Engineering applications of high strength alloys are often restricted due to their poor tensile elongation or ductility . Alloys with high yield strength typically exhibit limited strain - hardenability (the difference between tensile and yield- strengths), leading to reduced tensile ductility. Deformation twinning, resulting in high strain hardenability, can lead to enhanced tensile elongation in single phase solid solutions, including high entropy alloys (HEAs). However, alloy systems involving a solid solution matrix strengthened with an intermetallic phase do not exhibit deformation twinning, thus limiting thei r tensile ductility. We have successfully exploited deformation twinning in a novel HEA, strengthened using nano -lamellar ordered multi- component intermetallic precipitates, leading to an exceptionally high yield strength ( ~1630 MPa ), good tensile ductility ( ~15% ), and an ultimate tensile strength ( ~1720 MPa ), higher than any other reported fcc based alloy . Exploiting deformation twinning in a two- phase metal -intermetallic system , offers a new paradigm for addressing the strength -ductility trade -off plaguing alloy design. Keywords: High Entropy alloys, Precipitation Strengthening, Ductility, Twins, High strength The ever -growing need for high strength metallic systems to support extreme environments for terrestrial as well as extra -terrestrial space require s radical changes in alloy development approaches. Alloys can be strengthened via several approaches, including grain- refinement, lattice straining, precipitation- strengthening employing a high- density of refined second phase (s), or multi- phase composite strengthening. The drawback is, in all cases a higher strength leads to a reduction in ductility and the formability of the alloy. Precipitation strengthening is an effective strengthening mechanism, often without substantial loss of ductility. Recent ly, Yang et al . developed a novel multicomponent intermetallic nanoparticle based strengthening show ing promise in developing high strength alloys that retain high ductility1. The idea of multiple princip al elements resulting in high entropy alloys (HEAs) has been accepted widely across the materials research community to develop compositionally unbiased chemically complex alloys for extended tunability of properties. HEAs with concentrated mixture s of transition elements like Co, Cr, Fe , and Ni can provide optimum combination s of properties like strength -ductility, toughness, and corrosion resistance2.The stacking fault energy (SFE) of these alloys , which is another intrinsic property for structural alloys, is much lower as compared to Ni- base superalloy s 3 however their specific strengths are typically much higher as compared to steels4, which makes them very attractive for advanced applications. The most commonly investigated fcc-based CoCrFeNiMn HEA exhibits unusually high ductility and has a superior damage tolerance under cryogenic conditions5. Additionally, metastability of the parent fcc-phase leads to its stress- induced transformation to hexagonal closed- packed ( hcp) phase in CoFeMnNi -base HEAs 6 leading to transformation induced plasticity , ensuring a high strain -hardenability thus enhanced tensile ductility. He et al . 7 studied effect of Ti and Al addition on the microstructure and mechanical properties of FeCoNiCr- based HEA by forming coherent nanosized L1 2 precipitates. The L1 2 ( γ’) precipitates offer the most promising precipitation hardening effects in fcc- based ( γ) alloys, which forms the basis of high temperature Ni- base, Co -base, 8 and Al -base high temperature alloys9. Most optimal property combination was in (FeCoNiCr) 94Ti2Al4 alloy with tensile yield - strength ~ 900 MPa and ductility of 40%7. So far it has been very challenging to a chiev e high tensile strength of >1.5 GPa with a reasonable ductility in these alloys. Interestingly , bcc based HEAs such as based on WNbTaMoHf elements can exhibit extremely high compressive strength but typically have poor tensile ductility10. Ductility has been much more difficult to model due to the complex interplay between various deformation mechanisms. Hierarchical deformation twinning, or multiple generations of twins, can accommodate more strain. Hence, deformation by twinning can be a potential mechanism to enhance strength -ductility synergy as increasing the number density of twin boundaries increases the strength as well as elongation , and strain hardenability of the material, via increasing dislocation storage capacity 11. Here w e used an iterative CALPHAD -based solution thermodynamic approach for alloy design , to create a precipitation -strengthened alloy based on the fcc-L12 architecture. The composition was optimized such that L1 2 phase is stable to 1150°C and volume fraction of L1 2 is greater than 35 % up to 1000°C. The starting point was the Al xCoCrFeNi HEA system , which has been widely investigated12. We varied each of the element s iteratively to increase the phase fraction and high temperature stability of L1 2 phase . Figs. 1(a -b) show the first and last step while the complete iterations are