because of the rapid development of industry and the increasingly application of NAB alloys, improving its mechanical properties has been paid much attention to in recent years[1, 4, 9-12]. Microstructure refinement is an effective way to enhance strength and ductility of the metallic material simultaneously. Among the methods to refine microstructures, plastic deformation technique is one of the most important [13, 14]. Many NAB alloys with excellent mechanical performance are prepared by various categories of plastic deformation techniques. Friction-stir processing (FSP) method can optimize the surface properties and prepare refined, homogenized microstructure, while the interior part is hard to be changed[12]. With ECAP (Equal channel angular pressing) methods, microstructures can be significantly refined and mechanical properties can be dramatically improved[10], while it is quite difficult to process large bulk samples[14]. Sarath. K et al[11] has prepared a series of NAB alloys with tensile strength of 450~650MPa, and strain to failure of 10~30% by hot rolling. Since the selection of rolling parameters led to grain growth and excessive decomposition of retained β, the strengthening effects were not significant. Lv et al[9] prepared series of sheet material with superior strength and ductility by isothermal cyclic hot rolling at 850°C. The authors found that the strength of rolled NAB alloy increases with degree of deformation degree. The grain sizes of α phases can be reduced to a few microns. The particle-stimulated nucleation (PSN) of grain recrystallization and martensite twins generated by rolling are also regarded as effective mechanisms. In recent years, a number of methods aimed at refining microstructures of metallic materials by rolling have been implemented and examined by our group [15-18]. All these attempts have achieved drastically improvement on mechanical properties of alloys. To enhance the effect of microstructure refinement and achieve a higher strength, a method of continuous heavy hot rolling without reheating is reported in present work. Heavy hot rolling (HHR) and warm rolling (HWR) are common methods to prepare ultrafine structure metallic materials with high-strength and satisfactory ductility. In medium carbon steel, refinement of microstructures during HHR can help to achieve a superb mechanical properties[16]. Micro-size lamella structures formed in rolling reportedly can generate strengthening effect[17]. In Al-Mg alloy, stronger cube texture developed by HHR was studied, which is attributable to the preferentially recovered structure of cube regions leading to the preferential nucleation of cube grains[19]. In NAB alloy, related studies are still rare. In this study, NAB alloy with ultrahigh strength and sufficient ductility was achieved by HHR and subsequent aging. Strength-plasticity balanced properties were realized by subsequent 700°C/5min heat treatment. Such mechanical properties have reached the level of the strongest copper alloy, beryllium bronze. The evolution of 2
microstructure and tensile properties was investigated by scanning electron microscope (SEM), transmission electron microscope (TEM) and electron backscatter diffraction (EBSD) and tensile tests, respectively. Contribution of various strengthening factors to strength were quantified and high density annealing twins were focused on as a beneficial factor to high ductility.
2. Experimental procedures
2.1 Materials and processing A hot forged ingot of a commercial NAB alloy was prepared in this study, which has a nominal chemical compositions listed in Table 1. Such high contents of iron and nickel can refrain the formation of the brittle phase γ′[9]. The 20mm thick slabs were cut from the ingot for rolling. Heavy hot rolling (HHR) were carried out respectively from 950°C and 850°C (850HR and 950HR). The samples were both kept at the starting rolling temperature for 1h before entering into the rolling mill. The samples were rolled with an accumulative reduction ratio of 90% by three successive passes. A quenching process was followed by the last rolling pass and no reheating treatment was implemented during the whole rolling process. The thickness reduction of three passes was respectively 8mm, 6mm and 4mm. The time spent for the whole rolling process was nearly 20s. For the parameters of the rolling mill, the rolls diameter is 180mm and rotation velocity is 33rpm. After heavily hot rolling, samples were subjected to two kinds of heat treatments, low temperature aging and mid-temperature short-time heating, and then quenched. 350°C for 2h and 700°C for 5min are reported as typical in this work. Table 1 The chemical composition of the experimental NAB aluminum bronze, wt.% Sn
Al
Zn
Mn
Fe
Pb
Si
Ni
P
Others
Cu
＜0.1
9.5-11
＜0.5
＜0.3
4.5~5.5
＜0.02
＜0.1
4.5~5.5
＜0.01
＜1.0
Bal.
2.2 Microstructure characterization Samples for microstructure characterization were cut directly from the rolled and subsequent heat treated slabs. For optical microscopy (OM) and scanning electron microscopy (SEM), the specimens were mechanically polished and etched by a solution of 5g FeCl3+10mL HCl+100mL H2O. OM and SEM observations was respectively conducted using a ZEISS optical microscope and a JEOL JSM-7800F microscope coupled with energy dispersive spectrometer (EDS). The misorientation angle and grain sizes of specimens were determined by electron backscatter diffraction (EBSD) method using the 7800F SEM. A JEM-2100F transmission electron microscopy (TEM) was used to explore the microstructure evolution. For TEM sample preparation, thin plates were first mechanically grinded and polished to around 50µm and then twin-jet electron-polished using a solution of 4% perchloric acid and 96% ethanol solution at -25°C. 3
Phase structures were identified at room temperature by a Rigaku D/max 2500 X-ray diffraction (XRD) instrument with Cu-Kα radiation. The scanning 2θ range is 20 to 100 degree and the scanning rate is 2 deg/min. Data of 2θ range of Cu(111) and Cu(222) were scrupulously scanned to calculate the dislocation density. The Warren-Averbach method was used[20]. Grain size and elastic strain were calculated by data fitting. Samples for tensile test were machined from the center region of the rolling slabs along the rolling direction. The dimension of the gauge is 15mm in length. Tensile tests were conducted on a Zwick/Roell Z020 testing machine under an initial strain rate of 5×10-4 s-1 at room temperature.
3. Results 3.1 Tensile Properties Fig. 1a represents the uniaxial stress-strain curve of as-received, HHRed and subsequent heat treated NAB samples. Specific mechanical properties data are shown in Table 2. When continuous heavy hot rolled, the yield stress (YS) and ultimate tensile stress (UTS) of the NAB alloys can be largely improved. When started rolling at 950°C, a UTS of over 1300MPa can be obtained. Such UTS can be further increased by aging at 350°C for 2h, while the aging will reduce the plasticity of the HHRed samples. Ultimate strength of over 1300MPa in NAB alloy has not been reported in other works. Furthermore, a 700°C/5min heat treatment to HHRed samples will lead to a synergetic strengthening effect, which has an elongation of nearly 14.5% and yield strength of over 780MPa for 950HR samples. It can be clearly seen that 950°C is a better starting rolling temperature than 850°C. For all kinds of processing technics, the improvements of yield strength are remarkable. A digest of mechanical properties of HHRed NAB alloys in this study is shown by red five-pointed stars in light red oval within Fig. 1b. Some properties of NAB alloys in other works mentioned above are shown in ovals with other colors. The large improvement relative to the hot forged state is obvious. A yield strength of nearly 1200MPa is the highest of all reported NAB alloys. Of particular note is that, the strength of HHRed NAB is even higher than a large portion of beryllium bronze[21, 22], a group of strongest copper alloys that contains poisonous, rare-beryllium.
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Fig. 1 Mechanical properties of rolled NAB aluminum bronze a) Engineering stress-strain curves of the samples in this study and b) summary of yield stress-elongation data of samples in other studies
Table 2 Tensile properties of the NAB alloy in this study As-re
950HR
850HR
950HR350°C2h
850HR350°C2h
950HR700°C5min
850HR700°C5min
TS(MPa)
806.5
1309.5
1114.8
1341.1
1190.6
865
827
YS(MPa)
395.2
1098.2
859.2
1187.3
960.3
784.5
684.8
EL (％ ％)
19.8
5.6
5.61
3.41
2.77
14.5
11.1
3.2 Microstructures Fig. 2 shows optical microstructures of the samples before and after rolling. Samples in all states were composed of α phase, β′ phase and many kinds of k phases. The light-etching areas of copper-rich FCC solid solution or α phase, dark-etching regions of β phase transformation products or β′ phase and the uniformly distributed blue-etching tiny intermetallic or k phases can be easily distinguished[23]. In actual rolling process, the β′ phase area shown in Fig. 2a and Fig. 2b should be β phase, which is a BCC solid solution and stable only at temperatures above the eutectoid temperature (approximately 800°C). The microstructure before 850HHR is shown in Fig. 2a, α phase is equiaxed with an average grain size of nearly 10µm. Most k phases distribute on phase boundaries. As seen from Fig. 2b, almost all α phases exhibit Widmanstatten morphology, which usually precipitated rapidly during the quenching process[4]. Therefore, it can be inferred that little α phase exists before rolling. After heavily rolling, microstructures of samples 850HR and 950HR were both refined. α phases were elongated along the rolling direction. The grain sizes of the α and k phases of 950HR were smaller than those of 850HR since the two phases in 950HR precipitated through eutectoid phase transformation (β→α+k) during the rolling process while they barely existed before rolling (Seen from Fig. 2c and Fig. 2d).