provided in supplementary information). Our alloy composition (Al0.2Ti0.2Co0.7CrFeNi 1.7) has a theoretical density of ~ 7.85 gm/cc (close to mild steel), expected maximum L12 phase fraction ~50 % and temperature stability ~1100°C (Fig. 1 ( c)). The low Co and high Fe concentration in the alloy reduces the raw material cost compared to conventional Ni - base superalloys. The raw material cost of our alloy is estimated to be ~$5.5/lb whereas the cost of most commercial Ni -base superalloy s lies in the range of $7-12$/lb based on current raw material cost13. Additionally , due to higher Fe concentration the SFE of the alloy is expected to be low er, thus enhancing the wor k hardenability and ductility of the alloy3,4. The optimized alloy composition was arc -melted and subsequently thermo- mechanically processed to achieve average fcc grain size s of 30± 0 µm, 1.1±0.24 µm, and 0.67±0.20 µm , for the three conditions, CR -HTA- 800, CR -800, and CR -600, respectively (Fig s. 2 (a -d)). Annealing the severely cold worked alloy directly at aging temperature s results in con current recrystallization of fcc grains and precipitation of the strengthening phase (L1 2), limiting grain growth. The high magnification image s in the inset of Fig s. 2 (a, b) show the size and morphology of the precipitates in CR- HTA -800 and CR -800 conditions , while Fig s. 2 (c -d) show the grain size and precipitates in case of the CR -600 condition. A clear change in L1 2 precipitate morphology is noticed in CR- 800 condition (Fig. 2 (b)), from lamellar near the grain boundaries to spherical within the grains , due to heterogenous discontinuous precipitation versus homogenous nucleation, respectively . The precipitation is completely lamellar in the refined fcc grains in CR- 600 condition (Fig. 2 (d) , highlighted by white arrow ). The engineering stress -strain plots of the three conditions (Fig. 2(e)) show s that CR -HTA- 800 condition ha s a yield strength (YS) of ~980 MPa , ultimate tensile strength (UTS ) of ~1420 MPa , and 22% plastic strain to failure . The CR- 800 condition exhibited a YS of ~1370 MPa, UTS of ~1540 MPa , and 12% ductility. The CR- 600 condition exhibited a rather remarkable strength - ductility combination with YS of ~1630 MPa, UTS of ~1720 MPa , and 15% ductility. The strength -ductility are both higher for CR -600 as compar ed to CR-800; overcoming the well - established strength -ductility trade -off paradigm (Fig. 2 (e)). All three conditions exhibit a distinct four-stage work -hardening response, which is very characteristic of a low SFE material (Fig . 2 (f))14. The constant work hardening rate in stage B (black arrows) ha s been often correlated to deformation twinning15. This alloy demonstrates an outstanding strength- ductility combination as shown in supplementary Fig . S4 where the YS vs elong ation has been plotted for this alloy in comparison with various super alloys. To understand the microstructural origin behind these excellent and wide- ranging mechanical properties, more detailed characterization was carried out on the CR -HTA- 800 and CR -600 samples. TEM diffraction patterns from all the three conditions confirmed that the precipitates were ordered L1 2 precipitates. Dark field TEM images highlighting the L1 2 precipitates in CR - HTA- 800 and CR -600 conditions, recorded using an ordered superlattice diffraction spot (encircled in the inset s) are shown in Figs . 3 (a -b), respectively . Further the APT analysis show s the three -dimensional (3D) morphology, distribution , and compositions of the se precipitates . An iso-concentration surface delineating the region with Al higher than 8 at. % has been used to demarcate the L12 precipitates in the 3D reconstructions shown in Fig s. 3(c-d) for CR -HTA- 800 and CR -600 conditions , respectively. The precipitates have a spherical morphology with an average diameter of ~50 nm in CR -HTA- 800 condition, while a nano- lamellar morphology with ~10 nm average width in CR -600 condition. The nano- lamellar precipitation in CR -600 is a result of concurrent recrystallization of fcc grains coupled with L1 2 precipitation. The low temperature recrystallization limited fcc grain -growth, and discontinuous L12 precipitation initiating from the grain boundaries16 dominates over the more commonly observed homogeneous L1 2 precipitation within fcc matrix grains (details in supplementary Fig. ST9) . The composition profiles across the fcc /L1 2 interface are shown as proxi mity histograms in Fig s. 3 (e -f) and the phase compositions have been listed in supplementary Table ST2. Post-deformation microstructures exhibited a high fraction of deformation -twins and shear -bands (SEM -EBSD analysis of CR -HTA- 800 sample in Fig. 4 (a)) . TEM from one such twinned region (highlighted by box) is shown in Fig s. 4 (b -d). The bright -field (BF ) TEM image in Fig. 4 (b) further shows the presence of secondary twins within the primary deformation -twin. Secondary twins are also observed within the annealing twins after deformation (refer supplementary Fig . S10). The SAD patterns in Figs. 4 (c -d) were recorded on the primary twin -matrix