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Fig. 2 Optical microstructure of samples before HHR (quenched) a)850HR b)950HR and after HHR c)850HR d)950HR The typical SEM image of 850HR and its EDS mapping results is respectively shown in Fig. 3a and Fig. 3b. Lamella structures of α phase and retained β or β′ phase can be seen (Fig. 3a). After HHRed, some β′ phases penetrated into α phases along the rolling direction, leading to many thin crevices within the elongated primary α phases. Such phenomenon has been reported in FSP process on NAB alloy[24]. From the EDS mapping results (Fig. 3b), varieties of k phases can be distinguished. k2 phases distribute in the Fe-rich area with globular morphology and grain sizes of 1~2µm. Most of the k2 phase precipitates along α/β′ phase boundaries. From Al K and Ni K results of Fig. 3b, Al, Ni-rich laths can be observed, which corresponds to the β transformation products regions. There are many k3 phases precipitating from β phase during the cooling process through the eutectoid transformation. k4 phases are hard to be identified in Fig. 3 since it is too small. k1 phase does not exist in the experimental alloy because the Fe content is not high enough[4]. Fig. 3c shows the magnified microstructure of 850HR after 350°C/2h. After 350°C aging, the shape of α and β′ phase does not change much while the phase boundaries become smoother. Some tiny lamella k3 phases appeared, which means parts of β′ phase decomposed during the heat treatment. The magnified microstructure of 950HR after 700°C/5min heating is shown in Fig. 3d. After 700°C/5min treating, equiaxed α grains grew rapidly through recrystallization. Lots of defects such as dislocations were remained in the samples after rolling, so the nucleation process of recrystallization was apt to take place. There are also many globular nanocrystalline k phases can be seen, which are regarded as mixture of k phases based on NiAl and Fe3Al. 6
Fig. 3 EDS map scanning results of 850HR and the magnified microstructures of rolled samples after heat treatment a)magnified microstructure of 850HR b)EDS map scanning results of a) c)magnified microstructure of 850HR after 350°C/2h heating d) magnified microstructure of 950HR after 700°C/5min heating
3.3 X-ray diffraction pattern analysis Fig. 4 shows the XRD patterns of samples after normalization processing in this study. The main peak, Cu (111), of all samples are normalized to the same height. Bragg diffraction peaks of Cu (111)/ β′ martensite, Cu (200), Cu (220), Cu (311) and Fe3Al/NiAl (110) can be well detected. The analysis method refers to Zhao′s work[25]. Since the main peak of the Cu based solid solution or α phase and β′ martensite are coincident, it is hard to characterize phase transformation between α and β′ phase exactly. As shown by black arrows in Fig. 4, the integral area of the diffraction peak of k phase becomes larger after heavily rolling, which indicates the precipitation of k phases during the rolling process. Broadening of Cu (200), Cu (220), Cu (311) peaks may suggest that the significant reduction of grain (domain) size and the increase of the elastic strains occurred. The three peaks above, apart from the normalized main peak Cu (111), are all decreased, indicating the increasing of the texture in HHRed samples. After 350°C/2h heating, almost no difference from HHRed samples can be observed. While after 700°C/5min heating, some indication of recovery and recrystallization can be found.
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Fig. 4 XRD patterns of samples after normalization processing in this study
3.4 EBSD Characterization Fig. 5 shows EBSD results of 950HR samples. Inverse pole figure (IPF) grain maps and corresponding grain-to-grain misorientation distributions of α phases are presented. k phases are marked with black, while the martensite β′ phases and β decomposition products are hard to be detected. From Fig. 5a, numerous deformed α grains can be observed. After heavy plastic deformation, the fraction of low-angle boundaries (below 15°) is over 60%. Many cells divided by low-angle boundaries within ultrafine grains can be observed. Grain boundaries are distortional and consist of steps. The average grain size is reduced to nearly 300nm. Many sub-grains are generated in the primary equiaxed grains by HHR. After 350°C/2h heat treatment (Fig. 5d), some Σ3 annealing twins (60°) can be observed while the low-angle boundaries still account for the majority. The misoriention angle distribution became more uniform. Recovery and boundary migration happened during aging. The grain sizes are still maintained at around 300nm. When 700°C/5min heat treated (Fig. 5g), recrystallization took place in a large scale. The fraction of Σ3 twin boundaries increased to nearly 30%. Thanks to these twin boundaries, there were still many coherent boundaries within recrystallization grains. It can also be seen that the quantity and average grain size of k phases are both reduced after 700°C heating.
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Fig. 5 EBSD results of 950HR a) and b) 950HR c) and d) 950HR after 350°C/2h heating e) and f) 950HR after 700°C/5min heating
3.5 TEM characterization Fig. 6a shows the dislocation network and dislocation cells in HHRed samples. High density dislocations from many different slip systems tangled with each other to form dislocation cells during severely plastic deformation[26]. The initial grains are divided into many tiny sub-grains by dislocation cells. It can be predicted that these sub-grains will rotate to accommodate further deformation during subsequent tensile processes. Then grains with large-angle misorientation will be formed, which can further refine the microstructures[27]. Fig. 6b shows the large quantities of 9
high density dislocation area, indicating that work hardening make a prominent contribution to the strengthening effect of HHRed samples[9]. Interaction between k precipitations and dislocations can also be observed in Fig. 6b. Dispersed k phases with grain size of around 100nm can be observed in HHRed samples from Fig. 6c. After 350°C ageing, more nanocrystalline k phases precipitate (Fig. 6d).
Fig. 6 TEM images of high strength samples a) and b)950HR c)850HR d) 850HR after 350°C/2h heating Fig. 6 shows the microstructures of the HHRed samples after heating at 700°C for 5min. Recrystallization grains and high density annealing twins are identified. Many recrystallization grains with grain sizes of several hundred nanometers can be observed from Fig. 6a. Some thickness extinction coutours shows non-equilibrium grain boundaries[14]. After heating at 700°C, almost all dislocations within grains disappeared and many nanocrystalline k phases were generated. The selected area diffraction pattern (SADP) embedded in Fig. 6a presents an annular form, which is the typical feature of polycrystalline materials. As seen from Fig. 6b, annealing twin boundaries distribute in high density. SADP embedded in Fig. 6c proves the identification of twins. Nano k precipitations distribute both in α matrix (Fig. 6d) and twin (Fig. 6d). Fig. 6f shows 10
the FFT results of the high resolution diffraction pattern of the annealing twin boundary.
Fig. 7 TEM images of strength-ductility balanced samples (950HR after 700°C/5min heating). a)Recrystallization grains b)high density precipitations and twin boundaries c)bright filed images of annealing twin and its diffraction pattern d) and e)dark images of α matrix and annealing twins respectively f)FFT results of the high resolution diffraction pattern of the annealing twin boundary
4. Discussion 4.1 Strengthening mechanism analysis For the ultrahigh yield strength of over 1100MPa of the samples 950HR (with an ultimate strength over 1300 MPa), ultrahigh density dislocations, ultrafine microstructures and the dispersed ultrafine or nanoscale k phases are regarded as dominant factors. For the subsequent improvement by 350°C aging, k phase precipitation (Fig. 5c and Fig. 6d) and the annealing twin generated (Fig. 5d) through heat treatment are supposed to be main reasons. To demonstrate contribution of various kinds of strengthening factors, the total yield stress was split by the two formulas (1) and (2) [11]. Yield strength of sample 950HR will be practically computed to verify the strengthening mechanisms. = +
( ) = + + +
(1) (2)
From no less than 20 optical microscope images, the fraction of α and β′ phase, fα and fβ , was 11
measured to be 16.6% and 83.4%, respectively. The contribution of the dislocation density to the yield strength can be estimated by the relationship (3)[28], where b is Burgers’ vector (0.256nm[11]), µ is the shear modulus (28.5GPa for Cu-Al alloys[29]), ρ is the dislocation density and Mα is a constant (0.6 for NAB alloy[28]). Calculated by data of X-ray results[25, 30], ρ=3.95×1014m-2 was calculated by Warren-Averbach method. Therefore, σDislocation=274.9MPa. σ =
!"#
%$(3)
The well-known Hall-Petch formula[31] as shown below can be used to estimate the contribution of grain refinement to the yield strength, where d is the grain size, σ0 is the intrinsic strength of the material in single crystal form, including solid solution and dispersion strengthening contributions[32]. +