boundary and primary twin -secondary twin boundary, respectively , and clearly show the respective fcc crystal rotation s resulting from twinning. Similarly , the deformed sample of the CR-600 condition (after strain to failure) also showed a high density of nano- twins (Fig. 4 (e)) within the sub-micron fcc grain s. Typically, twinning is not a dominant deformation mechanism in such fine -grained fcc microstructures 17. The diffraction pattern shown as an inset in Fig. 4 ( e) exhibits twinned reflections from both fcc matrix and L12, indicating that the deformation- twins pass through both. An atomic resolution TEM image of a twin plane is shown in Fig. 4 ( f) and based on the image shown in Fig. 4 ( g), the thickness of these twins can be assessed to be ~6 nm. The excellent yield strength and strain hardening of this HEA stems from a combination of solid solution strengthening and grain -boundary strengthening (or Hall -Petch strengthening) within the compositionally concentrated fcc-matrix, coupled with homogenous precipitation of L1 2 phase. In the CR-HTA- 800 condition this leads to a YS of ~950 MPa, a high strain -hardening rate resulting from hierarchical twinning resulting in a UTS of ~1400 MPa, and ~ 20% tensile ductility . The high degree of deformation twinning indicates that the SFE has been substantially lowered in the matrix due to preferential partitioning of high SFE elements, Al , Ti, and Ni, into the L1 2 precipitates4,18. The further enhanced YS (~1630 MPa) in CR-600 condition can be attribu ted to finer size of recrystallized fcc grains (Hall- Petch strengthening) coupled with the nano- lamellar nature of the L12 precipitates. The Hall Petch and precipitation strengthening are calculated separately for the alloy (refer Supplementary Fig. S13 and Table S T4) and k HP (Hall Petch coefficient) is estimated to be ~ 516 MPa(µm)0.5. The twinning induced plasticity (TWIP) is important for high strain hardening rates observed in a variety of steels 4, titanium alloys 19, and more recently in single phase high entropy alloys5,20. However, typically twinning is not a dominant deformation mode at room temperature in fcc+L1 2 precipitation -strengthened alloys 21. In this study, deformation twinning , resulting in the TWIP effect, was observed in both coarse grained and ultra -fine-grained fcc+L12 conditions . Additionally, i n the CR-600 condition, electron diffraction evidence indicates twinning in both the matrix as well as in the L1 2 precipitates. Such a TWIP effect in microstructures containing intermetallic precipitates has not been reported so far in room temperature deformation and observing it in a HEA indicates that the L1 2 precipitates in these alloys, could potentially exhibit enhanced tendency for twinning. Previously, SF fringes and isolated faults have been reported to form in L1 2 precipitates under high temperature creep deformation 21,22. While the underlying rationale requires further investigation, it can be speculated that the complex site occupancies in these L1 2 precipi tates within HEAs presumably leads to changes in the nature of bonding 23 and lowering of the anti-phase boundary (APB) energy, facilitating growth of extended dislocations into deformation twins 22,24. In conclusion, we designed a new alloy first optimizing the composition for high precipitate fraction and high temperature stability using equilibrium thermodynamics. The thermomechanical treatments results in a n ano-lamellar precipitation of L1 2 phase in a l ow SFE disordered fcc matrix. Our alloy offers a low -cost material with an extremely high specific strength of 218 MPa/g/cc. The specific strengths of Ni -base superalloys are <150 MPa/g/cc and most steels are <200 MPa/g/c25. These new-generation HEAs or complex concentrated alloys could lead to superior structural properties, which are of both great fundamental and applied importance for advanced engineering systems. Our strategy can also be feasibly applied to many other alloy syst ems, such as steels, nickel or cobalt -base superalloys, and HEAs, to achieve enhanced properties. Acknowledgements The work was supported by cooperative agreement between the US Army Research Laboratory and the University of North Texas [grant number W911NF -16-2-0189], and the US Air Force Office of Scientific Research [grant number FA9550 -17-1-0395] . The authors acknowledge the Materials Research Facility (MRF) at the University of North Texas for use of microscopy facilities Figure captions Figure 1: CALPHAD based alloy design strategy (a, b) isopleths showing phase stability as a function of temperature and (a) Al, (b) Ti in base alloy. (c) The phase fraction vs temperature of the optimized composition. Figure 2: (a) SEM images from (a) CR -HTA- 800, (b) CR -800 (c -d) CR- 600 conditions. (e) The tensile en gineering stress vs strain and (f) w ork-hardening plots for the three conditions. Figure 3: Dark Field TEM micrograph showing L1 2 precipitates in (a) CR-HTA- 800 and (b) CR - 600 conditions. The APT results from (c, e) CR- HTA- 800 and (d, f) CR -600 conditions Figure 4: Post deformation results from (a -d) CR -HTA- 800 and (e -h) CR- 600.