σ& = σ' + () *,
(4)
Approximating with Cu-15at. pct Al alloy, σ0=27.5MPa and k=750MPa µm1/2 are determined according to Varschavsky’ s work for α phase, and σ0=68.6MPa and k=705MPa µm1/2 of Cu-19at. pct Al alloy for β′ phase[31]. When dα=300nm and d??=2µm is taken, σy of α phase and β′ phase will be 1396.9MPa and 567.2MPa, respectively. When multiplied by a weighting efficient or volume fraction, the total contribution of two phases σy=705.75MPa. Here Ashby’ s theory[33] , formula (5), is used to predicted the dispersion hardening contribution to the yield strength. Where b is Burgers’ vector; µ is the shear modulus; r and l are the particle radius and spacing, respectively. The adjustable parameter r0 is taken equal to 4b[11]. σ =
...01 4 23
2
56
7
(5)
The size (2r) and spacing (l) of globular k4 in α phase and lamella k3 in β′ phase were measured from over 20 high magnification SEM images, which is 97nm and 312nm for α phase and 262nm and 502nm for β′ phase, respectively. After weighting arithmetic, σPrecipitation=15.12MPa. For contribution of solid solution, process of calculating will not be shown in detail, while a range of 20~30MPa can be speculated for σSolid Solution [11]. In conclusion, the σtotal is calculated as 274.9+705.75+15.12+25=1014.3MPa, nearly 70MPa lower than the practical yield stress. The reason of the deviation can be attributed to the error of the model, or some other ignored factors, such as synergistic strengthening effects among α, β′ and k phases. It is sure that the remarkable microstructure refinement is most important factor of the ultrahigh yield stress. 4.2 High yield strength with considerable ductility after 700℃/5min heat treatment When aged at 700°C for 5min, high yield strength with considerable ductility can be achieved, as shown in Table 3. From Fig. 5e and Fig. 5f, after 700°C heating, most distorted α grains generated by HHR recovered and recrystallized to be equiaxed grains, while high density low-angle grain boundaries still existed in some distorted grains. Remarkable quantities of Σ3 twin boundaries were generated in the recrystallization grain interior, which accounted for nearly 30% of whole boundaries. Such a composite structure with relative coarse grains embedded with 12
nanotwin bundles would exhibit a good combination of high strength and ductility[34]. It is believed that these twin boundaries benefits the improvement of both yield stress and ductility[35]. The outstanding comprehensive mechanical properties can also be attributed to the promotion of the dislocation storage capability of the recrystallization grains. Different from deformation twin observed in other studies[9, 25], twins were induced by heat treatment after deformation in this study. These annealing twins exhibit more coherency than deformation twins, facilitating the gliding of dislocations along twin boundaries[14]. In conventional studies, grain size is an important factor that determines annealing twin density[36, 37]. Pande[38] proposed a model that shows the twin density is negative related with grain sizes. The large number of annealing twins emerged in this study (Fig. 7b) may be due to the ultrafine grain sizes generated by HHR. There are multiple strengthening mechanisms of yield strength in regard to annealing twin. The increase of boundaries can divide a primary grain into many parts, which effectively reduces the average grain sizes[39]. In Cu-Al alloys, when using inter-twin spacing as the effective grain size, the calculated yield strength by Hall-Petch relation will be more in accordance with the true value[40]. Moreover, it can be learnt that dislocations need higher applied stress to glide on annealing twin boundaries than other incoherent boundaries during gliding by molecular dynamics simulations[41]. The outstanding properties of annealing twin that can boost the strength and ductility have been reported by Lu[34, 42]. Ductility can be improved with the improvement of containing capacity of dislocations due to twin boundaries[34]. From Fig. 7b, many k precipitations can be observed in the high density twin area, which can have a cooperative effect on blocking dislocations sliding and improving the yield strength. During the recrystallization process under 700°C heat treatment, the quantity of dislocations were rapidly reduced. Dislocations absorption by grain boundaries and boundary migration may be one kind of recovery mechanism. For the fracture elongation of 14.8%, recrystallization grains without dislocations is the main factor, which has a considerable capacity to accommodate dislocations. It is noteworthy that the grain sizes still stay at ultrafine scale. Both high density annealing twin and k precipitates (Fig. 7a) can block the grain boundary migration, which will suppress the grain growth. To sum up, the contribution of the four conventional metallic material strengthening factors to the ultrahigh yield strength of the 950HR samples were respectively calculated. The results show grain-refinement strengthening is the dominant strengthening mechanism, and dislocation hardening is an important factor. In this study, the emergence of the high density annealing twin is a remarkable phenomenon, which is beneficial to the improvement of the strength and can effectively stabilize the grain sizes of the samples. For further research, some other feasible strengthening mechanisms should also be paid attention to, such as the synergistic effects among phases or the strengthening effect of the lamella structure.
5. Conclusions 1. Ultra-strong NAB alloys with a yield strength over 1100MPa, an ultimate tensile strength over 1300MPa and a fracture elongation nearly 5% have been prepared by continuous heavy hot rolling. 13
2. According to strengthening model analysis, dislocations, grain-refinement, precipitations and solid solution all contribute to the strengthening effect. Among these components, grain-refinement strengthening is dominant strengthening mechanism in HHRed samples. 3. Synergetic mechanical properties can be achieved by a 700°C/5min heat treatment after rolled from 950°C, leading to a yield strength of 784MPa and a fracture elongation of 14.8%. High density annealing twins generated by heating are beneficial to the improvement of strength and ductility.
Acknowledgments This work was funded by the National Natural Science Foundation of China (No. 51775258) and financially supported by National Major Science and Technology Project of China (No.2014ZX07214-002). The author (L. Fu) is grateful to the financial support from Startup Fund for Youngman Research at SJTU. (SJTU.18X100040023).
Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.
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Highlight • • •
Nickel aluminum bronze alloys with an ultimate strength of over 1300MPa were prepared. Significant microstructure refinement was achieved by continuous heavy warm rolling. Strength model analysis was carried out to evaluate the contribution of four kinds of strengthening factors.
•
High density annealing twins were characterized and their strengthening mechanisms were discussed.
1
Conflict of Interest
There are no conflicts of interest.
Liming Fu and Aidang Shan
School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai, 200240, People’s Republic of China Tel: 86-21-54748974 Email: [email protected]
Fig. 1 Graphic showing: (a) the tensile engineering stress-strain curves of the NAB alloy for different processing conditions and (b) comparison of the tensile properties in this study with those of other existing results in other studies.
Fig. 2 Typical optical microstructures of the NAB alloy specimens. (a) 850°C annealing and quenched before HHR; (b) 950°C annealing and quenched before HHR; (c) 850HR; (d) 950HR.
1
Fig. 3 SEM analysis of the HHRed NAB alloys. (a) 850HR; (b) chemical element distribution of the EDS map scanning of the image (a); (c) 850HR and 350°C/2h aging; (d) 950HR and 700°C/5min annealing.
2
Fig. 4
XRD patterns of the experimental specimens for different processing conditions.
3
Fig. 5 EBSD images including the inverse pole figure maps (IPF/RD) plus boundary distribution map (black line represents misorientation angle θ＞15° and red line represents 2°≤θ≤15°) and the statistical chart of misorientation angle: (a),(b) 950HR; 350°C/2h; (e),(f) 950HR and 700°C/5min annealing.
4
(c), (d) 950HR and
Fig. 6
Typical TEM images of the high strength NAB alloys. (a),(b) 950HR; (c) 850HR; (d) 850HR and 350°C/2h aging.
5
Fig. 7
TEM observation of the strength-ductility balanced specimens with 950HR and
700°C/5min annealing treatment. Typically graphic showing: (a) the recrystallization ultrafine grains , (b) high density precipitations and annealing twins, (c) bright field image of annealing twins and its corresponding diffraction pattern, (d) and (e) dark field images of α matrix and annealing twins, respectively , and (f) FFT results of the high resolution diffraction pattern of the annealing twin boundary.
6